UNSW Photovoltaics Annual Report 2012

Transcription

Never Stand Still
Faculty of Engineering
ARC Photovoltaics Centre of Excellence
ARC PHOTOV
P
OVOLTAICS
CS
CENTRE OF EXCELLENCE
ANNUAL REPORT
2012
12
Annual Report 2012
(Final Draft)
UNSW
The University of New South Wales
UNSW Sydney NSW 2052
Australia
Tel
Fax
61 2 9385 4018
61 2 9662 4240
http://www.pv.unsw.edu.au
CONTENTS
Page
1. DIRECTORS’ REPORT
1
2. HIGHLIGHTS
4
3. STAFF LIST
13
4. RESEARCH
4.1. INTRODUCTION TO RESEARCH
19
4.2. FACILITIES AND INFRASTRUCTURE
24
4.3. FIRST GENERATION: WAFER-BASED PROJECTS
37
4.4. SECOND GENERATION: SILICON, ORGANIC
92
AND OTHER “EARTH ABUNDANT” THIN-FILMS
4.5. THIRD GENERATION: ADVANCED CONCEPTS
169
4.6. SILICON PHOTONICS AND DEVICE CHARACTERISATION
225
5. ORGANISATIONAL STRUCTURE
241
6. PUBLICATIONS
243
1.
DIRECTORS’ REPORT
Photovoltaics involves the direct conversion of light, normally sunlight, into electricity when falling
upon devices known as solar cells. Silicon is the most common material used to make these
photovoltaic cells, similarly to its key role in microelectronics.
The Australian Research Council (ARC) Photovoltaics Centre of Excellence commenced at the
University of New South Wales (UNSW) on 13th June, 2003 with funding initially until 31st
December 2007. Upon review in 2005 and 2006, this funding was extended until 31st December 2010.
Criteria for such extension included “progress of the Centre towards becoming independent of the
ARC Centres of Excellence Scheme for funding”. With the rapid growth of the photovoltaics industry
since the Centre’s commencement and the Centre’s strong and seminal links to this industry, the
Centre was clearly in a much stronger position for independent operation beyond 2010 than many
others in its funding cohort. Accordingly, while remaining under the ARC Centres of Excellence
umbrella beyond 2010 in recognition of the quality of its research, the Centre is now dependent upon
industry-related funding for research with near-term outcomes and upon more academicallyorientated ARC and international schemes for its long-term research. In recent years, project funding
provided by the Australian Solar Institute has been particularly important.
The Centre‘s mission remains to advance silicon photovoltaic research on three separate fronts, as
well as to apply these advances to the related field of silicon photonics. The educational activities of
the former Key Centre for Photovoltaic Engineering are also integrated into the Centre and into the
School of Photovoltaic and Renewable Energy Engineering (SPREE), which now hosts the Centre.
Over the period of funding, photovoltaics has become the world’s most rapidly growing energy
source, with markets increasing by over 40 times over this period. The electricity generating capacity
of new photovoltaic product installed in 2006 exceeded new nuclear power capacity for the first time,
with the gap widening significantly every year since. In Europe, photovoltaics has been the largest
source of new electricity generating capacity for each of the last two years, ahead of wind generators
and gas turbines. Reducing prices, at least partly traceable to past Centre initiatives, means the
technology has reached “retail grid parity” in many parts of the world, where the cost of electricity
generated using photovoltaics can compete at the point of use with normal retail electricity prices in
many parts of the world including most of Australia. This is expected to lead to a period of selfsustaining growth with photovoltaics positioned to become one of the world’s major industries of the
present century.
Most present photovoltaic sales are of “first-generation” solar cells made from silicon wafers, similar
to the wafers used in microelectronics. The Centre maintains its world-leadership with these “firstgeneration” devices, with international records for the highest-performing silicon cells in most major
categories (including the outright highest-performing cells and modules). First-generation Centre
research addresses the dual challenges of reducing cost and further improving efficiency. The rapid
growth of the industry is generating widespread interest in ongoing innovations of the Centre’s first
generation technology with several distinct technologies now in large-scale production with annual
sales of technology under licence now approaching several billion dollars.
Silicon is quite brittle so silicon wafers have to be reasonably thick, presently more than a tenth of a
millimetre, to be sufficiently rugged for processing into solar cells with reasonable yield. Without this
mechanical constraint, silicon would perform well even if very thin, even 100 times thinner than
present wafers. Centre researchers have pioneered an approach where very thin silicon layers are
deposited directly onto a sheet of glass with the glass providing the required mechanical strength. This
“second-generation” approach gives enormous potential cost savings. Not only are the costly
processes involved in making wafers no longer required, but also there is an enormous saving in
silicon material. Cells also can be made more quickly over the entire area of large glass sheets. The
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Centre is at the forefront of international research with such “second-generation”, silicon based
approaches, with commercial product from “spin-off”, CSG Solar, powering several megawatt fields
in Europe. The Centre is now investigating electron-beam evaporation of silicon onto the glass
substrate, a much quicker process than the plasma-enhanced, chemical deposition processes used to
date, and also diode laser processing of the deposited films.
Steps to broaden “second generation” activities from silicon to both silicon and carbon were boosted
by the appointment of a new academic staff member, Dr Ashraf Uddin, in 2009, with this organic
solar cell work supported by an ARC Discovery Grant. The activities were further broadened in 2010
by the award to Centre Postdoctoral Fellow Dr Xiaojing Hao of an Australian Solar Institute
Fellowship to develop “earth abundant” CZTS (copper-zinc-tin-sulphide) solar cell technology. By
combining these four relatively benign and abundant elements in appropriate proportions, a
tetrahedrally co-ordinated “synthetic silicon” compound can be produced with some advantages over
silicon, such as stronger light absorption and more potential for bandgap control by alloying with
similar compounds.
The “second generation” thin-films have a clear potential cost advantage over the silicon wafer-based
approach, due mainly to reduced material costs. In large enough production volumes, even these
reduced material costs will dominate thin-film costs. This has led to the Centre’s interest in advanced
“third-generation” thin-film solar cells targeting significant increases in energy-conversion efficiency.
Higher conversion efficiency means more power from a given investment in materials, reducing
overall power costs. The Centre’s experimental program in this area is concentrating on “all-silicon”
tandem solar cells, where high energy-bandgap cells are stacked on top of lower-bandgap devices.
The silicon bandgap is controlled by quantum-confinement of carriers in small silicon quantum-dots
dispersed in an amorphous matrix of silicon oxide, nitride or carbide. Cells based on “hot” carriers are
also of great interest since they offer the potential for very high efficiency from simple device
structures. Although their implementation poses daunting challenges, considerable progress on
addressing these continued to be made during 2012.
The final Centre research strand involves silicon photonics where the emphasis is upon using our
experience with solar cells, using light to produce electricity, to the reverse problem of engineering
silicon devices that use electricity to produce light. The Centre holds the international record for the
light emission performance from bulk silicon, in both electroluminescent and photoluminescent
devices. Emphasis is now upon exploiting our expertise in silicon light emission to develop new
techniques for silicon wafer and cell characterisation. A Centre “spin-off” company, BT Imaging,
remains the premium equipment supplier internationally using this approach.
In addition to these four research strands, the activities of the former Key Centre for Photovoltaic
Engineering have been integrated into the ARC Centre of Excellence. The ninth year of students from
the Bachelor of Engineering (Photovoltaics and Renewable Energy) program graduated during 2012.
This program has been enormously successful, attracting some of the best and brightest students
entering the University and providing the human resources to fuel the recent growth of the industry.
The seventh year of students have now graduated from the Centre’s second undergraduate program,
leading to a Bachelor of Engineering (Renewable Energy).
A significant relevant development during 2010 was the blossoming into full scale operation of the
Australian Solar Institute, of which the Centre was a founding member given its key role in the
Institute’s establishment. The Institute has announced the outcome of several rounds of funding with
some notable successes for the Centre. The Institute was absorbed into the Australian Renewable
Energy Agency (ARENA) at the end of 2012, with indications the increased resources of this agency
will help accelerate the local development and deployment of photovoltaics.
We thank all those who contributed to the Centre’s success during 2012. The present continues to be a
very exciting time for photovoltaics. More are recognising the possibility of a future where solar cells
provide a significant part of the world’s energy needs, without the environmental problems and
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escalating costs associated with the traditional approaches. The recent transition to “grid parity” is
expected to make the coming decade a time of rapid change for the industry, with an accelerated pace
of adoption of newer, lower cost technologies pioneered by the Centre.
………………………………………………
………………………………………………
Professor Martin A. Green,
Executive Research Director
Professor Stuart R. Wenham,
Director
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2.
HIGHLIGHTS
Joint Australian-US research Institute Led by UNSW
A new joint Australian–US research institute led by UNSW aims to foster rapid development of “over
the horizon” photovoltaic technology and establish Australia as the solar cell research and education
hub of the Asia-Pacific region. Federal Minister for Resources and Energy Martin Ferguson
announced the $35 million Australia-US Institute for Advanced Photovoltaics (AUSIAP) on 13
December 2012 as part of a raft of new programs and projects through the United States–Australia
Solar Energy Collaboration. The funding is part of an $83 million commitment, one of Australia’s
biggest solar research investments.
The new Institute is led by UNSW’s world record-holding photovoltaic researchers. It provides a
pathway for research collaboration between Australian and American centres and agencies and will
drive innovation and commercialisation.
“We welcome such a significant government investment in renewable energy research and we are
pleased that UNSW’s world-leading solar researchers have been recognised in this way,” said UNSW
Vice-Chancellor Professor Fred Hilmer.
Program leader Professor Green said: “The Institute (AUSIAP) will establish Australia as the
photovoltaic research and educational hub of the Asia-Pacific region. It combines our expertise with
America’s world-class facilities and creates a tangible pipeline to ‘over the horizon’ photovoltaic
technology”. “The Institute will also be fundamental to the training of the next generation of
photovoltaic research scientists and engineers,” he said.
The USAIAP will work with the recently announced National Science Foundation/US Department of
Energy Research Center for Quantum Energy and Sustainable Technologies (QESST). QESST is
based at Arizona State University, but involves other key US groups including Caltech, MIT and
Georgia Tech.
Other US collaborators include Stanford University, University of California – Santa Barbara, and
Sandia National Laboratories. Local partners include the Australian National University, University
of Melbourne, Monash University, University of Queensland, the CSIRO, and the NSW, Victorian
and Queensland governments. Also playing key roles are Australian and US-based companies – large
entities Suntech Australia, BT Imaging, Trina Solar Energy, BlueScope Steel and smaller start-ups.
The AUSIAP is supported by both the Australian Solar Institute and the US National Renewable
Energy Laboratory.
The Centre will also take the lead in another research project announced by the Federal Government
through the USASEC Open Funding Round, a $6.7 million project to produce Low cost, high
efficiency copper-zinc-tin-sulphide (CZTS) on silicon multi-junction solar cells in partnership
with Suntech R&D Australia, the National Renewable Energy Agency (NREL), and the Colorado
School of Mines.
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Professors Martin Green, Staurt Wenham and Suntech Win Collaborative Innovation Award
Athena Prib from NewSouth Innovations and Scientia Professor Stuart Wenham
accepting the Collaborative Innovation Award on behalf of the collaborative team.
Centre-developed technology has been recognised with a Collaborative Innovation Award at the
Cooperative Research Centres Association Conference in Adelaide on 16 May 2012.
Scientia Professors Stuart Wenham and Martin Green at the ARC Photovoltaics Centre of
Excellence at UNSW developed the Pluto technology in close collaboration with Suntech, the world’s
largest solar cell manufacturer, with funding from the Australian Solar Institute.
The Pluto technology is based on the UNSW-developed PERL cell, which in 2008 set a new worldrecord for performance with 25 per cent efficiency.
The Pluto cell derivative is more economically viable to produce on a large scale, and recently
surpassed the 20 per cent efficiency barrier for converting the sun’s radiation into electricity, setting a
new benchmark for low-cost, mass produced silicon solar cells.
The cell features a unique texturing process that improves sunlight absorption, even in conditions of
low and indirect light, which enables higher efficiency.
“We recently broke through the 20 per cent target for solar cell efficiency, which many experts
thought was impossible and we’ve significantly lowered the costs compared to other technologies,”
says Green. “We believe Pluto technology has struck the ideal balance between conversion efficiency
and manufacturing costs to create a truly viable alternative for electricity production.”
“As we continue to refine the Pluto technology and push up the conversion efficiency, we have no
doubt that it will capture an increasing share of the global solar market,” says Professor Wenham,
who noted the importance of the partnership with Suntech, which allowed the cell to move beyond a
lab-scale device and become a commercial reality.
Dr Richard Corkish, who heads the School of Photovoltaic and Renewable Energy Engineering at
UNSW, says Australia is in a unique position to capitalise on its expertise in solar photovoltaics, and
could become a type of solar energy ‘Silicon Valley’ for research and development.
“Australia is at the ends of the Earth geographically – so you have to be really special to be at the
Centre of something from here, and we have that special status,” he says.
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ARC Linkage Grant Awarded For Buried Contact Cell Research
One of the world's most successful photovoltaic technologies is being "refined" for commercial
release by the University of New South Wales.
“The UNSW Buried Contact Solar Cell (BCSC) had over US$1 billion of product manufactured by
one of UNSW's licensees, BP Solar", said A/Prof CheeMun Chong. "Over the last twenty years we
have made some incredible improvements to this robust technology and now the Australian Research
Council has awarded us funding to take these improvements into production”.
The BCSC was implemented by BP Solar as their “SATURN” technology and reached efficiencies of
17% in commercial production in 2003. This is a significant result considering the production
efficiency of competing screen-printed solar cells was 1-2% (absolute) lower at that time.
The BCSC technology is listed amongst Australia’s Top 100 Inventions of the 20th Century as
determined by the Australian Academy of Technological Sciences and Engineering. Over the past
thirty years, screen printed cells silicon solar cells have dominated the PV market. Despite their large
market share, several fundamental limitations of screen printed solar cells have limited their
conversion efficiency. World-leading solar cell research at UNSW has shown that these limitations
can be overcome. The BCSC is one technology where improvements in solar cell structure have been
successfully transferred to the large scale production of large area, low cost commercial cells.
Structure of the UNSW Laser Grooved Solar Cell
(inset: A/Prof CheeMun Chong holding a laser grooved cell).
“UNSW has been awarded an Australian Research Council (ARC) Linkage Grant to support research
in improving the performance of the BCSC. Key additional areas of focus include the adoption of a
passivated rear surface as well as adapting the technology to suit new world-leading low cost silicon
wafers”.
“This project aims to increase the average commercial silicon solar cell efficiency from the current
17-19% to 23%. This improvement means that we will be able to generate 20% more electricity for
about the same cost and this is expected to be worth hundreds of billions of dollars per year if adopted
by the majority of industry. There is also the flow-on effect of significantly reducing the balance of
system costs as the higher efficiency solar products require less area, mounting structures, wiring and
transportation. The successful development of high efficiency, low cost solar cells will fast-track the
path for PV to reach grid parity in most countries world-wide”, said A/Prof Chong.
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Professor Allen Barnett named one of “The 50 Most Influential Delawareans of the Past 50 Years”
Professor Barnett (second from
left, front row) with his students
at UNSW.
Professor Allen Barnett was named one of “The 50 Most Influential Delawareans (US State of
Delaware) of the Past 50 Years” in 2012. He joined the School of Photovoltaics and Renewable
Energy Engineering as Professor of Advanced Photovoltaics in September 2011. At UNSW, his
research is focused on new high-efficiency solar cell modules, thin crystalline silicon (20+percent),
and tandem solar cells on silicon (30+percent). He joined University of Delaware (UD) in 1976 as
director of the Institute of Energy Conversion and Professor of Electrical Engineering. He left UD in
1993 to devote full time to AstroPower Inc., which became the largest independent solar cell
manufacturer and the fourth largest in the world. He returned to UD in 2003 and was executive
director of the Solar Power Program; research professor in the Department of Electrical and Computer
Engineering and Senior Policy Fellow in the Center for Energy and Environmental Policy.
Professor Barnett currently supervises a large group of doctoral students at UNSW. He is a fellow of
the Institute of Electrical and Electronic Engineers (IEEE) and received the IEEE William R. Cherry
Award for outstanding contributions to the advancement of photovoltaic science and technology and
the Karl W. Böer Solar Energy Medal of Merit. He has more than 280 publications, 28 U.S. patents,
and seven R&D 100 Awards for new industrial products. He actively consults for government
agencies, institutional investors, and private companies.
Move To Tyree Energy Technology Building
The Tyree Energy Technology Building (TETB) was completed in early 2012 on the lower campus to
house the Centre’s main laboratories. Centre staff moved to the new building soon after with
laboratories moving more gradually from their present site during 2012 and 2013. The new $155
million building was supported by Sir William Tyree as well as by a grant of $75 million from the
Australian Government’s Education Investment Fund. This grant was awarded after a presentation by
Centre researchers, Professors Stuart Wenham and Martin Green as well as by then doctoral student,
Nicole Kuepper, to the Selection Committee emphasising the past achievements of the Centre and the
impact of the new building on boosting Centre activities. Co-located in the building are other UNSW
energy-related activities.
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The new Tyree Energy Technologies Building (TETB) in operation in January 2012
to house Centre research and teaching (inset: interior of TETB).
Professor Martin Green Appointed a Member of the Order of Australia
Scientia Professor Martin Green appointed
member of the Order of Australia (AM)
Australia Day 2012.
Professor Martin Green's efforts in solar photovoltaics has been recognised on Australia Day, with
the professor being made a Member of the Order of Australia for service to science education as an
academic and researcher, particularly through the development of photovoltaic solar cell technology,
and to professional associations.
"We are all very proud of Professor Green", said Dr Richard Corkish (Head of School) speaking to his
staff. "This high honour is well deserved and I am sure you will join me in congratulating Martin and
thanking him for his continuing contribution and leadership."
The Order of Australia recognises outstanding achievement and service in our nation. Nominations to
the Order of Australia come directly from the community, which are then considered by the 19member Council for the Order of Australia.
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Professors Martin Green, Stuart Wenham and Dr Muriel Watt Inaugural Members of the
Australian Solar Hall of Fame
Scientia Professors Martin Green and Stuart Wenham were honoured with the unveiling of the
Australian Solar Hall of Fame on 6 December 2012, an initiative of the Australian Solar Council.
"Australia has long led the world in solar research and development, and the Solar Hall of Fame
honours the brightest lights in our solar system", said John Grimes, Chief Executive of the Australian
Solar Council.
"The inaugural inductees in the Solar Hall of Fame include some of Australia's greatest scientists,
innovators and entrepeneurs and collectively represent an extraordinary contribution to tackling
global climate change".
"Professors Martin Green and Stuart Wenham of the University of New South Wales have
revolutionised solar research and development and produced the world's most efficient solar cells".
"Dr Muriel Watt is honoured for her solar career and contributions to solar advocacy through the
Australian PV Association".
Dr Zhengrong Shi was also one of the twelve Australia’s leading solar practitioners inducted into the
Australian Solar Hall of Fame on this day.
"Dr Shi Zhengrong, the founder and Executive Chairman of Suntech - the world's largest producer of
solar panels - is a dual Australian/Chinese citizen, often referred to as the Sun King."
Martin Green Invited To Speak At Nobel Symposium
Professor Martin Green was an invited member of the Scientific Advisory Committee and invited
speaker at the Nobel Symposium on “Nanoscale Energy Converters” in Sweden in August 2012. The
symposia are prestigious events organised by the Nobel Foundation in areas of science where
breakthroughs are occurring or deal with topics of primary cultural or social significance. They are
“taken as an indication of where future Nobel Prizes may be heading”.
Professor Green was invited to talk on “Inorganic Nanophotovoltaics”.
Participants at the Nobel Symposium on “Nanoscale Energy Converters”.
Speakers (front row) included (from far left) Professors Eli Yablonovitch, Alan Heeger,
Mildred Dresselhaus, Martin Green (centre), Harry Atwater and Michael Graetzel (far right).
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Spree Student Wins Dean's Award for Excellence in Post-Graduate Research
Dean’s Award winner :Yuanjiang Pei.
Yuanjiang Pei from the School of Photovoltaic and Renewable Energy Engineering won the Dean's
Award for Excellence in Post-graduate Research, 2012. “We are all very proud of this achievement,"
said Dr. Richard Corkish, Head of School. "Yuanjiang has worked very hard and it is wonderful for
him to get this recognition.”
“Comprehensive Renewable Energy” Awarded the PROSE Award 2012
The “Comprehensive Renewable Energy” in which a book chapter contributed by Dr Supriya Pillai
and Professor Martin Green has been awarded the PROSE Award 2012, Reference Work: Best
Multivolume Reference/Science, by the American Association of Publishers.
S. Pillai and M.A. Green, “Plasmonics for Photovoltaics”,
In: ”Comprehensive Renewable Energy”, A Sayigh (Ed.),
Vol. 1, pp. 641-656, Oxford: Elsevier Ltd., 2012.
High Profile Visitors
Chinese Communist Party's Secretary in Guangdong Province, Mr Wang Yang, visited UNSW on 7
June 2012 for meetings with the Vice-Chancellor, Professor Fred Hilmer, followed by a visit to the
School of Photovoltaic and Renewable Energy Engineering, meeting the Dean of Faculty of
Engineering, Prof. Graham Davies, and Acting Head of School, A/Prof. Alistair Sproul. UNSW has
an ongoing cooperative relationship with South China University of Technology, and Sun Yat-Sen
University in the Guangdong Province. “It was an immense honour to meet Wang Yang and to show
him through some of SPREE’s photovoltaics facilities at the new Tyree Energy Technologies
Building”, said A/Prof Alistair Sproul - Acting Head of School. “Many of our PhD students are from
China and they were all especially excited about this visit”.
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Mr Wang Yang with A/Prof Alistair Sproul
viewing Suntech manufactured PV
on the TETB roof at UNSW.
Mr Lian Weiliang, Deputy Commissioner, National Development and Reform Commission (NDRC),
China'a top planner, and his delegation visited SPREE on 17 September 2012.
Mr Lian Weiliang with Dr Richard
Corkish, Head of School, SPREE.
On 11 December 2012, State Councillor, Her Excellency Madame Liu Yandong from the People’s
Republic of China, together with her more than 25 people delegation visited the ARC Photovoltaics
Centre of Excellence and met the UNSW Executive Team led by the Vice-Chancellor, Prof. Fred
Hilmer to better understand UNSW’s approach to research collaboration; , external partnerships and
the student experience.
Madam Liu’s laboratory tour at TETB
hosted by Prof. Martin Green.
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On 19 December 2012, Mr. Baoping Qiao, Executive Vice President of China Guodian Corporation
and Secretary of Party Leadership of China Guodian Corporation, together with his team of senior
executives, visited the ARC Photovoltaics Centre of Excellence, UNSW and met with Professor Les
Field, Professor Martin Green, Dr Richard Corkish, Dr Xiaojing Hao and a group of students.
“UNSW has established collaboration with many Chinese industries and happy to see further strategic
collaborations between Guodian and SPREE in the future” said Prof. Field. The meeting concluded
with a tour of the new TETB and laboratories inside the building.
Guodian (China Guodian Corporation) is one of the largest state-owned power generation groups in
China. Our Centre has ongoing collaboration with Guodian New Energy Research Institute in Beijing,
an institute formed by Guodian Corporation, since early 2012 on several mutually beneficial projects
including one supported by the Australian Research Council and one by ASI. Since the collaboration
commenced, this was the first official visit to the Centre by Mr Qiao.
Mr Baoping Qiao and Professor Les Field (Deputy Vice-Chancellor Research, UNSW)
were exchanging gifts at the meeting.
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3.
STAFF LIST
Director
Stuart Ross Wenham, BE BSc PhD UNSW, FTSE, FIEEE, FIEAust
(Scientia Professor)
Executive Research Director
Martin Andrew Green, AM, BE MEngSc Qld., PhD DSc(Hon) McM, DEng UNSW, FAA, FTSE, FIEEE
(Scientia Professor)
Deputy Directors
Gavin Conibeer, BSc MSc London, PhD Southampton
Thorsten Trupke, PhD Karlsruhe
Ashraf Uddin, BSc (Hons) MSc Dhaka U, Bangladesh; PhD Osaka U, SMIEEE
Sergey Varlamov, PhD Moscow
Business, Technology & Operations Manager
Mark D. Silver, BE UNSW, GMQ AGSM
Undergraduate Co-ordinator
Ivan Perez-Wurfl, PhD Uni. of Colorado, Boulder
Undergraduate Industrial Training Co-ordinator
Ashraf Uddin, BSc (Hons) MSc Dhaka U, Bangladesh; PhD Osaka U, SMIEE
Postgraduate Co-ordinator
Alistair B. Sproul, BSc (Hons) Sydney, PhD UNSW
Professors
Allen Barnett, BSEE, MS U Illinois, PhD Carnegie-Mellon U
Gavin Conibeer, BSc MSc London, PhD Southampton
Associate Professors
Chee Mun Chong, BSc BE (Hons) PhD UNSW
Evatt Hawkes, BSc (Hons) BEng (Hons) W. Aust., PhD Cantab.
Alistair B. Sproul, BSc (Hons) Sydney, PhD UNSW
Thorsten Trupke, PhD Karlsruhe
Ashraf Uddin, BSc (Hons) MSc Dhaka U, Bangladesh; PhD Osaka U, SMIEE
Senior Lecturers
Stephen Bremner, BSc (Hons), PhD UNSW
Richard Corkish, BEng (Com) RMIT, PhD UNSW
Alison Lennon, BSc (Hons) Sydney, PhD Sydney, PhD UNSW
Santosh Shrestha, PhD UNSW
Geoffrey J. Stapleton, BE UNSW
Ashraf Uddin, BSc (Hons) MSc Dhaka U, Bangladesh; PhD Osaka U, SMIEEE
Muriel Watt, BSc N.E., PhD Murdoch
Lecturers
Anna Bruce, BE(Hons) PhD UNSW
Merlinde Kay BSc (Hons), PhD UNSW
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Associate Lecturers
Emily Mitchell, BE PhD UNSW
Long Seng To, BE (Hons) BA UNSW
Senior Research Fellows
Anita W.Y. Ho-Baillie, BE PhD UNSW (returned April 2012)
Mark Keevers, Bsc (Hons) PhD UNSW (since January 2012)
Dirk König, Dipl. Ing. (EE) Dr. rer. nat. (Ph) Chemnitz
Sergey Varlamov, PhD Moscow
Research Fellows
Ian Brazil, BA BAI, Dubl., MEng DCU, PhD UNSW, MIEI (since October 2012)
Patrick Campbell, BSc BE PhD UNSW
Matthew Edwards, BE (Hons 1) PhD UNSW
Xiaojing Hao, BE, MSEE Northeastern Uni, China
Shujuan Huang, BE ME Tsinghua Uni., PhD Hiroshima Uni.
Henner Kampwerth, MSc Freiburg, PhD UNSW
Emily Mitchell, BE PhD UNSW
Supriya Pillai, B. Tech (India), PhD UNSW
Post-Doctoral Fellows
Pasquale Aliberti, BEng, MEng University of Salerno, Italy; PhD UNSW (until July 2012)
Jose Ignacio Bilbao, BEng (EE) MEngSc UC Chile, PhD UNSW (since September 2012)
Jessie Kai Copper, BEng (Hons), PhD UNSW
Hongtao Cui, BEng(Hons) Dalian U of Technol., PhD UNSW (since October 2012)
Armin Dehghan, PhD, U Western Ontario, Canada
Dawei Di (March – September 2012)
Jialiang Huang, BSc, PhD UNSW
Craig Johnson, BSc EE (Hons) Georgia Inst. Tech., PhD UNSW (since August 2012)
Siva Karaturi (since February 2013)
Sammy Lee, BSc (Hons), PhD UNSW (since October 2012)
Ly Mai, BE UNSW, PhD UNSW
Robert Patterson, B.Sc. (Hons) Queens U CAN, M.Sc. U Alberta CAN; PhD UNSW
Binesh Puthen-Veettil, BTech India, PhD UNSW
Ishwara Thilini Wijesingha Seneviratna, B.Sc. (Hons) Colombo U, Sri Lanka; PhD IC London, UK
Jing Rao (VC’s Post-Doctoral Research Fellow)
Budi Tjahjono PhD UNSW, MCom UNSW, BEng UNSW
Xi Wang, BEng Qingdao University, China; PhD Georgia Institute of Tech, USA (since Mar 2013)
Xiaoming Wen (since December 2012)
Yang Yang, PhD UNSW (since October 2012)
Adjunct Fellow
Dider Debuf, BE ME PhD UNSW
Visiting Professors/Fellows
Robert Bardos, BSc (Hons), PhD Melbourne
Yukiko Kamikawa, JSPS Fellow, AIST, Tsukuba, Japan (until March 2012)
Robert Opila, Professor, sabbatical from the University of Delaware (since Feb 2013)
Darcy Wentworth, BSc BE Sydney U., MSc Colorado Sch. of Mines, MSc UNSW (retired Dec 2012)
Fumin Xu (until June 2012)
Zhiping Zheng (until September 2012)
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Visiting Researcher
Daniel Hiller, Dipl. Phys. [U Halle-Wittenberg], Dr. Ing. [Freiburg U], Germany (Oct-Nov 2012)
Professional Officers
Patrick Campbell, BSc BE PhD UNSW
Kian Fong Chin, BE QUT, MEngSc QUT, MEngSc UNSW
Mark Griffin, BE UNSW
Yidan Huang, BSc Anhui
Omak Kuzuo
Robert Largent, AS USA
Hamid R. Mehrvarz, PhD UNSW
Tom Puzzer, BSc PhD UNSW
Nancy Sharopeam, BE U Sudan for Sci. & Tech., Sudan, ME UNSW (since June 2012)
Nicholas Shaw, BE UNSW, PhD UNSW
Lawrence Soria, AssocDipCompAppl W’gong
Bernhard Vogl
Alan Yee, BE UTS
Research Assistants
Malcolm Abbott
Vincent Allen
Christopher Antonini
Avantika Basu
Nino Borojevic
Catherine Chan
Mitchell Eadie
Hua Fan
Shahla Flynn
Brett Hallam
Jianshu Han
Florian Hess
Martin Hidas
Pei Hsuan
Jialiang Huang
Xiao Jin
Edwards Law
Sammy Lee
Raymond Leung
Yang Li
Dong Lin
Xiaolei Liu
Zhang Liu
Zhong Lu
Nitin Nampalli
Ben Noone
Yuanjiang Pei
James Rudd
Ming See
Nicholas Shaw
Lei (Adrian) Shi
Adeleine Sugianto
Mohsen Talei
Anthony Teal
Valantis Vais
Zhenyu Wan
Xi Wang
Gangqi Xu
Xaiohan Yang
Yu Yao
Sharon Young
Casual Staff
Joshua Allen
Catherine Davey
Tanya Gately
Wei Li
Wendi Peng
Alexandra Ryan
Erdinc Saribatir
Melissa Tang (until July 2012)
Matthew Wright
Manager, TETB Admin Unit
Kimberly Edmunds, MEdLead UNSW
Financial Officer
Julie Kwan (until February 2012)
Elena Estrina (since March 2012)
- 15 -
Administrative Staff
Danny Lin Chen, BSc UNSW
Mathew Cheung
Sue Edwards
Joyce Ho
Caroline Latimer
Jill Lewis
PhD Students
Xinrui An
Yameng Bao
Jose Ignacio Bilbao (until September 2012)
Nino Borojevic
Wenkai Cao
Kristen Casalenuovo (until February 2013)
Alexander Chan
Catherine Chan (since August 2012)
Kah Chan
Jian Chen
Simon Chung (since August 2012)
Brianna Conrad (since February 2013)
Jessie Copper
Hongtao Cui (until September 2012)
Jie Cui
Xi Dai
Dawei Di (until March 2012)
Martin Diaz
Claire Disney
Jonathon Dore
Hua Fan
Yu (Jack) Feng
Sishir Goel
Owen Grace
Morgan Green
Neeti Gupta
Brett Hallam
Phillip Hamer
Jianshu Han
James Hazelton
Pei-Chieh Hsiao
Andy Hsieh (until November 2012)
Pei Hsuan
Jessica Yajie Jiang
Xuguang Jia
Craig Johnson (until August 2012)
Mattias Juhl
Miga Jung
Shahram Karami
Dongchen Lan (since March 2013)
Sammy Lee (until August 2012)
Dun Li
Hongzhao Li (since February 2013)
Hua Li
Wei Li
Yang Li
Zhongtian Li
Yuanxun (Steven) Liao
Dong Lin
Rui Lin
Shu Lin
Ziyun Lin
Xiaolei Liu
Ziheng Liu
Doris Lu
Zhong Lu
Ibraheem Al Mansouri
Zichao Meng
Bernhard Mitchell
Nitin Nampalli
Amir Nashed
Mohd Zamir Pakhuruddin
Shyam Pasunurthi
Yuanjiang Pei
Binesh Puthen-Veettil (until July 2012)
Fang Qi
John Rodriguez
Fatemehsadat Salehi
Chao Shen
Rui Sheng (since February 2013)
Suntrana Smyth
Anastasia Soeriyadi (since February 2013)
Lihui Song
Ning Song
Anthony Teal
Peinan Teng
Alexander To (since February 2013)
Kai Wang
Kee Soon (Darryl) Wang
Lu Wang
Pei Wang
Qian Wang
Alison Wenham
Ned Western
Sanghun Woo
Lingfeng Wu
Matthew Wright
Guangqi Xu
Jiazhuo Xuan
Hongze Xia
Bo Xiao
Chaowei Xue
Chien-Jen (Terry) Yang
Xiaohan Yang
Yang Yang (until September 2012)
Yao Yao
Yu Yao
Jae Sung Yun
- 16 -
Pengfei Zhang
Haixiang Zhang
Haoyang Zhang
Tian Zhang
Xin Zhao
Masters Students
Chaho Ahn
Xin Ge
Mark Griffin
Kyunghun Kim
Taekyun Kim
Jongsung Park
Li Wang
Taste of Research Summer Scholars
Yu Jiang
Daniel Lambert
Albert Ng
Nathan Tam
Bo Wang
Jianzhou Zhao
Undergraduate Thesis Students
Tomas Adams
Aiman Albatayneh
Vincent Allen
Ashleigh Atkins
Isabel Baudish
Zyra Jane Bernadino
Joanne Borkowski
Hugh Bromley
Aden Brown
Flairy Anne Caragay
Deyuan Chen
Jiahao chen
Ran Chen
Jiahao Chen
Yiren Cheng
Rushil Chibber
Yuan Chu
Daniel Chung
Toby Coates
Jack Colwell
Yanting Deng
Mitchell Eadie
Michael Fairburn
Shu Fang
Hannah Farrow
Jade Fennell
Winnie Fu
David Godwin
Paul Gonzales
Joseph Grisold
Ricky Gu
Jinyi Guo
Thomas Harrison
Jai Hashim
Xiaoran He
Kenneth Hee
Florian Hess
Phillip Hodgkinson
Amnon Holland
Daniel Hooge
David Huang
Daniel Huynh
Matthew James
Aloysious John
Hannah Jones
Rowan Katekar
Nicholas Keene
Roland Kiel
Hee Chul Kim
Jin-Woo Kim
Dongchen Lan
Edward Law
Andrew Lem
Hongzhao Li
Junhan Li
Hongzhao Li (until November 2012)
Jessica Lilley
Shu Lin
Chao Liu
Kung Ting Lo
Mikael-Lu Lu
Christopher Mann
Luke Marshall
Thomas Milthorpe
Masahiro Miwa
Vicki Mo
Poojan Modi
Wangari Murchiri
Joanna Murphy
Anna Nadolny
Christopher Nheu
Benjamin Noone
Samrach Nov
Benjamin Pace
Jaga Park Ross
Guillaume Pauvert
Aobo Pu
Jonathan Pye
- 17 -
Thanh Vu
Jason Wu
Wei Wu
Qiyun Wu
Hons Wyn
Hang Xing
Dean Xue
Yiteng Yao
Jeeyeon Yeo
Chen Yi
Mark Yii
Wing Yip
Jinyang Yu
Quanzhou Yu
Qiuyang Zhang
Shuaijie Zhang
Xiaoyang Zhang
Yao Zhang
Yi Zhang
Yu Tian Zhang
Siyuan Zhao
Shichao Zheng
Leonard Quong
Matin Raupp
Mustaqim Raupp
Dean Redman
Alexander Robinson
Jake Saunders
Wei Chen Seah
Li (David) Shen
Joshua Shephard
Cheng Shu
Anastasia Soeriyadi (until November 2012)
Yi Sun
Xiaohe Sun
Andrew Sung
Desiree Surjadi
Anton Szilasi
Louise Temple
Ashwin Thomas
Alexander To (until November 2012)
Caitlin Trethewy
John Turle
Joshua Tang
- 18 -
4.
RESEARCH
4.1 Introduction to Research
Photovoltaics, the direct conversion of sunlight to electricity using solar cells, is recognised as one of
the most promising options for a sustainable energy future, with the corresponding industry poised to
become one of the world’s largest of this century. The ARC Photovoltaics Centre of Excellence
commenced in mid-2003, combining previous disparate strands of work supported under a variety of
programs, into a coherent whole addressing the key challenges facing photovoltaics, as well as “spinoff” applications in microelectronics and optoelectronics. The Centre was funded by the ARC until
December 2010 and has since been funded under a variety of other schemes, with much recent
funding provided no project grants by the Australian Solar Institute (ASI).
The Centre’s photovoltaics research is divided into three interlinked strands addressing near-term,
medium-term and long-term needs, respectively. The present photovoltaic market is dominated by
“first-generation” product based on silicon wafers, either single-crystalline as in microelectronics
(Fig. 4.1.1) or a lower-grade multicrystalline wafer. This market dominance is likely to continue for
at least the next decade. First-generation production volume is growing rapidly, with a huge effort
directed to streamlining manufacturing to reduce costs while, at the same time, improving the energy
conversion efficiency of the product. Also important is the reduction of the thickness of the starting
silicon wafer without losing performance or decreasing yield, to save on material use.
Figure 4.1.1: “First-generation” wafer-based technology (BP Solar Saturn Module, the photovoltaic
product that was manufactured in the highest volume by BP in Europe, using UNSW buried-contact
technology).
The Centre’s first-generation research is focussed on these key issues. Building upon the success of
“buried-contact” solar cell, the first of the modern high-efficiency cell technologies to be successfully
commercialised (Fig. 4.1.1), the Centre has developed several other high-efficiency processes in
commercial production or close to this, based on Centre innovations in laser and ink-jet processing.
Sales under licence are now approaching several billion dollars.
Although costs are reducing rapidly, wafers are still relatively expensive and need quite careful
encapsulation, since brittle and also thermal expansion mismatched to the glass coversheet, making
first-generation technology reasonably material-intensive. Several companies worldwide are
attempting to commercialise “second-generation” thin-film cell technology based on depositing thin
- 19 -
layers of the photoactive material onto supporting substrates or superstrates, usually sheets of glass
(Fig. 4.1.2). Although materials other than silicon are of interest for these films, silicon avoids
problems that can arise with these more complex compounds due to stability, manufacturability,
moisture sensitivity, toxicity and resource availability issues.
Figure 4.1.2: Example of “second-generation” thin-film technology (module fabricated on CSG
Solar’s German production line, based on thin-films of polycrystalline silicon deposited onto glass,
again UNSW-pioneered technology).
The Centre is developing a thin-film approach based on the use of the same high quality silicon used
for first-generation production, but deposited as a thin layer onto glass. The main emphasis is on the
development of lower-cost deposition and processing approaches (such as deposition by evaporation
rather than PECVD). The Centre also commenced activities on carbon-based, organic solar cells
during 2009 and a program on other “earth abundant” materials, particularly CZTS (Cu2ZnSnS4) in
2010.
At the present time, such second-generation thin-films are having difficulty entering the market, due
to rapidly decreasing first-generation cell costs. Large-scale commercialisation of thin-film product
leads to a different manufacturing cost structure composed to the wafer-based product. However,
costs again increasingly become dominated by material cost as production increases, for example, by
the cost of the glass sheet on which the cells are deposited.
More power from a given investment in material can be obtained by increasing energy-conversion
efficiency. This leads to the possibility of a third-generation of solar cell distinguished by the fact that
it is both high-efficiency and low cost. To illustrate the cost leverage provided by efficiency, Fig.
4.1.3 shows the relative cost structures of the three generations being studied by the Centre. This
figure plots efficiency against manufacturing cost, expressed in US$/square metre. First-generation
technology has relatively high production cost per unit area and moderate likely efficiencies at the
module level (14-20%). The dotted lines in Fig. 4.1.3 show the corresponding cost/watt, the market
metric. Values below US$1/watt are now common for large manufacturers.
- 20 -
Figure 4.1.3: Efficiency and cost projections for first-, second- and third-generation photovoltaic
technology (wafers, thin-films and advanced thin-films, respectively).
Second-generation thin-film technology has a different cost structure as evident from this figure.
Production costs per unit area are a lot lower, since glass or plastic sheets are a lot less expensive than
silicon wafers. However, likely energy-conversion efficiencies are lower (8-15%). Overall, this tradeoff produces costs/watt estimated as lower than those of the wafer product, in large production
volumes.
The third-generation was initially specified as a thin-film technology, which therefore would give
manufacturing costs per unit area similar to second-generation, but is based on operating principles
that do not constrain efficiency to the same limits as conventional cells (31% for non-concentrated
sunlight for the latter). Thermodynamic limits for solar conversion are much higher (74% for nonconcentrated light, giving an idea of the scope for improvement). If a reasonable fraction of this
potential can be realised, Fig. 4.1.3 suggests that third-generation costs could be lower than secondgeneration by another factor of 2 to 3.
Of the third-generation options surveyed by Centre researchers, “all-silicon” tandem cells based on
bandgap-engineering using nanostructures was selected as the most promising for relatively near-term
implementation (Fig. 4.1.4). This involves the engineering of a new class of mixed-phase
semiconductor material based on partly-ordered silicon quantum-dots in an insulating amorphous
matrix. The general CZTS material system may also have some potential here due to the range of
bandgaps accessible. Rapid recent progress in reducing silicon wafer costs means that even tandem
cells built on silicon wafers could fall into the targetted cost zone in Fig. 4.1.3. This realisation has
resulted in the Centre initiating several programs investigating such silicon wafer-based tandem cells
during 2012. Photon up-conversion as a way of “supercharging” the performance of relatively
standard cells forms a second line of research. A third is the investigation of schemes for
implementing hot-carrier cells. Both of the latter may prove to be particularly well suited to organic
solar cells.
- 21 -
Figure 4.1.4: Conceptual design of an all-silicon tandem cell based on Si-SiO2 (or Si-Si3N4 or SiSiC) quantum dot superlattices. Two solar cells of different bandgap controlled by quantum dot size
are stacked on top of a third cell made from bulk silicon.
The fourth Centre strand of silicon photonics draws upon elements of all three of the photovoltaic
strands. A by-product of this work has been the development of techniques based on silicon lightemission for characterising both completed devices, particularly solar cells, as well as silicon wafers
at different stages of processing (Fig. 4.1.5). Developing this approach to its full potential has formed
an increasingly large part of the Centre’s photonics program.
- 22 -
Figure 4.1.5: Schematic representation of a PL imaging system. An external light source is used to
illuminate the silicon wafer or solar cell homogeneously. The luminescent emission (red arrows) from
the sample is captured with a sensitive CCD camera. This technology was commercialised by Centre
“spin-off” BT Imaging in 2008.
- 23 -
4.2
Facilities and Infrastructure
The ARC Photovoltaics Centre of Excellence is located at the Kensington campus of the University of
New South Wales (UNSW), about 6 km from the heart of Sydney and close to its world famous
beaches including Bondi, Coogee and Maroubra (Fig. 4.2.1).
Figure 4.2.1: Centre of Excellence location in Sydney.
Organisationally, the Centre of Excellence is located within the School of Photovoltaic and
Renewable Energy Engineering (SPREE) within the Faculty of Engineering. The Centre of
Excellence has a large range of laboratory facilities (Fig. 4.2.2). These include the Bulk Silicon
Research Laboratories, the Device Characterisation Laboratory, the Optoelectronic Research
Laboratories, the Thin-Film Cell Laboratory, the Industry Collaborative Laboratory, Inkjet/Aerosol
Processing Laboratory and Organic Photovoltaic (OPV) Laboratory. Off campus the Centre has a
Thin-Film Cleanroom facility at Botany, 5km south-west of the main campus. Other important
resources are the Semiconductor Nanofabrication Facility (SNF) and Australian National Fabrication
Facility (ANFF) jointly operated by the Faculty of Science and the Faculty of Engineering.
Additional equipment commonly used for solar cell work is found at the new MicroAnalytical
Research Centre (MARC) and elsewhere on the University campus. Included in this category are
TEM/SEM electron microscopes, focused ion beam (FIB) specimen preparation equipment, X-ray
diffraction, Raman spectroscopy, AFM and surface analysis equipment. TEM, ellipsometry and a
femtosecond time resolved photoluminescence measurement system are also regularly accessed at
Sydney University.
During 2012, the development and acquisition of laboratory equipment and infrastructure continued
with the first Centre laboratories integrated into UNSW’s new flagship energy efficient Tyree Energy
Technologies Building (TETB), a building with a 6-Star rated Green Star design. Specific details
about significant additions in 2012 are found under the laboratory headings which follow.
- 24 -
Figure 4.2.2: Layout of laboratories and other facilities within the Electrical Engineering Building,
Kensington Campus. The OPV laboratory is in the Chemical Sciences Building (not shown). In
March 2012 the Device Characterisation and Optoelectronic Laboratories were relocated to Level 1 of
the new the Tyree Energy Technologies Building (TETB).
The Centre of Excellence has three computer networks. Both Kensington and Bay St Botany each
have a research/admin network. The third network is for students enrolled in the Undergraduate
Degree Courses in Photovoltaics and Renewable Energy. In 2011 the Kensington Research and
Administrative network grew by 9% to consist of 3 file servers, 1 terminal server, 1 intranet server, 2
Internet web-servers (hosted on Central IT infrastructure), 1 centrally hosted licence server and over
(196) client workstations. At Kensington, an additional 35 computers are dedicated to the computer
control of laboratory and other equipment. In addition a web-based lab-equipment booking system is
in place for the equipment in Rm 140 TETB. The Centre also has a 3 server computer controlled
SCADA and PLC network for equipment control and monitoring the research laboratories and related
infrastructure. Bay St Botany has 11 dedicated equipment controllers and their research/admin
network supports 7 workstations.
During 2012, the shared drive migration onto Central IT infrastructure, which commenced in 2011,
was completed. Also during 2012 a new server was commissioned to gather data from the grid
connected solar arrays on the TETB roof. The data will be used for PV grid connect power systems
research.
The computer resources are used for measurement, modelling/simulations, equipment control,
document control, laboratory design support, Internet access, general administrative purposes and
maintaining the Centre’s presence on the Internet.
Each postgraduate research student is allocated a dedicated personal computer and has access to
shared computer resources.
The Centre upgraded Linux (Beowulf) 64bit cluster has an estimated computational power of 4.4
Teraflops to support the demanding Density-Functional-Hartree-Fock and molecular dynamical
computations. A second 544 core, 4.6 Teraflop, cluster has been installed for complementary
investigations of energy efficiency and renewable energy via computational fluid dynamics
techniques.
- 25 -
In 2012, the SPREE student computer teaching laboratories were moved into two larger laboratories
in the TETB. The number of workstations increased from 34 dedicated machines to 84 machines
shared with other building stakeholders. At the same time domain control, file and print services were
migrated to UNSW IT central control and SPREE’s undergrad access to the CSE teaching laboratories
was no longer required and discontinued. Students enrolled in the Undergraduate Degree Courses use
these teaching computers for computer-related coursework and Internet access. There is also an
Internet capable web-server that gathers and displays data collected from solar arrays on the roof of
the Electrical Engineering building. These data can be viewed using a web browser and can be made
available for Internet access.
The Laboratory Development and Operations Team develops and maintains core Centre and
laboratory facilities. During 2012, the team, under the leadership of Mark Silver, comprised of an
additional 4 equivalent full-time and 10 casual + part-time employees, including: electrical engineers,
a computer/network manager, electronic/computer/laboratory technicians and administrative staff.
Bulk Silicon Research Laboratories
The Centre houses the largest and most sophisticated bulk silicon solar cell research facility in
Australia, incorporating both the High Efficiency and Buried Contact/Bulk Cell Laboratories. Total
laboratory processing space of over 580 m2 (including the Industry Collaborative, InkJet/Aerosol
Processing Development, Device Characterisation and Optoelectronic Research Labs) is located at the
Kensington Campus and is serviced with filtered and conditioned air, appropriate cooling water,
processing gases, de-ionized water supply, chemical fume cupboards and local exhausts. Another 480
m2 of combined roof space accommodates fixed PV arrays and 361m2 of accessible outdoor
experimental space. Off site, areas totalling 700 m2 are used for the storage of chemicals and
equipment spare parts along with a 100 m2 nickel sintering belt furnace facility at 78 Bay St Botany.
The laboratories are furnished with a range of processing and characterisation equipment including 23
tube diffusion furnaces, 6 vacuum evaporation deposition systems, a laser-scribing machine, 2 laser
doping machines, rapid thermal annealer, four-point sheet-resistivity probe, quartz tube washer,
silver/nickel and copper plating facility, visible wavelength microscopes, 3 wafer mask aligners, spinon diffusion system, photoresist and dopant spinners, electron beam deposition system, metallization
belt furnace, nickel sinter belt furnace, manual and automatic screen printers and a laboratory system
control and data acquisition monitoring system.
Figure 4.2.3: CNC Laser Scribe Tool.
- 26 -
The laser scribe tool, shown in Figure 4.2.3, has a 20 watt Nd:YAG laser for infrared operation (1064
nm) and an optional frequency doubler for green operation (532 nm). The work stage is CNC
controlled allowing 1 micron positional accuracy and table speeds approaching 25 cm/second across
an area of 15 cm by 15 cm. The tool is used primarily for Buried Contact and bulk silicon solar cell
fabrication, cutting 35-micron wide laser grooves as deep as 100 microns into silicon wafers and
scribing wafers in preparation for cleaving. It can also be used to cut other suitable materials, such as
stainless steel.
Device Characterisation Laboratory
The Device Characterisation laboratory was the first laboratory to be relocated to the TETB in March
2012.
It houses characterisation equipment including “Dark Star”, the Centre’s station for temperature
controlled illuminated and dark current-voltage measurements, the Centre’s Fourier-transform
infrared spectroscopy system (FTIR), frequency dependent impedance analyser, ellipsometer, Sinton
photoconductance lifetime equipment, wafer probing station, open circuit voltage versus illumination
measurement system (Suns-Voc), 4 point resistivity probe, spectral response system and
spectrophotometer with integrating sphere.
In 2008, a multi function wafer mapping tool, shown in Figure 4.2.4, was installed. This unit provides
state of the art capability for the measurement of carrier lifetime, bulk resistivity, emitter sheet
resistance and Light Beam Induced Current (LBIC) on wafer samples.
Additions in 2009 included a high speed commercial flash cell tester with 156 mm square wafer
capability as shown in Fig 4.2.5, a replacement UV/VIS/NIR spectrophotometer and a commercial
Luminescence Inspection System (LIS) from spin off company BT Imaging.
In 2011, a spectroscopic ellipsometer was acquired that allows non-destructive thin film
characterisation of samples at multiple wavelengths to determine thin film thicknesses and optical
constants. Other new additions, located in other Centre laboratories, include a surface profiler,
automatic 4 point probe and a front side grid contact resistance mapping tool.
In 2012, a new cryogenic micromanipulated probing station was commissioned for non-destructive
testing of devices on full and partial wafers up to 51 mm in diameter. The instrument can be used to
measure electrical, electro-optical, parametric, high impedance, DC and RF properties of materials
and test devices.
The system operates over a temperature range of 77 K to 475 K. The probe station provides efficient
temperature operation and control with a continuous refrigeration system using liquid nitrogen. A
control heater allows precise sample stage temperature control, and along with the radiation shield
heater, provides the probe station with fast thermal response.
The system has been configured with four micro-manipulated stages, each providing precise 3-axis
control of the probe position. Two probe arms are equipped with probes rated to 1GHz. Probe tips are
thermally linked to the sample stage to minimize heat transfer to the DUT. One probe arm is equipped
with an optical fiber that can be placed a few microns away from the device to be tested.
- 27 -
Figure 4.2.4: Wafer Mapping Tool.
Figure 4.2.5: Flash Cell Tester.
- 28 -
Optoelectronic Research Laboratories
This facility has six optical benches and several visible and near-infrared semiconductor diode lasers
along with other optical and electrical instrumentation. The facility is used for photoluminescence
(PL) and electroluminescence measurements in the visible and infrared spectral range up to
wavelengths of 2500nm, luminescence experiments with simultaneous two-colour illumination, quasi
steady state photoluminescence lifetime measurements and Sinton lifetime testing with the
conventional flash-light replaced by a high-power light emitting diode array.
This facility also houses the Centre’s world-leading first photoluminescence imaging system. Other
equipment includes a silicon CCD camera (for sensitive PL measurements), spectroscopic PL
systems, PL lifetime measurement unit, and LBIC measurement system. Areas separate from the
Device Characterisation Laboratory were necessary in order to meet stringent standards for safe laser
use. Cryogenic cooling equipment is shared with the Device Characterisation Laboratory.
Figure 4.2.6: Optical characterisation bench.
Thin-Film Cell Laboratory
This 40 m2 laboratory is equipped with a range of equipment for thin-film deposition and patterning,
including a plasma-enhanced chemical vapour deposition (PECVD) system, a sputtering system,
larger area plasma etcher, a reactive ion etcher (RIE), a resistively heated vacuum evaporator, helium
leak detector and an optical microscope with digital image acquisition system. Also used by the
laboratory is an electron-beam vacuum evaporator for silicon which is physically located within the
Bulk Silicon Research Laboratory. This Si evaporator is also equipped with an ionizer unit and a
sample heater, enabling fast-rate Si homoepitaxy at temperatures of about 500-600°C by means of
ion-assisted deposition (IAD). The IAD is also equipped with a residual gas analyser to permit real
time process monitoring. Other on campus equipment of use in thin-film projects is located within the
Semiconductor Nanofabrication Facility.
- 29 -
The PECVD system, shown in Fig. 4.2.7, has a 40 x 20 cm2 process platen and can handle large-area
silicon wafers as well as smaller pieces. Two types of plasma excitation (remote microwave and direct
RF) are available. The machine is used for the low-temperature deposition of thin dielectric films
(silicon nitride, silicon dioxide, silicon oxy-nitrides) and of amorphous silicon. The dual-cylinder,
remote microwave plasma source produces excellent-quality silicon nitride and silicon dioxide films,
with precise control over the stoichiometry at temperatures up to 500°C. Amorphous and
microcrystalline silicon films can also be deposited in the system.
In 2008, the thermal evaporator was upgraded to also support e-beam evaporation. This capability
greatly expands the range of materials which can be deposited and is of great interest to 3rd generation
researchers. In 2009, a new 4 point probing station was added to the laboratory. In 2010, the computer
support system for the optical microscope was upgraded to improve the useability of the digital
camera fitted.
Figure 4.2.7: Remote plasma PECVD machine.
Figure 4.2.8: Thermal and e-beam Evaporator.
- 30 -
Industry Collaborative Laboratory
This 120 m2 laboratory houses equipment needed for many of the industry-collaborative research
activities including the Buried-Contact + Bulk Solar Cell group. The laboratory equipment includes a
belt furnace; a state of the art laser micromachining tool; a PECVD deposition system (located in the
adjacent Thin Film Solar Cell Laboratory); a TiO2 spray deposition station; a high temperature
semiconductor muffle furnace, manual screen printer and a fully automatic production scale screen
printer. In 2008, a new 532nm green laser was added and customised in-house to increase the
capability for processing laser doped solar cells.
In 2009, new additions included a fast metallization firing furnace capable of processing 156 mm
square wafers, a light induced plating system, spin on doper and mirror steered scanning laser for
laser doping. In 2010 a new wafer spinner was added to the laboratory. Off site at 78 Bay St Botany a
nitrogen environment belt furnace capable of nickel sintering 156 mm square wafers and larger on its
25 inch wide belt was installed.
Inkjet/Aerosol Processing Development Lab
This laboratory houses the Centre’s inkjet printing development systems. The laboratory is used to
develop solar processing “inks” (chemical solutions) and for printing them onto a solar wafer under
computer control. The aim is to develop low cost processing techniques for creating fine structures in
solar cells. It is anticipated that the inkjet printing is capable of replacing processes such as laser
scribing and photolithography for forming fine patterns for contacting but at a cost of at least 10 times
cheaper. Equipment includes two ink jet material deposition printing systems, capable of depositing a
wide range of materials onto different substrates, a surface tension meter and viscometer.
In 2011, a new aerosol deposition system was commissioned. This has expanded our capability to
print and deposit a larger range of materials at typically smaller resolution than the inkjet printers. The
resultant reduction in material usage and loss is of great interest for production/manufacturing.
Figure 4.2.9: Aerosol Deposition and Inkjet development systems.
Semiconductor Nanofabrication Facility
The Centre also owns equipment within, and has access to, the Semiconductor Nanofabrication
Facility (SNF) at the University. This is a joint facility shared by the Faculties of Science and
Engineering and houses a microelectronics laboratory and a nanofabrication laboratory for e-beam
- 31 -
lithography. The SNF provides an Australian capability for the fabrication of advanced nanoscale
semiconductor devices and their integration with microelectronics. SNF research projects form an
integrated effort to fabricate innovative semiconductor nanostructures using the latest techniques of
electron beam patterning and scanning probe manipulation. A major applied objective of the facility is
the development of a prototype silicon nuclear spin quantum computer. The capabilities of this
facility were expanded to house the, NCRIS funded, Australian National Nanofabrication Facility
(ANNF). In 2010, the facility added a new large area e-beam lithography system and new e-beam
evaporator for non MOS-compatible metals deposition.
Thin-Film Clean Room facility in Bay Street, Botany
In 2003, the Centre added a 120 m2 cleanroom facility in Bay Street, Botany to its infrastructure,
greatly improving its experimental capabilities in the area of thin-films. This cleanroom is equipped
with several fume cupboards, two tube furnaces, an electron-beam vacuum evaporator, a thermal
vacuum evaporator, a glass washing machine, a rapid thermal processing (RTP) machine and a 5chamber cluster tool. The cluster tool presently features four plasma-enhanced chemical vapour
deposition (PECVD) chambers and one lamp-heated vacuum annealing chamber. The PECVD
chambers enable the low-temperature deposition of dielectric films (silicon oxide, silicon nitride, etc)
and amorphous silicon films (either n- or p-doped or undoped). Furthermore, samples can be
hydrogenated by PECVD using hydrogen plasma at substrate temperatures of up to 480°C. During
2004, the Centre purchased a low-pressure chemical vapour deposition (LPCVD) system, an infrared
NdYAG laser scriber and a box furnace for sample annealing. The LPCVD system is capable of
depositing doped crystalline silicon on glass and with its additional remote plasma source is currently
engaged in hydrogenation work. The NdYAG laser, located outside the clean room, is used for
scribing silicon films and other suitable metal and dielectric materials.
In 2006, a state of the art multi-target sputter machine was installed. In 2008 this vacuum tool was
upgraded from four to five separate targets. Each target is able to be operated independently of one
and other, allowing users to cosputter a thin film from more than one target and deposit multilayers
without breaking vacuum. Three power supplies are available with substrate biasing. The custom
made system can handle substrates up to 150mm x 150mm. Excellent film purity is assured as the
system incorporates a load-lock. Computer control can be used for most operations, including
substrate heating, allowing precise multilayers to be deposited repeatedly.
During 2007, a new computer controlled multi pocket, multi source UHV electron beam evaporator
system was delivered. This equipment provides the capability for high “industrial” rate silicon and
other thin film deposition.
In 2009, an optically assisted IV measurement system was installed for the determination of static hot
carrier distributions in the sample temperature range of 80 to 300 K. The system can also be used for
internal photo emission measurements. A dark-/oa-CV extension is planned for 2014/2015. The
LPCVD computer control system was also upgraded to provide enhanced throughput. A line beam
diode laser system was also installed, outside the clean room, for work on low thermal budget defect
annealing.
- 32 -
Figure 4.2.10: oa-IV setup (left).
In 2010, the first stage of a multi-chamber sputter system was ordered to allow our 3rd Generation
Group to carry out significantly different functions on a sample without exposing it to the atmosphere
between each process. Processes include sputter-depositing different dopants in separate chambers,
vacuum annealing, also PECVD layers. This one chamber purchase is “stage I”, a loadlock chamber
and other peripherals have been configured from decommissioned stock. Stage II will involve the
purchase of a second processing chamber and a transfer/loadlock chamber.
The addition of computer control to our hydrogenation tool provides more reliable process
repeatability and accuracy of set process parameters. The processor inputs all process parameters for
each “layer” defining each step before starting the process. This upgrade also frees processors’ time
in that they need not be in attendance to change settings when beginning each process layer.
2011 saw the additions of a split tube furnace and new sputter tool. The new furnace uses separate
quartz tubes for different samples to avoid cross contamination. The furnace is opened like a suitcase
to change processing tubes and each tube has its own endcaps, pressure gauge and pressure relief
valve. It presently has nitrogen purging and later will support processing under vacuum. The new
sputter tool will be used for intrinsic silicon based materials (ie no dopants) only.
In 2012, a sulphurisation furnace system was installed to support CZTS research. The furnace is a
dual chamber split tube type. Sulphur vapour is carried in N2 along the quartz tube from one chamber
where S is heated to the next where precursors absorb the vapour. All unabsorbed S vapour leaving
the chamber in the exhaust stream is condensed before reaching air, to avoid SO2 formation.
Figure 4.2.11: Bay Street Clean Room Sputtering Machines.
- 33 -
Figure 4.2.12: UHV Electron Beam Evaporator.
Organic Photovoltaic Laboratory (OPV)
The Organic Photovoltaic Devices group laboratory space in the Chemical Sciences Building was
further developed in 2012 for the processing and testing of organic photovoltaic devices and 3rd
Generation photovoltaic devices.
The laboratory occupies a space of 42m2 and consists of an acid-scrubbed fume cupboard, nitrogenpurged sample storage systems, nitrogen-atmosphere glovebox (comprising an internal liquid solution
spinner and hotplate for device processing within an oxygen-free environment) and a thermal
evaporator system for formation of metallic contacts to organic devices.
In 2012, additions to this area included a liquid source spinner with an exhausted work space hood, an
aerosol deposition system for application of TiO2 films for OPV devices and a lighted I-V
characterisation system.
For 2013, new thermal processing ovens will be commissioned for annealing processes of deposited
organic layer, and TiO2, films for OPV devices.
- 34 -
Figure 4.2.13: Nitrogen-purged, glovebox for spin-on and thermal processing of organic photovoltaic
devices in an oxygen-free atmosphere.
Figure 4.2.14: Exhausted workspace for liquid source spinner
- 35 -
Figure 4.2.15: New Light I-V Measurement tool for organic photovoltaic devices.
- 36 -
4.3
First Generation Wafer Based Projects
Research Team:
University Staff
S. R. Wenham
M. A. Green
Allen Barnett
Stephen Bremner
Gavin Conibeer
T. Trupke (Group Leader – Photoluminescence Imaging)
A. Lennon (Group Leader – Inkjet and Plating Technologies)
A. Ho-Baillie (Group Leader - High Efficiency Cells)
H. Kampwerth (Group Leader – Characterisation)
M. Keevers (Senior Research Fellows – Spectrum Splitting)
Postdoctoral Fellows
Ian Brazil (since October 2012)
M. Edwards (Tech Transfer)
Xiaojing Hao
L. Mai (Tech Transfer)
H. Mehrvarz (High efficiency Cells)
Ivan Perez-Wurfl
S. Pillai (High Efficiency Cells)
B. Tjahjono (Core Research)
Postgraduate Research Students and Research Assistants
Xinrui An
X. Bai
N. Borojevic
C. Chan
Jian Chen
Brianna Conrad
Hongtao Cui
Jie Cui
Martin Diaz
L. Gu
B. Hallam
P. Hamer
Z. Hamieri
Alex Han
Yajie (Jessica) Jiang
Mattias Klaus Juhl
Sammy Lee
M. Lenio
Dun Li
Hongzhao Li
Hua Li
Dong Lin
Ziheng Liu
D. Lu
B. Mitchell
Fang Qi
J. Roderiguez
- 37 -
Kenneth Schmieder
Chao Shen
Anastasia Soeriyadi
Lihui Song
A. Sugianto
Peinan Teng
Alexander To
K. Valliappan
D. Wang
Li Wang
Lu Wang
S. Wang
Alison Wenham
N. Western
Bo Xiao
G. Xu
Yang Yang
Y. Yao
Y. Yeung
L. Zhang
Xi Zhu
Research Associate
Wayne Zhenyu Wan (Research Associate)
Technology Transfer Team
Managers
 D. Jordan (Team leader)
 C.M. Chong (Deputy Team Leader – Devices)
 M. Edwards (Deputy Team Leader – Program Manager)
 S. Wenham (Technical Director)
Team Members
 X. Bai
 C. Chan
 C. M. Chong
 B. Hallam
 P. Hamer
 N. Kuepper
 A. Lennon
 L. Mai
 E. Mitchell
 A. Sugianto
 B. Tjahjono
NewSouth Innovations
 N. Simpson (until August 2012)
 A. Prib
 T. Montaldo
- 38 -
4.3.1
High Performance Cell Research
4.3.1.1
High Efficiency Silicon Cells
Further work on light trapping for high performance silicon solar cells has been carried out with
emphasis on optical characterization of textured front surface and various back reflectors.
Characterization of 2-D Reflection Pattern from Textured Front Surfaces of Silicon Solar Cells
Reflected light from textured front surfaces of a solar cell contains useful information about the
surface geometry as well as the optical properties of the cell. The reflected light patterns can be used
to extract details of surface morphologies and hence be used as a tool to fine tune and monitor
fabrication processes. In this work, the 2-D reflected light distributions from front surfaces of silicon
cells textured in various ways are characterized by a novel optical setup (Fig. 4.3.1.1.1). The
experimental results are compared to modelled results of conventional ray tracing. The impact of the
encapsulant’s refractive index on the amount of total internal reflection is also discussed for various
types of textured surface.
Figure 4.3.1.1.1: A cross-sectional schematic of the experimental setup for the recording of frontreflected light distribution. The pierced screen is either semi-transparent to allow digital imaging from
the back or consists of photosensitive photographic paper.
Figure 4.3.1.1.2a shows the expected reflected light distribution from the three rays A, B and C for a
regular inverted pyramid. Reflection angles of θa, θb and θc of the three rays relative to the vertical for
(111) planes as well as their corresponding weightings Ra, Rb and Rc (which are normalised to be unity
in sum) are shown in the inset of the figure. They are calculated by geometrical equations. Fig.
4.3.1.1.2b shows the measured spatial reflection distribution on photographic paper placed dt = 15.3
cm from the sample. The sample has mono-modal inverted pyramids without any antireflection
coating. The differences of escaping angles θa,b,c between theoretical and measured values are
attributed to the non-ideal average slope of the pyramid walls after each processing step, confirmed by
the SEM and FIB measurements in Fig. 4.3.1.1.4.
- 39 -
(a)
(b)
Figure 4.3.1.1.2: (a) A idealised reflection pattern for a regular inverted pyramid if the surface is hit
by a relatively small beam of light. (b) Reflected light distribution from a mono-modal inverted
pyramid texture on uncoated silicon surface.
Figure 4.3.1.1.3: Reflection patterns of silicon surface with inverted pyramids in a mono-modal
arrangement after (a) alkaline etch; (b) 1st acidic etch; and (c) 2nd acidic etch. (d)-(f) are the
corresponding top view SEM images of the inverted pyramids. (g)-(i) are the corresponding crosssectional FIB images.
Figure 4.3.1.1.4a,b,c show the reflection patterns for the random textured upright pyramids.
Compared to the regular, periodically arranged pyramids, these patterns are relatively feature-less due
to the randomness of the pyramid distributions. Data for R() are given as reflection weightings over
an angular range (every half degree) instead of for discrete angles, due to the continuous reflected
light distribution from random textures, see Eqn. 4.3.1.1.1. Figure 4.3.1.1.5 shows the reflection
weighting results (R) for the three selected random textures. We find that Texture-a with its relatively
large pyramids (~10 mm) shows the highest peak R. The peak R decreases for Texture-b and Texture-
- 40 -
c with smaller pyramids sizes (~5 mm and ~3 mm respectively). This can be explained by the
weakened effect of simple geometrical optics and the enhanced effects of scattering/diffraction when
the pyramids sizes become smaller.
(c)
(b)
(a)
Figure 4.3.1.1.4: Reflected light distributions captured on photographic papers for the corresponding
silicon surfaces with different sized random upright pyramids.
  0.5o
R  
 r   2  d  tan d
 r   2  d  tan d
t
90
(4.3.1.1.1)
o
t
0
Figure 4.3.1.1.5: Distribution of reflected light weighting R (over a half degree region) from the
three selected surfaces with random textures Texture-a, Texture-b and Texture-c.
The fraction of total internal reflection ftir indicates how much of light from the front-surface of the
cell is reflected back by the air-glass interface towards the cell again. The refractive index nenc of the
encapsulant is a dominating parameter determining ftir. In this work, ftir is defined as the intensity of
the integrated reflected light outside the escape cone of the module (θenc-cri to 90o) over the intensity of
the integrated reflected light from the cells surface (0o to 90o), as shown by Eqn. 4.3.1.1.2. The
relationship between ftir and nenc for surfaces with various different pyramids structures is summarised
in Fig. 4.3.1.1.6.


90 o
f tir 

enc _ cri
o
90
0
R( enc )d enc
(4.3.1.1.2)
R( enc )d enc
- 41 -
Figure 4.3.1.1.6: : The fraction of front reflected light that is totally internal reflected as a function of
the encapsulant’s refractive index.
Figure 4.3.1.1.6 indicates that for random textures, a conventional encapsulant with an index of nenc =
1.5 will result in only a fraction of total internal reflection of ftir = 5%. An increase to nenc = 1.6 will
increase ftir to 10% for random textures. Therefore a higher index encapsulant is needed to promote
total internal reflection for both random and regular textured surfaces. An alternative method to obtain
a high total internal reflection ratio is by developing novel surface textures that reflect light primarily
at oblique angles.
Angular reflection study to reduce plasmonic losses in the dielectrically displaced back reflectors of
silicon solar cells
The effective rear surface reflection of a solar cell will be determined by angular dependence if a
textured front surface is applied. This study investigates the angular rear reflection from various
planar reflectors experimentally. A novel optical setup (Fig. 4.3.1.1.7) has been developed that allows
the measurement of high resolution reflection characteristics over a broad angular range. The use of Si
hemispheres as test substrates (Figs. 4.3.1.1.8 & 4.3.1.1.9) for the reflector deposition enables the
analysis of the angular reflection properties of the back surface reflector over all incident angles
without the restriction caused by refraction at the Si-air interface.
(a)
(b)
Figure 4.3.1.1.7: (a) A simplified schematic of the experimental setup for the angular rear reflection
measurement. The inset indicates the rotating process during the measurement process. (b) A picture
of the realised setup mounted on an optical bench.
- 42 -
Figure 4.3.1.1.8: (a) A schematic showing the measurement limitations by using a flat wafer. The
angle of incidence is restricted to within a 17o cone inside Si due to the difference in the refractive
index between air and silicon. (b) A schematic of the proposed measurement with a Si hemisphere
allowing light illumination and detection at any angle between 0-90o.
Figure 4.3.1.1.9: Silicon hemispheres used as the substrates for the reflector deposition. Different
colours on the flat sides indicate different thicknesses of the thermally-grown SiO2.
The measured absolute reflected light intensities Rb (Φi) of the three rear reflector schemes are shown
by a half-polar chart in Fig. 4.3.1.1.10 for an oxide thickness of 366 nm. Measurements are carried
out on both sides of the hemisphere. Rb remains constant at angles below the critical angle of Si-SiO2
interface (Φc=24o) for rear schemes with displaced metal whereby most of the light is reflected back
into silicon with a small portion absorbed by the metal. Rb approaches unity at angles > 30o where
total internal reflection occurs. As light angle approaches Φc, Rb decreases and is at its minimum at an
angle of around 24.5o. This angle corresponds to the resonance angle Φr of the propagating surface
plasmon polaritons (SPP). Ag as a noble metal has slightly higher reflection than Al at angles < 18o,
and clearly higher reflection at angles near the resonance angle Φr where SPPs are excited. This
reflection difference comes from the different imaginary parts of the refractive index of the two
metals which indicates the amount of associated light absorption. The reflection dip at the resonance
Rb (Φr) for the Si/dielectric/metal structure indicates a loss due to excitation of SPPs, which should be
avoided in the solar cell. This can be achieved by optimising the front texture and/or by appropriate
choice of rear metal and rear dielectric with optimised thickness so that the refracted light at the front
interface is incident on the rear surface at an angle away from Φr.
- 43 -
Figure 4.3.1.1.10: Measured absolute reflected light intensities Rb (Φi) as a function of incident angle
(Φi) for three fabricated rear schemes: Si/SiO2/Al; Si/SiO2/Ag; and Si/SiO2/air.
For the Si/SiO2/Al scheme, the effect of SiO2 thickness on the angularly dependent rear reflection is
examined experimentally and by using W-VASE simulation as shown in Fig. 4.3.1.1.11. The
experimental data are normalised (whereby the maximum measured light intensity is set to equal
100%) to enable comparison with simulated results. As seen from Fig. 4.3.1.1.10, measured Rb is
slightly lower than the simulated Rb possibly due to the existence of a thin AlxOy layer at the SiO2/Al
interface which is not simulated in the model or due to differences in the experimental Al optical
constants from the values used in the simulation. Nevertheless, the measured Rb and the predicted Rb
show almost identical trends. The results have implications on the choice of dielectric thickness at the
rear for normal light incidence ie. for a planar front surface where Rb will fall within the escape cone
of the cell structure. As expected, Rb remains high at oblique angles where TIR occurs and is less
sensitive to dielectric thickness.
Figure 4.3.1.1.11 also indicates that front regularly-textured solar cells passivated by thin rear SiO2
are inferior to cells with thicker rear SiO2 due to i) increased parasitic absorption in the metal
(especially when most of the normally incident incoming light is deflected to an angle of 42o by the
textures of inverted pyramids as in Passivated Emitter Rear Locally diffused (PERL) cells fabricated
at UNSW) and ii) poorer surface passivation quality. Hence we can conclude that rear SiO2 with
thickness < 200 nm is not ideal for high performance solar cells with dielectrically displaced metal
back reflectors.
- 44 -
Figure 4.3.1.1.11: Rear Si reflectivity for unpolarised incident light from structure Si / SiO2 / Al with
various SiO2 thicknesses over the angle range of 2o to 56o (a) measured experimentally or by (b)
simulated.
References
4.3.1.1.1
4.3.1.1.2
4.3.1.1.3
4.3.1.1.4
4.3.1.1.5
Y. Yang, S. Pillai, H. Mehrvarz, H. Kampwerth, A. Ho-Baillie, M. A. Green (2012),
Enhanced Light Trapping for High Efficiency Crystalline Solar Cells by the Application
of Rear Surface Plasmons, Solar Energy Materials and Solar Cells, Volume 101, June
2012, Pages 217–226.
Y. Yang, M. A. Green, A. Ho-Baillie, H. Kampwerth, H. Mehrvarz, S. Pillai (2012),
Characterisation of 2-D Reflection Pattern from Textured Front Surfaces of Silicon Solar
Cells, submitted for publication in the Solar Energy Materials and Solar Cells.
Y. Yang, S. Pillai, M. A. Green, A. Ho-Baillie, H. Mehrvarz, H. Kampwerth (2012),
Surface plasmon enhanced rear light trapping schemes for c-Silicon solar cells – Effects
on electrical and optical properties, submitted for publication in the Solar Energy
Materials and Solar Cells.
Y. Yang, H. Kampwerth, S. Pillai, M. A. Green, H. Mehrvarz, A. Ho-Baillie (2012),
Angular reflection study to reduce plasmonic losses in the dielectrically displaced back
reflectors of silicon solar cells, submitted for publication in the Solar Energy Materials
and Solar Cells.
Y. Yang, S. Pillai, H. Mehrvarz, H. Kampwerth, A. Ho-Baillie, M. A. Green (2012), The
Effect of Rear Surface Passivation Layer Thickness on High Efficiency Solar Cells with
Planar and Scattering Metal Reflectors, Proceedings of the 38th IEEE Photovoltaic
Specialists Conference, Austin, USA, June 2012.
- 45 -
4.3.1.2
Thin, Light Trapped, Crystalline Silicon Solar Cells for High Efficiency at Low Cost
4.3.1.2.1 Overview
Epitaxial layer transfer technology is a promising way to make ultra-thin high performance crystalline
silicon solar cells. An efficiency of 20.6% was reported for a 35 um solar cell attached to resin and
fiber carrier [4.3.1.2.1] and 19.1% was reported for a 43 um free-standing solar cell [4.3.1.2.2]. This
paper reports initial fabrication results for 20 um silicon solar cells, electrically integrated with
bonded steel substrates and based on epitaxial layer transfer. The advantages of fabricating the thin
silicon solar cells on steel are robust during and after processing and matched thermal coefficient
between Si and steel. The fabricated cells had a rear p type emitter that was locally contacted through
a passivated surface structure and Al mirror known as PERC structure [4.3.1.2.3]. Uniform shallow
texturing was used with the width and height of the textured pyramids being ~1 um. Silicon
oxynitride (SiON) [4.3.1.2.4], which was deposited by PECVD on the front n type surface, was
patterned by laser doping [4.3.1.2.5] and a metal grid was formed by Ni/Cu self-aligned light induced
plating [4.3.1.2.6]. Based on measurements from test structures and calculations, a Voc of this cell
design can exceed 700mV and a Jsc of 38.3 mA/cm2 can be achieved without considering front grid
shading. Our best initial parameters are Jsc of 35.75 mA/cm2 on a textured solar cell with a nonoptimized AR coating, Voc of 623 mV and FF of 76.3%. The product of the best parameters indicates
17.0% efficiency. Our highest cell efficiency to date is 14.3% measured on a 1.21 cm2 cell.
4.3.1.2.2 Purpose of this work
The purpose of this work is to develop high performance 20um crystalline silicon solar cells
integrated with bonded steel substrate.
4.3.1.2.3 Approach
1.
2.
3.
4.
5.
6.
Develop an integrated design that will enable the thin silicon solar cell on steel to allow the
separate optimization of the key parameters;
Isolate the key parameters needed to achieve the potential and then develop test structures to
independently test them;
Separately optimise voltage through the device design (bulk, surface and contact
recombination);
Develop and integrate the optics for current increase through light trapping (adding texture
and back mirror);
Use epitaxial layers on FZ material for the surface and contact recombination optimization;
Epitaxial layer transfer technology;
4.3.1.2.4 Scientific innovation and relevance
1.
2.
3.
Fabrication of 20 um silicon solar cells on robust steel substrates;
Shallow texturing integrated with a back surface mirror is demonstrated realising Jsc values
of 35.8 mA/cm2;
Laser doping and light induced plating can be applied to these 20um silicon solar cells to
achieve FFs of 76%;
4.3.1.2.5 Results and Conclusions
Epitaxial silicon layers grown on porous silicon using reduced pressure chemical vapour deposition
(RPCVD) at AmberWave Inc. achieved reasonable lifetime. Best Photoluminescence (PL) counts
[4.3.1.2.7] of thin silicon on steel samples with SiON passivated front surface is 18000 per second,
which corresponds to a lifetime of ~100us according to correlation between PL counts and lifetime.
- 46 -
As stated in the introductive summary the highest Voc recorded to-date are 623 mV measured on a
non-textured cell and 620mV measured on a textured cell, and the highest Jsc is 35.75mA/cm2
measured using quantum efficiency on a textured solar cell with a non-optimized AR coating.
The highest efficiency recorded was 14.3% on a textured cell. The measured Jsc was 35.3mA/cm2
(which corresponded to a value of 37.3mA/cm2 after ARC correction), the Voc and FF were 602.3mV
and 74.8% respectively. To improve the efficiency, the AR coating and front contacts will be
optimized, which should result in increased Jsc and FF. The Voc is currently limited by the rear
surface passivation, and improved passivation schemes will be employed in the future.
4.3.1.2.6 Further Explanation
A typical planar solar cell structure, before laser-doping and plating, is shown in Figure 4.3.1.2.1.
SiON ARC
1 um n+ (1e18)
18um n 5e15
1um p+ 1e18
SiO2
Bonding stack w/ Al mirror
Steel substrate
Figure 4.3.1.2.1: Structure of planar 20um crystalline silicon solar cells on steel substrate.
Fabrication Steps
1.
2.
3.
4.
5.
6.
7.
8.
Epitaxial layer growth on porous silicon (AmberWave)
Rear p+ layer surface passivation and contact formation (AmberWave)
Epitaxial layer bonding and transferring onto steel substrate (AmberWave)
Front n type surface shallow texturing (UNSW)
AR coating and passivation (UNSW)
Laser selective doping and patterning (UNSW)
Self-aligned light induced plating (UNSW)
Laser edge isolation (UNSW)
Optics Measurement and Jsc Prediction
Light trapping is essential for indirect band gap thin silicon solar cells and effective light path length,
P, can be calculated using the following equation [4.3.1.2.8];
2
1
1
1
where W is the physical thickness of the cell, R is the average reflectivity at the rear surface, and f is
the fraction of light coupled out each time the light strikes the top surface, f= for textured surfaces
(n is the refractive index of silicon). Our measured rear surface reflectivity on test structure is 96.1%
at 1200 nm leading to an effective light path length 32 times the physical thickness of the solar cell.
Shallow texturing has been developed, and size of most pyramids is ~ 1um. Figure 4.3.1.2.2 shows
the SEM image of the textured surface and reflectance measurements with single layer AR coating.
- 47 -
Figure 4.3.1.2.2: SEM image of shallow textured surface (left), and two measured reflectance curves
with single layer AR coating (right).
The maximum Jsc for a 20 m solar cell was estimated to be 38.3mA/cm2 (without front-contact
shading) by assuming: (i) an optical thickness of 640um which is 32 times the thickness of our ultrathin silicon solar cells; (ii) a rear surface reflectivity of zero so that the light effective path length
equals to its physical thickness; (iii) a front surface reflectivity the same as measured and shown in
Figure 4.3.1.2.2; and (iv) all absorbed photons are collected.
QE and Reflection
100
90
QE or Reflection %
80
70
EQE corrected ARC
60
R corrected ARC
50
IQE
40
EQE non‐optimized ARC
30
R non‐optimized ARC
20
10
0
300
400
500
600 700 800 900
Wavelength (nm)
1000 1100 1200
Figure 4.3.1.2.3: QE and reflectance curves before and after correction on a textured sample.
As previously mentioned, the Jsc for the most efficient cell, which had a non-optimized AR coating,
was 35.3 mA/cm2. The IQE, EQE and reflectance curves that were measured for that cell are shown in
Figure 4.3.1.2.3. The cell’s reflectance curve is shifted from the ideal curve which is labelled as ‘R
corrected ARC’. The corrected EQE, which was calculated from the measured IQE and ‘R corrected
ARC’ data, was then used to estimate a Jsc of 37.3mA/cm2 for the cell (with a corrected ARC). This
result represents a significant increase from our previous reported results without light trapping where
the maximum Jsc was 17.6mA/cm2 [4.3.1.2.9].
- 48 -
Voc Analysis
The best Voc measured so far is 623mV which is much less than the expected 700mV. Figure
4.3.1.2.3 shows a good IQE response middle wavelength for a textured device, from which it can be
concluded that the minority carrier lifetime in the base is consistent with the lifetime estimated using
photoluminescence. The cell’s blue response is also reasonably good suggesting that front surface
recombination is probably not the reason for the lower than expected voltage either.
In order to investigate what was limiting the Voc, surface passivation was investigated for different
passivation layers and dopant densities (see Fig. 4.3.1.2.4). The results suggested that SiO2 was less
effective than SiON at passivating p-type surfaces in the dopant density 5e17cm-3. Silicon oxynitride
is also an effective passivation layer for n-type silicon. These results point the way to optimizing the
emitter, a less heavily doped p+ surface is preferred so as to achieve good surface passivation, SiON
is a good p+ surface passivation candidate at the doping density of 5e17cm-3.
p+ and n+ surface passivation by SiO2 or SiON
800
700
Lifetime (us)
600
500
400
p+ SiO2
300
p+ SiON
200
n+ SiON
100
0
0.E+00
5.E+17
1.E+18
2.E+18
2.E+18
Doping Density (cm‐3)
Figure 4.3.1.2.4: p+ and n+ surface passivation by SiO2 or SiON. It shows that SiO2 is a poor
passivation layer for p+ surface, SiON is a good passivation layer for p+ surface when doping density
is 5e17cm-3, and SiON is a good passivation layer for n+ surface.
References
4.3.1.2.1
4.3.1.2.2
4.3.1.2.3
4.3.1.2.4
Yahoo Finance 2012, Solexel Achieves Record 20.62% Thin-Silicon Cell Efficiency
Utilizing Applied Nanotech's Proprietary Aluminum Metallization Material, accessed 24
Feb
2013,
<http://finance.yahoo.com/news/solexel-achieves-record-20-62141500207.html>.
J. H. Petermann , D. Zielke , J. Schmidt , F. Haase , E. Garralage Rojas and R.
Brendel "19%-efficient and 43 μm-thick crystalline Si solar cell from layer transfer
using porous silicon", Progr. Photovoltaics: Res. Appl., 2012; 20:1-5.
Green MA, Blakers AW, Zhao J, Milne AM, Wang A, Dai X. Characterization of 23percent efficient silicon solar cells. IEEE Transactions on Electronic Devices 1990; 37:
331–336.
Brett Hallam, Budi Tjahjono, Stuart Wenham, “Effect of PECVD silicon oxynitride film
composition on the surface passivation of silicon wafers”, Solar Energy Materials &
Solar Cells 96, 173-179 (2012).
- 49 -
4.3.1.2.5
4.3.1.2.6
4.3.1.2.7
4.3.1.2.8
4.3.1.2.9
Hallam B, Wenham S, Sugianto A, Mai L, Chong C, Edwards M, Jordan D, Fath P.
Record large area p-type CZ production cell efficiency of 19.3% based on LDSE
technology, Proc. of the 36th IEEE Photov. Spec. Conf, Seattle, USA, 2011.
Alison Lennon, Yu Yao, Stuart Wenham, “Evolution of metal plating for silicon solar
cell metallisation”, Progress in Photovoltaics: Res. Appl. (2012), doi: 10.1002/pip.2221.
T. Trupke and R.A. Bardos, “Photoluminescence imaging of silicon wafers”, Applied
Physics Letters 89, 044107 (2006).
Campbell P. and M. A. Green, “Light trapping properties of pyramidally textured
surfaces”, Journal of Applied Physics 62(1): 243-249 (1987).
Allen Barnett, Ruiying Hao, C. Paola Murcia, Anthony Lochtefeld, Christopher Leitz,
Andrew Gerger, Michael Curtin “Independent approaches to increase voltage and current
in thin crystalline silicon solar cells”, 26th EUPVSEC, Hamburg, Germany, 2011.
- 50 -
4.3.1.3
III-V on Silicon Cells
With the 25% UNSW lab result providing a likely cap on what can be achieved in production with
conventional approaches, much higher efficiency is possible if tandem stacks of cells are used, with
the ultimate thermodynamic efficiency potential increased by a factor of 2-3. Silicon is a clean and
low-cost substrate and it is possible to grow high bandgap material on it. The limiting efficiency for
tandem solar cell using Si as bottom cell is quite close to the efficiency limit for tandem solar cell
with unconstrained material choice, which indicates Si is an excellent choice for the bottom cell.
Equally impressive to the progress with Si technology has been the rapid improvement in Group III-V
cell efficiency, driven by the use of these much more expensive cells in space and in systems with
sunlight focused 300-500 times. Efficiency has increased rapidly recently with 36% efficiency
demonstrated under standard sunlight, increasing up to 44% with strongly concentrated sunlight. The
preferred substrates for these cells are Ge wafers. Even though much cheaper than III-V
semiconductor wafers, Ge wafers are over 100 times more expensive than Si, providing the
fundamental limit on achievable III-V cell costs.
Tandem stack III-V technology provides a significant opportunity to improve efficiency but is
expensive in its usual form since Germanium (Ge) is a perfect match to the III-V group, but is over
100 times more expensive than Si. If the lattice spacing difference between Ge and Si is taken up near
the interface within a few atomic planes with a thin, essentially perfectly crystalline Ge layer formed
above this interfacial layer, the Si wafer can be converted into a “virtual Ge” substrate, suitable for the
subsequent deposition of a III-V cell stack.
Si Solar Cell Design for the III-V on Si Tandem Structure
In this work, we studied the design of the silicon device illuminated by a truncated solar spectrum due
to the optical filtering of the upper layers in the stack. Mainly, we have looked at the effect of the
junction depth and the doping concentration with different interface conditions in the silicon sub-cell.
Both modeling and experimental work have been performed using planar p-type Passivated Emitter
Rear Locally-diffused (PERL) high efficiency cell structure with high bulk lifetime substrate showing
good agreement.
To assist in the design the Si bottom solar cell, the solar cell simulation program (PC-1D), developed
at UNSW, has been used [4.3.1.3.1]. The emitter surface concentration range is between 1×1017 cm-3
and 1×1020 cm-3, while the junction depth is in the range of 0.1 μm to 10 μm. For implementing low
interface recombination, the model for thermally grown silicon dioxide passivation on n-type emitter
that has been developed in [4.3.1.3.2] is being used. Unlike the single Si junction solar cell where the
recombination at the emitter surface is minimized through the use of thermally grown oxide, the
ability to obtain very low interface recombination for the Si/buffer interface is not assured and so high
interface recombination conditions cannot be ruled out.
Figure 4.3.1.3.1 shows the modeled open circuit voltage (Voc) in volts as the emitter surface doping
level and emitter junction depth are varied for (A) low interface recombination and (B) high interface
recombination. As can be seen, at low doping levels, varying the junction depth does not impact the
voltage greatly; however, Voc drops with increasing junction depth at high emitter concentration. In
contrast, high interface recombination is showing totally different tendency; heavily doped emitter
profile is more preferable to attain high Voc where high interface recombination is expected.
- 51 -
1E20
0.4575
0.4675
0.4775
0.4875
0.4975
0.5075
0.5175
0.5275
0.5375
0.5475
0.5575
0.5675
0.5775
0.5875
0.5975
0.6075
0.6175
0.6275
0.6375
0.6475
0.6575
0.6675
0.6775
Doping Concentration (cm )
-3
-3
Doping Concentration (cm )
1E20
1E19
1E19
1E18
1E18
1E17
1E17
0.1
1
10
junction Depth (m)
(A)
0.1
1
junction Depth (m)
(B)
10
Figure 4.3.1.3.1: The modeled open circuit voltage in volt as a function of the emitter doping
concentration and the junction depth with (A) thermally grown oxide and (B) high interface
recombination velocity.
The sensitivity of emitter profiles with different interface conditions on the short circuit current (Jsc)
can be reliably studied using the modeled External Quantum Efficiency (EQE). Figure 4.3.1.3.2
illustrates the EQE for cases of lightly-doped and heavily-doped emitter profiles, the extreme two
cases, at low and high interface recombination. It can be seen that at high recombination, the lightlydoped emitter performs much better than the heavily doped. However, since the blue response is not
essential, as this part of spectrum is absorbed by the top junctions in multijunction solar cell, only a
small discrepancy is predicted. This is shown after the green line in Figure 4.3.1.3.3 which represents
the 1.5 eV junction above the Si subcell. This clarifies the small percentage change in the Jsc using the
modeled light IV measurement since all profiles have similar EQE at high wavelength.
100
90
80
70
EQE(%)
60
Lightly Doped Low FSRV
Lightly Doped High FSRV
Heavily Doped Low FSRV
Heavily Doped High FSRV
1.5 eV Junction
50
40
30
20
10
0
300
400
500
600
700
800
900
1000
1100
1200
Wavelength (nm)
Figure 4.3.1.3.2: The modeled external quantum efficiency for two emitter profiles at low and high
FSRV.
The experimental Voc of three emitter profiles, lightly-doped, intermediately-doped and heavily
doped, (A) with and (B) without SiO2 passivation are shown in Figure 4.3.1.3.3. The lightly-doped
emitter profiles have the highest Voc whereas the heavily ones have the worst at low interface
recombination. The opposite is observed with high recombination. Therefore, same results are
obtained as expected from the modeling for the two interface cases.
- 52 -
0.635
Lightly-doped
Mid-doped
Heavily-doped
0.68
Open Circuit Voltage Voc (V)
Open Circuit Voltage Voc (V)
0.69
0.67
0.66
0.65
Lightly-doped
Mid-doped
Heavily-doped
0.630
0.625
0.620
0.615
0.64
0.610
TP10-1
TP10-2
TP10-3
TP10-1
TP10-2
TP10-3
(B)
(A)
Figure 4.3.1.3.3: The experimental Voc for the lightly-doped (red), mid-doped (green) and heavilydoped (blue) emitter profiles (A) with (B) without SiO2 passivation.
The experimental spectral response measurement is shown in Figure 4.3.1.3.4. The straight lines
represent the case where SiO2 is passivating the emitter whereas the dash lines represent the high
interface recombination. Additionally, the green lines are for the lightly-doped profile, the red lines
are for the mid-doped profile and the blue lines are for the heavily-doped profiles. Only two samples
have been chosen showing the best and the worst scenarios in each profile. This perfectly agrees with
the modeling as well.
70
60
EQE (%)
50
40
TP10-1-4 no SiO2
TP10-2-6 no SiO2
30
TP10-3-6 no SiO2
20
TP10-1-4 SiO2
TP10-2-6 SiO2
10
TP10-3-6 SiO2
1.5 eV Junction
0
400
600
800
1000
1200
Wavelength (nm)
Figure 4.3.1.3.4: Experimental EQE of the three emitter profile with SiO2 passivation and after
removing the passivation. Green: lightly-doped, Red: mid-doped, Blue: heavily-doped, Orange:
1.5 eV junction, Straight line: with SiO2, and Dash line: without SiO2.
Further work is to be done on comparing modeled results for n-type PERL, p-type rear emitter
Passivated Emitter Rear Totally-diffused (re-PERT) and n-type re-PERT cells with experimental
results.
Low Cost Atomically Abrupt Si/Ge Transition for Virtual Ge Substrate
In this work, a novel approach of growing Ge on Si, where the lattice mismatch will be taken up
within a few atomic layers, deemed to be thermodynamically feasible because of the low energy
configuration, is investigated. Another unique feature with this work is the use of sputtering, an
environment friendly, economical, high throughput and simple deposition technique. Sputtering
avoids the requirement of a UHV system, the associated low deposition rates and the use of toxic
germane. Due to the thinness of the virtual Ge substrate, it can either be used as a passive
interconnection or as an active device, sharing current with the Si cell through the use of lighttrapping, see Figure 4.3.1.3.5. This light-trapping, pioneered at UNSW during the development of its
- 53 -
25% cell makes an “out of sequence” tandem feasible, since all low energy light does not have to be
absorbed by the Ge layer on its first pass. By trapping this light in the cell, most of it eventually will
be absorbed in the Ge layer, since not able to be absorbed by the other higher-bandgap layers present.
This allows the Ge layer to be very thin while still generating its share of current. This is a feature not
previously used in tandem cell design.
Figure 4.3.1.3.5: Possible cell stack grown on Ge coated Si wafer.
In the past year, experimental work has been focused on investigating the effects of sputter
conditions, thermal anneals and surface patterning on the crystallinity and dislocations on the virtual
Ge substrate and is summarized below.
Experimental Work Investigating Sputtering Conditions for Virtual Ge Substrate
Figure 4.3.1.3.6 shows the XRD pattern by the Ge layers sputter deposited at various temperatures.
The increased XRD peak intensity and the narrowing of the peaks suggest that higher temperature
favors the crystallinity. However, surface roughness is compromised at a higher temperature, see
Figure 4.3.1.3.7. This is possibly due to the changes in the island growth mode from 2 dimensional to
3 dimensional with increasing temperature. This effect can be minimized by the incorporation of
hydrogen, see Figure 4.3.1.3.8, where root means squared surface roughness has been reduced from 3
nm to 1nm for a Ge layer sputtered at a temperature of 450°C. This is possibly due to the presence of
the Ge-H cluster that serves to reduce diffusion barrier resulting in higher diffusivity & surface
mobility for the Ge growth.
Figure 4.3.1.3.6: XRD pattern of the virtual Ge substrates sputtered at various temperatures.
- 54 -
3D AFM
7
6
RMS(nm)
5
4
350
3
2
1
0
300
350
400
450
500
550
600
o
Substrate Temperature( C)
(a)
550
(b)
Figure 4.3.1.3.7: Atomic force microscopic (AFM) (a) calculations of the root mean squared
roughness in nanometers and (b) images surface morphologies of the virtual Ge substrates sputtered at
various temperatures.
(a)
(b)
Figure 4.3.1.3.8: AFM images of the surface morphologies of the virtual Ge substrates sputtered (a)
without and with (b) hydrogen.
- 55 -
In-situ vs Ex-situ Deposition Thermal Anneal
Different thermal anneal techniques have been employed and compared to study their effects on
reducing the threading dislocations in the Ge layer. Figure 4.3.1.3.9 shows the XRD patterns of two
Ge layers. Both samples under-went an initial 10nm sputtering deposition under the same conditions
at 500°C. One sample is then in-situ annealed at 700°C for 10 mins while the other had an ex-situ
anneal in a furnace at 900°C for 10 mins. A final 300nm Ge layer was then deposited in the same
conditions at 500°C on both samples. The better crystalline quality in the ex-situ annealed sample is
due to a higher temperature anneal resulting in reduced dislocation density (Figure 4.3.1.3.10) as well
as the formation of a SiGe alloy at the interface blocking many defects within the Ge layer, see Figure
4.3.1.3.11. It is anticipated that in-situ annealing using a newly acquired AJA sputter tool that is
capable of reaching a substrate temperature of 900°C will possibly produce better results to those
achieved by the current furnace annealing technique.
Rapid thermal annealing is another attractive option in reducing the dislocations of Ge film by
diffusing defects from Ge layer into Si wafer leaving behind a low dislocation density surface, see
Figure 4.3.1.3.12. Further work is required to avoid the formation of SiGe alloy throughout the entire
layer. Energy-dispersive X-ray spectroscopy (EDS) result shows that this is possible as the Si
concentration is seen to decrease towards the top surface, see Figure 4.3.1.3.13. Hence it is anticipated
a modified cooling process will likely result in a pure Ge layer with low dislocations at the top
surface.
Figure 4.3.1.3.9: XRD pattern of the virtual Ge substrates that underwent ex-situ and in-situ anneal.
- 56 -
Ge
SiO
Si wafer
(b)
(a)
(c)
Figure 4.3.1.3.10: TEM images of the virtual Ge substrates (a) before (b) after in-situ and (c) after
ex-situ anneal.
Si wafer
Ge film
Ge film
Interface
interface
Si wafer
(a)
(b)
Figure 4.3.1.3.11: TEM images of the virtual Ge substrates showing the Si-Ge interface after (a) insitu and (b) ex-situ anneal.
Figure 4.3.1.3.12: TEM image of a virtual Ge substrate deposited at 300°C followed by rapid
thermal anneal at 800°C for 3 minutes.
- 57 -
Pt
SiGe
Si
(a)
(b)
Figure 4.3.1.3.13: TEM image and EDS line scan of a virtual Ge substrate after a rapid thermal
anneal.
4.3.1.3.4 Identification of Surface Dislocation Density
To identify surface dislocation density of the virtual Ge substrate, a simple and fast technique of
preferential etch using a solution of 0.4M CrO3 and 49% HF at low temperature ~2°C [4.3.1.3.4] has
been employed. Figure 4.3.1.3.14 reveals defects on the surface of a 300 nm Ge layer at 2
magnifications. The cross-sectional TEM image in Figure 4.3.1.3.15 of the corresponding Ge sample
confirms that the original sample has not been over-etched, thus the defects revealed by the
preferential etching process were only those on the surface. Both the plan view (Figure 4.3.1.3.16)
and the cross-sectional TEM images, which are projection images, identifies the threading dislocation
density (TDD) within our Ge layers to be in the order of 109 to 1010 /cm2 while the TDD on the
surface of the Ge layer revealed by preferential etching as seen in the SEM image in Figure 4.3.1.3.15
was found to be 6108/cm2. This shows the potential of our low cost Ge growth technique to be
capable of achieving low defect surface.
(a)
(b)
Figure 4.3.1.3.14: Microscopic image at (a) low and (b) high magnifications showing virtual Ge
substrate surface dislocations revealed by preferential etching.
- 58 -
Figure 4.3.1.3.15: Cross sectional TEM of the same sample shown in Figure 4.3.1.3.14.
Figure 4.3.1.3.16: Plan view TEM of a virtual Ge substrate deposited at similar conditions as that
shown in Figure 4.3.1.3.15.
References
4.3.1.3.1
4.3.1.3.2
D. A. Clugston and P. A. Basore (1997), PC1D version 5: 32-bit solar cell modeling on
personal computers in Proc. of 26th IEEE Photovoltaic Specialists Conference, 1997,
p. 207.
A. Cuevas, G. Giroult-Matlakowski, P. A. Basore, C. Dubois, and R. R. King (1994),
Extraction of the surface recombination velocity of passivated phosphorus-doped silicon
emitters in Proc. of 24th IEEE Photovoltaic Specialists Conference, 1994, p.1446.
- 59 -
4.3.1.3.3
4.3.1.3.4
A. Langdo, et. al. (2000), High quality Ge on Si by epitaxial necking. APL, 2000.
76(25): p. 3700-3702.
J. Wernera, K. Lyutovich, and C. P. Parry (2004), Defect Imaging in Ultra-thin SiGe
(100) Strain Relaxed Buffers, Eur. Phys. J. Appl. Phys. 27, 367–370 (2004).
Publications
Ziheng Liu, Xiaojing Hao and Martin A. Green (2012), Influence of Substrate Temperature on
Epitaxial Growth of Ge on Si by Sputtering, In Proc. of 27th EUPVSEC, Sep 2012, Frankfurt
Germany. pp 298-300.
Ziheng Liu, Xiaojing Hao and Martin A. Green (2012), Investigation of Thin Epitaxial Ge films on Si
Substrates, 22th PVSEC, Nov 2012, Hangzhou China.
- 60 -
4.3.1.4
High Performance, High Voltage GaAsP on Silicon Solar Cells
Project Summary
The development of a cost-effective high voltage solar cell that can be grown directly on a siliconbased solar cell can lead to a significant increase in efficiency over the present best monocrystalline
silicon device. This achievement could lead to a significant reduction in the projected cost of solar
generated electricity. Devices composed of groups III, IV, and V of the periodic table have been
grown on silicon substrates. Each iteration, beginning with the single-junction windowless GaAsP
solar cell and concluding with a GaAsP/SiGe tandem device, is analyzed and reported. GaAsP
bandgap-voltage offsets achieve 0.54 in single-junction devices while dual-junction devices
demonstrate a device toward 15.2% efficiency (AM1.5) with AR-correction and rear texturing
4.3.1.4.1 Introduction
The development of a cost-effective high voltage solar cell that can be grown directly on a siliconbased solar cell can lead to a significant increase in efficiency over the present best monocrystalline
silicon device. This achievement could lead to a significant reduction in the projected cost of solar
generated electricity.
III-V materials are an excellent top solar cell candidate due to their proven performance near the
radiative limit [4.3.1.4.1]. However, the development of III-V solar cells on silicon (Si) substrates has
historically proven to be challenging.
The leveraging technology of this initiative is a metamorphic silicon:germanium (SiGe) buffer
[4.3.1.4.2] between the Si substrate and the active device layers. As developed by AmberWave Inc.,
it provides a low-dislocation interface for III-V nucleation and a high-quality bottom cell grown by
reduced pressure chemical vapor deposition (RPCVD) [4.3.1.4.3]. A basic layout of the material
structure is provided in Fig. 4.3.1.4.1.
Figure 4.3.1.4.1: Basic material structure layout.
4.3.1.4.2 Motivation
4.3.1.4.2.1
Single-junction detailed balance
Figure 4.3.1.4. 2 is yielded via detailed-balance methodology for single-junction efficiency limits at
AM1.5 [4.3.1.4.4]. Also plotted are some notable experimental record efficiencies. The two devices
that are closest to detailed-balance limits are gallium arsenide (GaAs) and Si. GaAs is predicted to
reach 33.2% and has resolved 28.3% efficiency in practice. Si is predicted to reach 33.4% and has
reached 25% efficiency in practice.
- 61 -
Figure 4.3.1.4.2: Detailed-balance single-junction solar cell efficiency.
Figure 4.3.1.4.3: Experimental percentage of detailed-balance maximum efficiency.
Figure 4.3.1.4.3 presents this data from another perspective, demonstrating a device’s percentage of
detailed-balance maximum efficiency. GaAs and Si are currently at 85% and 75%, respectively, of
their maximum predicted efficiencies.
However, other intrinsic limitations exist in Si, such as Auger recombination, which will limit the
efficiency beyond that predicted by detailed-balance. Accounting for these limitations will bring the
theoretical maximum efficiency of Si to 29.8% [4.3.1.4.5]. Therefore, Si is currently 84% of its
maximum predicted efficiency.
4.3.1.4.2.2
Dual-junction theoretical efficiency
Due to the impressive accomplishment of already reaching 84% of the theoretical maximum
efficiency of a Si solar cell, it is expected that the greatest gains from Si will not be from the singlejunction approach, but instead from the development of a tandem in which Si is the substrate.
Dual-junction models have been reported under the assumption of a top cell performing at the
radiative limit and a bottom cell limited by a bandgap-voltage offset (Woc) of 0.4 [4.3.1.4.6]. The
result suggests that with a bottom Si cell, a top cell of 1.72 eV can be applied to generate a two-
- 62 -
terminal (2T) efficiency of 42%. If a boundary condition is applied in which the III-V top cell must
be lattice-matched to a SiGe device, an ideal combination can be reached with gallium arsenic
phosphide (GaAsP) at 1.58 eV/0.84 eV and 40% tandem efficiency. Due to the mature SiGe graded
buffer technology available to this research, it is clear that a GaAsP/SiGe device has vast potential for
high-efficiency.
4.3.1.4.3 Background
Prior research on the SiGe single junction solar cell yielded devices with a 3E5 cm-2 threading
dislocation density (TDD) and an effective optical path length increase of 17x via a textured back
reflector [4.3.1.4.7]. The devices without light trapping reached 350 mV open-circuit for SiGe with
88% Ge concentration.
Afterward, GaAsP single-junction structures, detailed in Fig. 4.3.1.4.4, without a window layer were
grown and fabricated. Those results reached a Woc of 0.55 [4.3.1.4.8]. Using predictive modelling
developed concurrently, external quantum efficiency (EQE) of the device was fit with modelled EQE
to yield a front surface recombination velocity of 4E5 cm/s and a GaAsP/GaInP rear interface
recombination velocity of 8E4 cm/s.
Figure 4.3.1.4.4: Generation 1 GaAsP device layout.
4.3.1.4.4 Experimental Methodology
New structures were developed targeting voltage and efficiency improvements from the previously
reported windowless generation 1 devices. The second generation device includes a window layer
and targeted higher bandgap combinations of 1.662 eV GaAsP and 0.88 eV SiGe. Most recently, the
third generation device includes tunnel junctions and a bottom SiGe cell. The second and third
generation devices are provided in Figs. 4.3.1.4.5 and 4.3.1.4.6.
Figure 4.3.1.4.5: Generation 2 GaAsP device layout.
- 63 -
Figure 4.3.1.4.6: Generation 3 GaAsP/SiGe tandem device layout.
All devices have a 1.5 micron GaAsP base and 100 nm GaAsP emitter. The BSF is 100 nm and all
III-V growth is completed in less than 2.5 microns of epitaxy. The generation 1 contact layer is kept
very thin, 40nm, as it will not be etched off. Generations 2 and 3 have a 200nm contact layer which is
etched off after front metal evaporation. The generation 3 tandem has 30nm highly doped tunnel
junction layers.
The III-V devices were grown on a Veeco K475 As/P tool at temperatures ranging between 600C and
650C. Dopants of Si and Te are used for n-type layers while Zn and C are used for p-type layers.
GaAsP and GaInP were grown under high V/III ratios of 45 and 122, respectively.
All devices have 1 micron of Al for rear metal contact. Front surface contacts are unannealed Ti/Au
(20 nm/200 nm) or Ni/AuGe (5 nm/200 nm). Devices are developed for numerous cell sizes: 1 mm2,
9 mm2, and 1 cm2. Devices are isolated using standard selective and non-selective III-V etchants.
The front metal grid is responsible for shadowing 5% of the front surface. No AR coating has been
deposited at this time.
No textured back reflector has been applied to the generation 3 device as the focus here is on the
development of the top cell and tunnel junctions. Previous results [4.3.1.4.7] for SiGe solar cell
textured back reflectors suggest that the SiGe device will not be current limiting with the 1.662 eV top
cell. In the absence of texturing, the SiGe solar cell in this work is not isolated so as to prevent
current limiting behavior from the bottom cell and provide a more complete analysis of the III-V
layers and device.
Devices have been tested indoors under an ELH lamp calibrated for 1-sun measurement using an
outdoor-tested calibration cell. All reported results are for 2T device tests.
4.3.1.4.5 Device Results
4.3.1.4.5.1
Generation 2
From predictive modelling [4.3.1.4.8], it is expected that the inclusion of a window layer will reduce
the Woc by 10 mV. In practice, this was found to be the case, where the device bandgap is 1.662 eV
and the Voc is measured to be 1.125 V, yielding a Woc of 0.54. Table 4.3.1.4.1 provides measured and
modelled comparison with generation 1 devices.
Table 4.3.1.4.1: Reported results and modelling expectations of generation 1 and generation 2
devices.
Device
Gen. 1
Gen. 2
Eg
Experimental
Modeled
Voc
Woc
Voc
Woc
1.560 1.006 0.55 1.095 0.47
1.662 1.125 0.54 1.206 0.46
- 64 -
A Jsc-Voc plot is generated as Fig. 4.3.1.4.7 in order to isolate diode characteristics independent of the
series resistance (Rs) [4.3.1.4.9]. This is also beneficial since the device was developed for quick Voc
testing and was thus not optimized for efficient device operation.
Figure 4.3.1.4.7: Generation 2 device Jsc-Voc plot.
Figure 4.3.1.4.7 suggests the capability of a 76% FF for a device with no limitations from Rs.
4.3.1.4.5.2
Generation 3
With the inclusion of tunnel junctions and a bottom cell, it is expected that the voltage will be boosted
significantly. Experimentally, Voc reached as high as 1.32 V on some devices, surpassing the single
junction result by almost 0.2 V. The device which yielded the highest efficiency was 11% efficient
with a Voc of 1.3 V, Jsc of 12 mA/cm2, and FF of 70.5%. The IV curve is provided in Fig. 4.3.1.4.8.
When correcting for the inclusion of an AR coating and adding the textured back reflector, the
efficiency is expected to reach 15.2%.
Figure 4.3.1.4.8: Generation 3 tandem device IV and Jsc-Voc plot.
From the Jsc-Voc curve, a device with no Rs is expected to have a FF of 78%. The Rs will likely be
improved in future device generations after tests to develop optimized metallization schemes and
anneal recipes. The multijunction diode ideality factor was extracted [4.3.1.4.10] to be 4, which
suggests roughly n=2 at each junction.
The tandem device EQE is reported in Fig. 4.3.1.4.9. From EQE, the GaAsP device has a Jsc
equivalent to 12.1 mA/cm2, slightly higher than the Jsc registered from IV sweep under the 1-sun
- 65 -
calibrated lamp. This suggests, unsurprisingly, that some error exists in the indoor measurement of
devices under an ELH lamp. The EQE also confirms that the bottom cell would be limiting if its
junction area was defined as the top cell’s. The Jsc of the SiGe is resolved to be 8.1 mA/cm2.
Figure 4.3.1.4.9: Generation 3 device external quantum efficiency.
In Fig. 4.3.1.4.10, the 1.56 eV windowless generation 1 GaAsP solar cell is compared with the 1.66
eV generation 3 solar cell with a window layer. Results show significant boost in high-energy photon
collection efficiency with the inclusion of a 50 nm GaInP window layer.
Figure 4.3.1.4.10: EQE comparison between generation 1 & 3.
The GaInP/GaAsP interface recombination velocity extracted from generation 1 device modelling
[4.3.1.4.8] has been applied in generation 3 modelling to the window/emitter interface and the
base/bsf interface. Experimental EQE results show significant agreement with simulation. The EQE
step height change in generation 3 modelling at roughly 590 nm is an artifact of the simplifications
made in estimating the GaInP window layer absorption coefficiencts to facilitate their use over a wide
range of GaInP compositions.
Table 4.3.1.4.2 compares generation 1 and 3 devices by evaluating their average collection
efficiencies. Since the devices have different bandgaps, a direct comparison of Jsc does not adequately
compare the devices. Therefore, a simple method is used to consider the average collection efficiency
via the ratio of AR corrected current collection (Assuming 38% increase in transmission from AR
coating) to the current collection assuming 100% EQE for all photons with energy greater than or
equal to the bandgap. This method has some drawbacks, since experimental devices will collect some
photons at energies below their bandgap, but as these devices are direct bandgap with sharp dropoffs
in collection efficiency below their bandgap energy, the method is expected to provide an excellent
comparison.
- 66 -
Table 4.3.1.4.2: AR corrected average collection efficiency.
Device
Eg
Jsc
Gen. 1
Gen. 3
1.56
1.66
12.7
12.1
AR
Corr.
Jsc
17.5
16.7
Jsc for
100%
EQE
27.0
23.7
Average
Collection
Efficiency
64.9%
70.5%
Results in Table 4.3.1.4.2 show greater than 5% absolute increase in collection efficiency with the
inclusion of a 50nm GaInP window layer.
4.3.1.4.6 Conclusions
Results show a progression of better results with successive device generations. Generation 2 single
junction devices with window layers show an improved Woc of 0.54, a 10 mV better offset than
generation 1 devices previously reported. This 10 mV improvement in offset was also predicted by
simulations.
Generation 3 devices successfully add the bottom cell and tunnel junctions for operation at 1-sun.
Current collection has been improved greater than 5% absolute over generation 1 windowless solar
cells. Device voltages have been tested at a maximum of 1.32 V, which is a roughly .2 V boost from
the generation 2 device. This voltage increase is lower than expected, and increased attention must
therefore be paid to understanding the Voc in future work.
Acknowledgement
This research was, in part, funded by the Australian Solar Institute and United States-Australia Solar
Energy Collaboration.
References:
4.3.1.4.1
4.3.1.4.2
4.3.1.4.3
4.3.1.4.4
4.3.1.4.5
4.3.1.4.6
4.3.1.4.7
4.3.1.4.8
4.3.1.4.9
4.3.1.4.10
M. Green et al., “Solar cell efficiency tables (version 39),” Prog. Photovolt: Res. Appl.,
20, 12-20 (2012).
E. A. Fitzgerald et al., “Totally relaxed GexSi1-x layers with low threading dislocation
densities grown on Si substrates,” Appl. Phys. Lett., 59, 811 (1991).
M. Erdtmann et al., “Growth and characterization of high-Ge content SiGe virtual
substrates,” in Proc. Electrochemical Society, 11, 106-117 (2003).
W. Shockley and H. Queisser, “Detailed balance limit of efficiency of pn junction solar
cells,” J. Appl. Phys., 32, 510-519 (1961).
T. Tiedje et al., “Limiting efficiencies of silicon solar cells,” IEEE T. Electron Dev., 31,
711-716 (1984).
K. J. Schmieder et al., “Analysis of tandem III-V/SiGe devices grown on Si,” in Proc.
38th IEEE Photovoltaic Specialists Conf., Austin (2012).
Y. Wang et al., “Design and demonstration of light trapping for Ge:Si solar cell below Si
solar cell in a multi-junction solar cell system,” in Proc. 38th IEEE Photovoltaic
Specialists Conf., Austin (2012).
K. J. Schmieder et al., “GaAsP/SiGe on Si tandem solar cell performance predictions and
experimental results,” in review (2012).
R. A. Sinton and A. Cuevas, “A quasi-steady-state open-circuit voltage method for solar
cell characterization,” in Proc. 16th European Photovoltaic Solar Energy Conf.,
Glasgow (2000).
R. R. King et al., “Band gap-voltage offset and energy production in next generation
multijunction solar cells,” Prog. Photovolt: Res. Appl., 19, 797-812 (2011).
- 67 -
4.3.2
4.3.2.1
INDUSTRY COLLABORATIVE RESEARCH AND COMMERCIALISATION
Introduction
The Centre’s Technology Transfer team (TTT) was established in 2008 to accelerate the development
of the Centre’s commercially significant technologies and to carry out technology transfers to
industry. This includes assisting companies to establish and optimise the large scale production of
UNSW photovoltaic technology. With continued high demand for UNSW technology and the
inclusion of two new industry partners, the activities of the TTT team continued to expand. Figure
4.3.1 shows the TTT in June 2010 with members from left to right being D. Jordan, A. Sugianto, S.
Wenham, A. Lennon, B. Hallam, L. Mai, B. Tjahjono and N. Kuepper. More recent members
Catherine Chan, Chee Mun Chong, Phil Hamer and Danny Chen are not shown, while Dr Matt
Edwards was away on the date of the photo.
Figure 4.3.2.1: TTT photo June 2010 with Matt Edwards overseas at the time and new comers
Catherine Chan, Chee Mun Chong, Phil Hamer and Danny Chen about to join the team.
4.3.2.2
Industry Collaborators and Licensees of Technology
The Centre has a large number of collaborators and licensees including many cell manufacturers
working with the Centre to improve or develop new solar cell technology for commercialisation.
Some of these fund the Centre quite generously and are seen as being partners to NewSouth
Innovations (NSi), the commercial arm of the University that manages the entire intellectual property
portfolio for the Centre. Interest from companies wanting to become commercial partners with the
Centre or license technology grew rapidly during the last two years despite the international financial
downturn in late 2008 and subsequent international financial turmoil. This appears to be due in part to
the perceived world leadership of UNSW in this area, but also due to the growing long-term
importance placed on photovoltaics by a world struggling to deal with climate change, diminishing
resources and increasing energy requirements. Collaborators with the Centre in the first generation
photovoltaics area come predominantly from China, Germany, the United States, Australia, South
Korea and Taiwan and include Suntech-Power, Roth and Rau, Guodian Solar, Centrotherm, Tianwei
New Energy, Optomec, Sunrise Energy, Sunergy, CSG Solar, E-ton Solar, Spectraphysics, Hyundi
Heavy Industries, China Light, Corum Solar, Advent, Shinsung, Apollon Solar, LG Electronics,
Dopont, Heraeus and Yunnan Tianda. Several of these companies are amongst the world’s largest
solar cell and equipment manufacturers.
- 68 -
4.3.2.3
4.3.2.3.1
Commercially Relevant Technologies
Introduction
With screen-printed solar cells continuing to dominate commercial manufacturing with well over 50%
market share, the broad aim of this work has been to develop, in conjunction with several industry
partners, the next generation of screen-printed solar cell. In particular, the fundamental limitations of
the conventional screen-printed solar cell that have limited its performance for the last 30 years have
been identified, and innovative approaches to redesigning the emitter and front metal contact have
been devised, developed and analysed in this work. In addition to overcoming the current and voltage
limitations imposed by the design shown below, a further aim of this work has been to retain
compatibility with existing equipment and infrastructure currently used by our industry partners for
the manufacture of screen-printed solar cells.
150-200
120-150
m m
3 mmm
m
2-3
patterned metal contact
phosphorus
bulk of wafer
n ++
rear metal contact
p-type
p+
metal
Figure 4.3.2.2: Standard screen-printed solar cell.
Despite the dominance of this technology, this solar cell design shown in Fig. 4.3.2.2 has significant
performance limitations that limit the cell efficiencies to well below those achievable in research
laboratories around the world. In particular, the front surface screen-printed metallisation necessitates
a heavily diffused emitter to achieve low contact resistance and also to achieve adequate lateral
conductivity in the emitter since the metal lines need to be widely spaced compared to laboratory cells
to avoid excessive shading losses. Such cells therefore typically have emitters with sheet resistivities
in the range of 50-60 ohms per square, which inevitably give significantly degraded response to short
wavelength light. Good work in recent years by partners Dupont and Heraeus has improved paste
formulations, allowing ohmic contacts to be formed to more lightly doped emitters in the vicinity of
60 ohms/square while simultaneously lowering the bulk resistivity of the fired paste. To further raise
this sheet resistivity to above 100 ohms per square, as required for near unity internal quantum
efficiencies for short wavelength light, serious resistive losses are introduced, both in the emitter and
through the contact resistance at the metal to n-type silicon interface. Furthermore, the metal/silicon
interface contributes increasing dark-saturation current densities as the emitter doping is increased
when using a homogeneous emitter.
Furthermore, the conventional design for screen-printed solar cells has quite poor surface passivation
in both the metallised and non-metallised regions. Even if good ohmic contacts could be made to
more lightly doped emitters, the large metal/silicon interface area would significantly limit the
voltages achievable due to the high levels of recombination in these regions and the corresponding
contribution to the device dark saturation current. These voltage limitations are not of major
- 69 -
significance at the moment due to the limitations imposed by the substrates. However, in the future as
wafer thicknesses are reduced to improve the device economics and improved rear surface passivation
is introduced, the cells will have the potential for improved open circuit voltages, but only provided
the surfaces, including under the metal, are well passivated.
4.3.2.3.2
Semiconductor Finger Solar Cells
To accommodate a top surface emitter sheet resistivity of at least 100 ohms per square, metal fingers
need to be spaced no more than 1mm apart to avoid excessive sheet resistivity losses. Due to the large
width of screen printed metal lines or 100 microns or more, such a close spacing is not possible
without shading well over 10% of the cell surface. The concept of semiconductor fingers is therefore
introduced as shown in Figure 4.3.2.3. These semiconductor fingers are formed by laser doping the
silicon surface while simultaneously patterning the dielectric layer to expose the heavily doped silicon
surface. Developmental work with the laser doping process allows sheet resistivities as low as 1
ohm/square to be achieved while simultaneously forming laser doped lines of only 8 microns width.
Spacing such lines 0.8mm apart avoids significant resistive losses within the lightly doped emitter
which can be diffused to 100 ohms/square, ensuring excellent response to the short wavelengths of
light. In the laser doped regions, approximately half of the incident light is lost due to absorption
within the heavily doped semiconductor fingers.
This is defined as a 50% effective shading loss of the semiconductor fingers. Using this design, the
effective shading loss of the semiconductor fingers is only 0.5%, while the effective emitter sheet
resistivity of the emitter is 50 ohms/square. This latter figure results from the emitter sheet resistance
of 100 ohms/square in parallel with the laser doped lines which cover 1% of the area with a sheet
resistivity of 1 ohm/square and therefore also effectively contribute 100 ohms/square. This allows the
screen-printed lines to retain their normal spacing, but with 99% of the emitter being lightly doped
and therefore able to achieve near unity internal quantum efficiencies for short wavelengths of light.
In addition, the passivating dielectric layer not only passivates the lightly diffused surface so as to
give near unity internal quantum efficiencies for short wavelength light, but it also isolates the metal
from these same regions to minimise the device dark saturation current. Importantly, the silicon is
only exposed at the semiconductor fingers, with the screen-printed metal having been shown to make
excellent ohmic contact to the heavily phosphorus diffused silicon in these regions. Both thick oxides
and silicon nitride layers, when used with appropriate pastes, appear to provide adequate protection to
the lightly diffused surface regions, preventing the screen-printed metal from contacting the silicon.
Figure 4.3.2.3: Screen-printed fingers running perpendicular to the heavily diffused semiconductor
fingers where electrical contact is made. A dielectric/AR coating passivates the top surface and
isolates the metal from the lightly diffused top surface.
- 70 -
A typical cell design based on this concept uses semiconductor fingers 8 microns wide of sheet
resistivity 1-2 ohms/square and spacing of 0.8mm. The top surface sheet resistivity is typically 100
ohms/square, with the effective sheet resistivity in the direction parallel to the semiconductor fingers
typically 50-60 ohms per square. The screen printed metal lines are printed perpendicular to the
semiconductor fingers with a width of 100 microns and spacing of 2.2mm. The overall effective
shading loss of the combination of the semiconductor fingers and the screen-printed metal lines (not
counting interconnect strips) is therefore 5%, comprising 4.5% from the metal plus 0.5% from the
semiconductor fingers. This is typical of conventional screen-printed cells, but with the greatly
enhanced short wavelength response of these cells giving about a 4% advantage in short circuit
current.
This design also has the advantage of being the equivalent of a self-aligned selective emitter design
with the metal only contacting the silicon in heavily doped regions. This gives the cell top surface
structure much higher voltage capability although improved rear surface passivation will be necessary
to capitalise on this. The biggest challenge in turning this into a robust technology for large scale
production is the very small metal/silicon interface area which is less than 0.1% of the top surface.
This is sufficient for suitably low contact resistance to facilitate high fill factors provided everything
works properly. However, despite 18.6% efficiency being achieved in pilot production, variability in
this contact resistance is a weakness in the design causing efficiencies in production of the technology
to vary from 17 to 18.5%, with the average being below 18%.
The processing sequence for this technology is easily retrofitted onto a standard screen-print line and
is as follows:
1.
2.
3.
4.
5.
6.
7.
Surface texturing
Emitter diffusion (100 ohms per square)
Rear surface etch plus edge isolation
SiNx deposition by PECVD (top surface)
Semiconductor fingers by laser doping
Rear metal plus front metal screen-printing
Firing of metal contacts
Only step 5 deviates from standard homogeneous emitter screen-printed solar cell fabrication, with
the overall fabrication appearing to be simpler and shorter than those sequences proposed for
introducing selective emitter designs for screen-printed solar cells.
The best semiconductor finger solar cells have fill factors of 79%, demonstrated with this structure on
large area devices of approximately 150cm2, verifying the effectiveness of this contacting scheme for
minimising resistive losses. These cells also have near perfect response to short wavelength light as
shown below, leading to Jsc values of 36-37mA/cm2. Even though commercial p-type substrates are
not capable of voltages above about 640mV, Voc values approaching this have been achieved in pilot
production with this technology, with corresponding efficiencies as high as 18.6%.
- 71 -
120
100
%
80
EQE
60
Reflection
40
IQE
20
0
300
400
500
600
700
800
900 1000 1100 1200
Wavelength (nm)
Figure 4.3.2.4: Spectral response of semiconductor finger cell.
This concept of semiconductor fingers does not appear to have been used in large-scale commercial
solar cell production, and has considerable appeal as it facilitates good conductivity within the
emitter, but without the normal trade-off found in screen printed cells. Normally, such regions of good
emitter conduction are located at the top surface and therefore degrade the cell spectral response and
current generating capability due to the corresponding extremely short minority carrier diffusion
lengths in such regions. This technology is particularly well suited to multicrystalline silicon wafers
that normally degrade when exposed to prolonged thermal treatments.
Another possible implementation of this technology would appear to be the incorporation of the laser
chemical processing (LCP) techniques developed by the Fraunhofer Institute in Germany in
conjunction with the laser company Synova. This would allow the superior performance of the laser
grooved semiconductor finger solar cell to be combined with the simplicity and low cost of the laser
doped semiconductor finger solar cell. Both institutions have intellectual property that would appear
to provide significant benefits in combination.
4.3.2.3.3
Advanced Semiconductor Finger Solar Cells
As part of the core research of the Centre of Excellence funded by the ARC and UNSW, an enhanced
version of the semiconductor finger solar cell was designed and developed to overcome the
limitations of the standard semiconductor finger solar cell described in section 4.3.2.3.2. This involves
plating 2 microns of silver (or Ni/Ag or Ni/Cu/Ag) to the screen-printed metal at the end of
processing, which simultaneously plates a similar thickness to the semiconductor fingers when using
light induced plating as shown in Figure 4.3.2.5.
The application of this plated metal has several benefits. Firstly, it makes good ohmic contact to both
the heavily doped silicon and also the screen-printed metal, overcoming the high contact resistance
sometimes experienced by the standard semiconductor finger solar cell. This makes high fill-factors
routinely achievable while solving the yield problems of the original implementation. Secondly, the
increased conductivity of the semiconductor fingers facilitates increasing the screen-printed metal line
spacing to typically 2-3cm. Although the corresponding width has to be increased to 400 microns, it
has the benefit of allowing both tapering and increased height to above 50 microns. This reduces the
shading loss of the screen-printed metal fingers from 5% down to about 1%, with the corresponding
Jsc increase taking efficiencies to above 19%. Interestingly the effective shading loss of the
semiconductor fingers was expected to increase by 0.5% although in reality, the plated laser doped
surface appears rough enough to scatter the reflected light so that almost half is totally internally
reflected at the glass/air interface and returned to the solar cell surface. Therefore the effective
- 72 -
shading loss of the semiconductor fingers remains virtually unchanged. Importantly, by avoiding the
contact between the screen-printed silver and the lightly doped silicon, the top surface design of this
solar cell behaves very differently to standard screen-printed contacts and appears to have the
potential to achieve open circuit voltages approaching 700mV.
Figure 4.3.2.5: Photo of a hybrid screen-printed and plated solar cell showing the three levels of
metallisation. Each of the two busbars effectively collect current from 8 small solar cells, each with a
tapered screen-printed line that carries the current from the narrow plated lines to the respective
busbar/interconnect.
The development of this technology has formed part of the ASI project involving Suntech. It is
particularly well suited to multicrystalline silicon wafers due to the avoidance of prolonged high
temperature thermal processes. Based on small area test devices, the inclusion of rear surface
passivation is expected to take efficiencies on multi material to over 18% with corresponding open
circuit voltages in excess of 650mV while on mono, efficiencies of 21% are anticipated with open
circuit voltages exceeding 680mV.
4.3.2.3.4
4.3.2.3.4.1
Advanced N-type Screen-printed Solar Cells
N-type Screen-Printed Cells with Homogeneous Top Surface Diffusion
Increasing interest is being shown in n-type CZ material as a means for avoiding the widely reported
defects associated with the high boron and oxygen concentrations in p-type CZ material. In particular,
screen-printed aluminium has been used as a simple and cost-effective way to create an Al-alloyed
rear emitter for such n-type CZ material, especially in the n+np+ cell design with rear junction.
However low voltages for such structures reported in the range 617-627mV appear to be a severe
limitation of this approach since n-type CZ wafers should be capable of achieving much higher open
circuit voltages.
Discontinuities in the p+ layer have been identified as the main cause for such performance
degradation of this device. These discontinuities, as seen in Fig.4.3.2.6 (a), are isolated points where
the junction fails to form, usually created by non-uniform wetting of the silicon by the Al during the
- 73 -
alloying process. Although their presence can be minimized by optimizing the firing process so as to
allow uniform wetting of the surface to occur, they cannot be completely avoided. In small quantities,
such non-uniformities have almost negligible influence on the performance of the back surface field
in conventional p-type cells. However, they can significantly degrade the quality of the Al-alloyed
emitter in n-type wafers by allowing Al to locally bypass the p+ region and directly contact the n-type
bulk via a Schottky barrier causing a non-linear shunt. A new and modified firing process has
therefore been developed to avoid the damage from such non-uniformities. In this method, a patented
low temperature solid phase epitaxial growth process is employed after the conventional standard
spike firing to minimize the impacts of these junction discontinuities so that a uniform and good
quality junction as illustrated in Fig.4.3.2.6 (b) can be achieved.
(a)
(b)
Figure 4.3.2.6: Cross-sectional SEM photos show: (a) discontinuities in the Al-doped p+ layer; (b) a
deep and uniform Al-doped p+ layer.
These improvements have facilitated a 15-20mV increase in Voc relative to those reported in the
literature and efficiencies over 17% for standard screen-printed solar cell technology with
homogeneous n-type emitter (front surface field) applied to n-type wafers.
A common problem with the manufacture of conventional screen-printed solar cells is minute
amounts of aluminium paste accidently coming into contact with the cell front surface. This can
happen when wafers are face down during the printing of the rear or during wafer transfer/handling.
Aluminium paste contamination of such surfaces and transfer belts happens relatively easily in a
production environment such as through a broken wafer that has been Al printed, operators’
contaminated gloves, tiny holes in the screen-printing screen etc. However such contamination with
the current rear junction n-type technology creates unusual photoactivated shunts that are not present
in the dark. This is because a p-n-p transistor structure is formed when the unwanted aluminium on
the top surface is fired into the n-type surface (with the second p-n junction of course being at the rear
of the device). In this phototransistor, the lightly doped wafer forms the base of the transistor, and
despite the base being very thick, approaching 200 microns, the high lifetime of the CZ n-type
material allows the transistor to achieve moderate gain levels and therefore conduct large currents
when illuminated by light that generates the necessary base current for the transistor. The unusual
consequence is that the apparent shunt resistance of the cell is very poor (low) when illuminated
brightly, increasing to high values in the dark or even low illumination levels. The ramifications of
this for cell efficiency and fill-factor are that values fall with reducing light intensity as is normally
the case for shunted cells, but with values increasing again for low illumination levels as the shunt
problems disappear as the phototransistors are deactivated.
- 74 -
Even without such phototransistor shunting, analysis of such homogeneous emitter n-type devices
indicates that even higher voltages are potentially achievable if not for the large dark saturation
current contribution from the heavily doped phosphorus diffused top surface. This emphasises the
importance of moving to the equivalent of a selective emitter design for the front surface to facilitate
both improved short wavelength response as well as increased device voltages.
4.3.2.3.4.2
N-type Screen-Printed Cells with the Equivalent of a Selective Emitter
Until recently, the Centre held the world record (jointly with Stanford University) for the most
efficient n-type silicon devices with 22.7% efficiency. The cell design was based on the inverted form
of the PERL (Passivated Emitter and Rear Locally diffused) solar cell developed at UNSW and is
shown below. In this work, the cell design has been adapted to accommodate the use of low cost
screen-printed solar cell processes involving the alignment of the screen-printed front metal lines to
the heavily doped n+ regions to form the equivalent of a selective emitter on the front surface and the
use of a screen-printed aluminium grid pattern on the rear to form the localised p+ regions during the
spike firing of the Al.
"inverted" pyramids
finger
double layer
antireflection
coating
p+
n+
p+
rear contact
n
thin oxide
(~200Å)
+
p
p
0.9 -cm
n-silicon
p
oxide
Figure 4.3.2.7: Inverted form of the PERL (Passivated Emitter and Rear Locally diffused) solar cell
developed at UNSW based on the use of N-type silicon.
This approach enables the achievement of 18% efficiency but requires a quite complicated processing
sequence.
For comparison purposes, the screen-printed front contacts were replaced with laser doped contacts.
With a top surface homogeneous emitter diffused to about 200 ohms per square, the self aligned
metalisation with heavy doping beneath the metal contact was formed by melting the silicon through
the silicon nitride anti-reflection coating in the presence of an n-type dopant source. This laser doping
process automatically damages the overlying silicon nitride layer, facilitating direct plating of metal to
the heavily doped regions as a self-aligned top surface metal contact. The cell structure is shown
below. This approach appears to have significant advantages over the screen printing technique such
as self alignment to the locally diffused top surface, narrower metal lines and corresponding
significantly lower shading losses. Despite being a significantly simpler process than the screen-print
selective emitter design, excellent cell efficiencies of 19.0% have been achieved on large cell areas of
148.6cm2 using commercial grade CZ n-type wafers. Commercial partners have demonstrated even
higher efficiencies, exceeding 19.2% while simultaneously using only industrial equipment.
- 75 -
Figure 4.3.2.8: Schematic of a laser doped Al-alloyed rear junction n-type solar cell.
In this work, the cells were fabricated using industrial sized (125x125mm) phosphorus-doped CZ ntype wafers of 3ohm-cm resistivity and ~180m thickness. A texturing process was performed in a
NaOH/Isopropanol based solution to form random pyramids. A thin phosphorus diffused n+-layer
with a sheet resistivity of 200 ohm/sq serving as the front surface field was created by a thermal
diffusion at 850oC using liquid POCl3 source in a conventional tube furnace. A chemical etch
containing HF and HNO3 was then applied to the back of these wafers to remove the unwanted n+layer from this surface. A silicon nitride layer with a refractive index of 2.05 and thickness of 75nm
was subsequently deposited using a commercial Roth & Rau remote plasma PECVD system to
simultaneously form an anti-reflective coating and provide passivation for the front surface as well as
the bulk material. After screen-printing with a selected Aluminium paste on the entire back surface
with a gap of around 2mm from the edges, the wafers underwent an alloying process at 860oC in a
conventional conveyor belt furnace to produce the rear Al-alloyed emitter. A phosphorus dopant
source was then spun onto the front surface, followed by the laser doping process using a 532nm
wavelength Q-switch diode laser to create locally heavily diffused lines. The wafers were rinsed and
submerged in 1% HF solution for 30sec to remove the dopant source and any native oxide from the
laser doped lines. Lastly, light induced plating (LIP) was subsequently performed to deposit Ni, which
was sintered at 375oC, followed by Cu plating to form the front contacts.
The performance improvement relative to the selective emitter screen-printed counterparts arises
primarily from the reduced shading losses by the top surface metalisation although slightly higher
voltages as high as 650mV, fill-factors and yields are also achieved, apparently due primarily to the
reduced metal/silicon interface area and superior alignment with the laser doped contacts.
A spectral response measurement was performed to investigate different regions in the cell. A very
high value of more than 95% was maintained for internal quantum efficiency (IQE) from 580nm to
960nm. However, there was a slight drop in the short wavelength range, indicating that the front
surface passivation can be further improved.
- 76 -
100
90
80
70
60
IQE
(%) 50
EQE
40
Reflectance
30
20
10
0
360
480
600
720
840
960
1080
Wavelength (nm)
Figure 4.3.2.9: Spectral response of a typical laser doped n+np+ solar cell.
4.3.2.3.5
Buried Contact Solar Cells
Despite this technology being commercialised more than a decade ago, it remains a key technology
for collaborative research with industry and continues to do well commercially with close to $1
billion of product now deployed in the field and many new companies interested in its commercial
potential. New developments at the Fraunhofer Institute in conjunction with the laser company
Synova make the groove formation and doping a much simpler and lower cost process than
incorporated into the original technology implementation. The German company Rena is offering a
turn-key production line for this technology although it appears at this stage that significant laser
induced damage is making the achievement of high efficiencies difficult.
Never-the-less, this is expected to create significant new interest in the Buried Contact technology,
particularly with the apparent growing acknowledgement of the benefits of the technology over
existing screen-printed solar cell technology. In the 2006 European Inventor of the Year Awards, the
Buried Contact technology contributed to Green and Wenham receiving a Top 3 Ranking (out of
more than 200,000 inventions world-wide in the period 1990-2000) for inventors outside Europe. A
further distinction for this technology is its listing amongst Australia’s Top 100 Inventions of the 20th
Century as determined by the Australian Academy of Technological Sciences and Engineering. The
original design for this solar cell has the buried metal contacts on the top surface as shown in Figure
4.3.2.10 although in more recent years the interdigitated rear contact design applied to n-type wafers
shown in Fig. 4.3.2.11 has become particularly popular and a focus of UNSW research. Efficiencies
of 20% have been demonstrated on small area devices with efficiencies of 21% eventually believed
achievable on large area CZ material.
A newer implementation of this technology is under development in a collaboration between UNSW
and Tianwei in China. Efficiencies approaching 19% have already been demonstrated by using the
laser to simultaneously pattern the SiNx layer, ablate some silicon to form the groove and dope the
remaining side walls of the groove heavily with phosphorus. To study the performance potential of
this front surface design, cells were fabricated with a photolithographically defined PERL rear surface
with 20.8% efficiency being achieved.
- 77 -
Figure 4.3.2.10: Buried contact cell with buried contacts on both front and rear surfaces.
Figure 4.3.2.11: Buried contact cell with interdigitated rear contacts.
4.3.2.3.6
Laser Doped Selective Emitter Solar Cells
The benefits of a selective emitter have been well known and quantified for many years. The benefits
of heavy doping beneath the metal contacts contribute significantly to the high performance levels
achieved by technologies such as the Buried Contact Solar Cells, the semiconductor finger solar cell,
the point contact solar cells and the world-record holding Passivated Emitter and Rear Locally
diffused (PERL) solar cell. The heavy doping not only facilitates reduced contact resistance between
the metal and the silicon, but probably more importantly it shields the high recombination velocity
metal/silicon interface from the active regions of the cell. In addition, by restricting the heavily doped
material to the immediate regions beneath the metal contact, little light absorption takes place in such
regions thereby avoiding problems with carrier collection from heavily doped regions where the
minority carrier diffusion lengths are very short.
Despite the commercial success of the Buried Contact Solar Cell described above, a potentially more
effective and simpler way of achieving a selective emitter is by using laser doping to produce the
heavily doped regions beneath the metal contacts as shown in Fig. 4.3.2.12. Following top surface
- 78 -
emitter phosphorus diffusion to about 100 ohms per square and silicon nitride deposition, an n-type
dopant source is applied or can even be incorporated into the silicon nitride layer. A 532 nm NdYAG
laser is used to melt the silicon to a depth in the vicinity of a micron while simultaneously releasing
the n-type dopants into the molten region. The molten silicon subsequently regrows epitaxially,
heavily doped with phosphorus. Just as importantly, the overlying silicon nitride layer is removed
from the silicon surface in isolated regions, facilitating direct plating to the exposed n++ surface.
Electroless plating of Ni and Cu provides a particularly effective self-aligned metalisation scheme to
provide metal lines wherever the laser doping was effected. The laser doped regions are typically 12
microns wide, leading to metal lines of only 20 microns width after plating.
Figure 4.3.2.12: Laser doped selective emitter solar cell with self aligned plated metal contacts.
Innovative light induced plating techniques have been developed that facilitate the achievement of
very high aspect ratios for the plated Nickel and Copper. Using conventional electroless plating
techniques which tend to plate conformally with a uniform plating rate in all directions, the 10 micron
wide laser doped line plates to a width of about 30 microns for a plating height of about 10 microns as
shown in Figure 4.3.2.13 (a). In comparison, the new light induced plating techniques achieve line
widths of only 24 microns for even greater metal height as shown in Fig. 4.3.2.13 (b).
- 79 -
Figure 4.3.2.13: (a) electroless conformal plating of Cu onto a laser doped region and (b) Improved
aspect ratio for metal plated by the new light induced plating techniques where the height is 12
microns while the width is 24 microns. (c) FIB photo of Cu plating using the new light induced
plating techniques as for (b).
Contact resistances below 0.001 ohmcm2 have been demonstrated, leading to fill factors as high as
80% being achieved on devices of approximately 150cm2 in area. This has facilitated the achievement
of efficiencies above 19% on commercial-grade CZ p-type silicon. Of particular importance has been
the defect generation accompanying the laser doping process, particularly in conjunction with
dielectric-coated, textured silicon surfaces. Planar surfaces present minimal challenge in terms of
achieving near defect free regions in the vicinity of the laser melted regions. Textured surfaces
however, particularly in conjunction with dielectric coatings of significantly different thermal
expansion coefficient, have provided a significant challenge to match the low defect densities
achievable with planar devices. Important processing parameters in the optimization of the laser
doping process for textured surfaces have included laser pulse envelope shape, pulse duration, pulse
frequency, laser light frequency, laser power, beam focus as well as the type of dielectric and dopant
source being used.
The TT team has had significant success helping various companies get this technology into pilot
production, often achieving higher efficiencies and on larger area commercial wafers, than has been
achieved at UNSW. For example, 19.0% efficiency or higher has been achieved in pilot production at
several companies, four of which have chosen to do joint publications with UNSW documenting the
achievements including Sunrise Global Energy in Taiwan, Roth and Rau in Germany, Centrotherm in
Germany and Shinsung in Korea.
The preferred implementation of the laser doping selective emitter (LDSE) technology also uses an
equivalent laser doping/plating combination on the rear surface using a boron doping source. In this
cell design, the majority of both the front and rear surfaces is well passivated using silicon nitride
although the preferred rear surface passivation of the undiffused p-type surface uses somewhat
different deposition parameters for best results. Implied Voc values above 730mV at one-sun
demonstrate the near perfect passivation achieved with such surfaces. Following laser doping of the
rear surface, even though plated contacts can be used similarly to on the front surface, the preferred
rear contact is achieved through depositing aluminium over the entire surface followed by a low
temperature sinter used to form good ohmic contact with the boron laser doped regions. In this cell
design, the aluminium layer provides an excellent rear surface reflector. Even with full-size standard
commercial grade p-type CZ wafers, based on pilot production, this technology has achieved
independently confirmed efficiencies of 20% with pseudo efficiencies exceeding 23% therefore
showing the future potential of the technology once parasitic resistive losses are minimised.
Impressively, open-circuit voltages as high as 686mV have been achieved by these devices.
- 80 -
The simpler version of the LDSE technology equivalent to the pilot production established at Sunrise
with aluminium alloyed rear surface, is far simpler to retrofit onto existing screen-printed production
lines, with 90% of existing equipment retained. For several manufacturers such as Shinsung in South
Korea, this new technology whereby the laser doped selective emitter combined with light induced
plating of the metal contacts are used to replace the front surface screen-printed metal is achieving
significantly higher efficiencies well in excess of 19% compared to 18% for conventional screenprinted cells with minimal if any cost increase per cell. Roth and Rau have also secured the rights
from UNSW to develop and sell turn-key production lines based on the LDSE technology world-wide
This is expected to make it far simpler for many companies, particularly new manufacturers, to take
up the new LDSE technology in large scale production.
Two other popular implementations of the laser doping technology with industry collaborators are the
bifacial structure using laser doped self-aligned contacts of opposite polarity on both surfaces and the
interdigitated rear surface laser doped contacts for rear junction n-type devices. Collaborative research
projects based on these cell designs have been established with the aim of developing the technologies
to take cell efficiencies on n-type CZ also to above 20% in large scale commercial production at some
stage in the future.
Another use of the laser doping technology was described in Section 4.3.2.3.4.2 as a replacement for
screen-printed contacts with n-type CZ wafers in conjunction with aluminium alloyed rear contact and
junction. As reported above, excellent efficiencies in the vicinity of 19% have also been achieved
with this cell design on 148.6cm2 n-type CZ wafers. A simpler patented version of this technology has
also been developed which requires no diffusion processes, no edge junction isolation and no thermal
processes above 450 degrees Celsius except for the aluminium alloying process for several seconds.
Never-the-less, the technology is still able to achieve efficiencies above 18% on full-sized commercial
substrates. The processing sequence used is:
1.
2.
3.
4.
5,
wafer texturing
silicon nitride deposition
aluminium rear surface printing and firing
laser doping front surface
Ni/Cu light induced plating
4.3.2.3.7
4.3.2.3.7.1
Inkjet Technology for Solar Cell Fabrication
Resist Based Method
Inkjet technology has been an area of rapid development over the last decade, particularly for
printing. In recent years, its application has been spreading to other fields, but as yet has had only
minimal impact in photovoltaics. The first company apparently to commercialise a photovoltaic
technology incorporating inkjet technology was CSG Solar who during 2006 commenced production
of a thin-film technology that uses inkjet printing of a corrosive material to etch patterns in a resist
layer to facilitate metal contacting to the underlying silicon. In the present work, the use of inkjet
technology has been expanded to encompass a range of solar cell fabrication processes including
texturing, grooving, patterning of dielectric layers for metal contacting, localized diffusions, etc. The
techniques developed to carry out these processes are uniquely different to those used before. A noncorrosive plasticizer is inkjet printed onto a low cost resist layer, altering the chemical properties of
the resist layer in these localized regions to make them permeable to etchants such as hydrofluoric
acid (HF). This facilitates the patterning or etching of underlying dielectrics or semiconductor
material to facilitate a range of semiconductor processes. Importantly, the change in resist
permeability is a reversible process making it feasible to return the resist to its original state after
carrying out processes on the underlying material. This also opens the option to partially reverse the
permeability to reduce the hole or feature size produced in the underlying material.
- 81 -
A particular exciting application of this inkjet technology work is for very high efficiency silicon
solar cells. The University of New South Wales has held the world record for silicon solar cell
efficiencies for the last 15 years, initially with the Passivated Emitter solar cell (PESC) and more
recently with the Passivated Emitter and Rear Locally diffused (PERL) solar cell. Despite the
performance and achievements of these two technologies, neither has been used commercially, apart
from for space cells, primarily due to the sophistication, cost and complexity of the processes
involved. The photolithographic based processing is probably the main contributor to this. In this
work, inkjet printing techniques have been developed for patterning low cost resist layers as a simple,
much cheaper alternative to photolithographic based processing. These new approaches appear
capable of achieving similar device performance levels but with the greatest challenge being to match
the dimensions of features in the resist patterning achievable with photolithography. Test devices to
date based on inkjet technology have achieved feature dimensions such as holes of 30-40 microns
diameter as shown in the matrix of holes below. However, these dimensions need to be reduced to
about 10 microns diameter to fully match the performance levels demonstrated with photolithographic
based processing. New and innovative inkjet printing techniques have been recently developed for
further reducing the resist patterning dimensions. This work is being greatly assisted by the recent
availability of the new 1 picolitre inkjet heads.
Figure 4.3.2.14: 30-40 micron diameter holes formed by ink-jetting.
The described dielectric patterning capabilities via inkjet patterning can not only be used for defining
localized diffusion regions and locations for metal contacts, but also for the formation if textured
surfaces and light trapping schemes. Examples of the latter are shown below as well as corresponding
reflectance curves.
- 82 -
Figure 4.3.2.15: (a) Optical Microscope photograph of ink jet patterned holes in a SiO2 layer and (b)
Scanning Electron Microscope photograph of the inverted pyramids formed in the silicon surface
following KOH etching of (a).
Figure 4.3.2.16: (a) Optical Microscope photograph of ink jet patterned lines in a SiO2 layer and (b)
Scanning Electron Microscope photograph of the grooves formed in the silicon surface following
KOH etching of (a).
Figure 4.3.2.17: Reflectance of silicon wafers with different surface finish (i) inverted pyramids
formed by photolithography, (ii) inverted pyramids patterned by ink jet printing and (iii) v-grooves
patterned by ink jet printing.
- 83 -
80
25
70
20
R e fle cta n ce (% )
60
50
15
10
400
550
700
850
1000
40
Random upright pyramids
30
I.J. inverted pyramids
I.J. v-grooves
20
10
0
300
450
600
750
900
1050
1200
Wavelength (nm)
Figure 4.3.2.18: (a) Reflectance of silicon wafers with different surface finish (i) laboratory
fabricated random, upright pyramids, (ii) inverted pyramids patterned by ink jet printing and (iii) vgrooves patterned by ink jet printing and (b) SEM of (iii).
Significant interest is also being shown in developing high efficiency approaches for multicrystalline
silicon wafers. Record performance multicrystalline silicon cells have in the past benefited from
photolithographically defined “honeycomb” texturing using an acidic isotropic etch. Such structures
have also been demonstrated recently using inkjet technology as shown in Figure 4.3.2.19. Such
techniques appear to have significant commercial appeal.
(a)
(b)
Figure 4.3.2.19: Honeycomb texturing (a) and macrogroove texturing (b) of multicrystalline silicon
wafers through the use of inkjet technology.
4.3.2.3.7.2
Direct Etching Via Aerosol Jetting
The key to high-efficiency silicon solar cells is the ability to form small-area metal contacts to the
silicon through dielectric passivating layers. In this work, a new method for patterned etching of SiO2,
SiOxNy and SiNx dielectric layers has been developed and patented. The method uses Optomec’s
M3D aerosol jetting device to deposit a solution containing fluoride ions, according to an etching
pattern, onto an acidic water-soluble polymer layer formed over the dielectric layer as indicated in the
schematic of Figure 4.3.2.20. The deposited solution reacts with the polymer layer, at the locations
where it is deposited, to form an etchant that etches the SiO2 and SiNx under the polymer layer to
form a pattern of openings in the dielectric layer. After the pattern of openings is formed, the acidic
water-soluble polymer and the etch residue are easily removed by rinsing in water. The method
- 84 -
involves fewer steps than photolithography and is safer than existing immersion etching techniques in
that the corrosive etchant HF is only formed in-situ on the surface to be etched. Furthermore, the
method uses small amounts of inexpensive chemicals and produces significantly less hazardous
fluoride waste than existing immersion etching methods.
Figure 4.3.2.20: Schematic showing the principle of “Direct Etching” for patterning dielectric layers.
Aerosol jet printing is a new deposition technology being pioneered by Optomec, Inc. The technique
enables the finely-controlled deposition of an aerosol, which is generated from a liquid, by using a
sheath gas to constrict the aerosol into a fine jet which is directed to the substrate. The technique has
been previously used in applications such as printed electronics, fuel cells and displays.
The aerosol etching method has been used to etch groove structures in SiO2 which have been
thermally grown on polished silicon wafers. By varying the aerosol and sheath gas flow rates the
geometry and depth of etching can be varied. For example, grooves (see Figure 4.3.2.21) were cleanly
etched in a 260 nm thick SiO2 layer using aerosol and sheath gas flow rates of 21 and 50 cm3/min
respectively and depositing 10 layers of aerosolized 10% NH4F solution. These grooves were ~35
microns wide at the surface but only ~ 20 microns wide at the Si/SiO2 interface.
(a)
(b)
Figure 4.3.2.21: (a) Optical image, and (b) Dektak profile of a groove etched in a 260 nm thick
thermal oxide grown on a polished Si wafer by deposition of 10 layers of aerosolised 10% (w/v)
NH4F solution onto a polyacrylic acid polymer layer of thickness ~ 2.2 microns. The aerosol and
sheath gas flow rates were 21 and 50 cm3/min, respectively. The process velocity of the stage was 10
mm/s and the platen was heated to 45ºC. The groove width at surface ~35 microns.
- 85 -
Even narrower structures can be etched by reducing the aerosol flow rate and increasing the sheath
gas flow rate. Figure 4.3.2.22 shows the formation of central, more deeply, etched grooves which are
200 nm deep and ~10 microns at the Si/SiO2 interface. It is anticipated that the use of an even lower
flow rate for the first, wetting layer will largely eliminate the wider, shallow etched region at the
surface thus enabling very narrow grooves to be etched to a depth of at least 200 nm.
(a)
(b)
Figure 4.3.2.22: (a) Optical image of grooves etched in a 260 nm thick thermal oxide grown on a
polished Si wafer by deposition of 25 layers of aerosolised 10% (w/v) NH4F solution onto a
polyacrylic acid polymer layer of thickness ~ 2.2 m. The aerosol and sheath gas flow rates were 18
and 55 cm3/min, respectively. The process velocity of the stage was 10 mm/s and the platen was
heated to 45ºC. (b) Dektak profile of a section of the grooves shown in (a).
We have also been able to etch grooves 20-25 microns wide in 75 nm thick SiNx layers on
chemically-textured silicon surfaces. The etched grooves can be used to form front or rear metal
contacts for silicon solar cells using metal plating, screen printing or evaporation deposition methods.
Compared to current inkjet implementations of this direct etching method, aerosol jet etching results
in faster etching (fewer layers required to be deposited), smaller etched features and less variation in
etched groove width over large etching patterns. Although the current etched feature sizes are
sufficiently small for many current front-contact and rear-contact silicon solar cells, it is anticipated
that the feature sizes can be further reduced in the near future with refinement of the jetting
parameters and modification of the surface polymer layer composition. Replacing the thermal oxide
or SiNx layer with SiOxNy also has beneficial consequences for the dimensions of the feature sizes
formed such as shown by the etched lines of 15-20 microns width shown in Figure 4.3.2.23. These
further reductions in etched feature size may enable high-efficiency silicon solar cell designs, such as
the PERL cell, to be commercially realised.
Figure 4.3.2.23: Aerosol jet-etching of a silicon oxynitride layer to form lines of only 15-20 microns.
- 86 -
This area of work has attracted significant interest and funding from industry, including being
awarded a 3-year ARC Linkage grant for the design and fabrication of high efficiency solar cells
through the use of the described patented inkjet technology from UNSW.
4.3.2.3.8
PLUTO Technology
The Pluto technology is a low cost implementation of the UNSW PERL (Passivated Emitter and Rear
Locally diffused) solar cell developed at UNSW. The PERL technology currently holds the worldrecord for silicon solar cell performance with 25% efficiency being recorded as the new world record
for silicon solar cell efficiency during 2008. Through collaborative research with Suntech, two
generations of Pluto technology have been developed. The first uses the standard front surface design
of the PERL cell but with a screen-printed rear aluminium alloyed rear metal contact that simplifies
the technology and allows it to be easily retrofitted onto existing screen-printed solar cell lines. In
comparison to standard screen-printed solar cells made on the same production line, Pluto cells
achieve an increase in performance of more than 10% by taking average efficiencies from about 18%
to 19.6%. These efficiencies have been independently confirmed by the Fraunhofer Institute in
Germany. The main contributors to these increased efficiencies are: reduced shading loss (2-3%);
improved short wavelength response due to selective emitter (2-3%); improved Voc due to lightly
diffused emitter and well passivated surfaces (2-3%); and improved fill-factor due to reduced resistive
losses (2%).
In comparison to the semiconductor finger technology, many of the losses associated with the screenprinted metal contacts (shading, contact resistance, metal resistance, dark saturation current from the
metal/silicon interface, etc.) have been minimised through the Pluto technology, facilitating an
efficiency increase by about 1% in absolute terms. The spectral response and corresponding
performance spread in production are shown in Figures 4.3.2.24 and 4.3.2.25 respectively.
Figure 4.3.2.24: Spectral response curve of a 19% efficient Pluto solar cell.
- 87 -
Figure 4.3.2.25: Solar cell efficiencies of commercially manufactured Pluto solar cells (a) in
chronological order and (b) as a histogram.
In comparison to the laser doping technology, the main advantage is the elimination of most of the
defects and associated recombination that exists in the resolidified laser doped regions or in locations
immediately adjacent to these. Consequently, Pluto cells achieve reduced junction recombination and
therefore higher fill factors and Vocs leading to a performance increase of about 5%.
This technology is unique to Suntech with some of the intellectual property being owned solely by
Suntech. Intellectual property arising from the collaborative research is jointly owned by UNSW and
Suntech, but with the latter having the right to use it in its own production. Suntech has publicly
announced that it has increased the manufacturing capacity of the Pluto technology to 0.5GW during
2011. This was planned to be further scaled during 2012 with the full 2GW production capacity to be
retrofitted to the Pluto technology during 2013. However the severe down-turn experienced by the
industry during 2012 made it near impossible for any companies to expand the production of any
technologies. Figure 4.3.2.26 shows a photo of a Pluto cell juxtaposed to a standard screen-printed
solar cell.
Figure 4.3.2.26: Pluto production cell (right) with significantly reduced reflection losses compared to
standard screen-printed solar cells (left).
The second generation of Pluto technology uses the full PERL cell design by eliminating the rear
screen-printed aluminium layer and replacing it with a localised metal contact. An independently
confirmed world-record efficiency for commercial grade p-type CZ wafers of 20.3% was achieved
during 2012. Further already identified mechanisms for reducing losses is expected to take these
efficiencies to the 21-23% range over the next 1-2 years. It does however avoid the use of vacuum
- 88 -
deposition processes and photolithographic techniques, thereby avoiding the expensive fabrication
costs associated with most laboratory technologies. This technology is expected to follow the first
generation of Pluto technology into large scale production about 1-2 years later.
The PLUTO technology has been also successfully applied to multicrystalline silicon wafers and
implemented into large scale production with efficiencies of 17-18% also independently confirmed by
the Fraunhofer Institute. In mid 2009, Suntech beat the 15.5% multicrystalline silicon module worldrecord efficiency held by Sandia National Laboratories in the US for 15 years. This was followed later
in the year by Suntech’s achievement of another world-record of 16.5% (aperture area) for a multi
module, this time beating the recently achieved ECN world record in the Netherlands. During 2010,
this module efficiency was further increased to 17% and again confirmed by the Fraunhofer Institute.
Further improvements during 2012 focussed on improving the rear surface passivation with
efficiencies in the range of 19-20% expected in the near future.
4.3.2.3.9
Photoluminescent Imaging for Device and Material Characterisation
With the current growth in the photovoltaic industry and the trends towards higher cell efficiencies,
generally achieved on lower quality and thinner silicon wafers, there is an increasing demand in
research laboratories and in industrial manufacturing for fast and easy to use characterisation tools.
Recent research at UNSW has established photoluminescence (PL) techniques as extremely fast and
useful tools for the characterisation of silicon wafers and of silicon solar cells.
In particular, this work has shown that luminescence imaging techniques give two dimensional high
resolution images of electrical parameters such as localised series resistance values and the minority
carrier lifetime shown below for a multicrystalline silicon solar cell. The data acquisition times for
such measurements are of the order of typically only one second per wafer, which is orders of
magnitude faster than any competing experimental techniques. High resolution images of other
electrical parameters such as the shunt resistance are also feasible.
Figure 4.3.2.27: Luminescence image of a completed multicrystalline silicon solar cell.
A prototype bench-top system established at the Centre three years ago was in such heavy demand by
various UNSW research groups and for consultancy projects that a new state-of-the-art system was
purchased recently from BTI. It has helped identify a variety of unexpected processing problems and
in developing new processing sequences much more quickly than previously possible. In many cases
the two dimensional information contained in PL images has given easy to interpret clues about the
origin of specific problems. The ability to measure a large number of samples in a short time is very
beneficial in this context.
- 89 -
Collaboration with various industrial partners has also identified various processing problems
previously unknown to the manufacturers. In addition, recent research at UNSW has shown how the
luminescence imaging technique could be used for in-line process control, for example to remove the
influence of shunts. Other potential industrially relevant in-line applications for luminescence imaging
that are under investigation include crack detection, spatially resolved series resistance monitoring,
quality control of raw material and process control of individual key processing steps such as the
emitter diffusion. The collaborative work with various industry partners and several journal and
conference publications have resulted in significant interest worldwide in the technique, specifically
in a commercial PL imaging system.
A start-up company BT Imaging Pty Ltd has been formed by Centre researchers Robert Bardos and
Thorsten Trupke in 2007 with commercial sales of the system taking place during 2008. The company
has developed commercial off-line and in-line luminescence imaging tools, with healthy sales to
research groups and cell manufacturers throughout the world during 2012.
In summary, photoluminescence techniques have proven particularly useful for the characterisation of
various aspects of the screen printed solar cell process. Photoluminescence imaging was applied to the
study of belt furnace processing, while photoluminescence imaging with current extraction was
applied to the characterisation of n-type solar cells with printed metallisation to identify failure
mechanisms in the metallisation. The accuracy, high speed and non-destructive nature of PL
techniques make it an attractive candidate for use in high-throughput screen printed solar cell
production lines.
4.3.2.4
Pilot Production
Many of the above technologies have been implemented or are in the process of being implemented
into pilot production. The following table summarises the performance of each of the technologies in
pilot production using standard commercial wafers with commercial processes and equipment.
ARC
Jsc
(mA/cm2)
Voc
(mV)
FF
(%)
Effic
%
Area
(cm2)
Wafer
type
1 BCSC (Tianwei)
SiN
37.8
630
78.5
18.7
148.6
(p-type CZ)
1 BCSC (Tianwei)
SiN
35.2
620
78.4
17.1
156
(p-type multi)
2 IBBC
SiN
37.1
675
79.4
19.8
46
(n-type CZ)
3 Screen-print
SiN
37.1
632
78.8
18.5
148.6
(p-type CZ)
3 Screen-print
SiN
33.3
614
75.5
16.0
156
4 SCF
SiN
37.6
641
78.5
18.9
148.6
(p-type CZ)
5 LDSE (Sunrise)
SiN
37.8
638
79.0
19.0
225
(p-type CZ)
5 LDSE (UNSW)
SiN
38.6
642
79.4
19.7
148.6
(p-type CZ)
5 LDSE (Centrotherm)
SiN
38.4
638
78.9
19.4
148.6
(p-type CZ)
5 LDSE (Roth & Rau)
SiN
37.8
635
78.4
18.8
225
(p-type CZ)
5. LDSE (Shinsung)
SiN
38.6
640
79.3
19.6
148.6
(p-type CZ)
5. DS-LDSE (UNSW)
SiN
40.1
680
75.4
20.3
148.6
(p-type CZ)
5 LDSE
SiN
35.4
628
77.9
17.3
156
5 DS-LDSE (a-Si rear)
SiN
40.3
670
75.8
20.3
148.6
(n-type CZ)
6 Pluto
SiN
38.4
638
79.5
19.5
148.6
(p-type CZ)
6 Pluto 2
SiN
40.9
670
74.2
20.3
148.6
(p-type CZ)
6 Pluto
SiN
35.7
633
78.6
17.7
156
(p-type multi)
7 PERL
ZnS/MgF2
41.7
704
82.5
24.2
46
(p-type FZ)
Technology
- 90 -
(p-type multi)
(p-type multi)
4.3.2.5
Core Research
The distinguishing feature of core research is that it is funded by UNSW, the ASI and the Centre of
Excellence or other Government funding schemes. The priorities for the work are therefore set by the
Centre rather than by industry. Any IP is therefore wholly owned by UNSW, but with any such
developments made available at no cost to industry partners of the Centre such as those funding
collaborative research programs or else licensing Centre technology. Such core research is therefore
of great importance in the initial stages of development of a new technology to take the development
to the point where industry interest is gained and industry is willing to fund the ongoing development.
An example of this is the Advanced Semiconductor Finger Solar Cell described in Section 4.3.2.3.3
where core research has facilitated the development and demonstration of the concept, to the point
where test structures can show that efficiencies above 19% are achievable, making these the most
efficient screen-printed cells world-wide on standard commercial grade p-type CZ wafers. This work
is now to continue as an industry funded project in conjunction with the ASI, with the industry partner
being Suntech.
Other important areas of the core research are where problems are tackled and the solutions can be
used for the benefit of all or most industry partners. Examples include improvements in the light
induced plating techniques, reducing the laser induced defects during the laser doping process,
reducing feature sizes during inkjet patterning of dielectric layers, improving surface passivation
techniques, adapting processes to suit low cost silicon etc. The Centre has a large number of very
talented PhD students contributing to this core research, the achievements from which are
significantly accelerating the progress for most industry collaborative research projects. This
represents a major benefit to companies that engage in funding photovoltaic research at UNSW.
Some core research however is conducted in conjunction with companies, but where the companies
are not funding the work at UNSW and have no claim over any resulting IP from the work at UNSW
and cannot exercise any restriction regarding its use. An example of this is the collaboration with
Centrotherm to evaluate the compatibility between the LDSE technology and the standard
Centrotherm screen-printing technology. Cells from a standard Centrotherm screen-printing line for
producing homogeneous emitter solar cells of 18% efficiency were fabricated without front silver
metal. The emitter sheet resistivity was the only processing variation, being raised to 120 ohms per
square. After laser doping and plating at UNSW, the finished 5 inch cells were sent to the Fraunhoffer
Institute in Germany where they were confirmed to be 19.4% efficiency. This is believed to be the
highest ever independently confirmed for industrial solar cells fabricated on standard commercial
grade p-type CZ wafers. This record may however be shortly broken through collaborative work with
Shinsung where similar devices fabricated using similar wafers have been measured to be 19.6%
efficiency.
- 91 -
4.4
Second Generation: Silicon, Organic and other “Earth Abundant”
Thin-Films
4.4.1
Silicon Thin Films
University Staff:
Dr Sergey Varlamov (group leader)
Prof. Martin Green
Research Fellows:
Dr Jing Rao
Dr Jialiang Huang (from 10/2012)
Technical Staff:
Dr Patrick Campbell
Mark Griffin
Anthony Teal (from 10/2012)
Postgraduate Research Students:
Peter Gress (PhD, till 3/2012))
Bonne Eggleston (PhD,till 3/2012)
Hongtao Cui (PhD, till 8/2012))
Wei Li (PhD)
Qian Wang (PhD)
Kyung Hun Kim (Masters)
Anthony Teal (PhD)
Jae Sung Yun (PhD)
Chaowei Xue (PhD)
Miga Jung (PhD)
Jonathon Dore (PhD)
Jongsung Park (Masters)
Chaho Ahn (Master, from 3/2012)
Mohd Zamir Pakhuruddin (PhD, from 9/2012)
4.4.1.1
Summary
The silicon thin-film group conducts research on polycrystalline silicon (poly-Si) thin-film solar cells
on glass. The major focus in 2012 was shifting from solid-phase crystallised (SPC) silicon films to
liquid phase crystallised films. During previous years of research, which was funded by ARC Linkage
Grant with CSG Solar (now Suntech Research and Development, Ausralia, SRDA), it was found that
the performance of SPC poly-Si solar cells is inherently limited by the high defect density. A new
approach to overcoming this limitation has been identified which is liquid-phase crystallisation by
line-focus diode laser. A new project on poly-Si solar cells made of so-called “liquid-phase
crystallised silicon on glass” (LPCSG) commenced in 2012 in collaboration with SRDA and funded
by ASI/ARENA to facilitate development of this promising technology. The record breaking opencircuit voltage (Voc) up to 557 mV have been achieved, the highest Voc reported for poly-Si thin-film
cells on glass. The best short circuit current (Jsc), Fill Factor (FF) and the efficiency are 27.8 mA/cm2,
62.3% and 8.4% respectively, which are also higher than the values for the best UNSW e-beam SPC
poly-Si cells fabricated from e-beam films.
The performance of the best UNSW e-beam SPC poly-Si thin-film solar cells has peaked at the
efficiency of 7.1% (Voc 458 mV, Jsc 26.6 mA/cm2, FF 58%).
- 92 -
Significant progress was achieved in transient heating defect anneal of poly-Si films by a diode laser.
The laser annealed poly-Si cells have demonstrated Voc of 496 mV, even higher than Voc of 467 mV
for reference cells processed by conventional rapid thermal annealing.
Outstanding results were obtained in the area of light-trapping in poly-Si thin-film solar cells. A
simple and effective Si film etch-back texturing process was developed to improve light-trapping in ebeam poly-Si cells on planar glass, which allowed achieving a record current of 26.6 mA/cm2, the
highest ever reported for evaporated poly-Si solar cells. Optimisation of plasmonic poly-Si thin-film
cells led to Jsc enhancement of 53% and the cell efficiency of 5.95%, the highest reported efficiency
for plasmonic thin-film solar cells.
A concept of performance limiting recombination in SPC poly-Si has been developed. According to
the concept it is intragrain defects, such as dislocations that are responsible for limiting the minority
carrier lifetime in poly-Si material and thus the cell performance. In the practical range of dopant
concentrations in the quasi-neutral region of the cell absorber (> ~5E15 cm-3), the lifetime limiting
recombination occurs via shallow levels linked to dislocations. At dopant concentrations lower than
~5E15 cm-3 the lifetime is limited by recombination via deep levels introduced by charged impurities
at dislocation and possibly grain-boundaries. The dislocation density in poly-Si thin-film cells was
estimated from the TEM images and found to be of an order of 1E10 cm-2, which is consistent with
experimental voltages of about 500 mV.
4.4.1.2
Introduction
The technology of poly-Si thin-film solar cells on glass is developing steadily and it is expected to
become cost-competitive with other thin-film technologies. [4.4.1.1, 4.4.1.2]. The presently best
developed poly-Si on glass PV technology has been commercialised by CSG Solar AG where an
approximately 1.5 m thick PECVD amorphous silicon (a-Si) precursor diode is solid-phasecrystallised to form a poly-Si film, which is annealed and hydrogenated to activate dopants, remove
and passivate defects, and then processed into metallised modules [4.4.1.3]. The highest achieved
efficiency for 94 cm2 mini-modules produced by CSG technology is 10.5% [4.4.1.4], and full scale
7-8% efficient modules were manufactured in Thalheim, Germany in 2006-2010.
Research in the thin-film group at UNSW focuses on critical cell fabrication processes and explores a
range of advanced approaches to improve the performance and manufacturability of poly-Si thin-film
cells. The group’s work over the past years has led to innovative solutions to the key steps of the cell
fabrication process, including glass and Si film texturing [4.4.1.5, 4.4.1.6, 4.4.1.7, 4.4.1.8] and
localised surface plasmons (SP) in metal nanoparticles [4.4.1.9, 4.4.1.10, 4.4.1.11] for enhanced light
trapping, silicon deposition by high rate e-beam evaporation [4.4.1.12], defect anneal and passivation
[4.4.1.13, 4.4.1.14], and cell metallisation and interconnection [4.4.1.15]. The major focus in 2012 has
been on developing a new technology of poly-Si thin-film solar cells obtained by laser induced liquidphase crystallisation [4.4.1.16, 4.4.1.17, 4.4.1.18, 4.4.1.19]. The following sections summarise the
thin-film cell research projects and their results.
4.4.1.3
Solar cell fabrication sequence
A schematic representation of a poly-Si thin film solar cell is shown in Figure 4.4.1.1 and Figure
4.4.1.2 describes the cell fabrication sequence.
- 93 -
Figure 4.4.1.1: Structure of a
p-type poly-Si solar cell on
planar glass (layer
thicknesses and grain size not
to scale).
Glass cleaning/texturing
SiN deposition (PECVD, Sputtering)
a-Si precursor diode (PECVD, e-beam evaporation)
Figure 4.4.1.2: Process
sequence of the four types of
poly-Si thin-film on glass
solar cells under development
at UNSW. All cells are
designed for the superstrate
configuration (i.e., the
sunlight enters the cell
through the glass).
Solid Phase Crystallisation (SPC, 600ºC)
Defect Annealing/Dopant Activation (RTA, 1000ºC)
Defect Passivation (hydrogenation, 600ºC)
Cell metallisation
Cell characterisation
The substrate used for all cells is 3.3 mm thick borosilicate glass (Schott Borofloat33). Both planar
and textured glass substrates are investigated: typically e-beam and hybrid cells are fabricated on the
planar glass while PECVD cells are fabricated on textured glass. The texture on the Si facing glass
surface aims to reduce the reflection losses and to enhance the light trapping in a weakly absorbing
poly-Si thin-film. The texture is prepared by the UNSW-developed AIT method [4.4.1.20] described
below and consisting of an irregular rough array of sub-micron sized dimples. The next step is the
deposition of an approximately 80 nm thick SiNx film (refractive index ~2.1 at 633 nm) by PECVD at
about 400°C or by reactive sputtering at 200°C. The SiNx film acts both as an AR coating and a
barrier layer for contaminants from the glass. An a-Si n+/p/p+ precursor diode is then deposited by
either e-beam evaporation or PECVD on planar or AIT glass respectively, followed by SPC at 600°C.
The resulting as-crystallised poly-Si diodes have Voc of only about 150-180 mV. A high temperature
treatment is then performed at ~1000°C for 30-60 s to activate the dopants and to anneal defects,
which increases Voc up to 220-250 mV. The final step of the material preparation is hydrogenation in
a remote hydrogen plasma to passivate the remaining defects resulting in Voc of 450~500 mV typical
for the finished cells.
The cells are metallised using a combination of photolithographic patterning with dielectric and metal
deposition processes to create an interdigitated electrode structure on the rear (Si side) of the cells.
Typically, the BSF electrode of the PECVD cells covers the whole rear cell surface and the cells can
- 94 -
only be illuminated from the glass side [4.4.1.21, 4.4.1.22]. The BSF electrode of the e-beam cells
consists of thin Al fingers covering only 4% of the rear surface, thus allowing bifacial (i.e.
illumination from both glass and Si side) operation of the cells [4.4.1.23, 4.4.1.24, 4.4.1.25].
The hybrid cells are made on planar glass with the SiN layer and the emitter deposited in the SRDA
large area KAI PECVD tool and the absorber and BSF layers in the UNSW e-beam evaporator. The
films are processed into either individual 2 cm2 large solar cells using the UNSW bifacial metalisation
or solar mini-modules of 36 cm2 area using the CSG Solar metallisation technology described
elsewhere [4.4.1.4].
4.4.1.4
Glass Texturing
An advanced glass texturing method was developed at UNSW and it is referred to as “aluminium
induced texture” (AIT) [4.4.1.5, 4.4.1.6, 4.4.1.20]. A thin sacrificial Al film is deposited onto planar
glass, followed by annealing at about ~600°C in an inert atmosphere. The anneal initiates a red-ox
reaction whereby Al is oxidised to Al2O3 and SiO2 from the glass is reduced to silicon, as follows:
4 Al  3 SiO2  3 Si  2 Al2O3
A surface texture is imparted to the glass according to the nucleation conditions provided during this
anneal. Subsequent wet-chemical etching removes the reaction products from the glass surface and
reveals the glass texture (Figure 4.4.1.3a). Further improvement in the AIT process conditions led to
development of so-called “sub-micron AIT” (SAIT) glass with smaller and steeper feathers (Figure
4.4.1.3b), which ensure even better light absorption in a poly-Si film prepared on such texture (Figure
4.4.1.4) [4.4.1.26].
Typical atomic force microscope (AFM) images of the bare AIT and SAIT glass surfaces with the
RMS roughness of 800 and 150 nm respectively and an SEM image of the AIT glass coated with a
2 μm thick poly-Si film are shown in Figure 4.4.1.3.
a
b
- 95 -
Figure 4.4.1.3: AFM images of bare AIT (a) and
SAIT (b) glass and SEM image of 2 μm thick
poly-Si film on AIT glass (c).
c
As shown in Figure 4.4.1.4, the absorption in the 2 μm thick poly-Si films on the best AIT and
particularly SAIT glass in the 500~1000 nm wavelength interval is very close to the calculated
random scattering absorption limit, based on Monte Carlo ray tracing,. It should be noted that the
significantly higher absorption than the theoretical limit at 1000 nm and above can be attributed to the
enhanced parasitic absorption in the glass and/or measurements errors as discussed elsewhere
[4.4.1.21].
100
Planar
AIT
SAIT
Lambertian
90
80
Absorption, %
70
60
Figure 4.4.1.4: Calculated
random scatter absorption
limit(black dots) and
measured optical absorption in
2 μm thick poly-Si thin films
on planar, AIT, and SAIT
glass with air as the BSR.
50
40
30
20
10
0
350
450
550
650 750 850 950 1050 1150
Wavelength, nm
A useful model has been developed based on surface decoupling calculations that allow estimation of
light-trapping potential of a particular textured substrate without depositing a Si film on it [4.4.1.27].
The model first predicts the topography of the Si film surface based on the topography of the
underlying textured glass. It then calculates a correlation coefficient (CC) between two surfaces, a
lower CC corresponding more surface decoupling. It is assumed that the more decoupling (lower CC)
between two Si interfaces (one of which is experimentally characterised glass surface, another is the
predicted Si surface), the more efficient the light-trapping. A satisfactory agreement is found between
CC of Si films and their potential Jsc (an integral of a product of absorption and the AM1.5G solar
spectrum), which confirm the validity of the approach.
4.4.1.5
PECVD Cells
The poly-Si thin-film cells fabricated by PECVD are the most technologically advanced and have
achieved the highest efficiencies compared to other cell types [4.4.1.4].
- 96 -
The cells are made on 3.3 mm thick textured Schott Borofloat glass, have about 2-2.5 μm thickness,
an area of 4.0 cm2, and a structure of glass/SiN/n+p-p+. The standard processing sequence includes
SPC at 600oC for ~ 15 hrs; rapid thermal defect annealing and dopant activation at ~1000ºC for ~ 1
min; remote hydrogen plasma defect passivation at 600ºC for ~30 min; followed by metallisation
using the interdigitated contact scheme shown in Figure 4.4.1.5 and described elsewhere [4.4.1.12].
rear metal
Figure 4.4.1.5: Schematic representation of
the interdigitated metallisation scheme of
PECVD poly-Si thin-film solar cells on glass
(not to scale).
Si film
glass
sunlight
front metal
One of recent advances in improving the PECVD cell performance was the development of an
optimised metallisation scheme with a superior BSR made of a combination of a thin silica layer and
evaporated Al. The rear Al electrode makes a contact to the cell BSF layer through small vias
photolithographically defined in the silica film (Figure 4.4.1.6). When the PECVD cells are fabricated
on the AIT glass and metallised using such a point contacted rear-surface structure they possess
enhanced light-trapping and achieve both high Jsc of 29 mA/cm2 and record efficiency of 9.3%
[4.4.1.28].
Figure 4.4.1.6: Schematic
cross section of a poly-Si
thin-film solar cell with a
point-contacted rear surface
(not to scale).
The measured I-V curve of the record cell together with the other important cell parameters is
illustrated in Figure 4.4.1.7. The measurements were apertured to 4.0 cm2 area and a NREL calibrated
cell was used as a reference.
- 97 -
Figure 4.4.1.7 Measured I-V curve
of a 9.3% efficient PECVD poly-Si
solar cell made on AIT textured
glass (silicon thickness 2.3 µm, cell
area 4.0 cm2)
140
C u rren t [m A ]
120
100
80
Voc 510mV
Jsc 29.0mA/cm2
FF 62.6%
Eff 9.3%
60
40
20
0
-200
-100
0
100
200
300
400
500
600
Voltage [mV]
4.4.1.6
E-Beam Evaporated Si Cells
E-Beam cell structure, fabrication, and performance
Deposition of Si using an e-beam evaporator can be performed at a very high rate of up to 1 m/min.
Other advantages of the method include absence of toxic gases, good Si source material usage and
compatibility with a continuous in-line deposition mode. Thus, if performed in a non-UHV
environment (base pressure > 1×10-8 Torr, pressure during Si evaporation > 1×10-7 Torr), e-beam
evaporation is a cost effective Si deposition method for the thin-film PV applications (Figure 4.4.1.8).
Figure 4.4.1.8: Schematic the of e-beam Si
evaporation tool with boron and phosphorus
effusion cells for in-situ doping.
The thin-film group at UNSW has been developing e-beam Si thin-film cells since 2004, and
significant progress has been achieved in advancing the cell technology from a few square millimetre
area mesa devices at the start to recent fully functional 2 cm2 area 7.1% efficient cells [4.4.1.7].
Figure 4.4.1.21 illustrates the progress with the efficiencies of the e-beam cells and Table 4.4.1.1
summarises design features of typical e-beam cells fabricated at UNSW.
- 98 -
Figure 4.4.1.9: Evolution of the
PECVD and e-beam cell
efficiencies at UNSW.
Table 4.4.1.1: Typical design parameters of the UNSW E-beam Si cells.
Parameter
Details
Glass
3.3 mm planar Borofloat33
AR coating
SiN (~75 nm, n ~2.1)
Emitter
n+ (~100 nm, P, 5e19~1e20 cm-3, 400~500 /sq)
Base
p (~2 µm, B, 1e16~1e17 cm-3)
BSF
p+ (~150 nm, B, up to 1~5x1019 cm-3, ~1000 Ω/sq)
RTA
10~20 s at 1000°C or ~4 min at 900°C
Hydrogenation
15~20 min at ~600°C, remote plasma
Metal
0.5-1.5 μm thick interdigitated Al fingers on front & rear
Back reflector
Diffuse while paint
During previous years the focus of the e-beam cell research was on improving the poly-Si material
quality and optimising the cell structure and post deposition treatments (RTA, hydrogenation) to
achieve better voltages. Respectable Voc in the range up to 500 mV were measured by the Suns-Voc
technique on the non-metallised cells. More recently the focus shifted to developing a working
metallisation scheme, improving the cell currents and fill factors (FF), i.e. producing functional ebeam cells with appreciable energy conversion efficiencies.
E-Beam cell metallisation
The e-beam cell metallisation scheme, illustrated in Figure 4.4.1.10, was developed in 2008 and allow
bifacial illumination and versatile back surface reflector (BSR) study. It features interdigitated Al line
contacts for both heavily doped layers, the emitter and BSF. The BSF electrodes only cover about 4%
of the area. The emitter electrode is formed in a way that avoids contact with the absorber layer
entirely. Instead of plasma etching the grooves through the whole Si thickness, a timed etch stop is
used such that a thinned emitter layer still remains at the bottom of the grooves. Then, the glass-side
comb-like electrode is centred in the grooves leaving significant space (>10 μm) between the edges of
the electrodes and the groove sidewalls. After metallisation the entire cell rear surface is then coated
with a layer of a white paint as a diffuse BSR. The typical cell performance parameters obtained with
the described metallisation are listed in Table 4.4.1.2.
- 99 -
Figure 4.4.1.10: (a) Schematic (left) and actual (right) views of a metallised e-beam cell (white paint
BSR not shown).
Table 4.4.1.2: Performance parameters of 1.8 µm thick planar e-beam cells with aligned bifacial
metallisation ([4.4.1.14, 4.4.1.15]). The shunt resistance of the cells is > 2000 Ωcm2 and has a
negligible effect on the performance.
Voc (mV)
479
Jsc (mA/cm2)
20.1
FF (%)
66
Eff (%)
6.3
Rs (Ωcm2)
4
E-Beam cell absorber doping density
The electronic properties of the absorber layer, mainly the effective minority carrier diffusion length
and life-time, are expected to depend strongly on the layer doping level, which thus can have a large
effecperform optimisation of absorber doping. All cells in the experiment were intended to be
nominally identical except boron concentration in the absorber layer, which was determined using the
Z-analysis technique [4.4.1.29]. This technique yields the electrically active dopant density, which is
most relevant to the cell performance. Cells characterisation results are shown in Table 4.4.1.3. To
account for the spectral mismatch between the light I-V tester and the AM1.5G solar test spectrum,
which causes measurement artifacts, a conservative approach was used where Jsc during the J-V
measurements was adjusted to match those determined from EQE.
Table 4.4.1.3: Performance parameters of e-beam cells with different absorber doping densities.
Doping (cm-3) Voc (mV) FF (%) Jsc (mA/cm2) Rs (Ωcm2) Eff (%) pEff (%)
1.7×1015
413
64.1
19.2
2.0
5.09
5.5
15
5.7×10
418
61.9
17.6
3.4
4.67
5.3
1.0×1017
449
62.5
6.5
7.9
2.16
2.5
8.3×1017
403
54.4
4.0
19
1.04
1.3
The major observed effect of the doping density is a sharp reduction in Isc with increased doping. It
dominates a weaker effect on the cell Voc, which rises slightly up to dopant density of about
1e17/cm3 before falling sharply at higher densities. The best cell efficiencies are obtained at the
lowest active dopant concentrations in the range between 1e15/cm3 and 5e15/cm3. Figure 4.4.1.11
shows the effect of the doping densities on the key cell parameters.
- 100 -
Figure 4.4.1.11: E-Beam cell
performance parameters functions of
the absorber doping. Jsc is calculated
from EQE results. Jsc(EQE) and pEff
were obtained under two different
back
surface
reflector
(BSR)
schemes: air (black closed squares)
and white paint (red open squares).
Role of back surface filed
The BSF is very important for the poly-Si solar cell. Besides transporting the holes to the metallic
contacts, it also repels the electrons from the rear cell surface thus reducing recombination at this
interface. In fact, this effect is crucial for the poly-Si thin-film device as shown in Figure 4.4.1.12.
Without the BSF the Voc of the cell is reduced from 435 mV to only 270 mV. Increasing the BSF
thickness from 30 to 90 nm decreases the sheet resistance by a factor 3 (as expected) but does not
impact the Voc of the device. Increasing the deposition rate of the BSF has a positive effect on the
Voc. The Voc increases by 40 mV when the deposition rate increases from 5 A/s to 10 A/s.
- 101 -
Figure 4.4.1.12: Voc of 2 µm
thick evaporated poly-Si solar
cells with various BSF layers
thickness and deposition rate.
2 m thick film
480
BSF deposition rate
5 [A/s]
10 [A/s]
2.5 [A/s]
465
Voc [mV]
450
+40 mV
435
420
+165 mV
405
270
255
0
30
60
BSF thickness [nm]
90
The best efficiency reported previously for the poly-Si cell with the identical metallisation was 5.2 %
with a voltage of 435 mV [4.4.1.23]. The development of the high rate BSF clearly improved the
voltage. The I-V curves data of solar cells with and without a BSR are presented in Figure 4.4.1.13.
The poly-Si solar cells with a white paint BSR has a conversion efficiency of 6.3% with Voc 479 mV,
FF 66% and Jsc 20.1 mA/cm2. The BSR increases the Jsc by 5.3 mA/cm2 and decreases the FF by 2%
due to the increase of Jsc and (increase of ohmic losses in the metallisation).
Figure 4.4.1.13: I-V curves of 2
µm thick planar poly-Si thin
film solar cells with and without
white paint as back reflector.
20.0
17.5
2
Jsc[mA/cm ]
15.0
12.5
10.0
7.5
5.0
No BSR
482 mV 68% 14.8mA/cm2 4.9%
White paint 479mV 66% 20.1mA/cm2 6.3%
2.5
0.0
0
50 100 150 200 250 300 350 400 450 500
Voc [mV]
- 102 -
Light- trapping in e-beam poly-Si thin-film cells
Poly-Si thin-film solar cells need effective light-trapping to compensate for the moderate absorption.
This is typically achieved by glass substrate texturing for PECVD cells. However, evaporated poly-Si
cells are not compatible with textured glass due to low density material grown on textured substrates
as shown in Figure 4.4.1.14. A cross-sectional FIB images of the evaporated Si films on the textured
glass reveals noticeable morphological differences from the films on the planar glass. On relatively
smoother textures the Si film has apparently vague areas of lower material density, or high defect
density, extending from the inflexions in the texture features. Such defective areas further develop
into microcracks or voids either after the thermal treatment (SPC, RTA) or on rougher textures. When
the glass texture has vertical or overhanging features, extended discontinuities in the evaporated Si
film are found.
Figure 4.4.1.14: FIB cross section of an evaporated Si film on textured glass substrate. (a) is a greyscale image, and (b) is the same image but set to a black and white mode.
Most described defects are believed to be related to the columnar microstructure, which is typical for
evaporated films. Clear experimental evidence for the columnar microstructure and the low density
void network present in the SPC poly-Si films was recently reported [4.4.1.30, 4.4.1.31]. Such a
defective structural morphology is formed due to two inherent characteristics of the evaporation, high
directionality of the arriving atomic flux and the adatom low surface mobility. The deposition only
takes place on the surface areas in direct line-of-sight of the evaporation source resulting in
discontinuities in the areas shaded from the source either by the surface topographical features or by
the previously deposited atoms themselves. The earlier arrived adatoms cast “shadows” for
subsequently arriving adatoms, thus creating a network of parallel high density columns surrounded
by lower density material, as it is schematically shown in Figure 4.4.1.15. Besides, in contrast to
PECVD deposition, the evaporated adatoms have a sticking coefficient close to unity; they do not
have or receive the excessive energy to move in any way from the place of their initial landing to find
a more thermodynamically favorable position, thus leading to morphological and possibly electronic
defects.
- 103 -
Figure 4.4.1.15: Schematic representation of formation
of low density material and voids during evaporation,
where already adsorbed adatoms cast shadows for newly
arriving adatoms thus forming parallel channels filled
with no or low density material.
Possible approaches to reducing an extent of the defect formation during deposition by evaporation is
using smoother glass texturing and/or to texture the Si film surface instead of, or in addition to smooth
glass textures or a completely different approach, surface plasmon (SP) enhanced light-trapping
(described in section 4.4.1…). The respective cell structures are shown in Figure 4.4.1.16.
Figure 4.4.1.16: Schematic representation different light trapping scheme. A) planar, B) textured
glass, C) back texture, D) plasmonic nanoparticles
The silicon etch-back texture can be formed by etching a poly-Si film in a solution of KOH, or
H2O2/NH4F, or NH4F only [4.4.1.7, 4.4.1.8]. Examples of the etch-back textured Si film surface and
its roughness are shown in Figure 4.4.1.17. The Si surface roughness depends on the etching solution,
the process temperature and time and it varies from a few tens to a few hundred nanometres.
- 104 -
Figure 4.4.1.17: AFM and SEM images of etch-back textured poly Si film surface and its roughness
as a function of removed Si thickness.
RMS surface roughness first increases with the etched Si thickness, then stabilises at a certain level
depending on the etchant. A larger RMS surface roughness corresponds to steeper surface feature
angles (Figure 4.4.1.18). For 30% KOH etching, the RMS roughness of 120 nm develops after
removing about 6 µm of sacrificial Si, the maximum of the texture angle distribution is 7-9° while the
angle distribution tail extends up to 40°. For NH4F:H2O2 etching, the RMS roughness exceeds 250 nm
after removing 6 µm of Si, the maximum of the texture angle distribution is 12-15°, and the
distribution tail is up to 45°. The most effective etchant is a NH4F-only solution, for which the RMS
roughness exceeds 250 nm after etching off only 2 µm of Si, the angle distribution maximum is 2022°, and the largest texture angles reach 60°. All studied etch-back textures are compatible with the
metallised poly-Si thin-film cell structure and operation and very effective at enhancing light-trapping
and light absorption. After implementing the KOH etch-back texture into a 3.6 µm thick poly-Si thinfilm cell the record of 26.6 mA/cm2 was demonstrated, 21% higher than the Jsc for the reference cell
with the white paint diffuse reflector (Figure 4.4.2.19) [4.4.1.7]. Even higher Jsc enhancement of 27%
was achieved for a rougher NH4F etch-back texture but the Jsc itself was slightly lower due to shorter
diffusion length in the cell absorber.
Percen in bin
18%
16%
14%
12%
10%
8%
6%
4%
2%
0%
Figure 4.4.1.18: Texture
angle distributions and
RMS roughness for etchback Si films after
chemical etching
KOH, rms 95 nm
NH4F:H2O2, rms 220 nm
NH4F, rms 256 nm
0
10
20
30
40
Angle, degree
50
- 105 -
60
70
Figure 4.4.1.19: EQE curve
of the 2.5% KOH etch-back
textured poly-Si cell (red) as
compared to the reference
cell (black).
90
80
70
EQE, %
60
50
40
30
20
Planar, 21.9 mA/cm2
10
2.5% KOH textured 26.6 mA/cm2
0
350
450
550
650 750 850
Wavelength, nm
950
1050 1150
4.4.1.7 Plasmonic Poly-Si Thin-Film Solar Cells
The plasmonic light-trapping relies on plasmonic scattering and a local field enhancement effect
introduced by Ag nanoparticles (NP) as shown in Figure 4.4.1.20. Such NP can be relatively simply
formed by depositing a 5-30 nm thin Ag precursor film on already metallised and characterised polySi cells followed by annealing at 200-230ºC leading to formation of 50-200 nm wide irregularly
shaped NP. The AFM profile and SEM image of the Ag NP on a poly-Si film are shown in Figure
4.4.1.20.
- 106 -
Figure 4.4.1.20: Light-scattering by surface
plasmons (top-left); SEM image of Ag
nanoparticles formed from 16 nm thin
precursor film (top-right); AFM micrograph of
Si cell surface coated with Ag nanoparticles
(left).
Jsc enhancement of 45% due to the plasmonic light-trapping was previously demonstrated using a Ag
nanoparticle array formed from 16 nm thick precursor Ag film directly on a poly-Si cell rear surface
and then overcoated with a 350 nm thick MgF2 cladding layer with a white paint as BSR on the top
(Figure 4.4.1.21) [4.4.1.9].
(a)
(b)
Figure 4.4.1.21: (a) - Schematic structure of plasmonic poly-Si thin-film solar cell; (b) EQE of the
plasmonic cell with 45% Jsc enhancement (green line) compared to the original cell (black dots).
Recent focus in plasmonic cell research has been on optimisation of the plasmonic structure and its
fabrication, i.e. dielectric layer refractive index (RI) and thickness, nanoparticle formation conditions,
and two-layer NP arrays. The best experimental Jsc enhancement has been improved from 45% to
50.2% [4.4.1.32] and the highest achieved plasmonic thin-film solar cell efficiency was 5.32% and
5.95% without and with a back reflector respectively [4.4.1.33].
Optimization of dielectric coating
Nanoparticle absorption and scattering cross-sections strongly depend on the RI of the surrounding
medium. Thus, by optimising the surrounding medium, the light-trapping efficiency by NP can be
enhanced. Systematic study of the effects of the RI and thickness of the surrounding dielectric
medium on the light absorption and the Jsc of poly-Si thin-film solar cells was conducted and the cell
performance enhancement was further improved. This is particular important for randomly distributed
NP arrays which are simple to fabricate and which provide a broad plasmonic resonance peak but the
particles size and shape cannot be precisely controlled. Different dielectrics with varying RI (MgF2 –
1.4; Ta2O5 – 2.2; TiO2 – 2.4) were studied as the NP overcoating layers.
- 107 -
Experimentally, it was found that higher RI dielectric coatings result in a redshift of the main
plasmonic extinction peak and higher modes were excited within the spectral region of interest to
poly-Si thin-film solar cell application. The optical characterisation shows that NP coated with the
highest RI dielectric TiO2 provide highest absorption enhancement of 75.6%. However, from the EQE
measurements (Figure 4.4.1.22) the highest Jsc enhancement of 45.8% was achieved by coating the
NP with the lowest RI MgF2. By further optimising the thickness of MgF2 the highest Jsc
enhancement of 50.2% was achieved for the 210 nm thick layer.
Figure 4.4.1.22: EQE enhancement vs different
dielectric coating MgF2 – 1.4; Ta2O5 – 2.2; TiO2
– 2.4.
Optimization of nanoparticle fabrication
The NP fabrication process parameters (Ag precursor film thickness, annealing temperature and time)
can affect the particle size, shape, and coverage, and therefore, the plasmonic light-trapping in the
cells. Comprehensive optimisation of a NP nanoparticle fabrication process was performed. The
thickness of the precursor film was 10, 14 and 20 nm; annealing temperature was 190, 200, 230 and
260°C; and annealing time was varied between 20 to 95 min. Figure 4.4.1.23 shows SEM images of
NP formed from 10, 14 and 20 nm Ag precursor films. These three samples were annealed at 260°C
and for approximately 25 min. Ag precursor films of different thicknesses (tAg) yield nanoparticles of
different size and coverage, as shown in Figure 3.3.3.1. With increasing tAg, the average particle size
increases: the average particle size of (a), (b) and (c) are 27, 167 and 233 nm while the NP coverage
decreases from 45.7% to 41.5% and then to 37.1% respectively.
(b)
(a)
- 108 -
Figure 4.4.1.23: SEM images of
nanoparticles formed from (a) 10 nm,
(b) 14 nm, and (c) 20 nm thick Ag
precursor films and annealed at 260°C
for ~25 min.
(c)
An important finding was also that the Ag precursor film thickness has the largest effect on NP
properties, the average NP size most of all, and hence on plasmonic Jsc enhancement. Similar NP
arrays with similar light-trapping performance can be produced from the Ag film of a particular
thickness at different combinations of the annealing temperature and time. NP from 10 nm Ag film
are too small and NP from 20 nm Ag film are too large for optimum performance.
Performance enhancement due to light-scattering by nanoparticles was calculated by comparing
absorption, Jsc and energy conversion efficiency in solar cells with and without NP formed under
different process conditions. Nanoparticles formed from 14 nm thick Ag precursor film annealed at
230°C for 53 min result in the highest absorption enhancement in the 700-1100 nm wavelength range,
and in the highest Jsc enhancement of 33.5%. The plasmonic cell efficiency of 5.32% was achieved
without a back reflector and 5.95% with the back reflector, which is the highest reported efficiency
for plasmonic thin-film solar cells (Figure 4.4.1.24) [4.4.1.33].
Figure 4.4.1.24: Light I-V curves
for best performing plasmonic cell
with the highest Jsc and efficiency
enhancement with and without
back reflector.
NP fabrication parameters: 14 nm
Ag film, 230C, 53 min.
Double nanoparticle structure
A motivation for a two-layer NP array (Figure 4.4.1.25) is two-fold. Firstly, not all incident photons
can interact with a single-layer-nanoparticle structure and therefore, adding an extra nanoparticle
layer can increase the interactions between the incident photon and metallic nanoparticle. Secondly, a
second layer of the nanoparticles can result in higher scattering angles and longer light path in the
active layer.
- 109 -
Back diffuse reflector
MgF2(n=1.4)
Silicon
Ag
Silicon
AR coating
Glass
Incoming light
Figure 4.4.1.25: Schematic sketch of a thin- Figure 4.4.1.26: shows the electrical field
film poly-Si th-film cell with a two-layer- distribution at 600 nm wavelength of the simulated
nanoparticle plasmonic structure.
structures of (a) a single-layer-NP structure without
a back reflector and (b) a two-layer-NP structure
without a back reflector; (c) a single-layer-NP
structure with a back reflector and (d) a two-layerNP structure with a back reflector.
Ag films with a thickness of 14 nm were deposited on the rear side of the cell by thermal evaporation
and annealed at 230C for 53mins to form NP. After application of MgF2 coating the NP process was
repeated to form the second NP level. It is found that without a back reflector, the two-layer-NP
structure results in a slightly higher Jsc enhancement. With a back reflector, two-layer-nanoparticle
gives 33% Jsc enhancement, compared with 40% enhancement for a single-layer-nanoparticle
structure. Simulated results (Figure 4.4.1.26) support the experimental finding and give the insights of
the loss mechanism in the two-layer structure. Two-layer-NP structure can provide two–layer
scattering objects but at the same time, it increases the parasitic absorption and also leads to the light
confinement within the dielectric layer. Modifying the RI of the dielectric layer, the NPs size or
coverage may reduce the light confinement within the dielectric layer and thus provide potentially
higher Jsc enhancement.
4.4.1.8 Diode laser treatments of poly-Si films
Millisecond defect annealing
As was mentioned in section 4.4.1.3, the defect annealing in poly-Si films after SPC is a very
important process for improving cell Voc. Finding the best conditions for this process is complicated
due to conflicting trends. On one hand, higher annealing temperatures and longer times are preferred
because of better dopant activation and more complete defect dissolution. On the other hand, the same
conditions cause excessive dopant smearing and glass distortion, which need to be avoided. Typically,
for any given annealing temperature, there is an optimum time balancing favourable processes with
unfavourable ones. For example, at 950ºC such an optimum time is about 3 min, while a similar Voc
effect can be achieved at 1050ºC for only 50 s. Following this trend, it can be estimated that, at
temperatures near the Si melting point (1414ºC), the optimum annealing times can be of order of a
few milliseconds, because of which this type of processes are generally called “millisecond
annealing” [4.4.1.34]. Very importantly, for such short times the thermal diffusion length in glass is of
an order of tens of microns only [4.4.1.35, 4.4.1.36], which means the bulk of the glass stays cold
resulting in less distortion and lower thermal budget for the annealing process.
Two recently available techniques, which allow heating a few micron thick Si film up to its melting
point within a few millisecond long time are diode lasers and flash lamps [4.4.1.34, 4.4.1.36]. With
- 110 -
the diode lasers in particular, the exposure time can be easily controlled in the CW mode from about
1 ms to 100 ms by the laser beam scanning speed, while the degree of Si film heating is controlled by
adjusting the laser power.
Process development for millisecond annealing of poly-Si cell structures have been conducted using a
diode laser (LIMO450-12x0.3) with a line focus beam with the dimensions of 12 mm length and
0.175 mm FWHM (Figure 4.4.1.27).
Figure 4.4.1.27: LIMO line
beam diode laser.
The laser beam laser scans the surface of silicon rapidly heating it up to near the melting point for a
few milliseconds. Typical film temperature profiles during diode laser annealing are shown in Figure
4.4.1.28 for different laser exposures and doses [4.4.1.13, 4.4.1.14]. It is instructive to also plot
relative temperatures as a function of a distance from the beam centre for a range of exposures and
laser powers but below the melting threshold, which are shown in Figure 4.4.1.29. Despite very large
changes in both exposures and powers, the relative temperature profiles are very similar. It suggests
that a poly-Si film reaches an effective steady state at its every point under the laser beam and also
that the shape of the temperature profile is determined by the beam shape, rather than its speed and
power. It means that if a particular process requires a specifically different temperature profile, it can
only be achieved by modifying the beam shape in the short, scanning direction.
Figure 4.4.1.28: Temperature profiles in poly-Si Figure 4.4.1.29: Normalised temperature profiles
film on glass for different scanning speeds and in poly-Si film on glass for different exposures
doses: 0.075 m/s (2.3 ms, 21 J/cm2); 0.100 m/s and doses.
(1.8 ms, 17 J/cm2); 0.133 m/s (1.3 ms, 12 J/cm2;)
It is demonstrated that the laser treatment anneals defects in a SPC poly-Si film and activates dopants.
Both a reduction in the boron-doped (BSF) sheet resistance and an increase in Voc are observed with
- 111 -
increasing laser dose, for a wide range of scanning speeds (Figure 4.4.1.30). The maximum achieved
Voc is 492 mV resulting from 4 ms exposure at 50 J/cm2 laser dose, which is an improvement of 32
mV over the voltage after optimised rapid thermal annealing. The highest substrate temperature
required during the process sequence is reduced from 960C to 620C (the peak temperature during
hydrogenation).
Figure 4.4.1.30: Voc of laser annealed
poly-Si thin-film cells as a function of
laser dose and exposure time. The peak
intensities are normalised.
Diode laser dopant diffusion
A poly-Si thin-film solar cell structure is typically created by introducing dopants during deposition.
However, not always it is most convenient or even feasible. Etch-back textured poly-Si cells require
BSF diffusion after texturing and LPCSG cells (section 4.4.1.10) require emitter diffusion after
crystallisation. Until recently both BSF and emitter diffusion were typically done by RTA. New and
more technologically advanced processes have been developed where dopant diffusion is achieved by
using the line-focus diode laser treatment [4.4.1.14]. Firstly, a spin-on-dopant (SOD) source,
phosphorous or boron, is applied to the poly-Si film surface and baked. Then, it is scanned by the line
laser beam that heats silicon with SOD above 1000C causing relatively shallow, a few tens of
nanometres, dopant diffusion. Laser conditions are optimised for each process to achieve cell
performance, which is at least comparable to the cells with conventional thermal dopant diffusion or
the cells with the as-deposited structure. Both sheet resistances and cell voltages similar to such for
the standard solar cells can be obtained by laser diffusion at doses 40~90 J/cm2. Dopant concentration
profiles after laser diffusion can be estimated by spreading resistance analysis, as shown in Figure
4.4.1.31.
Figure 4.4.1.31: Boron
concentration profiles a in
poly-Si film after laser
diffusion as measured by
spreading resistance analysis.
- 112 -
4.4.1.9
Recombination Processes in Thin-Film Si on Glass Solar Cells
The current output of poly-Si thin-film solar cells is approaching the so-called Lambertian limit,
which serves as a gauge for quantum efficiencies. However, there is little understanding about how
high the Voc could be – the lowest foreseeable practical limit is the Voc of multicrystalline Si wafer
based solar cells. Thus, the voltage of poly-Si cells as an academic topic is still a research frontier
where the fundamental limits are unknown and speculations abound. In previous years many
researchers have attributed the lower lifetimes of poly-Si solar cells to their small grain size, a view
which holds merit if the grain boundaries are not well passivated [4.4.1.37], but in more recent years
it is becoming clear that intragrain defects are also potent enough to limit the lifetime to the
nanosecond range [4.4.1.38]. While the grain boundary recombination can be successfully mitigated
by enlarging the grain size or optimising the passivation of dangling bonds within the grain
boundaries, so far there has not been an effective method to circumvent the detrimental effects of
intragrain defects. The goal of research at UNSW is to elucidate the dominating recombination
pathway in the poly-Si thin-film solar cells so that the attainment of higher Voc becomes a more
tractable problem.
As a result of a comprehensive study of the poly-Si on glass solar cell characteristics (Suns-Voc,
EQE) and film properties over a range of temperatures and absorber dopant concentrations, a concept
consistent with experimental observations, which describes the dominant recombination in poly-Si
has emerged [4.4.1.39~43].
According to the concept, the carrier recombination in poly-Si on glass solar cells can be modelled as
a superposition of two processes. The first one involves the electronic transition between shallow
states which are 0.05-0.07 eV below the conduction band and 0.06-0.09 eV above the valence band,
respectively, which are consistent with the shallow bands at silicon dislocations (Figure 4.4.1.32a).
The second process occurs via deep levels at charged defects. The shallow band recombination
dictates the solar cell properties over the entire practical absorber doping range of 5E15~1E17 cm-3;
the deep level recombination originates from either charged dislocations or grain boundaries and only
becomes influential at dopant concentration lower than ~5E15 cm-3 (Figure 4.4.1.32b).
qΦ
Ec
EF
1
2
Ev
a)
b)
Figure 4.4.1.32: (a) The dominant recombination pathways in poly-Si solar cells involving shallow
defect levels: 1) direct transition between shallow bands; 2)transition between deep level and shallow
band; (b) Simulation of the lifetime at Voc vs dopant concentration. The short dashed line represents
shallow bands recombination, and the long dashed line represents deep level recombination at grain
boundaries.
In the ideal sc-Si the diode saturation current J01 follows the Arrhenius law with the activation energy
Ea equal to the Si bandgap when extrapolated to 0K. The short-circuit current density, Jsc, in sc-Si is
only very weakly dependent on the temperature (due to the effect of band-gap variation). However,
- 113 -
poly-Si solar cells exhibit different characteristics [4.4.1.39]. The Arrhenius plots for J01 of poly-Si
solar cells have distinctively a smaller slope, which reflects a lower Ea associated with J01 by 0.150.18 eV. As a consequence, the Voc plots versus T extrapolated to 0K have a lower intercept in the
range 1.0~1.1 eV as compared to 1.27 eV for sc-Si solar cells (Figure 4.4.1.33a, b). The Arrhenius
plots for Jsc of poly-Si cells whose absorber diffusion length is shorter than the absorber thickness
have steeper slopes indicating stronger temperature dependence of the poly-Si cell current (Figure
4.4.1.33c).
a
c
b
Figure 4.4.1.33: Arrhenius plots for the diode saturation current J01/(JL/I) (a) and for short-circuit
current Jsc (c) for poly-Si and c-Si solar cells; b) Voc plots extrapolated to 0K for poly-Si and c-Si
solar cells.
Although the smaller Ea and the intercept could in principle be due to the narrower band gap in polySi material as compared to c-Si, it cannot explain the temperature sensitivity of Jsc. On the other
hand, both effects are consistent with existence of shallow sub-bandgap states acting as either
minority carrier traps or recombination centres.
Further confirmation of the shallow-level related recombination in poly-Si solar cells comes from the
study of the bulk lifetimes as a function of temperature which allows a more detailed understanding of
the behaviour of these shallow levels in the device quasi-neutral, bulk regions. Minority carrier
diffusion lengths in the absorber layers of various poly-Si thin-film cells were extracted by fitting the
device EQE simultaneously with its optical reflectance, a procedure which has become a routine for
poly-Si thin-film cells whose short lifetimes precludes quantification by the standard techniques such
as the quasi steady state photoconductance (QSSPC). The EQE measurement and fitting procedure
were repeated for each sample at different temperatures ranging from 120K to 320K. Figure 4.4.1.34
superposes the front-side EQE curves of the same cell at different temperatures (note that the heights
of the curves are reduced slightly due to additional reflectance by a quartz window on the cryostat in
which the sample is placed). Clearly the quantum efficiency of the cell varies strongly with
temperature.
- 114 -
Figure 4.4.1.34: Typical sequence of
EQE curves at different temperatures
ranging from 120K to 320K.
70
60
EQE (%)
50
40
Increasing T
30
20
10
0
300 400 500 600 700 800 900 1000 1100 1200
wavelength (nm)
Figures 4.4.1.35 a) and b) plot the extracted lifetime τ against q/kT for p- and n-type cells
respectively. Although the actual lifetime values at a given temperature differ greatly from cell to cell,
all the lifetime curves exhibit Arrhenius law with the activation energy (Ea) of 0.17-0.21 eV near
room temperature [4.4.1.40]. These activation energies are much larger than those reported for
plastically deformed Si and multicrystalline Si [4.4.1.44], where the dominant recombination is
thought to be via transitions between the shallow and deep levels. Assuming direct transitions
between shallow levels, the Ea is expected to be (Ec-Ede)+(Edh-Ev)+2.5kT=180~120 meV (where Ede
and Edh are shallow band energies near the conduction and the valence bands respectively), which is
in remarkable agreement with the experimental values.
a)
PECVD 3e15
PECVD 1e16
PECVD 2.4e16
PECVD 1e16
e-beam 1.56e17
e-beam 7.2e15
e-beam 6.9e16
e-beam 2.5e16
PECVD 2.2e15
PECVD 5.1e15
e-beam 6.4e15
ALICE 2e16
AIC
1.E-09
lifetime (s)
1.E-09
lifetime (s)
b) 1.E-08
1.E-08
1.E-10
1.E-11
AIC
ALICE 4.5e16
AIC
ALICE 5e16
1.E-10
1.E-11
a)
b)
1.E-12
1.E-12
30
40
50
60
70
80
30
q/kT
40
50
60
70
80
q/kT
4.4.1.35: Lifetime vs q/kT plots for a) p-type, and b) n-type poly-Si thin-film cells.
Furthermore, as a consequence of electron transitions between the shallow bands, the recombination
rate becomes proportional to the concentrations of both minority and majority carriers, as in the case
of radiative recombination, in contrast to recombination via deep levels, where the rate is proportional
to the minority carrier concentration only. It leads to minority carrier lifetime inversely proportional
to the majority carrier concentration, or roughly to the dopant concentration [4.4.1.40]. The
- 115 -
experimental confirmation of this effect is shown in Figure 4.4.1.36 which plots the minority carrier
lifetimes in a number of poly-Si solar cells versus the dopant density. There is a clear inverse
lif dopant density.
relationship between the lifetime and
1.E-09
Series1
ALICE (n-type)
AIC
e-beam (n-type)
PECVD (n-type)
e-beam (p type)
PECVD (p-type)
Power (Series1)
1.E-10
Figure 4.4.1.36:
Lifetimes at 232K for
various poly-Si thin-film
cells plotted against
dopant concentration.
-0.9619
y = 462576x
2
R = 0.8161
1.E-11
1.00E+15
1.00E+16
1.00E+17
1.00E+18
dopant concentration (cm-3)
Amongst the n-type cells in Figure 4.4.1.35b, three are made from small grained (~1 μm) poly-Si
(circle or square data points), and three are made from large grained (~5 μm) poly-Si epitaxially
grown on an Al induced crystallised (AIC) seed layer (triangle data points). The similarity in the
temperature behaviour of the lifetime, regardless of the grain size, indicates that large and small grain
materials have a common lifetime limiting mechanism. Also, in Figure 4.4.1.36 one readily sees that
the AIC cells (triangles) can take on the lifetime values either above or below the power law trend
line, indicating that they do not have consistently superior lifetimes just by virtue of their larger grain
size. It was found that the predominant grain orientation plays a greater role in determining the Voc of
AIC cells in fact, of the two rightmost triangle data points representing AIC cells at about 5E16 cm-3
doping in Figure 4.1.1.36, the higher lifetime point originates from (100) preferentially oriented
material and the lower lifetime point comes from (111) preferentially oriented material.
All the above evidence points at intragrain defects, as a likely limit on poly-Si thin-film cell lifetimes.
In particular, dislocations are the most likely candidates because they provide the strain fields
necessary to create split-off states from the band edges that can act as shallow levels [4.4.1.45]. To
estimate dislocation density in poly-Si solar cells a TEM study was conducted using the Weak-BeamDark-Field (DFWB) technique. An example image is shown in Figure 4.4.1.37 and the dislocation
density deduced by the image analysis is in the order of 1E10 cm-2. There is very rough correlation
between the dislocation density found in poly-Si solar cells and the cell Voc. The best cells with Voc
> 500 mV have dislocation density of 6~8x109 cm-2, while for the cell with very poor Voc < 450 mV
the dislocation density is about 1.5x1010 cm-2 [4.4.1.46].
- 116 -
Figure 4.4.1.37: Typical TEM DFWB image of a
poly-Si film on glass.
As the dopant density decreases, the shallow band recombination gradually becomes less and less
significant and eventually, at sufficiently low dopant densities, less than about 5x1015 cm-3 the
recombination involving deep levels starts to dominate (process 2 in Figure 4.4.1.32a). Such deep
levels are associated with defects concentrating in dislocations and grain boundaries. The defects are
usually charged to the same polarity as the majority carriers thus creating potential barriers which
attract the minority carriers. The lifetime in this case depends on the potential barrier heights, which
become greater with lowering dopant density.
The effect of recombination via deep levels at charged defects is to place an upper limit on the
minority carrier lifetime (thus the cell Voc). Figure 4.4.1.38 shows Voc simulation results of poly-Si
solar cells where two recombination processes in the absorber, both via shallow bands and via deep
levels, taking place in parallel. Above 1x1016 cm-3 doping, the Voc is fairly constant and approaches
the 485 mV limit corresponding to the case where only shallow band recombination occurs. Below
1x1016 cm-3 doping, deep level recombination becomes influential to the device performance eroding
the Voc by about 21 mV when the doping reaches 1x1015 cm-3. Thus, in the practically significant
dopant concentration range of 5x1015~5x1016 cm-3, an immediate challenge to consistently raising the
Voc above 500 mV is to reduce the dislocation density to below the 1x109 cm-2 level. If and when this
is achieved the next step to even higher Voc is minimisation of charged defects likely associated with
impurities.
Figure 4.4.1.38: Simulated Voc of a polySi solar cell (2.2 µm thick n-type absorber,
Jsc 27 mA/cm2) in which both shallow
bands recombination and deep level
recombination at grain boundaries (with
electrostatic fluctuations) occur.
- 117 -
4.4.1.10
Liquid-Phase Crystallised Poly-Si Thin Film Solar Cells
The conclusion from the previous section is that poly-Si thin-film solar cells obtained by SPC are
capable of achieving only modest Voc of around 500 mV due to inherently high defect density in the
solid-phase crystallised Si. However, in order to become a commercially competitive technology, the
voltages of poly-Si thin-film cells need to be significantly higher, which requires higher crystal and
electronic quality poly-Si material.
Such material is now made possible with a liquid-phase Si crystallisation process developed at
UNSW. A line-focused diode laser beam is scanned across the surface of the silicon, causing it to
melt. Upon directional re-solidification, continuous lateral growth can be achieved, in which the
crystal growth front is seeded by the preceding crystallised region, forming parallel grains several
mm long and several hundred µm wide in the direction of scanning (Figure 4.4.1.39). Such is the
quality of this silicon, that 10 µm thick Si may be used in solar cells without compromising collection
efficiency, whereas with the SPC process, no more than 2 µm could be used. This thickness
difference, together with the higher quality silicon, increases the potential for higher efficiency solar
cells.
Figure 4.4.1.39:
(a) Diode laser
crystallisation
process; (b)
optical
microscope
image of
resulting lateral
grains.
(a)
(b)
The grains have been found to be consistently of [101] orientation. The most common boundaries
between grains are defect-free, electrically inactive Σ3 twin-boundaries as shown in Figure 4.4.1.40.
- 118 -
Figure 4.4.1.40: (a) High
resolution TEM image of Σ3
twin-boundaries; (b) bond angle
frequency showing peaks at 0°
(bonds within a grain) and 60°
(bonds at a twin boundary); (c)
EBSD images showing twins
about a [101] grain orientation.
(a)
(c)
(b)
The grains contain regions with varying defect density. Most laterally grown grains are nearly defectfree on the TEM scale which translates into the defect density of 1E5 cm-2 or lower (Figure 4.4.1.41a),
while the most defect-rich grains can contain dislocations with the density up to 1E9 cm-2 (Figure.
4.4.1.41b). The Hall mobility varies between 300 and 470 cm2/V*s, for the carrier concentration of
~10E16 cm-3, and resistivity 1~3 Ω*cm. These values are similar to those measured for a reference cSi wafer: 414 cm2/V*s, 1.6E16 cm-3, 0.94 Ω*cm respectively, indicating a high, Si-wafer-like
electronic quality of the laser crystallised silicon film, which is far superior to the quality of a
reference SPC Si film with a similar boron concentration and a mobility of only 50-120 cm2/V*s.
Figure 4.4.1.41: Cross
sectional TEM images of (a)
a defect free region and (b)
a defect rich region.
(a)
(b)
During the silicon solidification, cracks usually form through the silicon thickness parallel to the scan
direction. An example is shown in Figure 4.4.1.42b. Such cracks may lead to mild short-circuit losses
in metallised devices and higher crystal defect density in the region surrounding the crack, as shown
in Figure 4.4.1.42a, and are also an aesthetically undesirable. Optimisation of the laser process
parameters has enabled a significant reduction in crack severity and ongoing work aims to eliminate
the problem entirely.
- 119 -
(a)
(b)
(c)
Figure 4.4.1.42: (a)
High defect density
shown in SEM after
defect etch adjacent to a
silicon crack; light
transmission images
showing (b) silicon
cracks caused by
standard crystallisation
process and (c) nearelimination of cracks
via optimised process.
Laser melting results in a material with uniformly distributed dopants throughout the whole film and a
p-n junction has to be introduced after material fabrication by phosphorus diffusion. Until recently
emitter diffusion has been done by rapid thermal annealing (RTA), which is undesirable in a
manufacturing environment, due to the expensive and time-consuming belt furnace required and
results in significant heating of the glass superstrate, which constrains the choice of superstrate
material. We have now demonstrated that this process can also be performed by using the same linefocused diode laser which is used for crystallisation, resulting in a faster process with significantly
lower glass heating. Firstly, a spin-on phosphorous source is applied to the poly-Si film surface and
baked. Then, it is scanned by the line laser beam, which heats the silicon above 1000C causing
relatively shallow (a few tens of nanometres) dopant diffusion. Cell performance is comparable to the
RTA process.
Two different metallisation schemes are under investigation. The simplest scheme is shown in Figure
4.4.1.43. First, the 1 cm2 cell area is defined by laser scribing. The silicon is then coated with a TiO2
impregnated isolating resist. Holes are etched through the resist by inkjet and then through the silicon
to expose the buried (absorber) layer. The resist is then exposed to solvent to cause it to reflow over
the walls of the emitter, but leave the absorber still exposed. A second round of holes is then etched
through the resist to expose the top (emitter) layer. The surface is coated with a thin layer of sputtered
aluminium, which makes contact with both the absorber and emitter layers, before laser scribing
separates the contacts, which are connected to tabs at either side of the cells. Finally, a 1-hour contact
bake at below 200C completes the process.
Figure 4.4.1.43: Cell metallisation scheme
for laser crystallised thin-film solar cells
Significant, but reversible performance degradation was found for these cells in the hours subsequent
to contact baking. The degradation was attributed to high contact resistance between the lightly doped
absorber and the aluminium.
- 120 -
A solution to this has been developed, which uses a selective p+ arrangement to improve the absorber
contacts. In contrast to the baseline scheme, the absorber contact holes are larger in diameter and
fewer in number. Additional boron is diffused into the floors of the holes, creating p+ regions and
thus allowing better contact between the silicon and the aluminium. Table 4.4.1.4 shows the cell
characteristics for cells with each scheme. For the selective p+ scheme, similar initial performance to
the baseline scheme has been achieved, while the degradation is eliminated.
Table 4.4.1.4: One-sun cell characteristics for Cell A, with the selective p+ contacting scheme and
Cell B with the simple scheme.
Cell
JSC [mA/cm2]
VOC [mV]
FF [%]
Cell A (initial)
Cell A (delayed)
Cell B (initial)
Cell B (delayed)
24.4
24.2
23.6
23.9
555
559
524
434
56.2
56.9
62.3
46.4
Eff
[%]
7.6
7.7
7.7
4.7
The role of the intermediate layer between the glass and the silicon has also been shown to have a
significant influence on the process with regard to wettability, diffusion of contaminants and dopants
and material VOC, while also determining the optical properties for the superstrate configured devices.
Silicon oxide (SiOx), silicon nitride (SiNx) and silicon carbide (SiCx), either alone or in multilayer
stacks are suitable candidates for the intermediate layer and each has various advantages and
disadvantages.
If the energy delivered to the silicon during laser crystallisation is too low, incomplete melting occurs,
resulting in the small-grained structure shown in Figure 4.4.1.44a. If the energy is too high, the silicon
film dewets the glass surface and balls up as shown in Figure 4.4.1.44b. The intermediate layer has a
great influence on how wide the process window between these two extremes is. SiCx has been found
to be the best wetting layer, allowing the largest process window, although a successful process is also
possible with SiOx and SiNx.
Figure 4.4.1.44: Lower
and upper boundaries
for laser energy during
crystallisation process,
resulting in (a) small
grain size, shown by
cross-sectional TEM
and (b) dewetting
shown by optical
reflection micrograph.
(a)
(b)
Reflectance data, measured after metallisation and shown in Figure 4.4.1.45a, demonstrate the benefit
of using a 70 nm SiNx, which results in much lower reflectance in the important wavelength region
around 600 nm. The sum of reflectance (measured after metallisation) and absorbance (measured
prior to silicon deposition) is included in Figure 4.4.1.45a to show the total lost photons in this region.
The SiCx has significant parasitic light absorption below about 700 nm wavelengths. This restricts the
use of the SiCx to such thin ~20 nm layers as used here, which is too thin to act as an ARC. The SiOx
is optically identical to the glass and therefore has no impact on reflection when used on its own or on
the glass side of a multi-layer stack.
- 121 -
The addition of a thin oxide layer on the silicon side to make SiOx/SiCx/SiOx (OCO) or
SiOx/SiNx/SiOx (ONO) triple-intermediate-layer stacks increases the optical thickness compared to
double-layer stacks, changing the interference properties and thus the reflectance. This is shown in
Figure 4.4.1.45b. The O/C/O is a much better anti-reflection coating than the SiCx alone. The O/N/O
stack is thicker than desirable, with the reflectance minimum pushed out to around 700 nm.
Figure 4.4.1.45:
Reflectance (solid
lines) and sum of
reflectance and
parasitic absorbance
(dashed lines) for cells
with (a) single or
double intermediate
layer stacks and (b)
triple layer stacks.
(a)
(b)
Figure 4.4.1.46a shows J/V curves for the same solar cells using each type of intermediate layer. The
SiOx is clearly best, with an active area efficiency of 7.7 %, while the SiNx, despite its better optical
properties, reaches only 3.2 %, with the SiCx is not much better at 3.3 %.
The J-V data shown in Figure 4.4.1.46b clearly demonstrate the benefit of the additional oxide layers
in triple-layer stacks. The SiCx based cell improved to 7.8 % active area efficiency, while the SiNx
based cell achieved 8.4 %.
Figure 4.4.1.46:
Reflectance (solid
lines) and sum of
reflectance and
parasitic absorbance
(dashed lines) for cells
with (a) single or
double intermediate
layer stacks and (b)
triple layer stacks.
(a)
(b)
The best efficiency achieved so far is 9.1 % and was made using an O/N/O intermediate layer stack as
described above. The VOC of this cell is 555 mV, making this one of many cells to consistently reach
VOC greater than 550 mV. This confirms the increased silicon quality compared to SPC films. The JSC
of the best cell is 26.8 mA/cm2. An even higher JSC of 27.8 mA/cm2 was measured for a similar cell
with optimised intermediate layer thicknesses. These cells were fabricated with no optical texturing
and the JSC values are expected to increase much further as this is developed.
The internal quantum efficiency (IQE) of the best cell is compared in Figure 4.4.1.47 to that of a
typical sample with only a SiOx intermediate layer and is considerably lower in the short-mid
wavelength region. This suggests that yet further gains to both current and voltage may be achieved
by optimising the intermediate layers to achieve both the anti-reflective benefit of the O/N/O stack
and the higher minority carrier lifetime when only the SiOx is used.
- 122 -
Figure 4.4.1.47: Internal
quantum efficiency of the
record 9.1 % efficient cell
(blue) employing an O/C/O
intermediate layer and a
typical cell (red) employing
a SiOx only intermediate
layer
Figure 4.4.1.48: Efficiency
improvement of record cell
over time.
4.4.1.11
4.4.1.1
4.4.1.2
4.4.1.3
4.4.1.4
4.4.1.5
4.4.1.6
Dec-2012
Sep-2012
Jun-2012
Mar-2012
Dec-2011
Sep-2011
Jun-2011
Mar-2011
10.0
9.5
9.0
8.5
8.0
7.5
7.0
6.5
6.0
Dec-2010
Record Efficiency [%]
Modelling a cell with the typical IQE suggests that the minority carrier lifetime is at least 260 ns and
that the front surface recombination velocity is at most 1900 cm/s. Modelling of simple device
optimisations consistent with typical poly-Si thin film solar cells suggests that cell efficiency of over
13 % is possible with the present material quality. Figure 4.4.1.48 shows the regular increase in record
efficiency since the beginning of this work. Continuing this trajectory would see such efficiency
realised within two years from now.
References
M.A. Green, “Polycrystalline silicon on glass for thin-film solar cells”, Applied Physics
A: v. 96, p. 153, 2009.
A. G. Aberle, “Fabrication and characterisation of crystalline silicon thin-film materials
for solar cells”, Thin Solid Films, v. 26, p. 511, 2006.
P. A. Basore, “Crystalline Silicon on Glass Device Optimization”, Proceedings of the
21st EUPVSEC, p. 544, Dresden, Germany, 2006.
M. Keevers, T.L. Young, U. Schubert, R. Evans, R.J. Egan and M.A. Green, “10%
Efficient Csg minimodules:, Proceedings of the 22nd EUPVSEC, p. 1783, Milan, Italy,
2007.
G. Jin, P. I. Widenborg, P. Campbell, S. Varlamov, “Lambertian matched absorption
enhancement in pecvd poly-Si thin film on aluminium induced textured glass
superstrates for solar cell applications”, Progress in Photovoltaics: Research and
Application, 2010, v. 18, pp. 482-589.
H. Cui, P. Campbell, “A glass thinning and texturing method for light incoupling in thinfilm polycrystalline silicon solar cells application”, Phys. Status Solidi RRL, 2012, v. 6,
no. 8, pp. 334–336, DOI 10.1002/pssr.201206266.
- 123 -
4.4.1.7
4.4.1.8
4.4.1.9
4.4.1.10
4.4.1.11
4.4.1.12
4.4.1.13
4.4.1.14
4.4.1.15
4.4.1.16
4.4.1.17
4.4.1.18
4.4.1.19
4.4.1.20
4.4.1.21
4.4.1.22
T. Soderstrom, J. Qian, K. Omaki, O. Kunz, D. Ong, S. Varlamov, “Light confinement in
e-beam evaporated thin film polycrystalline silicon solar cells”, Phys. Stat. Solidi, RRL
(A), 2011, v. 5, pp. 181-183.
Q. Wang, T. Soderstrom, K. Omaki, A. Lennon, S. Varlamov, “Etch-back silicon
texturing for light-trapping in electron-beam evaporated thin-film polycrystalline silicon
solar cells”, Energy Procedia, 2012, v. 15, pp. 220-228in press.
Z. Ouyang, X. Zhao, S. Varlamov, Y. Tao, J. Wong, S. Pillai, “Nanoparticles enhanced
light-trapping in thin-film silicon solar cells” Progress in Photovoltaics, 2011, v. 19, no.
8, pp. 917-926, DOI: 10.1002/pip.1135.
S. Varlamov, J. Rao, T. Soderstrom, “Polycrystalline silicon thin-film solar cells
with plasmonic-enhanced light-trapping”, J. Vis. Exp (JoVE), 2012, v. 65, e4092,
DOI: 10.3791/4092.
J. Rao, S. Varlamov, J. Park, S. Dligatch, A. Chtanov, “Optimization of
Dielectric-coated Silver Nanoparticle Films for Plasmonic-Enhanced Light
Trapping in Thin Film Silicon Solar Cells”, J of Plasmonics, 2012,v. 7, no. 4, pp.
1-7, DOI 10.1007/s11468-012-9473-y.
Z. Ouyang, O. Kunz, A. B. Sproul, S. Varlamov, “Influence of the absorber doping for ptype evaporated polycrystalline silicon thin-film solar cells on glass, J. Appl. Phys.,
2011, v. 109, 060104.
B. Eggleston, S. Varlamov, M. Green, “Large area diode laser defect annealing of
polycrystalline silicon solar cells”, IEEE Trans Electronic Devices (A), 2012, v.
59, no. 10, pp. 2838-2841, DOI: 10.1109/TED.2012.2208648.
S. Varlamov, B. Eggleston, J. Dore, R. Evans, D. Ong, O. Kunz, J. Huang, U.
Schubert, K. H. Kim, R. Egan, M. Green, “Diode laser processed crystalline
silicon thin-film solar cells”, SPIE Photonics West, 2-7 February 2013, San
Francisco, USA.
P.J. Gress, P.I. Widenborg, S. Varlamov, A.G. Aberle, Wire bonding as a cell
interconnection technique for polycrystalline silicon thin-film solar cells on glass”,
Progress in Photovoltaics: Research and Applications, 2010, v. 18, pp. 221-228.
J. Dore, R. Evans, B. Eggleston, S. Varlamov, M. Green, “Intermediate layers for
thin-film pc-Si solar cells formed by diode laser crystallisation”, MRS Spring
Meeting, San-Francisco, USA, 9-13 April 2012, v. 1426, DOI:
10.1557/opl2012.866.
B. Eggleston, S. Varlamov, J. Huang, R. Evans, J. Dore, M. Green, “Large
Grained, Low Defect Density Polycrystalline Silicon on Glass Substrates by
Large-area Diode Laser Crystallisation”, MRS Spring Meeting, San-Francisco,
USA, 9-13 April 2012, v. 1426, DOI: 10.1557/opl2012.1260.
J. Dore, R. Evans, U. Schubert, B. Eggleston, D. Ong, K. Kim, J. Huang, O.
Kunz, M. Keevers, R. Egan, S. Varlamov, M. Green, “Thin-film polycrystalline
silicon solar cells formed by diode laser crystallisation”, Progress in
Photovoltaics: Research and Application (A), 2012, DOI: 10.1002/pip.2282.
J. Dore, S. Varlamov, B R. Evans, Eggleston, D. Ong, O. Kunz, J. Huang, U.
Schubert, K. Kim, R. Egan, M. Green, “Performance potential of low-defect
density silicon thin-film solar cells obtained by electron beam evaporation and
laser crystallisation”, EPJ Photovolt., 2013, v.4, article no. 40301.
P.I. Widenborg, N. Chuangsuwanich, A.G. Aberle, Glass texturing, Patent
AU2004228064A1.
P.I. Widenborg, A.G. Aberle, “Polycrystalline Silicon Thin-Film Solar Cells on AITTextured Glass Superstrates”, Advances in Optoelectronics, 2007, article ID 24584.
P.I. Widenborg, S.V. Chan, T. Walsh, and A.G. Aberle, “Thin-Film Poly-Si Solar Cells
on AIT-textured Glass – Importance of the Rear Reflector”, Proceedings of the 33rd
IEEE PVSC, San-Diego, 2008.
- 124 -
4.4.1.23
4.4.1.24
4.4.1.25
4.4.1.26
4.4.1.27
4.4.1.28
4.4.1.29
4.4.1.30
4.4.1.31
4.4.1.32
4.4.1.33
4.4.1.34
4.4.1.35
4.4.1.36
4.4.1.37
4.4.1.38
4.4.1.39
4.4.1.40
O. Kunz, Z. Ouyang, S. Varlamov, and A.G. Aberle, "5% efficient evaporated solidphase crystallised polycrystalline silicon thin-film solar cells", Progress in Photovoltaics:
Research and Applications, v. 17, p. 567, 2009.
O. Kunz, Z. Ouyang, J. Wong, and A.G. Aberle, “Advances in evaporated solid-phasecrystallized poly-Si thin-film solar cells on glass (EVA)”, Advances in Optoelectronics,
2008, article ID 532351.
O. Kunz, J. Wong, J. Janssens, J. Bauer, O. Breitenstein, A.G. Aberle, “Shunting
Problems Due to Sub-Micron Pinholes in Evaporated Solid-Phase Crystallised Poly-Si
Thin-Film Solar Cells on Glass”, Progress in Photovoltaics, 2009, v. 17, pp.35-46.
H. Cui, P. Gress, P. Campbell, M. Green, “Developments in the aluminium induced
texturing (AIT) glass process”, Glass Technol., v. 53, no. 4, pp. 158-165.
H. Cui, M. Green, P. Campbell, O. Kunz, S. Varlamov, “A Photovoltaic light-trapping
estimation method for textured glass based on surface decoupling calculation”, Sol. Mat.
(A), 2013, v. 109, pp. 82-90.
P. I. Widenborg, G. Jin, P. J. Gress, and S. Varlamov, “High rate 13.56MHz PECVD aSi:H depositions for SPC poly-Si thin film solar cells”, Proceedings of the 24th
EUPVSEC, Hamburg, 2009.
A. Straub, R. Gebs, H. Habenicht, S. Trunk, R. A. Bardos, A. B. Sproul, and A. G.
Aberle, “Impedance analysis: A powerful method for the determination of the doping
concentration and built-in potential of nonideal semiconductor p-n diodes”, J. Appl,
Phys., 2005, v. 97, p. 083703.
Z. Ouyang, O. Kunz, M. Wolf, P. Widenborg, G. Jin and S. Varlamov, “Challenges of
Evaporated Solid-Phase-Crystallised Poly-Si Thin-Film Solar Cells on Textured Glass”,
Proceedings of the 18th International Photovoltaic Science and Engineering Conference
& Exhibition, Kolkata, India, 2009.
M. Werner, U. Schubert, C. Hagendorf, J. Schneider, M. Keevers, R. Egan, “Thin film
morphology, growth and defect structure of e-beam deposited silicon on glass”,
Proceedings of the 24th EUPVSEC, Hamburg, Germany, 2009.
J. Rao, S. Varlamov, J. Park, S. Dligatch, A. Chtanov, “Optimization of Dielectric-coated
Silver Nanoparticle Films for Plasmonic-Enhanced Light Trapping in Thin Film Silicon
Solar Cells”, J of Plasmonics, 2012,v. 7, no. 4, pp. 1-7.
Park, T. Kim, J. Rao, S. Varlamov, “Highest efficiency plasmonic polycrystalline silicon
thin-film solar cells by optimisation of plasmonic nanoparticle fabrication”, Plasmonics,
2013, in press.
P. Timans, J. Gelpey, S. McCoy, W. Lerch, S. Paul, “Millisecond Annealing: Past,
Present and Future”, Mat. Res. Soc. Symp. Proc., 2006, v. 912, 0912-C01-01.
K. Ohdaira, T. Fujiwara, Y. Endo, S. Nishizaki, H. Matsumura, “Formation of SeveralMicrometer-Thick Polycrystalline Silicon Films on Soda Lime Glass by Flash Lamp
Annealing”, Jap. J. Appl. Phys., v. 47, p. 8239, 2008.
K. Ohdaira, T. Fujiwara, Y. Endo, S. Nishizaki, H. Matsumura, “Explosive
crystallisation of amorphous Si films by flash lamp annealing” J. Appl. Phys. v. 6,
044907, 2009.
P. Altermatt and G. Heiser, “Development of a three-dimensional numerical model of
grain boundaries in highly doped polycrystalline silicon and applications to solar cells”,
J. Appl. Phys., v. 91, p. 2002.
L. Carnel, I. Gordon, D. Van Gestel, G. Beaucarne, J. Poortmans, A. Stesmans, “High
open-circuit voltage values on fine-grained thin-film polysilicon solar cells”, J. Appl.
Phys., v. 100, p. 063702, 2006.
J. Wong, J. L. Huang, O. Kunz, Z. Ouyang, S. He, P. I. Widenborg, A. G. Aberle, M.
Keevers, M. A. Green, “Anomalous temperature dependence of diode saturation currents
in polycrystalline silicon thin-film solar cells on glass”, J. Appl. Phys., v. 105, p. 103705,
2009.
J. Wong, J. Huang. B. Eggleston, M. Green, O. Kunz, M. Keevers, R. Egan, “Lifetime
limiting recombination pathway in thin-film polycrystalline silicon on glass solar cells”,
J. Appl. Phys. v. 107, 123705.
- 125 -
4.4.1.41
4.4.1.42
4.4.1.43
4.4.1.44
4.4.1.45
4.4.1.46
J. Wong, J. Huang, M. Keevers, M. Green, “Voc-Limiting recombination mechanism in
thin film silicon on glass solar cells”, Proc. 35th IEEE PVSC, Hawaii, USA, 2010.
J. Wong, J. Huang, S. Varlamov, M. Green. R. Evans, M. Keevers, R. Egan, “Structural
inhomogeneities in polycrystalline silicon on glass solar cells and their effects on device
characteristics”, Progress in Photovoltaics”, 2011, v. 19, no. 6, pp. 695-705.
J. Wong, J. Huang, S. Varlamov, M. Green, “The roles of shallow and deep levels in the
recombination behaviours of polycrystalline silicon solar cells”, Progress in
Photovoltaics, 2012, v. 20, pp, 915-922.
M. Kittler, W. Seifert, K.W. Schroeder, “Temperature dependent EBIC diffusion length
measurements in silicon”, Phys. Stat. Sol. (a), 2986, v. 93, pp.K191-K104.
V. Kveder, M. Kittler, W. Schroter, “Recombination activity of contaminated
dislocations in silicon: a model describing electron-beam induced current contrast
behaviour”, Phys. Rev. B, 63, p. 115208 2001.
J. Huang, J. Wong, M. Keevers, M. Green, “Cell performance limitation investigation of
polycrystalline silicon thin-film solar cells on glass by transmission electron
microscopy”, Proc. 25th EUPVSEC, Valencia, Spain, 2010, pp.3565-3567.
- 126 -
4.4.2
Organic Thin Films
University Staff
A/Prof. Ashraf Uddin
Prof. Martin Green
Prof. Gavin Conibeer
Senior Research Fellow
Dr Dirk König
Postdoctoral Fellow
Dr Shujuan Huang
Dr Thilini Ishwara
Research Assistant
Xiaohan Yang
Postgraduate Research Students
Xiaohan Yang
Fred Qi
Matthew wright
Pengfei Zhang
Rui Lin
Kah Howe (Alex) Chan
Rui Sheng
Undergraduate Thesis Students
Ran Chen
Anna Nadolny
Zhang Xiaoyang
Wenkai Cao
Shi Chao Zhang
Yao Zhang
Yuan Chu
Jinyang Yu
Hang Xing
Masahiro Miwa
Taste of Research Student
Yu Jiang
4.4.2.1
Introduction
An organic photovoltaic (OPV) research lab has been established at UNSW with world class research
facilities to work on organic semiconductor materials and devices. Today this lab is equipped with
facilities supported by an ARC discovery grant (DP10), a UNSW equipment grant, a grant from the
Green Family, a Faculty grant and from the support of ARC Photovoltaic Centre of Excellence,
UNSW. Currently, the research team consists of three academics, three research fellows, one research
assistant, seven PhD, and 10 honours students.
Our research has produced meaningful data over the last three years. Four journal papers and seven
conference proceeding papers have already been published. Several other papers are underway to
publish. Our research activities are mainly focusing on the following three areas:
- 127 -
(i) Morphology control of bulk heterojunction organic solar cells
(ii) Light trapping in organic solar cells
(iii) Organic/inorganic nanoparticles hybrid solar cells
The nanoscale texture or thin film morphology of the donor/acceptor blends used in most bulk
heterojunction OPV devices is a critical variable that can dominate both the performance of new
materials being optimized in the lab and efforts to move from laboratory-scale to factory-scale
production. Although OPV conversion efficiencies have improved significantly in recent years,
progress in morphology optimization still occurs largely by trial and error, in part because much of
our basic understanding of how nanoscale morphology affects the optoelectronic properties of these
heterogeneous organic semiconductor films has to be inferred indirectly from macroscopic
measurements. If the morphology could be controlled on a molecular scale, the efficiency of charge
separation and transport could be expected to be substantially higher. We are now working on
fundamental issues of morphology optimization for organic films for OPV devices such as control of
electronic structure at film interfaces, exciton dissociation and carrier transport for photovoltaic
operation to achieve our goals. Ab-initio (density functional theory – DFT) methods are a valuable
tool to obtain insight into the charge separation (exciton dissociation) and transport process as well as
the photochemical stability of organic molecules – a major challenge for organic PV. This helps us to
find robust and electronically suitable molecules which determine the morphology of OPV devices.
Furthermore, we have investigated the plasmonic enhancement of bulk heterojunction OPV cells with
an incorporated thin silver (Ag) film. Such films consist of plasmon-active and size-variable Ag
nanostructures. Incorporation of a plasmon-active Ag nano-material is shown to enhance light
absorbance in the photoactive layer. Consequently, enhancements of external quantum efficiency at
red wavelengths are observed. This plasmonic enhancement needs to be optimised to further improve
the photo conversion efficiency of OPV cells.
Hybrid solar cells are a mixture of nanocrystals (NCs) of both organic and inorganic materials. They
combine the unique properties of inorganic semiconductor nanocrystals with properties of
organic/polymeric materials. Inorganic NCs such as quantum dots, nano-tubes, etc. have also high
absorption coefficients and particle size induced tunability of the optical band-gap. Band-gap tuning
in inorganic nanocrystals with different sizes can be used for realization of device architectures, such
as tandem solar cells in which the different bandgaps can be obtained by modifying only one chemical
compound. A substantial interfacial area for charge separation is provided by nanocrystals, which has
high surface area to volume ratios. Thus, the organic-inorganic NCs hybrid concept for PV cells has
become increasingly interesting and attractive in recent years.
Figure 4.4.2.1: Schematic of the films that form the multilayer porous metal-oxide: polymer solar
cell.
As an example, a hybrid organic-inorganic solar cell combining a metal-oxide with a polymer is
shown in Fig. 4.4.2.1. The structure has a similarity to the solid state dye-sensitized solar cell but the
semiconducting polymer can be used for both light absorption and hole transport. Light absorption in
- 128 -
the polymer leads to the creation of an exciton. Following exciton diffusion to the organic-inorganic
interface, exciton dissociation results to produce separated charges. Electrons are transported by the
metal-oxide and the holes by the polymer. The nanostructure of the metal-oxide can be controlled
during growth and the same material can be formed in different structures to create various
morphologies. The device can be fabricated by depositing the metal-oxide layer on ITO coated glass
substrate. The polymer is then introduced into the nanoporous metal-oxide film, creating a bulk
heterojunction of a thickness ~250 nm. The top contacts (PEDOT:PSS and Au) are then applied.
Blocking layers of pure material are coated at both electrodes to avoid losses by shunting.
4.4.2.2
OPV Conversion Efficiency
Interest in OPV with increasing conversion efficiency has grown exponentially over recent years. The
key development in the OPV device is the bulk heterojunction cells as shown in Fig. 4.4.2.2. Both
packaging and cell efficiency had improved by 2006 to the stage where independent measurement of
cell efficiency was both feasible and warranted, with subsequent efficiency improvements well
documented. Konarka established 4.8% efficiency for a tiny 0.14 cm2 cell in July 2005, with Sharp
demonstrating 3.0% for a relatively massive 1 cm2 cell the following year. Konarka increased this to
5.15% in December 2006 with 5.24% posted for a 0.7 cm2 cell in July 2007. In the latter device,
P3HT was replaced by lower “bandgap” PCPDTBT. Further benefits have been claimed using C71PCBM rather than the usual C61–PCBM. Companies such as Plextronics have reported roughly
comparable results using similar materials with 5.4% efficiency confirmed in July 2007. Konarka
achieved 6.3% efficiency in a 0.7 cm2 device based on the use of PCZ (polycarbazoles) as the donor
material.
OPV device developer Solarmer Energy achieved the conversion efficiency record of a plastic
OPV champion cell—7.9% in December 2009. In November 2010 Konarka achieved their highest
OPV efficiency 8.3%. Mitsubishi reported their highest OPV efficiency around 9.3% in early 2011.
University of California has developed 10.6% efficient tandem polymer solar cells in the beginning of
2012 [4.4.2.1].
Figure 4.4.2.2: Recent improvements in independently confirmed efficiency for small area (1cm2)
organic solar cells (extracted from Green et al. [4.4.2.2, 4.4.2.3]).
- 129 -
The photovoltaic power conversion efficiency of nearly every single junction solar cell is
fundamentally limited by the Shockley-Queisser (SQ) limit [4.4.2.4]. The applicability of the SQ
theory is based on the assumption of a step function quantum efficiency, perfect extraction of carriers,
and radiative recombination as the only loss processes. This approach places appropriate upper limits
for many inorganic semiconductors used for photovoltaics. For organic bulk heterojunction cells, the
SQ limit is still an upper efficiency limit; however, some of the original assumptions of the SQ theory
are no longer fair to assume in organic solar cells. The radiative efficiency limits have the advantage
that hardly any internal properties of the photovoltaic energy conversion process have to be known.
Any charge carrier is assumed to be collected directly after photogeneration and every injected carrier
recombines radiatively leading to an emission defined by the absorptance of the device and the Planck
spectrum. The theoretical conversion efficiency limit of OPV cells is over 22% [4.4.2.5]. The current
experimental single-junction OPV cell efficiency is η ≈ 11%. This observation directly leads to the
question, where this factor of nearly 2 in efficiency drop originates. There are in principle four
reasons for this efficiency gap:
(1) Optical losses, meaning that light may be parasitically absorbed in solar cell layers that are
not active in charge carrier collection or that the current organic solar cells, have no effective
light trapping scheme as most inorganic solar cells.
(2) Exciton losses due to insufficient transport of excitons to the next donor-acceptor interface or
due to inefficient exciton dissociation at the interface.
(3) Recombination losses, meaning that only a small part of the carriers recombines radiatively,
but most of the carriers recombine nonradiatively at donor-acceptor interfaces or defects in
the absorber and at the contacts.
(4) Collection losses due to insufficient mobilities allowing only some of the free carriers to
reach the contacts. In order to quantify the importance of these loss mechanisms, we need to
have a simple device model capable of reproducing experimental current/voltage curves.
The key issues of polymer design include engineering the bandgap and energy levels to achieve high
JSC and VOC, enhancing planarity to attain high carrier mobility, materials processing ability and
stability. All of these issues are correlated with one another. In the ideal case, all factors should be
optimized in a single polymer, but this remains a significant challenge. The cell design improvement,
implementation of a range of new ideas relevant to carrier’s recombination reduction and improved
light trapping are needed for OPV cells to improve the efficiency. The exploration of the potential of a
range of “third generation” options also needs to be investigated to improve the efficiency of cells.
This will open the opportunity for fabricating optimised tandem stacks of OPV-based cells, with
much higher efficiency. Many new ideas and much new technology will be required before
competitive products eventuate. It is very important to identify the key degradation mechanisms to
improve the device lifetime and performance. Much work needs to be done, both in cell material
selection and in low cost encapsulation, before a level of durability required for mainstream product is
obtained.
4.4.2.3
Standard OPV structures
The fullerene based acceptor material (PCBM) is interspersed with a donor polymer, commonly
P3HT, with line-bond diagrams for these materials and the energy band diagram of standard “bulk
heterojunction” cell structure is shown in Fig. 4.4.2.3. From this baseline of device structure, the PV
Centre believes it has several strengths that will allow it to contribute to the ongoing cell efficiency
evolution documented in Fig. 4.4.2.2 and to the development of more stable and durable devices. The
challenge for OPV is that absorbing light in an organic donor material produces coulombically bound
excitons that require dissociation at the donor/acceptor interface. The energy-level offset of the
heterojunction (donor/acceptor interface) is also believed to play an important role for the dissociation
- 130 -
of bound excitons in OPV cells. An efficient exciton separation into free charge carriers at the
interface of donor/acceptor materials will increase the photocurrent of the solar cell. The transport of
excitons to the interface also limits the conversion efficiency. The details of the exciton transport
mechanism in OPV are still not understood. There is an urgent need for more systematic study to gain
in depth understanding of the mechanisms of light absorption, excitons dissociation and charge
transport, particularly their relationships with molecular and morphological structures of the materials as
well as the nano-scale architectural design of the devices. We believe that the conversion efficiency of
OPV can be increased to much higher than the reported values for commercial applications through indepth understanding of the light harvesting behaviour and mechanisms of organic materials and by
selecting suitable materials and new device architectures.
(a)
(b)
Figure 4.4.2.3: (a) A schematic diagram of a bulk-heterojunction organic solar cell. The molecular
structure of donor (P3HT) and acceptor (PCBM) is shown. (b) The schematic band energy structure of
the bulk-heterojunction organic cell is shown.
Figure 4.4.2.4 shows the two examples of disorder donor/acceptor blending design and systematic
alignment of donor and acceptor layers for bulk-heterjunction organic solar cell.
- 131 -
(a)
(b)
Figure 4.4.2.4: (a) A schematic random distribution of the disordered donor/acceptor morphology in
the blended bulk-heterojunction organic solar cell. (b) A schematic diagram of systematically aligned
and ordered donor/acceptor morphology for bulk-heterjunction organic solar cell.
An ideal donor/acceptor morphological structure is proposed to be an ordered and vertical wellaligned interpenetrating donor and acceptor (D/A) network with nanoscale domain sizes, i.e. 20-40
nm, as shown in Fig. 4.4.2.4(b) [4.4.2.6]. The domain sizes are comparable with exciton diffusion
length ( 10 nm) in pristine organic polymer. To achieve the ordered D/A morphology, researchers
have developed two techniques, namely wetting of free-standing anodic aluminium oxide (AAO)
template and nanoimprinting by pre-patterned mold techniques. However, both techniques are quite
costly and complicated.
We have developed an improved ordered D/A morphology control technique by introducing a
designed AAO mould to allow preparing ordered D/A networks. An important point of our research is
to develop a simple, cheap and less time consuming technique to fabricate a highly ordered AAO
nano-imprinting mould. The implementation of this nano-imprinting mould into the proposed selfaligned D/A morphological control technique needs more investigation.
4.4.2.4
Active Layer Annealing
Thermal annealing is an effective method to increase the crystallinity of polymers and to improve the
OPV device performance due to the change of donor/acceptor morphology of active layer into
nanoscale domains as well as the better intermolecular alignment of the P3HT and PCBM giving
higher hole mobility. Usually, as-cast P3HT:PCBM films have featureless and homogeneous surfaces,
while annealing produces a coarser textured surface with obvious phase segregation. In our studies,
we have found that annealing the device after the application of the cathode improves the efficiency
of the device much more significantly that annealing before the cathode deposition. The reason for
this is believed to be reduced mobility of the PCBM molecules under the increased confinement of the
cathode. In our studies we have explored a range of anneal temperatures and find that the optimum
annealing temperature depends upon the blend ratio, and in general lies in the range of 120°C to
150°C. The optimum annealing temperature for our blend appears to be ~130°C for 10 minutes in
nitrogen atmosphere. An improved transport property of active layer is also observed which leads to a
reduced series resistance and in turn a better device performance. Quantitative information about the
donor P3HT average crystal size ‘D’ is estimated after annealing at different temperatures from the
XRD spectra using Scherrer’s relation [4.4.2.7]:
D
0.9
 2 cos
(4.4.2.1)
- 132 -
where 2θ is the full width at half maximum (FWHM) of the XRD peak (in radians). The P3HT
crystallite size undergoes a significant increase – from 11.05 nm in the as-cast film to 18.95 nm –
after annealing at 180°C for 10 minutes. A linear increase in crystallite size is observed for increasing
annealing temperatures.
(a) (b)
Figure 4.4.2.5: (a) Absorption spectra of P3HT:PCBM film after annealing in nitrogen atmosphere
for 10 minutes at different temperatures. (b) Plot of P3HT crystallite size vs annealing temperatures in
the P3HT:PCBM films, estimated from the XRD spectra of P3HT [4.4.2.8].
However, annealing temperatures higher than 130C causes a reduction in electronic performance. It
was found that a trade-off governs the optimum thermal annealing temperature for a bulk
heterojunction device. Low annealing temperatures induce crystallinity in the polymer phase which
leads to an improvement in charge carrier conduction. Thermal annealing temperatures above the
optimum value ( 130C) cause an aggregation of PCBM molecules, which reduces the ability for
excitons to dissociate, consequently causing a reduction in electronic performance.
(a) (b)
Figure 4.4.2.6: Plot of EQE spectra of P3HT:PCBM devices after annealing at different temperatures
for 10 minutes in nitrogen atmosphere. (b) The plot of EQE peak values at around 510 nm wavelength
region as a function of annealing temperature for P3HT:PCBM devices. The fitted line is a Gaussian
function [4.4.2.8].
Figure 4.4.2.6 shows the EQE spectra of P3HT:PCBM devices as a function of annealing conditions.
The significant improvements of electrical performance of our OPV devices are observed as a result
of thermal annealing from 80 to 130°C due to the increase of crystallinity which leads to improve the
charge carriers transport properties. Annealing temperature above 130C led to reduced device
- 133 -
electrical performance due to the diffusion and aggregation of PCBM molecules, which limits the
dissociation of excitons due to the reduction of P3HT (donor) and PCBM (acceptor) interfacial area.
There is an optimal annealing condition which best compensates the trade-off between these two
effects. The plot of EQE peak value, at around 510 nm, as a function of annealing temperature is
shown in Fig. 4.4.2.6(b). This provides a very clear indication of two distinct trends of annealing
temperature on electrical performance of OPV devices. Initially, the electrical performance of device
rises with increasing annealing temperature, until a maximum is obtained at 130°C, after which,
electrical performance decreases. The behaviour of peak EQE with annealing temperature can be
described by the following expression:
Peak EQE = A*exp[-((Ta – To)/B)2] + C
(4.4.2.2)
where A = 33.64 is a constant, Ta is the annealing temperature in C, To = 127.5 C is the optimum
annealing temperature obtained from the experimental values, B = 36.79 is related to the FWHM of
peak EQE curve with annealing temperature and C = 5.05 is the intersection of peak EQE curve with
y-axis.
4.4.2.5
Anodic Aluminum Oxide (AAO) Nanoimprinting Mould for Ordered Morphology
Control of Organic Solar Cells
A simple experimental setup as shown in Figure 4.4.2.7 was utilised for the anodic oxidation process.
A piece of Nickel alloy plate worked as the cathode metal and the Al sample was clipped to the
anode. A constant DC voltage was applied between the two terminals. 0.3M H2SO4 was selected as
the electrolyte solution. The solution was continuously agitated by a magnetic stirring ball to maintain
a homogenous solution and to release hydrogen generated by the anodic reaction. The first
anodisation process could be just full anodisation lasting about 10 minutes or the full anodisation plus
an extended anodisation up to 3 hours. After that, samples were etched in 5 wt% H3PO4 solution to
obtain enlarged pore sizes.
Figure 4.4.2.7: A Schematic diagram of standard anodisation process.
The target of our research is to implementation of this prepared imprinting mold into the proposed
self-aligned D/A morphological control technique. The major process to prepare the AAO
nanoimprinting mold is to anodise high purity thin film aluminium evaporated on silicon substrate to
create a brief porous alumina surface. The anodisation and etching process improve the whole system
to an AAO mold with high regularity, which was characterised using high resolution scanning
electron microscopy (HRSEM Hitachi S900). By varying anodic conditions, such as applied DC
- 134 -
voltage, electrolyte solution and time of anodisation and each etching step, the pore sizes, pore
distribution and packing density can be well manipulated. A highly ordered AAO mould has been
successfully fabricated through simple anodisation and chemical etching steps as shown in Fig.
4.4.2.8. This simple, cheap and less time consuming technique provides another path to create wellaligned nanostructures as shown in Fig. 4.4.2.8. The AAO mould process parameters and steps are
listed in Table 1. Coupled with the future pattern transfer process (e.g. imprint the AAO mould on the
planar organic film), this technique has the great potential in achieving ordered D/A morphology
control of BHJ OPV devices. However, a feasible imprinting process needs to be further investigated
and developed.
Figure 4.4.2.8: SEM and TEM images of AAO sample on SiO2/Si with 20 minutes subsequent
etching.
Table 4.4.2.1: AAO thin film on Si wafer or SiO2/Si prepared in phosphoric acid (medium to large
pore size nanopattern) – A07-A11.
- 135 -
4.4.2.6
Light Trapping in OPV Devices
The use of a light trapping systems has been actively investigated to increase the total path length of
light into the active material without the need of increasing OPV cell physical thickness. The concept
is to induce sunlight to travel a longer distance in the active layer by a multi-pass path in order to
exploit most of the radiation. Many different light trapping schemes have been proposed to enhance
the quantity of light absorbed in OPV cells such as metal gratings [4.4.2.9], buried nanoelectrodes
[4.4.2.10] and multireflection structures [4.4.2.11]. However, the requirement that the feature sizes of
the scattering structures be comparable to the film thickness of the OPV cells, i.e., 50–300 nm, and at
the same time larger than the wavelength of light to effectively alter the photon propagation direction
is intrinsically contradictory. The integration of plasmonic structures with OPVs has been previously
examined in the context of metallic nanoparticles, which can be used to increase optical absorption.
Silver nanosphere (Ag NS) buffer layers were introduced via a solution casting process to enhance the
light absorption in poly(3-hexylthiophene) (P3HT) and [6,6]-phenyl-C61 butyric acid methyl ester
(PCBM) bulk-heterojunction organic solar cells as ahown in Fig. 4.4.2.9. In order to improve the
performance of P3HT:PCBM bulk-heterojunction organic solar cells, Ag NS buffer layers were
deposited in between ITO and PEDOT:PSS layer. The Ag NSs used for this investigation have
average diameters of 20 nm and 40 nm. A thin layer of Ag NSs was spin coated between ITO and
PEDOT:PSS hole-conducting layer. Physical film properties were analysed using TEM and optical
absorption, whilst electronic characteristics were examined using EQE measurements. These Ag NSs,
since they support surface plasmons, increase the optical electric field in the photoactive layer whilst
simultaneously improving the light scattering. As a result, this buffer layer improves the light
absorption of P3HT:PCBM blend and consequently improves the external quantum efficiency (EQE)
of organic solar cells as shown in Fig. 4.4.2.9. By incorporating plasmon-active Ag NSs, an enhanced
optical absorption and improved short circuit current Jsc is observed. The overall optical absorption of
P3HT:PCBM blend films in the spectral range of 350-650 nm were increased ~4 and 6 % as a result
of utilizing the 20 and 40 nm Ag NSs, respectively. The Jsc increases from ~3.93 mA/cm2 to ~4.76
and 4.89 mA/cm2 for 20 and 40 nm Ag NSs, respectively. These improvements are due to the light
scattering of Ag NSs and the high electric field generated by the excited surface plasmons.
Figure 4.4.2.9: A schematic diagram of OPV device with Ag nano-spheres (NSs) on an ITO anode
surface for the investigation of plasmonic effects on device performance. The TEM images of (a) 20
nm size Ag NSs; and (b) 40 nm size Ag NSs on ITO surfaces [4.4.2.12].
- 136 -
(a)
(b)
Figure 4.4.2.10: (a) Absorption spectra of P3HT:PCBM films with and without Ag NSs at room
temperature. (b) The EQE spectra of P3HT:PCBM blends OPV devices with and without Ag NSs.
There is a little effect of Ag NSs on the optical properties of OPV devices, but there is a significant
improvement of electrical performance of devices [4.4.2.12].
The scattering cross section is defined as:
| |
(4.4.2.3)
Here, α is the polarizability of the particle, given by α = 3V(ε-1)/(ε+2) for a small spherical particle in
vacuum, where V is the volume of the particle and ε is the permittivity of the metal. Scattering
depends on the size of the particles. As the size of the particles increases, scattering increases, this is
advantageous for light trapping. The optical absorption spectra of Ag NSs are shown in Fig. 4.4.2.11
for different particle sizes.
Absorbance (a.u.)
10nm Ag NSs (Peak: 408nm)
350
400
450
500
550
600
650
700
750
800
Wavelength (nm)
Figure 4.4.2.11: Optical absorbance spectra of 10 nm (green circle), 20 nm (red up-triangle), 30 nm
(blue down triangle) and 40 nm (brown square) Ag NSs. Inset: TEM image of 40 nm Ag NSs in
solution.
- 137 -
Figure 4.4.2.12 shows the uv-vis absorbance spectra of PEDOT:PSS and PCPDTBT:PC70BM blend
films with and without Ag NSs buffer layer using cell structure in Fig. 4.4.2.9. The secondary axis
displays the enhancements in absorbance, due to the addition of an Ag NSs buffer layer. For the
control blend, with no Ag NSs buffer layer (black dashed line), two main absorption peaks at ~400
nm and 760 nm are observed. Both PCPDTBT and PC70BM contribute absorption in the region of the
~400 nm peak, whilst absorption in the region of the ~760 nm peak is attributed only to PCPDTBT.
Absorption in the blend film is contributed to primarily by PCPDTBT, especially in long wavelength
region. PC70BM has an absorption in the spectral range of 300 nm to 650 nm. The concentration ratio
of 2:7 for PCPDTBT and PC70BM was selected, which is due to a balance between light absorption
and change transport within the active layer.
80
60
40
20
0
300
400
500
600
700
800
900
Absorbance Enhancement (%)
Absorbance (a.u.)
PEDOT:PSS + active layer
with 10nm Ag NSs
with 20nm Ag NSs
with 30nm Ag NSs
with 40nm Ag NSs
1000
Wavelength (nm)
Figure 4.4.2.12 Optical absorbance spectra of PEDOT:PSS and PCPDTBT:PC70BM active layer with
Ag NSs 10 nm (green solid), 20 nm (red solid), 30 nm (blue solid), 40 nm (brown solid) and control
blend (black dash) (primary axis). Enhancements of absorbance for the films with 10 nm (green
circle), 20 nm (red up-triangle), 30 nm (blue down-triangle) and 40 nm (brown square) Ag NSs buffer
layer (secondary axis).
Incorporation of the Ag NSs buffer layers slightly enhances the absorbance spectra. The optical
absorption enhancements of films with Ag NSs buffer layers have peaks around 400 nm, which
correspond to absorptions of Ag NSs. The absorbance displays continuous enhancement after 550 nm,
which benefits from plasmonic resonance of Ag NSs. The PEDOT:PSS buffer layer, as a low
refractive index media, causes the red-shift of the plasmonic resonance position. The overall
absorptions of PEDOT:PSS and PCPDTBT:PC70BM blend films are increased as a result of
incorporation of Ag NSs buffer layers. This is attributed to light scattering of the Ag NSs and the
increased electric field in the photoactive layer by excited localized surface plasmons around the Ag
NSs.
Figure 4.4.2.13 displays the EQE spectra of PCPDTBT:PC70BM devices with and without Ag NSs in
cell structure. The secondary axis displays the enhancements in EQE, due to the addition of an Ag
NSs buffer layer. For all compared devices, the EQE at wavelengths shorter than 335 nm is identical.
This indicates that the measurements were reliable; the condition of the tested devices was stable
during the entire study. Although our OPV devices have not been optimized, significant changes in
electronic performance as a result of localized surface plasmon resonance (LSPR) have been
- 138 -
observed. The EQE profile of the control device (black dashed line) primarily follows the profile of
optical absorbance of the PCPDTBT:PC70BM blend film, and has a maximum EQE value of ~40 % at
420 nm. With the presence of an Ag NSs buffer layer, the EQE values increase after 560 nm for
different size Ag NSs, this correlates with the plasmonic resonance position in the absorbance spectra.
The reason for the EQE enhancements around 340 nm is not fully understood. Similar results have
been observed in poly-Si thin film solar cells. It is suggested that the enhancement is related to a
higher order plasmonic resonance (quadrupole resonance) [4.4.2.13, 4.4.2.14]. Moreover, EQE
decreases due to the absorption of Ag NSs and red-shift of EQE peaks are observed around 400 nm.
The calculated maximum EQE enhancements in the spectral range 300-900 nm are 14 %, 18 %, 25 %
and 30 % for 10 nm, 20 nm, 30 nm and 40 nm Ag NSs, respectively. The larger-sized Ag NSs led to a
larger enhancement of EQE in red wavelength region.
90
Reference Cell
10nm Ag NSs
20nm Ag NSs
30nm Ag NSs
40nm Ag NSs
EQE (%)
30
75
60
45
20
30
10
15
0
0
300
EQE enhancement (%)
40
400
500
600
700
800
900
1000
Wavelength (nm)
Figure 4.4.2.13: EQE spectra of PCPDTBT:PC70BM devices with Ag NSs 10 nm (green solid), 20
nm (red solid), 30 nm (blue solid), 40 nm (brown solid) and control device (black dash) (primary
axis). Enhancements of EQE for the devices with 10 nm (green circle), 20 nm (red up-triangle), 30 nm
(blue down-triangle) and 40 nm (brown square) Ag NSs buffer layer (secondary axis).
Figure 4.4.2.14 displays the J-V characteristics of PCPDTBT:PC70BM devices with and without Ag
NSs. The control device exhibited a Voc of 551 mV, a Jsc of 10.16 mA/cm2, and a FF of 0.40, resulting
in a PCE of 2.24 %. After the addition of Ag NSs below the PEDOT:PSS layer, the values of Voc and
FF remained unchanged, suggesting that the interface between PEDOT:PSS and active layer was not
affected by the Ag NSs. Some portion of the Ag NSs surface was probably modified by PEDOT:PSS.
The photocurrents, however, were improved as the particle size was increased. For the devices
prepared with Ag NSs, the Jsc values were ~10.77, 10.97, 11.32 and 11.54 mA/cm2 for 10 nm, 20 nm,
30 nm and 40 nm Ag NSs, respectively. As a result, the PCEs increased to 2.35 %, 2.38 %, 2.45 %
and 2.56 %, respectively. This represents an increase of ~5 %, 6 %, 9 % and 14 % for 10 nm, 20 nm,
30 nm and 40 nm Ag NSs, with reference to the control device. The larger-sized Ag NSs led to a
larger enhancement of Jsc and PCE, which is thought to result from enhanced scattering interactions.
With the incorporation of Ag NSs, an additional photocurrent is created due to the enhancement of the
photogeneration of excitons associated with enhanced electric field intensity from the localized
surface plasmonic resonance. This plasmonic coupling of light is a major contributor to the improved
electronic transport through the blend films.
- 139 -
12
Reference Cell
10nm Ag NSs
20nm Ag NSs
30nm Ag NSs
40nm Ag NSs
Current Density (mA/cm2)
10
8
6
4
2
0
0.0
0.1
0.2
0.3
0.4
0.5
0.6
Voltage (V)
Figure 4.4.2.14: J-V characteristics of PCPDTBT:PC70BM devices with Ag NSs 10 nm (green
solid), 20 nm (red solid), 30 nm (blue solid), 40 nm (brown solid) and control device (black dash),
recorded under illumination at 100 mW/cm2.
4.4.2.7
Inverted structure bulk hetero-junction OPV Cells
The inverted organic BHJ solar cell consists of a sandwich structure, which incorporates a photoactive
layer deposited between two electrodes as shown in Fig. 4.4.2.15. Figure 4.4.2.16 shows the energy
band diagram of device in Fig. 4.4.2.15. The device fabrication process can be divided into four steps.
The first step is the substrate cleaning in which the ITO substrates are sonicated in di-water (15
minutes), acetone (15 minutes) and 2-propanol (15 minutes). Secondly, the sol-gel derived ZnO films
are prepared using zinc acetate (Zn(CH3COO)2•2H2O, Aldrich, 99.9%) and ethanolamine
(NH2CH2CH2OH, Aldrich, 99.5%) in 2-methoxyethanol (CH3OCH2CH2OH, Aldrich, 99.8%) as a
precursor solution. This solution was stirred for 1 hour in air. The precursor solution is then spun onto
the ITO substrate and subsequently treated at different annealing temperature. After the formation of
the sol-gel derived ZnO layer, the active layer (PCPDTBT:PC71BM, PTB7:PC71BM or other
polymers) which is stirred on a hotplate at 40 ˚C overnight in a N2 purged glove box is then cast onto
the substrate. Finally, a 10 nm buffer layer of MoO3 and 100 nm silver electrodes are thermally
evaporated to form the device. Morphology control techniques are used during this device fabrication
process, especially during the active layer solution preparation and the active layer drying process.
The effects of sol-gel derived ZnO layer upon the donor and acceptor morphology of the active layer
are also investigated.
- 140 -
Figure 4.4.2.15: Schematic diagram of inverted organic solar cell device structure.
Figure 4.4.2.16: Energy band diagram displaying the electronic structure of the inverted structure
solar cell (energy band values obtained from reference [4.4.2.15])
Figure 4.4.2.17: The EQE profile for PCPDTBT:PC71BM with different thickness sol-gel derived
ZnO layer, around 132 nm for 2000 rpm 30s, around 115 nm for 2000 ram 60s and around 13.5 nm
for 3000 rpm 30s
- 141 -
Figure 4.4.2.18: The EQE profile of the 2:7 PCPDTBT:PC71BM with and without ZnO buffer layer
with varied zinc acetate concentration ratio.
The EQE profile in Fig. 4.4.2.18 shows that the ZnO buffer layer is successfully formed and performs
well by increasing the zinc acetate dehydrate concentration from 0.02 g/ml to 0.1 g/ml. If we keep
increasing the concentration, the short current density will decrease mainly due to the increase of the
series resistance.
4.4.2.8
Organic/inorganic hybrid solar cells
The hybrid cells can be processed at room temperature and can be deposited on a large area and/or
flexible substrates. At present, one major challenge to hybrid solar cells remains their low conversion
efficiencies (up to 5.5 % [4.4.2.16]) and high cost of materials because of the limited use. Theoretical
efficiency of hybrid PV devices is limited up to 44% [4.4.2.17]. Hybrid solar cells are a mixer of
nanocrystals (NCs) of both organic and inorganic materials as shown in Fig. 4.4.2.19. Inorganic NCs
such as quantum dots (QDs), nano-tubes, etc. have also high absorption coefficients and particle size
induced tunability of the optical band-gap. When the inorganic NCs size become smaller than an
exciton (typically about 10 nm or smaller), their electronic structure will change than their bulk. The
electronic and optical properties of such small particles depend not only on the bulk material, also on
their size. Band-gap tuning in inorganic nanocrystals with different sizes can be used for realization of
device architectures, such as tandem solar cells in which the different bandgaps can be obtained by
modifying only one chemical compound. A substantial interfacial area for charge separation is
provided by nanocrystals, which has high surface area to volume ratios. Thus, the organic-inorganic
NCs hybrid concept for PV cells is getting interesting and attractive in recent years.
- 142 -
Figure 4.4.2.19: A Schematic structure of an organic-silicon NCs hybrid solar cell. Thin films of
PEDOT:PSS and Si NCs/P3HT were spun sequentially on transparent ITO coated glass substrate.
Aluminum metal electrode is deposited on the top. The energy band diagram of the structure is shown
in right [4.4.2.18].
There is a lack of information in the literature about the structure of the organic-NCs interface in
hybrid PV cells. The nature of the organic-NCs interaction and the extent of adsorption will affect
dispersion stability and also photoactive layer morphology. It will also play an important role in
charge transport across the organic-NCs interface. Good interfacial contact is essential in order to
maximise the interfacial area and minimise the series resistance across the cell. It is very important to
study on organic-NCs interface to enhance exciton dissociation and suppress carrier recombination.
Normally, free carriers are produced when the band offsets are greater than the exciton dissociation
energy. Controlling aggregation within organic/NCs/solvent mixture to obtain morphological
optimisation will improve the hybrid cells efficiency. Increasing the dispersion stability of
organic/NCs/solvent mixtures should improve the morphology of the photoactive layer and improve
the efficiency. A key challenge for organic-NCs hybrid PV cells is to control the aggregation within
the NCs phase in order to produce equilibrium morphologies to minimise the recombination losses.
The open circuit voltage (Voc) is plotted as a function of diagonal band-gap of different
inorganic/organic materials from different published works as shown in Fig. 4.4.2.20. The Voc is
directly proportional to the magnitude of the heterojunction diagonal band gap; however, there exists
an empirical loss factor related to the bulk heterojunction design. Detailed balance theory suggests
that maximal Voc will be obtained when recombination is exclusively radiative. Large luminescence
quenching in bulk heterojunction blend films implies that radiative recombination is only a small
fraction of total recombination, and thus, maximal values for Voc have not yet been obtained. The
experimental work suggests that the Voc is, in fact, related to the spectral position of the charge
transfer band, which is determined mainly by the energetics of the HOMO level of the donor and the
LUMO level of the acceptor in OPV device. This understanding, based on empirical observation, is
assumed to apply to hybrid solar cells, even though the constituent materials are different. A rigorous
exploration of this hypothesis, applied to hybrid solar cells, is currently lacking from literature.
Ideally, theoretical knowledge will be coupled with experimental data to build a comprehensive
model to explain the origin of Voc in hybrid solar cells.
- 143 -
Figure 4.4.2.20: Plot of open circuit voltage (Voc) as a function of diagonal band gap of
organic/inorganic materials. Doted points are for different inorganic nanocrystal materials in organic
donor materials.
4.4.2.9
TiO : polymer solar cell
2
A conjugated polymer is combined with the wide-bandgap semiconductor metal-oxide (TiO2 or ZnO)
to form the active layer in these devices. Light absorption is achieved in the polymer film while
electron transport is carried out by the inorganic layer. The device combines the advantages of the
organic and inorganic components, allowing low-temperature solution-processed fabrication,
improved stability compared to all-organic solar cells, good electron transport, and the ability to
control the nanostructure. Even though this device type exhibits great potential, performance levels
are low as this field is still young and further research effort is required for optimisation. The
knowledge gained through research of these devices is also of use for tandem solar cells (The world
record efficiency of 6.5% for a solution processed organic tandem solar cell combines similar hybrid
layers [4.4.2.19]).
Figure 4.4.2.21: Schematic band structure diagram of TiO : polymer solar cell.
2
- 144 -
4.4.2.10 Simulation Works
The Centre has purchased an upgrade of the Gaussian DFT program package to Gaussian09. We can
model larger systems in the so-called ONIOM model which consists of up to three model shells of
different accuracy. The highest accuracy is limited to the molecule of interest, while the lower
accuracy regions are bigger and cover the electronic environment of the species under investigation.
This can be neighbouring molecules for charge transfer as well as electrodes. An improved tool for
calculating infrared- and Raman-spectra is included as well, providing an important link to
experimental observations at the organic species. The present strengths include a dedicated ab-initio
molecular simulation capability (Fig. 4.4.2.22) presently rated at 650 gigaflops but with upgrading to
1.6 teraflops planned for the near-term. To date, this facility has been used for ab-initio modelling of
silicon quantum dots in a dielectric environment. The extended capability for excited state modelling
should allow the screening of new candidate OPV materials for both performance and durability.
Figure 4.4.2.22: Dedicated Linux 64 bit cluster for ab-initio molecular simulations presently rated at
1080 gigaflops.
Other strengths include past experience with a number of engineered vertical junction structures and
with light-trapping, not yet fully exploited in OPV devices. The Centre has pioneered light-trapping in
both wafer-based and silicon thin-film devices, as well as the use of plasmonics for light-trapping in
thin, plane-parallel photovoltaic structures [4.4.2.20].
The Centre is also particularly interested in hybrid organic/inorganic systems as a way of improving
both performance and stability. Some work in this area has already commenced as reported below.
Hybrid solar cells are a mixer of nanostructures of both organic and inorganic materials. They
combine the unique properties of inorganic semiconductor nanoparticles with properties of
organic/polymeric materials. Inorganic semiconductor nanoparticles or quantum dots may have high
absorption coefficients and particle size induced tunability of the optical band-gap. Band-gap tuning
in inorganic nanoparticles with different nanoparticle sizes can be used for realization of device
architectures, such as tandem solar cells in which the different bandgaps can be obtained by
modifying only one chemical compound. Thus, the organic/inorganic hybrid concept for photovoltaic
solar cells is getting interesting and attractive in recent years. The solubility of the n-type and p-type
components is an important parameter of the construction of hybrid solar cells processed from
solutions.
- 145 -
4.4.2.11 Ordered nanoparticle arrays for hybrid organic/inorganic solar cells
Experimental work on fabrication of highly ordered arrays of nanoparticles as absorber materials
originally for hot carrier solar cells has been initiated as a potential means to realise ordered
superlattice structures. In this work, the aim is to establish a fabrication system for depositing
sequential monolayers of nanoparticles with uniform shells.
We have installed a Langmuir-Blodgett (LB) system for fabrication of highly ordered nanoparticle
monolayers as shown in Fig. 4.4.2.23. The LB technique leads to the development of ordered
monolayers at an air-water interface while exploiting the self-organization mechanism of colloidal
dispersion. Compression of the monolayer is monitored via measurements of surface pressure and
then controlled by a feedback loop. This unique technique allows transfer of this ordered monolayer
onto a wide range of solid substrates such as glass or Si wafers. By controlling the interspacing
between adjacent particles, i.e. the shell thickness, by varying the molecular weight of capping
species, we can control the periodicity of the film - leading to new optical and electrical properties.
substrate
Figure 4.4.2.23: Schematic of the LB apparatus used to fabricate a monolayer of encapsulated
nanoparticles.
Silicon (Si) nanoparticles are being used as core materials. In order to control the interspacing
between the particles, the termination of Si nanopaticles is carried out using organosilanes of varying
alkyl chain lengths.
Progress to date is reported elsewhere [4.4.2.21]. Once the assembly approach is mastered, it should
be possible to build up device structures incorporating layers of quantum dots with the doping in each
layer individually controlled.
References
4.4.2.1
4.4.2.2
4.4.2.3
4.4.2.4
4.4.2.5
4.4.2.6
4.4.2.7
Martin A. Green, Keith Emery, Yoshihiro Hishikawa, Wilhelm Warta and Ewan D.
Dunlop, “Solar cell efficiency tables (version 39)”, Prog. Photovolt: Res. Appl. 20 (2012)
p12–20.
M.A. Green, K. Emery, Y. Hishikawa and W. Warta, “Solar Cell Efficiency Tables”,
Versions 26-34, Progress in Photovoltaics, 2006-2009.
Martin A. Green, Keith Emery, Yoshihiro Hishikawa, Wilhelm Warta and Ewan D.
Dunlop, “Solar cell efficiency tables (version 41)”, Prog. Photovolt: Res. Appl. 21 (2013)
p1–11.
W. Shockley and H. J. Queisser, “Detailed Balance Limit of Efficiency of p‐n Junction
Solar Cells”, J. Appl. Phys. 32 (1961) p510.
Thomas Kirchartz, Kurt Taretto, and Uwe Rau, “Efficiency Limits of Organic Bulk
Heterojunction Solar Cells” J. Phys. Chem. C113 (2009) p17958 – 17966.
Huang, S., et al. Fabrication of nanorod arrays for organic solar cell applications. Materials
for Photovoltaics. Symposium, 29 Nov.-2 Dec. 2004. 2005. Warrendale, PA, USA:
Materials Research Society.
T. Erb, U. Zhokhavets, G. Gobsch, S. Raleva, B. Stühn, P. Schilinsky, C. Waldauf, and C.
J. Brabec, Adv. Funct. Mater. 15, 1193 (2005).
- 146 -
4.4.2.8
4.4.2.9
4.4.2.10
4.4.2.11
4.4.2.12
4.4.2.13
4.4.2.14
4.4.2.15
4.4.2.16
4.4.2.17
4.4.2.18
4.4.2.19
4.4.2.20
4.4.2.21
Xiaohan Yang, Ashraf Uddin, and Matthew Wright, “Effects of Annealing on Bulkheterojunction Organic Solar Cells”, Adv. Sci. Eng. Med., 4 (2012) p1–7.
L.S. Roman, O. Inganäs, T. Granlund, T. Nyberg, M. Svensson, M.R. Andersson, J.C.
Hummelen, Advanced Materials 12 (3) (2000) p189–195.
M. Niggemann, M. Glatthaar, A. Gombert, A. Hinsch, V. Wittwer, Thin Solid Films 451–
452 (2004) p619–623.
K. Tvingstedt, V. Andersson, F. Zhang, O. Inganas, Applied Physics Letters 91 (2007)
p123514.
Xiaohan Yang , Ashraf Uddin, Matthew Wright, “Plasmon enhanced light absorption in
bulk-heterojunction organic solar cells”, Phys. Status Solidi RRL 6 (2012) p199–201.
P. Spinelli, C. van Lare, E. Verhagen and A. Polman, Opt. Express, 2011, 19, A303-A311.
F. Zhou, Z.-Y. Li, Y. Liu and Y. Xia, The Journal of Physical Chemistry C, 2008, 112,
p20233-20240.
Sun, Y., J.H. Seo, C.J. Takacs, et al., Inverted polymer solar cells integrated with a lowtemperature-annealed sol-gel-derived ZnO Film as an electron transport layer. Adv Mater,
2011. 23(14): p1679-83.
B.R. Saunders and M.L. Turner, “Nanoparticle-polymer phovoltaic cells” Adv. Coll.
Interface Sci., 138, 1(2008).
I. M-Sero, S. Gimenez, F. F. Santiago, R. Gomez, Q. Shen, T. Toyada and J. Bisquert,
“Recombination in quantum dot sensitized solar cells,” Accounts of Chemical Res., 42,
1848(2009).
C.Y. Liu, Z.C. Holman, and U.R. Kortshagen, “Hybrid Solar Cells from P3HT and Silicon
Nanocrystals” Nano Lett. 9, 449 (2009).
J.Y. Kim, et al., Efficient Tandem Polymer Solar Cells Fabricated by All-Solution,
Science, 317:222-225, 2007.
S. Pillai, K.R. Catchpole, T. Trupke, and M.A. Green, “Surface Plasmon Enhanced Silicon
Solar Cells”, Journal of Applied Physics 101 (2007) p093105.
L. Treiber, C. Bumby, S. Huang, G. Conibeer, 23rd European Photovoltaic Solar Energy
Conference, Valencia, Spain (2008).
- 147 -
4.4.3 Earth-Abundant CZTS Photovoltaic Devices
University Staff
Dr Xiaojing Hao (group leader)
Prof. Martin Green
Prof. Gavin Conibeer
Research Fellows
Dr Shujuan Huang
Dr Hongtao Cui
Dr Fangyang Liu
Postgraduate Research Student
Xiaolei Liu
Ning Song
Jian Chen
Chang (Eric) Yan
Lei (Adrain) Shi
Visiting Fellow
A/Prof Zhiping Zheng (Huazhong University of Science & Technology, China, till Dec 2012)
Dr Zhanxia Zhao (Shanghai University, China, from Sept 2012)
4.4.3.1
Strands Outline
All successfully commercialised non-concentrating photovoltaic technologies to date are based on
silicon or the chalcogenides (semiconductors containing Group VI elements, specifically Te, Se and
S). As indicated by Figure 4.4.3.1, the successful chalcogenide materials, CdTe and CuInSe2, can be
regarded as “synthetic silicon” where the balance between atoms in these materials provides the same
average number of valence band electrons as in silicon.
Figure 4.4.3.1: Synthetic silicon-Portion of the Periodic Table showing pathways to engineering
semiconductors with 4 valence electrons/atom.
Unfortunately Cd and Se are toxic “heavy metals” while Te and In are amongst the 12 most scare
elements in the Earth’s crust, factors that would seem to limit the long-term potential of the
established chalcogenide technologies. However, as indicated in Figure 4.4.3.1, by investigating more
deeply into the Periodic Table, the compound Cu2ZnSnS4 is uncovered with the same number of
valence band electrons on average but involving earth-abundant, non-toxic elements.
- 148 -
Although the potential of this material is relatively unexplored for photovoltaics, initial results have
been promising with a group at IBM reporting 11.1% efficiency for small cells based on related
materials (alloy of the above sulphide and the corresponding selenide). The appearance of the cells
involved (Figure 4.4.3.2) is quite close to the appearance of cells made using CIGS technology
(CuInSe2 plus CuGaSe2 alloy).
Figure 4.4.3.2: SEM cross-section of
CZTS-type cells [4.4.3.1].
The Centre’s work in this new strand of activity takes a different direction from most of the present
international work in this area. The starting point is CZTS thin film solar cells through a viable high
throughput manufacturing process. Efforts are to develop technology suitable for the desired process
rather than persevering with solution growth processes that have given the best laboratory results to
date. Materials selection would be guided by stable compositions formed in naturally occurring
minerals such as kesterite and stannites. Another key area that is to be investigated in detail is
epitaxial relationships to silicon to investigate the suitability for tandem cell stacks, including stacks
involving silicon as the lowermost layer.
4.4.3.2
CZTS Thin Film Solar Cells
The CZTS thin film research in 2012 is focused on separately
fabricating and optimizing materials for different functional layers
as shown in Figure 4.4.3.3. The investigation on thses functional
layers was summarized as follows.
4.4.3.2.1 Back contact-Mo
The properties of low resistivity and good adhesion to substrate are
pursued for back contact of CZTS solar cells. The strong adhesion
of Mo thin films to soda lime SLG (SLG) substrate is crucial for the
subsequent deposition of CZTS thin films. The adhesion tape test
complying with the standard ASTM D3359-09 was carried out for
such purpose. It was found that the adhesion of Mo film to SLG
substrat could be improved with increasing pressure. This is mainly
due to the transition from compressive to tensil stress in such Mo
Figure 4.4.3.3: Schematic of
thin film since better adhesion is favoured under tensile stress.
CZTS solar cells on SLG.
However, the higher pressure giving good adhesion would also
result in higher resistivity. This indicates that the Mo films deposited at a certain Argon pressure
cannot satisfy these two goals simultaneously. Mo Films deposited at higher Argon pressure show
good adhesion to SLG substrates but high resistivity. In contrast, Mo Films deposited at lower Argon
pressure exhibit low resistivity but suffer from poor adhesion. In order to obtain both advantages of
low resistivity and good adhesion to SLG, a double layer scheme was proposed and investigated,
where the bottom layer was deposited at higher Argon pressure for strong adhesion to SLG and the
top layer was deposited at lower Argon pressure for low resistivity.
- 149 -
The double layer Mo film was deposited on SLG substrates at room temperatire without intentional
substare heating. Argon pressure was varied during the deposition, with pressure of 10-20 mTorr at
the beginning and 2mTorr for the rest of the deposition. For further improving the resistivity, one of
the as-depsited Mo film was annealed at 550°C for 5 min for a comparison. As shown in Figure
4.4.3.4, all Mo double-layer thin films exhibite preferred orientation of (111). It indicated that it is the
original deposition conditions that largely determins the Mo film ctystallinity though annealing can
improve the crystallinity as well.
Figure 4.4.3.4: XRD spectra of Mo double layer
on SLG.
Figure 4.4.3.5: Resistivity of Mo double
layer on SLG.
Similarly the resistivity, shown in Figure 4.4.3.5, is significantly dependent on deposition conditions
but can be further decreased after post-deposition annealing. The lowest resistivity of 12.1 μΩcm was
obtained after annealing the double layer with the pressure combination of 20 mTorr and 2 mTorr for
bottom and top layers respectively, which might be due to the better crystallinity obtained under such
a condition. Moreover, the best adhesion category of 4B was achieved on this Mo sample, which
means less than 5 per cent of test area was removed. The surface morphology, also important for
subsequent CZTS absorber growth was examined by using AFM. The surface roughness of Mo film
increases with pressure applied in the Mo bottom layer deposition as shown in Figure 4.4.3.6. The Mo
double layer with the pressure combination of 20 and 2 mTorr for bottom and top Mo layers
respectively, meets the requirement for the back contact. However, the effect of the surface roughness
on the properties of subsequent growth of CZTS still needs further investigation.
a) 2/10 mTorr, Rq=1.96 nm
b) 2/20 mTorr, Rq=3.36 nm
Figure 4.4.3.6: Surface morphology of Mo double layer on SLG.
- 150 -
4.4.3.2.2 CZTS absorber
Fabrication of CZTS thin films by sulphurisation of SLG/Zn/Sn/Cu precursor layers
CZTS polycrystalline thin film can be formed by one-stage PVD method, i.e., reactive sputtering or
co-evaporation, or by two-stages of CZTS precursor (including CZT and CZTS) followed by
subsequent sufurisaton annealing.
As for the two-stages method, we started with the investigation of CZT in the structure of stacking
layers. The effects of the order of stacking layers and sulpurisation condition on the properties of asdeposietd and annealed films were investigated. Precursor stacking layers of SLG/Zn/Sn/Cu were
sputtered on the SLG substrates, followed by the sulphurisation annealing at various annealing
conditions. The chemical composition of the as-deposited and annealed films were analysed by LAICP, shown in Table 4.4.3.1. The desired composition for reported high efficiency CZTS solar cells is
listed as a reference. It indicates that the film suffers from a significant loss of Zn and Sn during postdeposition annealing process. A Cu poor as-deposietd stacker is necessary to compensate the loss of
Zn and Sn in order to obtaine the desired chemical compositon. In addition, S content in the anneald
film seems to be significantly lower than the desired amount. This might be explained by the
formation of Cu2S, where the ratio of S over Cu is 0.5. The presence of CuxS (1<x<2) can be
evidenced by the XRD pattern shown in Figure 4.4.3.7 and Raman spectra in Figure 4.4.3.8. In fact,
CuxS clusters can be found sitting on the surface of annealed sample. If focusing the laser beam on
such clusters, a strong CuxS peak is presented as shown in Figure 4.4.3.8.
o
550 C, 5min
60000
CuS (475)
50000
intensity (counts)
40000
30000
20000
10000
CZTS (337)
CZTS (287)
0
200
250
300
350
400
450
500
550
-1
Raman Shift (cm )
Figure 4.4.3.7: XRD spectra of CZTS thin films
sulfurisation-annealed at different conditions.
Figure 4.4.3.8: Raman spectra of CZTS thin films
sulfurization-annealed at different conditions.
Table 4.4.3.1: Chemical composition of as-deposited & annealed thin films.
Unit
As-deposited
Annealed SLG/Zn/Sn/Cu under S Desired
SLG/Zn/Sn/Cu
vapour at 550 C 30 min
ratio
Zn/Sn
Molar
2.445
1.313
1.2-1.3
ratio
Cu/(Zn+Sn)
Molar
0.6001
0.99
0.8-0.9
ratio
S/(Cu+Zn+Sn) Molar
0.0083
0.3888
1
ratio
Figure 4.4.3.9 shows the SEM micrograph of the surface of annealed stacked layer of SLG/Zn/Sn/Cu.
The surface is quite rough and many voids are presented. EDS mapping in Figure 4.4.3.10 indicates
that the clusters sitting on the surface are Cu-rich which is speculated as CuxS. In addition, Si rich
region presented in Figure 4.4.3.10c indicates the significantly high porosity. To understand the origin
of surface roughness and the voids formation mechanism, the stacking order was examined. Figure
4.4.3.11 shows the SEM micrograph of the surface of the stacked layer of Sn/Zn/SLG. We can see the
- 151 -
islans and spheroid of Sn on the relatively smooth surface of Zn, and the size of islans and spheroid
increases with the power applied on the Sn target. Many voids and rough surface are expected with
the deposition of Cu on the top of Sn/Zn/SLG. From these results of SEM observation, it was found
that the precursor order significantly affects the surface morphology of CZTS film. Particularly, the
surface roughness was found highly dependent on the Sn layer. From the results of these preliminary
examinations, the new type of stacking layer is proposed and under investigation.
Figure 4.4.3.9: cross section SEM of CZTS thin film
sulphurised at 550 °C for 30 min.
Figure 4.4.3.10: EDS mapping of CZTS thin films sulphurised at 550 °C for 30 min.
SLG/Zn
SLG/Zn/Sn(2.5W)
SLG/Zn/Sn(30W)
SLG/Zn/Sn(9W)
SLG/Zn/Sn(30W)/Cu
Figure 4.4.3.11: surface morphology of metallic stacking layers.
- 152 -
Simplified fabrication of CZTS thin film solar cells
Though various approaches have been applied to fabricate CIGS solar cell, CIGS with the highest
efficiency (up to 20.4% so far) have been achieved by the physical vapour deposition (PVD)
approach, with the first 10% efficient cell demonstrated in 1980, the first 15% cell in 1993 and the
first 20% cell in 2008. As a substitute for CIGS thin film solar cells, CZTS has many similarities to
CIGS as mentioned above. As such, it is logical that the full potential of high efficiency CZTS thin
film solar cells will be extracted using the PVD approach.
Though a major milestone was reached at the end of 2009 with the demonstration of 9.7% CZTS
efficiency (some Se also in device) by a research team at IBM by a “hydrazine-based solution”
approach [4.4.3.1], solution growth is believed inherently incapable of giving the material control
required for best-possible performance. Evidence for this arises in the similar case of CIGS where the
same IBM team held the record of 12.2% efficiency for solution-deposited CIGS (again hydrazine–
based) [4.4.3.2], while in contrast the PVD approach gives 20.3% efficiency.
Evaporation and sputtering are the two most important PVD techniques. The disadvantage of
evaporation is that it is very difficult to control the Cu evaporation source and the lack of
commercially available equipment for large area thermal evaporation. In contrast, magnetron
sputtering is a more attractive process because large area/high rate sputtering is a better developed
technique than evaporation. In addition, sputtering has commercially available equipment that gives
good uniformity over large areas. UNSW has considerable experience in PVD techniques, such as the
achievement of the first silicon quantum dot solar cell on quartz [4.4.3.3], a recent patent relating to
crystal Ge on Si wafer by sputtering technique and the polycrystalline silicon thin film solar cell on
SLG by an evaporation technique [4.4.3.4].
Almost all of the methods that have been attempted for CZTS fabrication involve two steps;
deposition of a metal/metal-sulfide precursor and a subsequent anneal at high temperature in a
hydrogen sulphide or elemental sulphur environment to incorporate more sulfur into the film and
evolve the correct phase [4.4.3.5-4.4.3.7]. It is reported that the H2S/sulphur anneal is a destructive
process that can potentially introduce defects into the film due to the out-diffusion of metals. In
addition, for large scale production of solar cells, a direct one-step, large area deposition process is
preferred.
Fortunately, reactive sputtering is a well suited technique because of the above-addressed feature of
precise control on the stoichiometry of deposited films over a large area with high uniformity at a
relatively low cost. The experience documented in our patent of “A Method of Forming A
Germanium Layer On A Silicon Substrate And A Photovoltaic Device Including A Germanium
Layer”, where the introduction of H2 gas into the in-situ deposited germanium thin film by magnetron
sputtering results in the significant improvement of the performance of germanium thin film, appears
highly relevant.
As such, the idea of one-step reactive sputtering by in-situ introducing H2S gas into the deposited
CZT (Cu2ZnSn) film is a logical step to attempt a better CZTS film quality. However, there have been
few attempts to date at such a simplified one-step process by introducing the sulphide into the CZTS
film during in-situ deposition. In addition, this process is not well understood and significantly more
effort is required in order to understand growth mechanisms involved in the reactive sputtering
process. Comparison of reactively sputtered films with 2-step sulphidisation processes will provide
the early emphasis of the CZTS strand’s work.
- 153 -
4.4.3.2.3 Cd-free buffer layer
Alternative Cd-free buffer by dry fabrication methods are required for reducing production and
recycling costs and the toxicity of the involved materials, as well as easily integrated into an in-line
process. Non-toxicity will ensure that the future progress of technology is not impeded by the
increasingly stringent legislation relating to the use and disposal of toxic material. Thus the potential
of Cd-free as an alternative to CdS buffer assumes greater significance under this scenario. Following
the trace of the development of Cd-free devices for CIGS solar cell, it flagged in 1992 with an
efficiency of about 9% has progressed to the current status of 18–19.9% efficiency. Comparing with
CdS buffer layer, the initial difference of 4% absolute has been erased to the current negligible level.
The primary function of a buffer layer in a heterojunction is to form a junction with the absorber layer
while admitting a maximum amount of light to the junction region and absorber layer. As thus, this
layer should have (1) large energy band gap for minimal absorption losses; (2) a suitable conduction
band line-up to the absorber and to the i-ZnO. Based on the understanding of record CIGS cells, a
moderate spike is preferred for conduction band offset for reducing interface cross-recombination; (3)
a very low defect density at the absorber/buffer interface by less lattice mismatch at the junction for
consideration for epitaxial or highly oriented layers; (4) desired doping density sufficiently exceeding
that of the absorber to confine the space charge region at absorber; (5) Position of the Fermi level at
the absorber/buffer interface close to the absorber conduction band, i.e., achieve type inversion of the
absorber surface. For industrial applications, further important criteria must be taken into account,
including production costs (e.g., deposition time, compatibility with other process steps and waste
management); technological feasibility (e.g., development of suitable large-area deposition systems);
reproducibility and long-term device stability. The production cost of compatibility with other process
steps mainly implies that the buffer layer process must be a dry vacuum-based process as discussed
above, particularly promising options are ALD, sputtering and evaporation. The promising Cd-free
buffer materials focused in our research are Zn related compounds, for example ZnS, Zn(Ox,S1-x) and
ZnxMg1-xO.
4.4.3.2.4 Al:ZnO (AZO) window layer
Among transparent conducting oxides, AZO is the most promising material for window layers in thin
film solar cells and modules because of its wide band gap, high conductivity, relatively low cost and
the non-toxicity of the material [4.4.3.8]. Many techniques have been employed to fabricate AZO
films. However, magnetron sputtering is the most promising, permitting deposition at low
temperature, and giving better adhesion on the substrate and higher film density than other methods
[4.4.3.9]. For its application in CZTS solar cells, AZO needs to be also highly transparent in UV
range and near IR range as well. The latter requires the AZO has high mobility to remain the low
resistivity since in the near IR range free carrier absorption dominants. In order to prepare AZO thin
films with low resistivity and high transparency for its application in CZTS thin film solar cells, the
influences of the thickness and substrate temperature on the structural, electrical, and optical
properties of AZO films were investigated.
AZO films deposited in Ar atmosphere
The AZO films were deposited on quartz substrates by sputtering a 2 wt.% Al2O3 doped ZnO target.
In order to minimize the plasma bombardment damage on the substrates, a lower RF power was
applied on the target. As the film is thicker than 140 nm, TEM and AFM images indicated that the
growth mode changed from vertical to lateral. A lateral growth mode is beneficial for forming high
quality films. The growth mode transition maybe due to the decreasing stress within the grains as the
film thickness increases. Bulk resistivity, carrier concentration and Hall mobility changed
accordingly, as displayed in Figure 4.4.3.12.
- 154 -
Bulk resistivity
Carrier concentration
Mobility
20
1.4x10
-3
3.5x10
20
1.2x10
-3
3.0x10
-3
2.5x10
20
1.0x10
-3
2.0x10
19
8.0x10
-3
1.5x10
6.0x10
-3
1.0x10
100
200
300
400
19
34
32
30
28
26
24
2
-3
Carrier concentration (cm )
4.0x10
Bulk resistivity (cm)
36
20
1.6x10
-3
Mobility (cm /Vs)
-3
4.5x10
22
20
18
500
Thickness (nm)
Figure 4.4.3.12: Bulk resistivity, carrier concentration and Hall mobility of AZO films as a function
of film thickness.
100
10
4x10
90
80
50
40
30
20
10
3x10
551 nm
443 nm
335 nm
224 nm
143 nm
67 nm
Substrate
a2 (cm-2)
Transmittance(%)
60
551 nm
443 nm
335 nm
224 nm
143 nm
67 nm
10
70
10
2x10
10
1x10
0
3.0
0
200 400 600 800 1000 1200 1400 1600 1800 2000
3.4
3.6
3.8
4.0
4.2
4.4
Hv(eV)
Wavelength(nm)
Figure 4.4.3.13: Optical transmission of AZO
films as a function of film thickness.
3.2
Figure 4.4.3.14: Square of the absorption coefficient as
a function of photon energy for AZO thin films with
different thicknesses.
- 155 -
Figure 4.4.3.15: TEM image of AZO film with thickness of 551 nm.
Figure 4.4.3.16: AFM images for films with different thickness: (a) 67nm (b) 335nm (c) 551nm.
The average transmission is higher than 85% for all AZO layers in the wavelength range of 390 to
1100 nm, as shown in Figure 4.4.3.13. The obvious drop in transmission of the IR region as film
thickness increases is mainly due to the increment of free carrier absorption. Such transmission drop
can be reduced by decreasing the free carrier concentration while increasing the mobility without
sacrificing the desired resistivity. Figure 4.4.3.14 shows that the Eg of AZO film increases from 3.4 to
3.55 eV with its thickness increasing from 143 to 551 nm, which can be explained by the BursteinMoss shift.
The deposion temperature of AZO is critical that it needs to be high enough to yield good crystallinity
and low resistivity of AZO while having no detrimental effect on the underlying CZTS absorber and
thereby formed heterojuntion. AZO films were prepared under different substrate temperatures, from
room temperature (RT) to 450°C. To ensure high transmission and low resistivity, the film thickness
of around 330 nm was chosen. Other deposition parameters were retained the same as previously
detailed. The minimum resistivity was obtained from the sample prepared at 250°C, as shown in
Figure 4.4.3.17. The electrical properties of the film deposited at RT were beyond the detection limit
and therefore not shown here. The electrical properties of the film are improved as the substrate
temperature increases from RT to 450°C. This may be explained by increased grain size and thus
reduced grain boundaries with an increase in the substrate temperature. However, the electrical
qualities deteriorate at temperature above 250°C. This can be expected since the grain boundary
barrier was known to increase with the substrate temperature above 200°C [4.4.3.10].
Figure 4.4.3.18 shows the optical transmission of AZO films prepared under various substrate
temperatures. The average transmission was higher than 85% for all AZO layers in the wavelength
- 156 -
range of 390 to 1100 nm. The lowest transmittance in the IR region is obtained from the sample
prepared at 250°C which was attributed to the highest free carrier concentration. Correspondingly, as
we can see in Figure 4.4.3.19, this sample had the largest Eg of 3.5 eV as a result of the BursteinMoss shift. Therefore, 250°C is close to the optimum substrate temperature that allowing the
synthesis of high quality AZO films and in the meanwhile will not decompose the underneath CZTS
absorber. However, the mobility needs to be improved to reduce the free carrier concentration to
allow for retaining the desired low resistivity.
1.4x10
-3
1.2x10
-3
1.0x10
-3
8.0x10
-3
6.0x10
20
4.0x10
20
3.0x10
19
2.0x10
19
1.0x10
150
200
250
300
350
400
32
28
24
20
2
Bulk resistivity ( cm)
20
-3
5.0x10
100
36
20
1.6x10
-3
6.0x10
Carrier concentration (cm-3)
-3
7.0x10
Mobility (cm /Vs)
Bulk resistivity
Carrier concentration
Mobility
16
450
o
Temperature( C)
Figure 4.4.3.17: Electrical properties of AZO films as a function of substrate temperature during
deposition.
100
o
22 C
o
150 C
o
250 C
o
350 C
o
450 C
10
8x10
10
6x10
60
-2
 cm 
RT
o
150 C
o
250 C
o
350 C
o
450 C
40
20
10
4x10

Transmittance(%)
80
10
2x10
0
400
800
1200
1600
0
2000
2.0
Wavelength(nm)
Figure 4.4.3.18: Transmission of AZO films
as a function of substrate temperature during
deposition.
2.5
3.0
3.5
4.0
4.5
5.0
hv (eV)
Figure 4.4.3.19: Optical band-gap of AZO films
as a function of substrate temperature during
deposition.
AZO films deposited in H2 & Ar atmosphere
The AZO films were deposited on quartz substrates by sputtering a similar 2 wt.% Al2O3 doped ZnO
target. The deposition processes were conducted in an atmosphere of Ar:H2=95:1. The substrate
temperature was retained at 200 °C. H2 added to the sputtering ambient could reduce absorbed O- on
ZnO grain boundaries by reacting with O- to form complexes. This can return carriers to conduction
band and thus increase the measured Hall mobility. H2 addition also enables environment to form Vo
and Zni. In ZnO, only H+ is stable, and hence hydrogen exclusively acts as a donor [4.4.3.11]. AZO
films with thickness of 283 nm and 120 nm were deposited to compare with that of NREL’s used in
their high efficiency CIGS solar cells. As shown in Table 4.4.3.2, both of our films have higher
mobility, lower carrier concentration and lower resistivity than that of NREL’s. High carrier
- 157 -
concentration results in significant IR absorption, therefore, high TCO conductivity is best achieved
by increasing mobility rather than carrier concentration for photovoltaic device application.
Table 4.4.3.2 Comparison with NREL’s Al:ZnO films used in high efficiency CIGS solar cells.
Thickness Resistivity
(nm)
(Ωcm)
Carrier
concentration
-3
Mobility
Bandgap
(eV)
57.9
Average
Transmission
(390nm1100nm)
86.9
43
87.9
3.72
2
(cm /Vs)
(cm )
UNSW
283
-4
2.7×10
-4
UNSW
123
5.7×10
NREL
120
8.5×10
20
4.4×10
20
2.6×10
-4
20
3.75
13
5.6×10
4.4.3.2.5 i-ZnO buffer layer
The Voc can be improved by having higher bandgap at the interface and by reducing shunt paths.
Anant H [4.4.3.12] reported improved Voc of CIGS cells by reducing shunts path via optimising the
thickness of i-ZnO layer. Too thin i-ZnO (including without i-ZnO) will have regions of high bandgap
discontinuity leading to overall low open circuit voltage. On the other hand, too thick i-ZnO could
weaken the built-in field by spreading the space charge region over the thickness of i-ZnO thus
reducing the efficiency. We choose a thickness of 50 nm for i-ZnO layer. The i-ZnO films were
deposited on quartz substrates by sputtering an intrinsic ZnO target. The deposition processes were
conducted in an atmosphere of Ar:H2=95:1. The substrate temperature was retained at 200°C. The
resistivity of this i-ZnO layer is 2.24×10-2 Ωcm and the transmission is 83.6% due to a low bandgap of
3.28 eV. The i-ZnO and AZO bi-layers were designed and deposited with parameters as
aforementioned. The resistivity increased from 3.7 ×10-4 Ωcm of 300nm thick AZO layer to 9.9×10-4
Ωcm of 50nm thick i-ZnO layer/300nm thick AZO bi-layers. Further investigation needs to be carried
out to optimise the i-ZnO and AZO bi-layers.
4.4.3.2.6 Matallization
Metallisation is important for high efficiency CZTS solar cells to minimise the resistive loss and
shadow loss and extract power from the solar cells. According to the procedure described in ref.
[4.4.3.13], the top contact of the CZTS solar cells has been designed as shown in Figure 4.4.3.20,
based on Jmp, Vmp values reported in the literature [4.4.3.14-4.4.3.16] and from UNSW’s
collaborator as well as the sheet resistance of UNSW’s AZO and evaporated metal contacts. The top
mask is for defining the cell area and the bottom mask is for defining the top contact.
Figure 4.4.3.20: Metallisation mask for CZTS
device fabrication.
- 158 -
4.4.3.3
CZTS/Si Multijunction Solar Cells
The highest efficiency solar cells achieved so far is by using single crystal III-V semiconductor
homojunctions grown on Ge substrates to form monolithic three-junction cells. All three junctions
must be approximately lattice-matched to minimize crystal imperfections, which will reduce minoritycarrier recombination times and hence cell efficiencies. The lattice matching requirement limits the
choice of bandgaps and thus the efficiencies, although efficiencies over 41% have been achieved
under high light concentrations. In addition, their high cost has limited the applications of these cells
to space and to high concentration photovoltaic systems. The high cost of high concentration PV
system itself has been a substantial drawback. In contrast, Si is an optimum choice as the bottom cell
for tandem application. For a single cell, Si record cell (25% at UNSW) has already achieved 86% of
the silicon limit (29%). Until now silicon wafer based cell still accounts for about 90% of the solar
cell market and will continue to dominate the PV market in the future. Any concept, allowing
significant boost to higher efficiency and lower cost, and fully making advantage of and being
compatible with commercialized Si wafer cell technology, would be a natural “add-on” to Si wafer
based PV. The concept of a tandem cell with Si as bottom cell offers the best opportunity to boost the
efficiency beyond 30% or even 40%. This suggests the use of an absorber system that is lattice
matched with Si and also flexible in tuning to the desired bandgap combination for tandem cells.
CIGS has been shown to have a large range of bandgap tunability. However, a lattice matched
substrate suitable for their growth acting as bottom cell is absent. On the other hand, even though
CIGS based chalcopyrite devices have been reported as the highest efficiency (20.4%) in the thin film
solar cell area, efficiencies are still lower as compared to single crystalline solar cells based on Si
(25% at UNSW) or GaAs (29% under unconcentrated light). The question arises how far the
polycrystalline nature of these materials limits their efficiencies. Single crystalline materials, obtained
by single crystal or epitaxial growth would bear the possibility to avoid the disadvantageous influence
of grain boundaries and interfaces in this compound. CZTS, as the most promising substituting
material for CIGS, is suitable for epitaxial growth on Si due to very low lattice mismatch and also
feasible for the tandem cell structure due to its large tunable bandgap range, spanning a wider range in
energy beyond 2.25 eV which is above the accessible range for the highest efficiency III-V top
junction materials. Therefore this will allow the exploration of both CZTS family/Si based tandem
cells by taking advantage of single crystalline CZTS and mature Si wafer cell technology.
4.4.3.3.1 CZTS on Si substrate
CZTS on Si with different orientations
Intensity (a.u.)
Hetero-epitaxial growth is a highly desirable approach for achieving high quality single-crystalline or
preferred-orientation materials. One-stage fabrication is necessary for such epitaxial growth. Si and
ZnS/Si were investigated as alternative
(112)
70000
substrates to engineer the crystalline
properties of CZTS.
CZTS on Si (100)
60000
o
Orientated CZTS film was obtained by using
CZTS on Si (100) 4
50000
CZTS on Si (111)
these substrates. This allows for alternative
o
CZTS on Si (111) 4
paths for forming high quality crystalline
40000
CZTS
on
Si
(110)
CZTS absorber with desired orientation.
30000
About 400 nm CZTS thin films were
20000
(101)
(220)
deposited on different orientated Si wafers by
(200)
(312)
10000
sputtering of a single CZTS target. The
(008) (332)
0
orientation of Si wafer used for substrates are
20
30
40
50
60
70
80
(100), (100)4, (111), (111)4 and (110). It
2deg
was found in Figure 4.4.3.21, that the XRD
Figure 4.4.3.21: XRD patterns of CZTS films
patterns of the thin films show the preferred
on Si substrates with different orientations.
orientations of the (112), (220) depending on
- 159 -
the substrate orientation. Here we use the intensity ratio of I220/I112 to relatively compare the
variation of preferred orientation upon various substrates. As shown in Table 4.4.3.3, the maximal
ratio of I220/I112 of CZTS thin films is obtained on Si (110) substrate while that of the minimal is on
Si (100) substrates, indicating CZTS grown on Si (100) yields preferred (112) orientation while that
on Si(110) yields (220) preferred orientation. This proves the concept that CZTS with desired
orientation could be achieved by surface engineering of the Si substrate and thus the increased CZTS
solar cell performance resulting from the high quality of CZTS absorber with desired crystallographic
orientation. The mechanism of the effect needs to be further studied.
Table 4.4.3.3: I220/I112 of XRD peaks of CZTS films on Si substrates with different orientations.
Substrate
orientation
(100)
(100)4
(111)
(111)4
(110)
I220/I112
0.082
0.115
0.223
0.286
0.298
CZTS 338
140000
CZTS on Si (100)
o
CZTS on Si (100) 4
CZTS on Si (111)
o
CZTS on Si (111) 4
CZTS on Si (110)
Intensity (a.u.)
120000
100000
80000
60000
40000
CZTS 288
20000
CZTS 370
Si 521
CZTS 257
0
0
200
400
600
800
1000
-1
Raman Shift (cm )
Figure 4.4.3.22: Raman spectra of CZTS films on Si substrates with different orientations.
Figure 4.4.3.22 presents Raman scattering spectra of CZTS thin films. The intensity of CZTS peaks
varies with the change in substrate orientations, indicating crystalline of CZTS thin films is also
affected by the orientation of the Si substrates.
Figure 4.4.3.23 shows the SEM micrographs of CZTS films. All CZTS films are columnar structured
but slightly different in crystal size and shape. This suggestes that the morphology of CZTS thin films
is also affected by the orientations of the Si substrates.
- 160 -
- 161 -
Figure 4.4.3.23: SEM images of CZTS films on Si substrates with different orientations. (a) CZTS
on Si (100)_cross section, (b) CZTS on Si (100) surface, (c) CZTS on Si (100)4_cross section, (d)
CZTS on Si (100)4 surface (e) CZTS on Si (111)_cross section, (f) CZTS on Si (111) surface, (g)
CZTS on Si (111) 4_cross section, (h) CZTS on Si (111) 4 surface, (i) CZTS on Si (110)_cross
section, (j) CZTS on Si (110) surface.
Figure 4.4.3.24: Concentration depth profile of CZTS film on Si (100).
- 162 -
Figure 4.4.3.24 shows the concentration depth profile of CZTS film on Si (100) substrate determined
by XPS. From the surface to the interface of CZTS and Si, the relative concentrations of each element
are almost constant. An obvious oxygen hump is presented in the Si/CZTS interface which might be
due to SiO2 formed on the surface of the Si wafer before CZTS was deposited on it. This is likely
because of long time pumping down required for achiving the desired base pressure due to the lack of
a load lock in the current sputter system. The average composition ratios estimated from XPS are
Cu/(Zn+Sn)=1.33; Zn/Sn=1.48; and S/(Cu+Zn+Sn)=0.75.
Sulphurisation annealing
Sulphurisation annealing was applied to increase the crystallinity of the as-deposited CZTS, the CZTS
on Si (100) films deposited by sputtering a brand-new target with the same deposition parameters as
aforemented were sulfurized for 0.5h in the MTI sulfurization furnace at 530°C, 540°C and 550°C,
respectively.
Figure 4.4.3.25: SEM surface image of CZTS thin films on Si (100): (a) as-deposited, (b) annealed at
530°C, (c) annealed at 540°C (d) annealed at 550°C.
The surface morphology (shown in Figure 4.4.3.25) and chemical composition (shown in Table
4.4.3.4) obtained from EDS measurement changes with the sulfurization temperature. As the
sulfurization temperature increases from 530°C to 550 °C, the surface becomes rougher and more
porous, and Sn-loss becomes more serious. In addition CuxS clusters were shown sitting on the
surface and porosity surrounding the CuxS clusters is significantly higher than the rest of the surface
area. This also explaines the more pronounced Si peak and more scattered Raman spectra with
increasing annealing temperature. The presence of CuxS was also confirmed from XRD pattern
(Figure 4.4.3.26) and Raman spectra (Figure 4.4.3.27).
Table 4.4.3.4: Composition ratios of CZTS thin films estimatd from EDS measurement.
Zn/Sn=
As-deposited 530°C 540°C 550°C
1.62
1.57
2.10
2.39
Cu/(Zn+Sn)= 0.89
0.91
0.97
0.92
S/metal=
0.76
0.82
0.82
0.74
- 163 -
CZTS on Si (100)
530C, 0.5h sulfurized CZTS on Si (100)
CZTS 338
Intensity (a.u.)
30000
20000
CZTS 370
112
20000
Intensity (a.u.)
40000
CZTS on Si (100)
530C, 0.5h annealed CZTS on Si (100)
15000
10000
101
CZTS 288
220
200
Si 521
CZTS 257
10000
303
5000
202 211
224
008
60
70
332
0
0
0
100
200
300
400
500
600
700
20
800
30
40
50
80
2deg
-1
Raman shift (cm )
Figure 4.4.3.26: Raman spectra of CZTS films on Figure 4.4.3.27: XRD patterns of CZTS films on
Si (100) sulphurised at different temperatures.
Si (100) sulphurised at different temperatures.
It was also found that at 550°C, in addition to the second phase of CuxS, ZnS is also presented in the
film, which indicates the decomposition of the CZTS films through the loss of SnS. All these data
suggested that the optimal temperature for obtaining the best quality of CZTS is around 530°C and
longer time at this temperature could be applied for improving the crystallinity and crystal size of
CZTS.
CZTS on ZnS-coated Si substates
A thin ZnS layer was deposited on HF dipped Si by thermo evaporation and thereafter the ZnS coated
wafer was loaded in the sputtering chamber for CZTS deposition. The deposition parameters were the
same as above. Figure 4.4.3.28 shows the TEM image of CZTS/ZnS/Si. ZnS and CZTS are both
polycrystal and in columnar structure.
Figure 4.4.3.28: TEM images of CZTS/ZnS/Si: (a) cross sectional bright field image, (b) high
resolution image of the interface of Si and ZnS, (c) high resolution image of the interface of CZTS
and ZnS.
- 164 -
Sulphurisation annealing temperature effect
Figure 4.4.3.29: SEM cross-sectional images of CZTS/ZnS/Si: (a) as-deposited (b) sulphurised at
550°C for 5 mins.
Figure 4.4.3.30: EDS line scan on CZTS/ZnS/Si after sulphurisation at 550°C for 5 mins.
CZTS/ZnS/Si (100)
530 C 5 min sulfurized CZTS/ZnS/Si (100)
550 C 5 min sulfurized CZTS/ZnS/Si (100)
140000
CZTS 338
120000
Intensity (a.u.)
100000
Cu2-xS 475
80000
60000
CZTS 288
Si 521
CZTS 370
CZTS 257
40000
20000
0
-20000
0
100
200
300
400
500
600
700
800
-1
Raman shift (cm )
Figure 4.4.3.31: Raman spectra of CZTS films on Si (100) sulphurised at different temperatures.
The effect of sulphurisation annealing temperature on the properties of CZTS/ZnS/Si films was
obvious. After annealing at 550°C for 5 min, CZTS thin film tends to decompose resulting in
significantly rougher surface as shown in Figure 4.4.3.29. EDS line scan on the big clusters revealed
- 165 -
that in addition to the large grains of CuxS sitting on the surface, Cu2SnS3 secondary phase might also
presented in the annealed film as shown in Figure 4.4.3.30. The rough surface and presence of
secondary phases was confirmed by its Raman spectra (Figure 4.4.3.31). Therefore, 550°C is too high
for sulfurization of CZTS/ZnS/Si (100).
Sulphurisation annealing time effect
From the results discussed above, 530°C is a favourable temperature for sulphurisation of
CZTS/ZnS/Si structure. To improve the film quality, we have extended the sulphurisation time from 5
min to 0.5h. From Figure 4.4.3.32, we can see that after 5 mins sulfuphisation, the intensity of (112)
increased and the ratio of I220/I112 changs from 0.717 to 0.244, indicating the improved crystallinity
of preferred orientation of (112). In addition, there is no significant change in the surface roughness as
shown in the Raman specra (Figure 4.4.3.33). With the annealing time extended to 0.5 hour,
I220/I112 continues to decrease to 0.092, which indicats the further growth of CZTS crystal with the
orientation of (112). The growth of CZTS crystal can also be evidenced by the significant increase in
the scattering of incoming laser beam indicating the much rougher surface resulted from bigger CZTS
crystal size. The underlying mechanism of the structure change caused by the annealing conditions
needs to be further investigated.
112
14000
CZTS/ZnS/Si (100)
530C, 5 min sulfurized CZTS/ZnS/Si (100)
530C, 0.5h sulfurized CZTS/ZnS/Si (100)
12000
10000
Intensity
101
220
8000
110
6000
303
200
4000
202 211
2000
224
008
332
0
20
30
40
50
60
70
80
2
Figure 4.4.3.32: XRD patterns of CZTS films on Si (100) sulphurised at 530°C within 5min and 0.5h.
80000
CZTS 338
CZTS/ZnS/Si
530C, 5min sulfurized CZTS/ZnS/Si (100)
530C, 0.5h sulfurized CZTS/ZnS/Si (100)
Intensity (a.u.)
60000
40000
CZTS 288
20000
CZTS 370
CZTS 257
Si 521
0
0
100
200
300
400
500
600
700
800
-1
Raman shift (cm )
Figure 4.4.3.33: Raman spectra of CZTS films on Si (100) sulphurised at 530 °C within 5min and
0.5h.
4.4.3.3.2 Bandgap engineering
The development of CZTS material for photovoltaic application is based on cation cross-substitution
(mutation) from zinc-blende II-VI and III-V binary semiconductors to form quaternary zinc-blende
- 166 -
superlattices. In this way, cation elements in CZTS can also be substituted by other elements to make
new photovoltaic semiconductor materials fulfilling required qualities. One application of this cation
substitution method is for bandgap engineering.
To promote kesterite device efficiency, multi-junction structure is being investigated. UNSW has a
world reputation in making high-efficiency crystalline silicon solar cell and also has pioneered the
research of silicon based tandem solar cells. It is therefore a straightforward idea to combine kesterite
cell and silicon cell into two-junction tandem cells. To get optimal efficiency for two junction cells,
we need to have optimal set of bottom and top cell band gaps. According to the calculation of De Vos
[4.4.3.17], the optimal band gap set for two-junction device is 1.0eV for bottom cell and 1.9eV for top
cell. Crystalline Silicon has a band gap of 1.1eV, which is close to the optimal bottom cell; while
CZTS has a band gap of 1.45eV, which needs to be modified for an optimal top cell.
To achieve bandgap engineering, an effective method is to make quaternary alloys. For example, by
increasing the Ag2ZnSnS4 concentration in the alloy of Ag2ZnSnS4/Cu2ZnSnS4 the band gap of CZTS
can be increased from 1.45 to 2 eV [4.4.3.18]. This alloy method can also be regarded as a way of
doping Ag into CZTS. Generally, the band gap modification of quaternary alloy follows the rule
[4.4.3.19]:
(4.4.3.1)
where A and B are two quaternary materials, b is the band gap bowing parameter. To obtain a top cell
with the band gap of 1.9 eV, alloys of Ag2ZnSnS4/Cu2ZnSnS4 and Cu2ZnGeS4/Cu2ZnSnS4 have been
proposed and focused in this strand. Actually, pure Ag2ZnSnS4 and Cu2ZnGeS4, which have band
gaps of 2.0 eV and 2.05 eV [4.4.3.20], can be excellent potential materials for the top cell. The
tandem cell structure is proposed as shown in Figure 4.4.3.34.
Lattice mismatch is another problem need to be considered when designing tandem cells. The lattice
constant modification of quaternary alloys follows Vegard’s law [4.4.3.21]:
(4.4.3.2)
Or
(4.4.3.3)
where δ, δ’, δ” are three constant coefficient. When the lattice constant difference between material A
and B is small, the first equation of Vegard’s law is applied; otherwise the second equation is applied.
Figure 4.4.3.34: Schematic of Si/kesterite
tandem cell structure.
- 167 -
Apart from calculating the alloy structure constants, these constants can also be directly simulated
from first-principles. By using the PBE (Perdew-Burke-Ernzerhof ) method, the lattice constant can
be simulated while by using HSE method, band gap can be simulated, as currently underway
4.4.3.4 References
4.4.3.1
4.4.3.2
4.4.3.3
4.4.3.4
4.4.3.5
4.4.3.6
4.4.3.7
4.4.3.8
4.4.3.9
4.4.3.10
4.4.3.11
4.4.3.12
4.4.3.13
4.4.3.14
4.4.3.15
4.4.3.16
4.4.3.17
4.4.3.18
4.4.3.19
4.4.3.20
4.4.3.21
T.K. Todorov, K.B. Reuter, D.B. Mitzi, Advanced Materials 2010, 22, E156.
W. Liu, D.B. Mitzi, M. Yuan, A.J. Kellock, S.J. Chey, and O. Gunawan, Chem.Mater.,
2010, 22(3), 1010-1014.
I. Perez-Wurfl, X.J. Hao, A .Gentle, D-H. Kim, G. Conibeer, M.A. Green, Applied Physics
Letters 95 (2009), 153506.
Z. Ouyang, S. Pillai, F. Beck, O. Kunz, S. Varlamov, et al., Applied Physics Letters, 2010,
96 (26), 261109.
J.J. Scragg et al., Thin Solid Films, 517 (2009), pp. 2481-2484.
R. Schuee et al., Thin Solid Films, 517 (2009) pp. 2465-2468.
Hironori Katagiri et al., Thin Solid Films, 480-481 (2005), pp. 426-432.
Cooray, N., et al., Optimization of Al-doped ZnO window layers for large-area Cu (InGa)
Se2-based modules by RF/DC/DC multiple magnetron sputtering. Japanese Journal of
Applied Physics, 1999. 38(11): p. 6213.
Kuo, S.Y., et al., Effects of RF power on the structural, optical and electrical properties of
Al-doped zinc oxide films. Microelectronics Reliability, 2010. 50(5): p. 730-733.
Kim, K.H., K.C. Park, and D.Y. Ma, Structural, electrical and optical properties of
aluminum doped zinc oxide films prepared by radio frequency magnetron sputtering.
Journal of applied physics, 1997. 81(12): p. 7764.
Van de Walle, C.G., Hydrogen as a cause of doping in zinc oxide. Physical review letters,
2000. 85(5): p. 1012-1015.
Jahagirdar, A.H., A.A. Kadam, and N.G. Dhere. Role of i-ZnO in optimizing open circuit
voltage of CIGS2 and CIGS thin film solar cells. in Photovoltaic Energy Conversion,
Conference Record of the 2006 IEEE 4th World Conference on. 2006. IEEE.
Green, M.A., Silicon Solar Cells: Advanced Principles & Practice. 1995, Sydney: Centre for
Photovoltaic Devices and Systems.
B. Shin, O. Gunawan, Y. Zhu, N. A. Bojarczuk, S. J. Chey, S. Guha*, Thin film solar cell
with 8.4% power conversion efficiency using an earth-abundant Cu2ZnSnS4 absorber, Prog.
Photovolt: Res. Appl. 2013 (21):72–76.
D. A. R. Barkhouse, O. Gunawan, T. Gokmen, T. K. Todorov, D. B. Mitzi*, Device
characteristics of a 10.1% hydrazine-processed Cu2ZnSn(Se,S)4 solar cell, Prog. Photovolt:
Res. Appl. 2012 (20):6–11.
T. K. Todorov , J. Tang , S. Bag , O. Gunawan , T. Gokmen , Y. Zhu, .D. B. Mitzi *,
Beyond 11% Efficiency: Characteristics of State-of-the-Art Cu2ZnSn(S,Se) 4 Solar Cells,
Adv. Energy Mater. 2012, DOI: 10.1002/aenm.201200348.
Vos, A.D., Detailed balance limit of the efficiency of tandem solar cells. Journal of Physics
D: Applied PhysicsEmail alert RSS feed, 1980. 13: p. 839.
Tsuji, I., et al., Novel Stannite-type Complex Sulfide Photocatalysts AI2-Zn-AIV-S4 (AI =
Cu and Ag; AIV = Sn and Ge) for Hydrogen Evolution under Visible-Light Irradiation.
Chemistry of Materials, 2010. 22(4): p. 1402.
Chen, S., et al., Compositional dependence of structural and electronic properties of
Cu_{2}ZnSn(S,Se)_{4} alloys for thin film solar cells. Physical Review B, 2011. 83(12): p.
125201.
Leon, M., et al., Optical constants of Cu[sub 2]ZnGeS[sub 4] bulk crystals. Journal of
Applied Physics, 2010. 108(9): p. 093502.
KT, J., R. Shubhra, and R. L, Vegard's law: a fundamental relation or an approximation?
International Journal of Materials Research, 2007. 9: p. 776.
- 168 -
4.5
Third Generation Strand – Advanced Concepts
University Staff:
Prof. Gavin Conibeer (group leader)
Dr Richard Corkish
Prof. Martin Green
Senior Lecturer:
Dr Santosh Shrestha
Senior Research Fellows:
Dr Dirk König
Lecturers:
Dr Ivan Perez-Wurfl
Research Fellows:
Dr Shujuan Huang
Dr Supriya Pillai (ASI Fellow)
Dr Xiaojing “Jeana” Hao (ASI Fellow)
Postdoctoral Fellows:
Dr Pasquale Aliberti (until June 2012)
Dr Robert Patterson (ASI Fellow)
Dr Binesh Puthen-Veettil (ASI Fellow)
Dr Xiaoming Wen (from Dec 2012)
Dawei Di (March to Sept 2012)
Craig Johnson (from Aug 2012) (ASI Fellow)
Dr Siva Karaturi (from Feb 2013)
Professional officers:
Dr Tom Puzzer (part time)
Dr Patrick Campbell (shared with Thin Film)
Mark Griffin (shared with Thin Film)
Research Associates:
Dr Didier Debuf (adjunct fellow)
Yidan Huang
Higher Degree Students:
Andy Hsieh (until Nov 2012)
Dawei Di (until March 2012)
Craig Johnson (until Aug 2012)
Sammy Lee (until Aug 2012)
Haixiang Zhang
Tian Zhang
Yu “Jack” Feng
Pengfei Zhang
Sanghun Woo
Yao Yao
Neeti Gupta
Chien-Jen “Terry” Yang
Yuanxun “Steven” Liao
Suntrana Smyth
- 169 -
Lingfeng Wu (from Mar 2012)
Xuguang Jia (from Mar 2012)
Ziyun Lin (from Mar 2012)
Hongze Xia (from Mar 2012)
Xi Dai (from Mar 2012)
Shu Lin (from Aug 2012)
Pei Wang (from Aug 2012)
Simon Chung (from Aug 2012)
Casual Staff:
James Rudd (part time)
Adrian Shi (Aug to Dec 2012)
Visiting Researchers:
Dr Yukiko Kamikawa (AIST, Tsukuba, Japan, until Mar 2012)
Undergraduate Students:
4th year Thesis Projects:
Dongchen Lan
Guillaume Pauvert
Aden Brown
Li Shen
Alexander Robinson
Abstract
There has been a significant increase in the work on hot carrier cells in 2012 with the start of both the
Australian Solar Institute (ASI) project and US-Australia Solar Energy Collaboration (USASEC)
projects. The Group IV nanostructure tandem cells project - the “all-Si” tandem cell – has also seen an
increase with the start of its own ASI project. There has also been further work on up-conversion,
photo-electrolysis and on plasmonics.
There has been increased understanding of the parameters required to grow good Si nanostructure
materials. Characterisation of materials is leading to better understanding of the defects which narrow
the effective bandgap and hence limit VOC. Improvements in VOC have been seen with lower
temperature measurements and in current with light trapping by various techniques. Modelling using
equivalent circuits is giving greater insight to device performance.
Hot carrier cells efficiencies are being modelled with greater input of real material properties and with
integration of more complex structures. Combined contact and absorber materials mdoels and those
including both multiple quantum well and phononic bad gap material have been developed. New
designs of contacts for rectification or for incorporation of energy selective contacts directly in
absorbers are being developed. A wide range of potential materials for phononic modulated structures
have been identified and are being developed. Growth of nanoparticle ordered arrays has progressed
with modelling of such superstructures also being further developed. Fully integrated devices are now
being designed and are expected to be fabricated with collaborators in the near future.
The up-conversion project now has good synthesis routes for Er containing phosphors and for PbS
downshifting quantum dots (QDs) used to sensitise Er to a wider wavelength range. Progress on
incorporating such materials with light trapping structures is expected to give improvements in upconversion for Si cells.
This progress in all the main Third Generation project areas is improving understanding and allowing
optimisation of modelling, structures and devices. Developments to come in the next year will see
significant advancement in these areas, with excellent prospects for good demonstration devices.
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4.5.1 Third Generation Photovoltaics
The “Third Generation” photovoltaic approach is to achieve high efficiency whilst still using “thin
film” second generation deposition methods. The concept is to do this with only a small increase in
areal costs and to use abundant and non-toxic materials and hence reduce the cost per Watt peak
[4.5.1]. Thus these “third generation” technologies will be compatible with large scale
implementation of photovoltaics. The aim is to decrease costs to well below US$0.50/W, towards
US$0.20/W or better, by dramatically increasing efficiencies but maintaining the economic and
environmental cost advantages of thin film deposition techniques (see Fig. 4.1.3 showing the three PV
generations) [4.5.1, 4.5.2]. To achieve such efficiency improvements such devices aim to circumvent
the Shockley-Queisser limit for single band gap devices that limits efficiencies to the “Present limit”
indicated in Fig. 4.1.3 of either 31% or 41% (depending on concentration ratio). This requires
multiple energy threshold devices such as the tandem or multi-colour solar cell. The Third
Generation Strand is investigating several approaches to achieve such multiple energy threshold
device [4.5.1, 4.5.3].
The two most important power loss mechanisms in single-band gap cells are the inability to absorb
photons with energy less than the band gap (1 in Fig. 4.5.1), and thermalisation of photon energy
exceeding the band gap, (2 in Fig. 4.5.1). These two mechanisms alone amount to the loss of about
half of the incident solar energy in solar cell conversion to electricity. Multiple energy threshold
approaches can utilise some of this lost energy. Such approaches do not in fact disprove the validity
of the Shockley-Queisser limit, rather they avoid it by the exploitation of more than one energy level
for which the limit does not apply. The limit which does apply is the thermodynamic limit shown in
Fig. 4.1.3, of 68.2% or 86.8% (again depending on concentration).
S
C
T
2
3
1
2
C
C
Figure 4.5.1: Loss processes in a standard solar cell: (1) non-absorption of below band gap photons;
(2) lattice thermalisation loss; (3) and (4) junction and contact voltage losses; (5) recombination loss.
In the Third Generation Strand, we aim to introduce multiple energy levels by fabricating a tandem
cell based on silicon and its oxides, nitrides and carbides using reduced dimension silicon
nanostructures to engineer the band gap of an upper cell material. We are aiming to collect photogenerated carriers before they thermalise in the “hot carrier” solar cell. Also we are investigating
absorption of two below bandgap photons to produce an electron-hole pair in the cell by upconversion in a layer behind the Si cell using erbium doped host materials. In order to optimise the
requisite properties, all these structures are likely to be thin hence maximising absorption of light in
thin structures through light trapping is very important. Hence we are also investigating localised
surface plasmon enhanced coupling of light into these Third Generation devices.
4.5.2 Si Nanostructure Solar Cells
Researchers: Shujuan Huang, Ivan Perez-Wurfl, Dirk König, Tom Puzzer, Xiaojing Hao, Dawei Di,
Zhenyu Wan, Sammy Lee, Yidan Huang, Lingfeng Wu, Xuguang Jia, Ziyun Lin, Tian Zhang, Terry
Yang, Gavin Conibeer, Martin Green
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4.5.2.1 The ‘‘all-Si’’ tandem cell
We are developing a material based on Si (or other group IV) quantum dot (QD) or quantum well
(QW) nanostructures, from which we can engineer a wider band gap material to be used in tandem
photovoltaic cell element(s) positioned above a thin film bulk Si cell, see Fig. 4.5.2.
Previously we have demonstrated the ability to fabricate materials which exhibit a blue shift in the
effective band gap as the QD or QW size is reduced, using photouminescence [4.5.4] and absorption
[4.5.5] data [4.5.6]. A thin film deposition of a self-organised QD nanostructure is achieved through a
sputtered multi-layer of alternating Si rich material and stoichiometric dielectric [4.5.4]. On
annealing the excess Si precipitates into small nanocrystals which are limited in size by the layer
thickness, thus giving reasonable size uniformity, as first demonstrated by Zacharias [4.5.7].
Demonstration of doping of these layers with both phosphorus and boron to create a rectifying p-n
junction has resulted in devices with a photovoltaic open circuit voltage of 490mV [4.5.6, 4.5.7,
4.5.8].
This image cannot currently be display ed.
4
Figure 4.5.2: A tandem photovoltaic cell using quantum
confined QDs or QWs to engineer the band gap of the top cell
and potentially also the lower cells. Short wavelength light is
absorbed in the top cell and longer wavelengths in lower cells,
thus boosting the overall voltage generated and hence efficiency.
Formation of Si (or Ge or Sn) QDs through layered thin film
deposition of Si rich material which crystallises into uniform
sized QDs on annealing.
A cell based entirely of Si, or other group IV elements, and their dielectric compounds with other
abundant elements (i.e. silicon oxide, nitride or carbide) fabricated with thin film techniques, is
advantageous in terms of potential for large scale manufacturability and in long term availability of its
constituents. Such thin film implementation implies low temperature deposition without melt
processing, it hence also involves imperfect crystallisation with high defect densities. Hence devices
must be thin to limit recombination due to their short diffusion lengths, which in turn means they must
have high absorption coefficients.
For photovoltaic applications, nanocrystal materials may allow the fabrication of higher band gap
solar cells that can be used as tandem cell elements on top of normal Si cells [4.5.11, 4.5.12]. For an
AM1.5 solar spectrum the optimal band gap of the top cell required to maximize conversion
efficiency is ~1.7 to 1.8eV for a 2-cell tandem with a Si bottom cell [4.5.13]. To date, considerable
work has been done on the growth and characterization of Si nanocrystals embedded in oxide [4.5.7,
4.5.14] and nitride [4.5.15, 4.5.16] dielectric matrices. However, little has been reported on the
experimental properties of Si nanocrystals embedded in SiC matrix [4,5,17]. These are of particular
interest for application in photovoltaic devices because of an expected significant increase in carrier
transport due to a decrease in the barrier height between adjacent nanocrystals [4.5.18]. As a result,
sufficient carrier mobility can be obtained to satisfy device fabrication requirements.
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4.5.2.2 Fabrication of Si QD nanostructures
Thin film techniques are used for nanostructure fabrication. These include sputtering and plasma
enhanced chemical vapour deposition (PECVD). The deposition is a variation of the multi-layer
alternating ‘stoichiometric dielectric / Si rich dielectric’ process, shown in Fig. 4.5.3, followed by an
anneal during which Si nanocrystals precipitate limited in size by the Si rich layer thickness [4.5.12,
4.5.7]. The most successful and hence most commonly used technique is sputtering, because of its
large amount of control over deposition material, deposition rate and abruptness of layers. A multitarget remote plasma sputtering machine with two independent RF power supplies as well as
additional DC power supplies is used in this work.
Si rich dielectric
2-5nm
Stoichiometric
dielectric, 2-5nm
Quartz or glass
Substrate
substrate
Figure 4.5.3: Multilayer deposition of alternating Si rich dielectric and stoichimetric dielectric in
layers of a few nm. On annealing the Si precipitates out to form small nanocrystals of a size
determined by the layer thickness. Nanocrystal or quantum dot size is therefore uniform.
RF magnetron sputtering is used to deposit alternating layers of SiO2 and SRO of thicknesses down to
2nm. [SRO refers to Si rich oxide, formed by co-sputtering Si and SiO2.] Deposition of multi-layers,
consisting typically of 20 to 50 bi-layers, is followed by an anneal in N2 from 1050 to 1150C.
During the anneal, the excess silicon in the SRO layer precipitates to form Si nanocrystals between
the stoichiometric oxide layers.
For Si QDs in SiO2 the precipitation occurs according to the following:
Precipitation of excess Si from Si rich dielectrics in SiNx and SiC follows a similar crystallisation
reaction as Si precipitates from the amorphous matrix. The techniques have also been applied to
growth of Sn and Ge QDs in SiNx in SiO2. Ge quantum dots can be precipitated at substantially lower
temperature, as discussed below.
4.5.2.3 Silicon nanocrystal devices on quartz substrates
Researchers: Ivan Perez-Wurfl, Xiaojing Hao, Dawei Dai, Adrian Shi, Lingfeng Wu, Xuguang Jia,
Ziyun Lin, Tian Zhang, Terry Yang
As previously reported in previous annual reports, devices have been fabricated using the SiQD in
SiO2 materials, with p-n junction formation using B or P of multilayers, respectively [4.5.9, 4.5.10,
4.5.19]. The fabricated p-n diodes consisted of sputtered alternating layers of SiO2 and SRO onto
quartz substrates with in-situ boron and phosphorus doping. The top B doped bi-layers were
selectively etched to create isolated p-type mesas and to access the buried P doped bi-layers.
Aluminium contacts were deposited by evaporation, patterned and sintered to create ohmic contacts
on both p and n-type layers. The fabricated interdigitated solar cells have an effective area of up to
0.12cm2. The devices exhibit rectification and a photovoltaic response with VOC up to 493 mV, but
with as yet very small currents and very bad fill factors. These are partly due to the very high
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resistance of the material and in particular to the relatively long lateral paths to contacts at the back
contact, necessitated by growth on an insulating quartz substrate. The device structure with
appropriate contacting is shown schematically in Fig. 4.5.4. This also shows the current crowding
effect which results from the high lateral resistance.
+
V m(Temp)
-
W
I
Al
contacts
SRO 4nm
SiO2, 2nm
B doped
SRO
bilayers
B doped
SRO
bilayers
B doped
SRO
bilayers
Wa
P doped SRO bilayers with
sheet resistance Rsheet
Non-uniform current
injection
Quartz substrate
Figure 4.5.4: Schematic diagram of the of the fabricated inter-digitated Si QD in SiO2 devices, with
B and P doping to create the p-n junctions and illustrating the effect of current crowding due to high
lateral resistance to the back contact.
4.5.2.4 Equivalent circuit model of nanocrystal devices on quartz substrates
In order to better understand the limitation imposed by the device on the solar cell performance, it is
necessary to find a good equivalent circuit model specific to our devices. The high resistivity of the
base layer makes it imperative to consider the two dimensional effects of current flow [4.5.20]. In
order to analyse the I-V characteristics of these diodes we first generalised the model to include any
number of diodes in series. The series connection of diodes is used to explain the ideality factors
higher than two normally observed in our structures. We believe this is a reasonable model as the
ideality factor observed is almost independent of the diode current. This type of behaviour has also
been observed in multi quantum well laser diodes [4.5.21], where it has been proven to arise from an
unintended series combination of diodes. Based on this series combination of diodes, it is possible to
obtain an expression relating temperature dependent I-V measurements to the band gap [4.5.9]. It is
further possible to linearise the expression around an average measuring temperature, Tavg, as follows:

i
kTavg
Eg i
 VD  IRs (T , I ) 
q
q

i
i

 I
kT 
 ni ln
q  i
 I oi
   Tavg  
q
  i 
   ln
 T   1
kTnom
   nom   i

i
Eg i 

q 
(4.5.1)
where k is Boltzmann’s constant, q is the electron charge, Tnom is the temperature at which the
saturation current, Ioi, is defined, ni is the diode ideality factor, and αi is the saturation current
temperature exponent. Notice that this equation shows that the I-V characteristics are related to a sum
of band gaps.
As the current flows from the base contact to the emitter, a linear voltage drop along the base and
under the diode isolation mesa causes an exponential change in the diode current. This crowding of
the current at the edge of the diode mesa, depicted in Fig. 4.5.4, can be modelled adding a current
dependence on the series resistance. This series resistance, RS, can be expressed as the sum of a
current independent, Rext and a current dependent series resistance Rint arising from current crowding:
Rs (T , I )  Rext  Rint 
Rsheet (T )Wa Rcont (T ) RsheetW ln cos( ) sin( )


2L
2L
8L
 tan( )
(4.5.2)
The value of this resistance can be found by numerically solving the following transcendental
equation:
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 tan  
1 q
W
IRsheet
8n kT
L
(4.5.3)
Within this mathematical framework it is then possible to extract the value of the series resistance,
remove its effect from the measured I-V characteristics to extract the actual ideality of the diode
current. We normally observe an ideality factor of 3. Based on the proposed model of at least two
diodes in series, without any assumptions, it is only possible to establish that the value of the band gap
extracted from temperature dependent I-V measurements corresponds to a sum of the band gaps (or
activation energies) of these diodes.
An interesting behaviour observed in these devices was the apparent lack of correspondence between
the dark and light I-V characteristics. The series resistance extracted from the diode dark I-V
characteristic is too small to explain the limited short circuit current as well as the low fill factor
measured under a simulated 1-sun condition. A more complete circuit model is necessary to explain
this discrepancy. We have proposed a model where the observed behaviour is due to two distinct
areas in the fabricated devices. The photocurrent is produced only in a small area of the device, this
area being proportional to a fraction of the normalized diode area. In a Spice circuit model this area is
given a value smaller than 1, that we denote as fraction. This will be a fitting parameter to reproduce
the measured dark and illuminated I-V characteristics. Figure 4.5.5 shows the circuit proposed.
Figure 4.5.5: Equivalent Spice circuit representation of the fabricated devices.
A relatively large percentage of the device area is responsible for the measured characteristics in the
dark. The current in this area is caused by the diffusion of minority carriers caused by the applied
voltage (V1 in Fig. 4.5.5) from the p or n side to the opposite region. This current is expected to be
large due to the low lifetime of the minority carriers (mostly recombination current in the depletion
region). The observed series resistance will be proportional to the total series resistance, Rtot, and
inversely proportional to the area, R2=Rtot/(1-fraction), where the diffusion current occurs. The dark
I-V behaviour is modelled by the top branch of the circuit depicted in Fig. 4.5.5. D_QD1 and D_Sch
represent the series connection of diodes whose band gaps are extracted using Eq. (4.5.1). The series
resistance, R2, has the temperature and current dependence detailed in Eqs. (4.5.2) and (4.5.3).
Only a small part of the diode area may have a large enough lifetime to produce a photocurrent. As
this photocurrent flows only through this fractional area, the series resistance is inversely proportional
to this fraction: R1=Rtot/fraction. Since the photocurrent, Iph, flows through R1, the illuminated I-V
characteristics are limited by a larger resistance than that observed in the dark condition, as long as
the fraction of the diode area is smaller than one half. The simulations depicted in Fig. 4.5.6 show the
reduction of ISC as the fraction, f, is varied from 99% to 1% of the total diode area.
- 175 -
f=0.01
f=0.03
f=0.1
f=0.31
f=0.99
Figure 4.5.6: Spice simulations of 1-sun I-V characteristics based on the circuit depicted in Fig.
4.5.5. The fraction, f, represents the normalized area of the diode where the photocurrent is produced.
With the circuit model described, it will be possible to extract complementary information from dark
and illuminated I-V measurements.
In view of these simulations, it is clear that the electrical characterisation of our devices needs to take
into consideration previously overlooked limitations. For example, great care should be taken when
interpreting the Quantum Efficiency extracted from a spectral response measurement as the assumed
condition of short circuit current may be incorrect even if the device is externally short circuited
(internally, the diode may be forward biased). Moreover, as the current is proportional to the photon
flux, and the flux is generally different at each wavelength tested, the internal bias of the device can
be different at each point of the spectrum investigated.
4.5.2.5 Modelling for device optimisation
Researchers: Ivan Perez-Wurfl, Ziyun Lin, Tian Zhang
When designing and implementing the Si-QD solar cell model the most important performance
limiting factor identified was the in-plane resistivity of the material. To take this into consideration a
2-D model needed to be implemented. This model takes into consideration the lateral flow of current
in the solar cell. By implementing this 2-D effect we have been able to recreate the current-voltage (IV) characteristics of diodes made with Si-QD on insulating substrates. The model takes into
consideration the temperature dependence of the resistivity and saturation current. The actual
behaviour of the resistivity and saturation current are modelled in terms of activation energies
obtained from temperature dependent measurements. The activation energy of the resistivity is related
to defects in the material which act as electronic traps. The activation energy of the saturation current
is related to an effective electronic bandgap that, in some cases, can be modified by the doping itself.
In crystalline semiconductors such as Si, doping does not affect the electronic bandgap unless extreme
levels of dopants are used (> 1x1019cm-3). However, in disordered or amorphous semiconductors,
doping the material will always affect the electronic bandagap. For example, doping hydrogenated
amorphous silicon (a:Si:H) increases the defects in the material which in turn act as traps. The end
result is a highly resistive material despite the very high doping levels used. Furthermore, the high
doping itself adds such a high density of defects that the electronic bandgap of the material is clearly
reduced. To minimize the effect, undoped materials are commonly used for the active PV layer of the
device. We have identified that the type of Si-QD films required for PV devices, has indeed a similar
behaviour as that of a-Si:H when it comes to doping. What is more, the native defects in the Si-QD
layers we produce are high enough to affect the electronic bandgap even when no intentional doping
is present.
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The realisation that the effective electronic bandgap is in fact the most important aspect limiting VOC
has clarified an optimisation plan to follow.
Efforts are now in place to establish optical methods to quantify the effective electronic bandgap.
Specifically, we are working on tailoring standard characterisation methods, Photo Thermal
Deflection Spectroscopy (PDS) and Photo Luminescence (PL), to specifically study Si-QD materials.
Both techniques offer the advantage of allowing evaluation of defects. These defects should be
minimised to reap the benefits of the expected enhanced bandgap and, with the aid of these
techniques, could be quantified without the need to fabricate devices. Once the quality of material has
been optimised according to minimised defects determined using these optical methods, this material
can then be used to fabricate devices.
4.5.2.6 Light trapping for increased current in devices
Researchers: Ivan Perez-Wurfl, Supriya Pillai, Lingfeng Wu, Ziyun Lin, Xuguang Jia
The photogenerated current in p-n unction Si nanostructure devices has been increased by more than
100% compared to a cell without any light trapping. The application of an Aluminium back-reflector
gave a current enhancement of up to 46%. Plasmonic silver nanoparticles gave an even more
impressive result increasing the current by up to 109%.
The work concentrated around devices fabricated on quartz substrates in the substrate configuration.
This means that the devices were designed such that illumination was applied from the top of the
structure. The substrate is used as a support for the device as well as the back-reflector as shown in
Fig. 4.5.7.
Figure 4.5.7: Schematic of the structure used to optimise photocurrent using an Al back-reflector.
The amount of light absorbed by these structures is limited by the thinness of the structure and the
relatively low silicon content. Furthermore, light trapping is not possible using a textured substrate as
the material is deposited by sputtering which is not as conformal as CVD or ALD, for example.
Furthermore, the thinness of the layer does not allow the typical chemical texturing used in standard
Si solar cells. The most straightforward way to increase light absorption, and therefore short circuit
current, is with the use of a metal back-reflector. The efficiency of a back-reflector can be further
enhanced (or substituted) with the use of plasmonic effects that could increase the optical absorption
path by scattering incoming light.
A sample was chosen such that the current was as high as possible for multiple devices contained
within the same insulating substrate. We will refer to this sample as SM2A from here on. The devices
showed a reasonable uniformity within the sample to allow a statistically significant comparison
between cells with and without a back-reflector. The resistivity of the base had to be chosen to be as
low as possible to reduce its effect on the current without compromising the quality of the diodes
fabricated.
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The results of the photocurrent enhancement are summarised in Table 4.5.1. Notice that the lowest
enhancement is 31.3%. The average enhancement across the sample was 40% with a maximum
enhancement of 46.1%.
Table 4.5.1: Summary of photocurrent enhancement achieved with an Aluminium back-reflector on
sample SM2A.
ISC without Reflector (A)
ISC with Reflector (A)
5mm2 device
Device A
18.7
27.3
% increase using Al (M2A)
46.1%
1mm2 devices
Device B
Device C
Device D
4.1
4.6
4.6
5.4
6.7
6.4
31.3%
45.1%
37.5%
Figure 4.5.8: Typical dark and illuminated I-V characteristics of a 1mm2 diode tested in sample
SM2A. Three curves are shown: Illuminated with Reflector (green), Illuminated with No Reflector
(grey) and in the Dark (blue).
Larger devices were also tested with the best result shown on a 5mm2 device.
The largest device showed the highest percent enhancement in photocurrent when tested with a backreflector or with silver nanoparticles. This particular enhancement is probably due to variations in the
device performance across the sample rather than having an actual dependence on the area. It is worth
pointing out that the largest device is closer to the centre of the sample which positively impacts the
quality of the device. An interesting feature observed in Fig. 4.5.8 is the larger VOC observed when
tested without a back reflector. This is probably due to heating of the sample while testing,
consequently, the temperature may not have been the same for all the measurements. The actual
temperature of the device is difficult to control due to the thermally insulating substrate used. Details
of the thermal management issues will be discussed below.
Silver nanoparticles have also been investigated. The enhancement obtained in the photocurrent was
significant. When tested on a 5mm2 solar cell, the current improved more than two fold. Specifically,
the current was increased by 109% compared to that obtained without any light trapping scheme.
The silver nanoparticles were deposited on the top of the cell. Illumination was then applied from the
side of the substrate after removing the Al back reflector as shown in the schematic in Fig. 4.5.9.
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Figure 4.5.9: Schematic of the structure used to optimise photocurrent using silver nanoparticles.
The nanoparticles were obtained by evaporating 18nm of Silver followed by a 1 hour anneal in a
nitrogen purged oven at 200oC. Figure 4.5.10 shows the comparison between the I-V characteristics of
the cell tested with an aluminium back-reflector or silver nanoparticles.
Figure 4.5.10: Illuminated I-V characteristics of a 5mm2 diode showing an enhanced photocurrent
due to Silver nanoparticles (green curve) or an Al backreflector (blue curves) compared to the same
cell without any light trapping (red curve).
4.5.2.7 Devices on conductive substrate
We have investigated three metals and one highly doped semiconductor thin film suitable for use as
conductive virtual substrates on a quartz supporting substrate.
The first all Si-QD devices demonstrated by our group were fabricated on quartz substrates. The main
reason for using quartz substrates is to demonstrate that any PV response is in effect due to the Si-QD
material and not the substrate itself. The disadvantage of using these substrates is the current
crowding that occurs due to the high resistivity of the Si-QD layers. This current crowding effect
limits the maximum current that can be extracted from the solar cell. Conductive substrates are
investigated to mitigate this detrimental effect.
We note that a thick substrate is not actually necessary. A thin, highly conductive film is sufficient to
reduce current crowding effects to a negligible level. The conductive film of choice will be deposited
on a quartz substrate to ensure compatibility with our current process.
The high temperature needed to precipitate the Si-QD from the silicon rich oxide layers, precludes the
use of low melting point metals such as Aluminium. To ensure compatibility with the unavoidable
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high temperature of our process, we chose three refractory metals with the highest melting points,
namely Tungsten (W), Tantalum (Ta) and Molybdenum (Mo). We also investigated the use of silicon
rich silicon carbide (SRC) layers. SRC can be deposited with a low resistivity and is perfectly
compatible with the sputtering tool used for depositing SRO materials.
4.5.2.7.1 Metals as back contacts
Researchers: Ziyun Lin, Ivan Perez-Wurfl, Xuguang Jia, Terry Yang
Sputtering is the most common technique used to deposit refractory metals. The deposition itself is
straightforward but much optimisation is needed to ensure the stress in the films is low enough to
avoid the films delaminating from the substrate. Early on in the optimisation process we identified
severe difficulties in obtaining good quality films with W and Ta. We concentrated our efforts then on
improving the quality of Mo thin films. It was found that a good indicator of the compatibility of the
film with a high temperature process was the minimisation of pinholes in the film. To accurately
quantify the density of pinholes we developed a technique based on the optical transmission of Mo
thin films. We have been able to use this characterisation technique to quantify the pinhole density
down a 1/100,000 total area of the film. Fig. 4.5.11 shows the measured optical transmittance of two
Mo films deposited on quartz.
Figure 4.5.11: Optical transmission of two Mo films on quartz.
The transmittance was measured using a Perkin Elmer Lambda 1050 spectrophotometer. The
measurement has been optimised by ensuring a baseline calibration compatible with measurements of
very low transmittance. Figure 4.5.11 shows the nominally zero transmittance (~10-6 %) of a 3mm
thick Al slab. This “zero” level shows a clear differentiation with a film transmitting only 0.001%.
This optical characterisation method is fast and accurate and allows us to pre-screen Mo films before
using them as a substrate.
The resistivity, surface roughness and the areal density of pinholes of the optimised Mo layer is
reported in Table 4.5.2.
Table 4.5.2: Characteristics of the best overall Mo film used as a substrate for a Si-QD p-i-n
structure.
Substrate
deposition
temperature (oC)
420
Film thickness
(nm)
Surface
roughness (nm)
Sheet resistance
(/square)
Area of pinholes
(% of total area)
350 ± 50
0.84 ± 0.2
2.35 ± 0.8
0.0018
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A 300 nm Mo film similar to that of Table 4.5.2 but with 50X higher pinhole density was sputtered on
quartz at room temperature. A bilayered structure consisting of 25 bilayers (4nm SRO:4nmSiO2),
nominally un-doped with 45% Si by Volume in the SRO, was simultaneously deposited on a quartz
substrate and the Mo film. Both structures were also simultaneously annealed at 1100oC for 1 hour in
a nitrogen purged quartz tube furnace. Visually, there was no apparent difference between the
structure deposited on Mo and that deposited on quartz after annealing. Comparing the PL emission
from the two samples also showed no difference as shown in Fig. 4.5.12.
Figure 4.5.12: a) Picture of the annealed 25 bilayer structure on Mo b) Comparison of the PL
emission from the multilayer on quartz and Mo.
This experiment proved that Mo is indeed compatible with the high temperature process and has no
obvious interaction with the multilayer structure.
The same Mo layer was then used as a back contact for a standard full p-i-n structure (70 bilayers
total, bilayer = 4nm SRO/2nm SiO2, SRO = 66% Si by volume). After the annealing process it was
clear that Mo had interacted with the p-i-n structure as can be seen in Fig. 4.5.13.
Figure 4.5.13: Picture of p-i-n structure after 1100°C anneal. The bottom left corner shows clearly
that a reaction occured between the p-i-n structure and the underlying Mo thin film.
Electrical measurements showed that a dead short was created across the full p-i-n structure
precluding the measurement of VOC or ISC. It is very likely that the formation of a Silicide caused the
shunt observed across the diode. The silicide formation is very sensitive to the presence of SiO2
between Si and Mo. The thicker oxide and lower Si content of the 25 bilayers may have been enough
to avoid a reaction with the underlying Mo thin film. In the case of the p-i-n structure, the high
concentration of Si in the SRO and the thin SiO2 barriers promote the formation of a Silicide. Mo
films are still under investigation. Further optimisation of these films is under intensive study. The use
of a silicide is being considered to limit the reaction of the metal and the p-i-n structure. Our findings
in this critical area will be reported in future reports.
- 181 -
4.5.2.7.2 Wide bandgap semiconductor as back contacts
Researchers: Dawei Di, Ivan Perez-Wurfl, Gavin Conibeer
A second option investigated, geared at reducing the lateral resistivity, is the use of a thin, highly
doped, wide bangap material.
Silicon rich carbide (SRC) layers have been investigated as the back contact for a p-i-n structure. A
significant advantage of using SRC rather than a metal conductive substrate is the large amount of
transmission of SRC as compared to zero transmission for metals. This allows a wider range of device
testing including rear illumination and is more compatible with a full tandem cell structure.
The SRC layers are deposited directly on quartz. The films are doped with phosphorous. After
1100oC annealing for 1 hour, the SRC layers show a sheet resistance an order of magnitude smaller
than that of bilayered SRO layers of similar thickness and with the highest silicon content of interest
(~66% Si by Volume).
Despite its greater transmission than metals SRC as a conductive substrate nonetheless does have an
unavoidable absorption loss caused by the layer itself. The SRC layer has to be as thin as possible to
reduce this loss. In the interest of investigating the effectiveness of SRC as a conductive substrate, we
choose a relatively thick layer that would be easier to handle and characterise. With enough evidence
to prove the effectiveness of SRC we can then optimise the thickness without significantly sacrificing
its conductivity.
A 237 nm thick phosphorous doped SRC layer was deposited on quartz, annealed and characterised
prior to using it as a substrate. The measured sheet resistance of this film was 5 k /square. The layer
was then used as a substrate for a standard p-i-n structure. The full stack was annealed once more,
devices were fabricated and characterised. It was possible to demonstrate a working device on this
structure. Dark and light I-V characteristics of a 1mm2 diode is shown in Fig. 4.5.14.
Figure 4.5.14: I-V characteristics of a standard p-i-n structure on and SRC substrate.
The measured ISC and VOC are lower than those of a standard p-i-n structure deposited directly on
quartz. The reason for this lower performance is mainly due to the absorption of light in the SRC
layer which is most probably not converted into current.
- 182 -
4.5.2.8 Increase in open circuit voltage
Researchers: Lingfeng Wu, Xuguang Jia, Ziyun Lin, Ivan Perez-Wurfl
The original intention to show a path towards improving the open circuit voltage of these devices
beyond 493mV (the previous record voltage) was to vary the silicon content in the SRO in order to
exploit a suspected trend relating it to the open circuit voltage. Our previous experience had shown
that by reducing the silicon content it is possible to increase the open circuit voltage sacrificing some
of the short circuit current due to the inherent increase in the resistivity of the material. This suspected
trend had to be revised after developing the proper techniques to reproduce material with the same
silicon content on a regular basis. Contrary to expectations, the change in silicon content of the SRO
has very little effect on the open circuit voltage.
An open circuit voltage of 503mV was measured on a 1mm2 solar cell at 250 K under an approximate
0.8 suns illumination condition. The demonstration of a voltage higher than 500mV was one of the
outcomes of this more structured line of work. The reason for these non-standard testing conditions
will be explained in detail in the following paragraphs. It is important to point out that the
demonstration of a voltage larger than 500mV was measured on a device fabricated on an oxidised Si
wafer. Such devices can absorb less than half the number of photons compared to a similar cell on a
transparent quartz substrate. The I-V curve of this solar cell is shown in Fig. 4.5.15.
Figure 4.5.15: Dark and illuminated I-V of cell with VOC = 503 mV, under an approximate 0.8 suns,
250 K.
Having developed the desired control on the composition and thickness, we deposited p-i-n structures
with variable silicon content (65.8, 63.7 and 60.9% Si by Vol.) with nominally the same size of SiQD. Using photoluminescence (PL) as a metric to compare the similarity of QD in three different
structures, we confirmed that indeed the Si-QD were very similar in the three cases. The PL spectra
measured on these structures is shown in Fig. 4.5.16.
- 183 -
Figure 4.5.16: Measured PL on three p-i-n structures with variable silicon content in the SRO (Si
percent by Volume: 65.8%, yellow; 63.7%, pink; 60.9%, blue).
As can be seen in Fig. 4.5.16, the peak emission of the three structures is nearly equal. The obvious
differences are a small red shift of the PL emission spectra for the 65.8% Si by Vol. sample and a
slightly higher tail in the high energy emission for the sample with the lowest Si content. The low
variation in the silicon content was intentional as these 5% variation represents a change in the lateral
resistivity of four to five orders of magnitude. The detailed relationship between Si content and
resistivity is currently under study.
The measured optical absorption of these three samples confirms that indeed, the sample with a higher
silicon content has the highest optical absorption becoming slightly less pronounced as the silicon
content is decreased.
The fabricated devices expected to highlight the relationship between VOC and Si content in the SRO,
do not in fact show any particular trend. The open circuit voltage measured ranges between 200mV
and 300mV without any apparent correlation to the silicon content. The lowest silicon content
investigated in fact did not produce any measurable devices, probably due to the extreme high series
resistance which in turn limits the device performance.
Having determined that the silicon content is not indeed determining the open circuit of these devices,
we set ourselves to the task of determining the discrepancy of these observations with our previous
experience showing an actual dependence. We have concluded that the lack of accurate control in the
material composition in the early days of this research at UNSW, along with the variation in the
deposition rate, lead to unintended variations in the deposited material which was interpreted as a
trend in VOC as a function of Si content in the SRO. We believe that there may much more important
factors affecting VOC besides the Si content. Investigation of other options to increase VOC led us to
consider devices not fabricated on our standard quartz substrates.
Our strongest proof that the substrates indeed affect the quality of the devices built on them, come
from an experiment performed by depositing single SRO layers on quartz substrates. An apparent
random variation in obtaining a PL signal from only a handful of the samples prepared was eventually
correlated to an absorption peak observed in the IR region of the spectrum, present only in the
substrates that showed no PL signal. This absorption peak (~2.7m) was later identified as an
indicator of OH groups. The higher concentration of OH groups caused out-diffusion of oxygen from
the substrate which oxidised the single layer deposited on them. The vendor of the quartz substrates,
however, does not differentiate between the substrates with and without the OH related peak. In the
same manner, there are other impurities that could be infiltrating the Si-QD structures which could
potentially limit their electronic and PV quality.
- 184 -
With this in mind we tested a p-i-n structure with an intended 64% Si by Volume in the SRO layers.
The structure was deposited on an oxidised Si wafer instead of quartz. The oxide on the wafer is
97nm thick and ensures that there is no possible electrical contact between the substrate and the SiQD devices.
The open circuit voltage measured in seven devices under 1sun illumination varied between 420 mV
and 440 mV. However, these devices are thermally isolated due to the thick oxide present on the
substrate. As the device is illuminated, it heats up due to the light absorbed in the film however, due
to the thermal insulation caused by the substrate, a temperature gradient is bound to exist with the
highest temperature always being on the side of the device. An advantage of measuring the I-V
properties of these devices as a function of temperature, is the possibility of determining the electronic
bandgap of the material used in the devices. The uncertainty in the actual temperature of the device
results in an uncertainty in the determination of this bandgap. Fig. 4.5.17 presents a plot of VOC vs.
temperature measured on a 1mm2 Si-QD solar cell on a thermally oxidised Si wafer.
Figure 4.5.17: VOC vs. Temperature measured on a 1mm2 Si-QD solar cell on a thermally oxidised Si
wafer. The bandgap extracted depends on the actual temperature of the device measured.
The bandgap can be estimated based on the intercept on the y-axis of the linear fit to the measured
data. This intercept is then the bandgap plus a correction factor in the order of 3kTaverage. Based on this
approach, we conclude that the effective electronic bangap of the Si-QD layer is somewhere between
0.84 and 0.91 eV. These values are well below the bandgap estimated from the PL peak emission
reported in Fig. 4.5.16 (~1.4 eV). The large discrepancy can be explained if the Si-QD material has a
very high density of defects bellow and above the conduction and valence band edges respectively.
We note that a lower electronic bandgap than anticipated will make it difficult to obtain open circuit
voltages as large as one would expect from a material with a bandgap given by the peak of the PL
emission. Having identified this issue we can now concentrate on reducing the density of electronic
defects rather than tackling the resistivity or varying the diameter of the Si-QD.
4.5.2.9 Other Si nanostructure materials studies
The use of matrices other than SiO2 can give advantages in carrier transport. Or the use of NCs other
than Si can give lower processing temperatures and different possibilities for bandgap engineering.
- 185 -
4.5.2.9.1 Carrier tunnelling transport in Si QD superlattices
Transport properties are expected to depend on the matrix in which the silicon quantum dots are
embedded. As shown in Fig. 4.5.18 different matrices produce different transport barriers between
the Si dot and the matrix, with tunnelling probability heavily dependent on the height of this barrier.
Si3N4 and SiC give lower barriers than SiO2 allowing larger dot spacing for a given tunnelling current.
SiO2
Si3 N 4
3.2 eV
SiC
c-Si
1.1 eV


8m *
Te  16 exp  d

E

2


1.9 eV
0.5 eV
c-Si
1.1 eV
c-Si
1.1 eV
0.9 eV
2.3 eV
4.7 eV
Figure 4.5.18: Bulk band alignments between silicon and its carbide, nitride and oxide. Tunnelling
probability between QDs separated by d depends exponentially on the square root of the barrier height
(ΔE1/2) multiplied by d. (e.g. [4.5.22] p244)
The results suggest that dots in a SiO2 matrix would have to be separated by no more than 1-2 nm of
matrix, while they could be separated by more than 4 nm of SiC. Fluctuations in spacing and size of
the dots can be investigated using similar calculations. It is also found that the calculated Bloch
mobilities do not depend strongly on variations in the dot spacing but do depend strongly on dot size
within the QD material [4.5.18]. Hence, transport between dots can be significantly increased by
using alternative matrices with a lower barrier height, ∆E. For the same tunnelling current the spacing
of QDs can increase for oxide to nitride to carbide matrix.
4.5.2.9.2 Hetero-Interlayers
Researchers: Dawei Di, Zhenyu Wan, Gavin Conibeer
Experiments have been carried out in which the layer between the Si rich layers is a different
dielectric. Si3N4 (nominally stoichiometric) has two advantages. The higher density of the nitride
restricts diffusion of Si to nucleating nanocrystals during annealing. Thus size control and a more
mono-disperse size distribution of the nanocrystals result. Secondly the lower barrier height of the
nitride, as compared to the oxide within the Si rich layers, leads to a higher tunneling probability in
the growth direction, whilst good quantum confinement is maintained in the nanocrystals planes.
Initial results on this approach show promising decreases in the vertical resistivity of such Si QD
nanostructures with SiNx interlayers [4.5.23]. Figure 4.5.19 shows I-V curves measured on vertical
samples on a conducting Si wafer substrate, for both SiNx or SiO2 interlayers, each doped with either
B or P. In each case the substrate used was of the same carrier type as the nominal type of the doped
layer on top. The most striking feature of the data is the much lower resistivity of the B doped nitride
interlayer samples as compared to the B doped oxide interlayer sample. Whilst for P doping the oxide
and nitride interlayer samples show almost no difference in resistivity.
Hence there is some evidence for significantly reduced resistivity for nitride as compared to oxide
interlayers, for the B doped samples. But the similar resistivities for the two interlayers for P doping
is not as first expected and requires further investigation. The explanation is related to the differing
effects on Si QD crystallisation of P and B.
- 186 -
1.0E+01
1.0E+02
P2O5 doped samples
nitride barrier
oxide barrier
2
Current density (A/cm )
nitride barrier
oxide barrier
1.0E+01
2
Current density (A/cm )
B doped samples
1.0E+00
1.0E+00
1.0E-01
1.0E-01
1.0E-02
1.0E-02
0.0
0.2
0.4
0.6
Bias voltage (V)
0.8
1.0
1.2
0.0
0.2
0.4
0.6
Bias voltage (V)
0.8
1.0
1.2
Figure 4.5.19: Logarithmic I-V curves for Si QD samples grown with either nitride or oxide
interlayers, and subsequently annealed in H2 forming gas at 350°C. (a) B doped samples: the nitride
interlayer shows a dramatically lower resistivity; and (b) P doped samples: the resistivities of the
nitride and oxide interlayer samples are similar. Insets show schematic diagrams of QD sizes and
spacings.
Figure 4.5.20 shows the XRD spectra for SiO2 and SiNx interlayers between SRO layers. Figure
4.5.20(a) shows that for the oxide interlayer (i.e. for the SiO2-SRO-SiO2 structure) the effect of P (B)
increasing (decreasing) the Si QD size is demonstrated in the narrowing (broadening) of the XRD
peaks for all of the {111}, {220} and {311} Si peaks. The fact that all these principal peaks are
narrowed (broadened) to the same extent for a given sample indicates that the effect is due to an
increase (decrease) in QD size allowing a narrower (broader) range of angles that satisfy the Bragg
condition. This is in contrast to inhomogeneous strain or a disordered structure in the nanocrystals
which, if it were present, would distort lattice spacing differently for different plane families. Hence
these QD sizes can be estimated using the Scherer equation for peak broadening to give 5.7nm and
14.8nm for B and P, respectively [4.5.23]. Thus indicating that P incorporation significantly increases
the size of Si QDs. But in Figure 4.5.20(b) for the nitride interlayer the QD sizes are much closer at
3.3nm and 5.0nm for B and P, respectively. This suggests that the nitride interlayers suppress the
growth of larger (smaller) nanocrystals for P (B) doping respectively.
(a)
(b)
Figure 4.5.20: XRD comparison of P and B doped multilayers for both (a) SiO2 and (b) SiNx
interlayers between SRO layers, all on quartz substrate.
The likely reason for this uniformity of QD sizes for the nitride interlayer with P and B doping, as
compared to the oxide interlayer, is that the presence of a strong gradient in N concentration at the
interface between the silicon rich oxide and silicon nitride layers will tend to promote heterogeneous
nucleation whether or not B or P are present. This will tend to suppress the differing QD sizes in P
and B material caused by their differing nucleation behaviour, thus reducing the difference in QD
spacing and its differing effects on resistivity.
This suggests a reason for the differing I-V behavior shown in Fig. 4.5.19 based on the opposite
effects on Si QD crystallisation of P and B. For the oxide interlayer samples the QDs are expected to
- 187 -
be smaller with B and larger with P incorporation. The large size of the QDs in P doped material are
such that they penetrate several layers and significantly reduce the number of tunneling events
required in the vertical current transport direction. This will tend to reduce resistivity. But in the
nitride interlayer the much more similar QD sizes for P and B doping shown in Fig. 4.5.20, mean this
effect will be significantly reduced. Hence for B doping in Fig. 4.5.19 (a) the QD sizes are very
similar and the presence of the lower barrier height nitride interlayer dominates in reducing resistivity.
Whilst for the P doping of Fig. 4.5.19 (b) the larger QDs in the oxide interlayer material reduce
resistivity to approximately the same extent as does the lower barrier height of the nitride in the
nitride interlayer material even though this latter has larger spacing between the smaller QDs.
4.5.2.9.3 Alternative group IV materials for tandem cells
Researchers: Sammy Lee, Shujuan Huang, Zhenyu Wan, Gavin Conibeer
Ge nanocrystals in SiO2, made by a similar phase separation from solid solution process, are being
investigated as an alternative material with a lower precipitation temperature and wider potential
range of effective band gap than Si nanocrystals. The temperature range employed is typically around
650°C, but with substrate heating during growth can be as low as 350°C [4.5.24]. These materials
demonstrate a weak p-type conductivity without intentional doping and can be made more p-type by
addition of Sb during growth, although the exact mechanism of doping is still under investigation.
Other analogues involving Sn nanocrystals in dielectric matrices have been fabricated. These require
even lower processing temperatures, around 250-300°C, and can be annealed in situ. However, Sn
nanocrystals formed in SiO2 matrix tend to undergo significant oxidation [4.5.25]. In a Si3N4 matrix
there is no oxidation but the tetragonal β form of Sn always forms, because the β-Sn allotrope is more
thermodynamically stable than the cubic α-Sn allotrope [4.5.26]. But as β-Sn is also semi-metallic as
compared to the semiconducting α-Sn, this also means that no semiconducting optical or electronic
properties of these Sn nanocrystal materials have yet been measured.
Other alternatives to group IV elements in dielectric matrices are alloys of group IV elements.
Several of these are attractive for controlling the band gap in a group IV tandem cell as summarized in
[4.5.27]. SiC alloys with excess Si have been investigated both as material with precipitated Si and
SiC nanocrystals and as amorphous unannealed alloys. Band gap control over the range of about 1.7
to 2.2eV is possible for the Si nanocrystals material by control of nanocrystal size [4.5.28]. Si:Ge
alloys are of course well known as a means of continually varying the band gap from 0.7 to 1.1eV.
But whilst they are suitable in a tandem cell for cells under a Si cell, they are not suitable for a tandem
cell element on top of a Si cell.
Alloying of Sn with small percentages of Ge can stabilize the cubic α-Sn phase [4.5.29]. The Ge rich
Ge:Sn alloy can also be grown. These alloys have band gaps from 0.2-0.7eV and are hence
potentially useful for cells to go under a Si cell in a tandem, but again are not suitable for cells with
higher band gap than Si. However, the Ge:Sn alloy can also be used to stabilize the α phase in
nanocrystals incorporated in a dielectric matrix. SiGeSn alloys are a variation of Sn:Ge alloys. The
addition of less than 1% of Si allows lattice matching of the alloy with Ge substrate [4.5.30] with
addition of more Si also allowing a higher range of direct band gap of 0.8-1.4eV [4.5.31]. The
advantage of the ternary alloy/compound is that the extra degree of freedom allows independent
control of band gap and lattice parameter, which in turn allows for better quality materials.
Ge:C and Sn:C alloys and GeC and SnC compounds are interesting as tandem cell materials [4.5.27].
Ge:C offers a potentially wide band gap range from 0.6 to 1.1eV and a fairly continuous range of
composition. Several of the material properties of GeC have been calculated [4.5.32] and GeC thin
films can be sputtered [4.5.33] with material approaching stoichiometric GeC for high methane gas
flows. In addition high sputter powers lead to a greater degree of sp3 hybridization of carbon bonds
thus leading to cubic diamond or zinc-blende structures as opposed to the graphitic structures for sp2
hybridization [4.5.34].
- 188 -
4.5.2.10 Transfer to Si QD growth using PECVD
Researchers: Tian Zhang, Ivan Perez-Wurfl
Collaboration with: Manuel Mogano, University of Milan; Birger Berghoff, RWTH Aachen
Sputtering is a high energy process in which many defects are incorporated in deposited material, it is
also relatively slow for thin film deposition. A transfer of the technology to plasma assisted chemical
vapour deposition (PECVD) has the potential advantages of improving material quality for Si QD
nanostructures in SiO2 matrix and of reducing long term material costs.
PECVD has been used to date primarily for growth of Si QD nanostructures in SiNx and in SiC
[4.5.35, 4.5.36]. It has now been extended to a SiO2 matrix.
Figure 4.5.21: XRR measurements of (a) as deposited PECVD multilayers of SRO/SiO2 on quartz
substrate (b) after annealing at 1100oC.
Figure 4.5.21 shows the X-ray reflectivity (XRR) measurements of multilayer SRO/SiO2 material as
deposited by PECVD and after annealing at 1100oC. They clearly show a good quality multilayered
structure as deposited with reduced quality after annealing. This is very consistent with the formation
of Si nanocrystals breaking up the multi-layer arrangement.
Figure 4.5.22: Raman spectroscopy of annealed multilayer PECVD sample (blue line) compared to cSi wafer.
Figure 4.5.22 shows Raman data on the annealed multilayer nanostructure. The clear asymmetric
broadening of the peak at around 510cm-1 strongly indicates formation of Si nanocrystals.
- 189 -
4.5.2.11 Summary of Group IV nanostructures for tandem cell elements
Modelling of the parameters of Si nanostrucrue growth and measured performance has indicated that
the band gap of these materials is dominated by defects. This is allowing a more systematic
investigation of the material quality in order to move towards higher quality material to give higher
VOC devices. Currents have been improved through light trapping using both rear reflection and
plasmonics. On measuring devise significant temperature increase is seen because of the high
resistivity, therefore lower temperature measurements of Voc have seen an increase to over 500mV for
devices that are in effect locally at room temperature. Reduced resistivity has been seen in structures
with nitride interlayers and other group IV elements have been incorporated into Ge QW and Ge:C
alloys. Further improvement of device parameters is expected as greater understanding of the
underlying material properties is gained with the current robust characterisation methodology.
4.5.3 Hot Carrier Cells
Researchers: Shujuan Huang, Santosh Shreshtha, Dirk König, Yukiko Kamikawa, Robert Patterson,
Pasquale Aliberti, Binesh Puthen Veettil, Ivan Perez-Wurfl, Andy Hsieh, Yu Feng, James Rudd,
Pengfei Zhang, Yao Yao, Hongze Xia, Neeti Gupta, Yuanxun Liao, Suntrana Smyth, Xi Dai, Pei
Wang, Simon Chung, Martin Green, Gavin Conibeer
Hot carrier solar cells offer the possibility of very high efficiencies (limiting efficiency above 65% for
unconcentrated illumination) but with a structure that could be conceptually simple compared to other
very high efficiency PV devices – such as multi-junction monolithic tandem cells. For this reason, the
approach lends itself to ‘thin film’ deposition techniques, with their attendant low costs in materials
and energy usage and facility to use abundant, non-toxic elements.
An ideal hot carrier cell would absorb a wide range of photon energies and extract a large fraction of
the energy to give very high efficiencies by extracting ‘hot’ carriers before they thermalise to the band
edges. Hence an important property of a hot carrier cell is to slow the rate of carrier cooling to allow
hot carriers to be collected whilst they are still at elevated energies (“hot”), and thus allowing higher
voltages to be achieved from the cell and hence higher efficiency. A hot carrier cell must also only
allow extraction of carriers from the device through contacts which accept only a very narrow range
of energies (energy selective contacts or ESCs). This is necessary in order to prevent cold carriers in
the contact from cooling the hot carriers, i.e. the increase in entropy on carrier extraction is minimized
[4.5.37]. The limiting efficiency for the hot carrier cell is over 65% at 1 sun and 85% at maximum
concentration – very close to the limits for an infinite number of energy levels [4.5.1, 4.5.38, 4.5.39].
Fig. 4.5.23 is a schematic band diagram of a hot carrier cell illustrating these two requirements.
h+
selective
energy contact E
Si
S
O
E
Q2
ua
(
(
L
O
TA
Ef(p)
(c)
((
s
Figure 4.5.23: Band diagram of the hot carrier cell. The device has four stringent requirements:
a) To absorb a wide range of photon energies; b) to slow the rate of photogenerated carrier cooling in
the absorber; c) To extract these ‘hot carriers’ over a narrow range of energies, such that excess
carrier energy is not lost to the cold contacts; d) to allow efficient renormalisation of carrier energy
via carrier-carrier scattering.
- 190 -
Modelling of Hot Carrier efficiencies has progressed with implementation of real material properties
to give more realistic efficiencies for InN which include Auger processes and more realistic contact
structures. Significant progress has been made on demonstrating resonance in double barrier selective
energy structures. Further work on triple barrier double Si QW structures has been carried out for
rectifying ESCs. This is complemented by improvements in 2/3D modelling of transport in these
ESC structures. For absorbers, modelling of nanocrystals superlattice arrays has been applied to real
material systems. The growth of such systems in both III-V QD superlattices with collaborators and
with colloidal Langmuir-Blodgett dispersion of Si nanocrsystals has produced structures which are
now being characterised for their modulation of phononic properties. Also measurement of carrier
cooling rates has been extended to other large phononic gap bulk materials including InN,
demonstrating the importance of material quality. Design of structures for hot carrier cells which
should be practical and realisable has developed, with the device properties more carefully specified
and plans for fabricating such structures in real devices.
4.5.3.1 Modelling of hot carrier solar cell efficiencies using optimised absorber and ESCs
Researchers: Yu Feng, Pasquale Aliberti, Hongse Xia, Gavin Conibeer
The performance of a hot carrier solar cell device has been evaluated taking into account the nonequilibrium distributions of electrons and holes as well as different rates of Coulomb scattering
between charge carriers. This includes processes of electron-electron (e-e) scattering, hole-hole (h-h)
scattering and electron-hole (e-h) scattering. Their effects on the carrier renormalization and the
electron-hole energy transfer have been quantitatively examined under the relaxation time
approximation (RTA). Other processes relevant to an operating device, including photo-generation,
radiative recombination, carrier extraction, carrier cooling (inelastic carrier scattering) and impact
ionisation/Auger recombination (AR-II), have also been incorporated with RTA.
The 1000-sun performance has been predicted for the cell configuration of an intrinsic InN absorber
(thickness: 50nm) with two symmetric ESCs, in accordance with previous work. The maximum
efficiencies have been computed with different relaxation time parameters, following the A and B
given in the following equation.
where
,
,
are the relaxation time of the respective processes (i.e. e-e/h-h/e-h scattering); N
is the carrier concentration.
The results are shown in Fig. 4.5.24. In the upper figure efficiency decreases with A, suggesting
insufficient rates of e-e and h-h elastic scattering could reduce the efficiency. This is because the
carriers with energies within the transmission window of the ESCs will be depleted unless the elastic
scattering is quick enough to supply new carriers with these energies. The curve at B = 108 is almost
flat, for e-h elastic scattering quick enough to prevent carrier depletion, noting that e-h scattering
allows both carriers to exchange their energies, providing possibilities of supplying carriers with the
energies of extraction. For the same reason, with negligible e-e and h-h scattering (A > 10−16) the
efficiency tends toward a constant value instead of dropping to zero. The lower figure in Fig. 4.5.24
shows that the efficiency decreases with the relaxation time of e-h scattering. It is partly due to the
carrier depletion especially with insufficient e-e and h-h scattering, at A = 10−14. With sufficient e-e
and h-h scattering, for example the equilibrium limit in which carriers are completely re-normalised,
the same trend still exists. In fact a long relaxation time of e-h scattering blocks the energy flow from
electrons to holes, allowing the holes to relax rapidly due to their fast thermalisation events compared
to electrons hence leaving the electron population hot. This would limit the hole flux out of the device
since only holes with an elevated energy and within a small energy window could be extracted. On
the contrary, the efficiency goes up with B if we replace the hole ESC with a normal contact, one that
- 191 -
is transmissive for all holes, since less e-h scattering means less energy loss from hot electrons. To
optimize the cell performance, the selection of the hole contact depends on the e-h scattering rate in
the absorber. A normal non-selective valence band contact is beneficial for slow scattering rates (B >
1011) while the ESC configuration adopted here has a greater advantage for rapid scattering rates (B <
1011).
Figure 4.5.24: Maximum efficiency with different A and B. The e-h isolation limit indicates no
scattering between electrons and holes. The equilibrium limit indicates ultrafast e-e scattering and h-h
scattering.
The energy-dependent distributions of electrons and holes are shown in Fig. 4.5.25, corresponding to
the Maximum Power Points (MPPs) with different A and B values. From left to right the figures show
the trend of increasing carrier temperature with e-e and h-h scattering rates, as with less carrier
depletion a faster extraction rate can occur reducing the retention time of carriers. From top to bottom
the figures demonstrate a larger deviation between electron distribution and hole distribution with less
e-h scattering. The carrier depletion within the transmission window of ESCs is illustrated in the
corresponding insets. The resupply of depleted holes is more rapid than electrons because of the
proportionality between the scattering rate and the carrier effective mass. With a small A, i.e. fast ee/h-h scattering rates, the supply of holes is sufficient for the carrier extraction, leading to a larger
current density and a higher overall efficiency. In this case the electron depletion could be even more
severe since the increase in current depletes electrons more quickly. This is observed with B = 1010
and 1012. Apart from the region of carrier depletion the distributions are close to Fermi functions,
which may be used for defining a “quasi-temperature” for each type of carrier.
- 192 -
Figure 4.5.25: Energy-dependent distributions of electrons (black solid line) and holes (red dashed
line) when operating at the Maximum Power Points (MPPs) with different A and B: A = 10−14 (left),
A = 10−19 (right); B = 1010 (upper), B = 1012 (middle), B = 1014 (lower). The depleted region within
the transmission window is enlarged in the corresponding inset.
The calculated current-voltage curves are shown in Fig. 4.5.26. With faster e-h scattering (smaller B)
the short- circuit current density is increased while the open-circuit voltage is reduced. It is difficult to
extract holes at the elevated energy level required by ESCs unless e-h scattering is quick enough to
resupply highly energetic holes. On the other hand the increase of voltage with slower e-h scattering
comes from the hotter electron distribution at the open-circuit condition, with less energy lost to the
thermalised holes. If a HCSC is designed with the narrow high energy transmission window for holes,
as adopted in this work, the maximum efficiency is always higher for quicker e-h scattering. The
drawback of slow e-h scattering would be pronounced for insufficient e-e and h-h scattering rates,
since the significant carrier depletion would inhibit the carrier extraction resulting in a colder carrier
distribution. Hence the quasi-temperatures of both electrons and holes are low at MPPs, if e-h
interaction is weak, while in the lower inset the contrast between them increases. The carrier quasitemperature insets also indicate a test of the validity of commonly used assumptions: electrons and
holes share the same temperature only for fast e-h scattering (B < 1010) while holes are completely
thermalized when e-h scattering rates are slow (B > 1014).
- 193 -
For real materials the choice of parameters A and B can be determined by computing the realistic
scattering rates between carriers. We have calculated the scattering relaxation times for bulk cubicInN, taking into account both the direct Coulomb potential and the exchange potential. The results
closely follow the relation shown in the equation above for relaxation times, with A = 2.5×10−19 and B
= 2.4×1011.
The results show that the maximum efficiency varies significantly with the carrier scattering
properties due to energy-dependent carrier depletion and asymmetric statistics of electrons and holes.
With different hole contacts the dominant loss mechanism switches: with an ESC at an elevated
energy the hole extraction is blocked until significant heat is transferred from hot electrons. Thus at
the hole side an ESC is more suitable for fast e-h scattering while a normal contact is beneficial for
slow e-h scattering rates. In addition, the RTA model suggested here provides the possibility of
analyzing any realistic HCSC device. It also serves to optimize the transmission properties of both
contacts. According to this work it heavily depends on the carrier scattering properties in the absorber.
4.5.3.3 Energy Selective Contacts (ESCs)
Researchers: Binesh Puthen-Veettil, Yu Feng, Andy Hsieh, Yuanxun Liao, Santosh Shrestha, Gavin
Conibeer
One practical implementation of the requirement for a narrow range of contact energies is an energy
selective contact (ESC) based on double barrier resonant tunnelling. Tunnelling to the confined
energy levels in a quantum dot layer embedded between two dielectric barrier layers, can give a
conductance sharply peaked at the line up of the Fermi level on the ‘hot’ absorber side of the contact
with the QD confined energy level. Conductance both below this energy and above it should be very
significantly lower. This is the basis of the current work on double barrier resonant tunnelling ESCs.
4.5.3.3.1 Modelling optimised materials for Energy Selective Contacts
Energy selective contacts (ESCs) for a HC solar cell can be implemented using different materials and
structures. The main property that needs to be achieved is precise energy selectivity in carrier
extraction. Different models have been developed to identify possible materials for ESCs and to
optimize their properties. Amongst the different possible material configurations, group IV and group
III-V double barrier resonant tunnelling structures look to be the most promising.
Figure 4.5.26: Schematic diagram of a double barrier structure. Silicon QDs are formed in a dielectric
matrix that may not be the same as the barrier material. The growth direction of the structure is along
the z axis.
Modelling of group IV materials has shown that the structure with the highest potential to be used as
ESC consists of double barrier structure (DBS) with quantum dots (QDs) embedded in a dielectric
matrix.
Modelling of electrical properties for the structure shown in Fig. 4.5.26 has been performed for Si
QDs in Si based dielectric materials. The reasons for focusing on Si in the first instance for group IV
- 194 -
materials are abundance and well established growth techniques. Structures consisting of Si QDs and
the following dielectric barriers have been modelled: SiO2, Si3N4 or SiC. Results show that optimal
energy selectivity properties are obtained with higher lateral confinement within the middle layer of
the structure and higher vertical conductivity. Amongst the material taken into account in the
simulations, structures consisting of Si QDs in a SiO2 with SiC barriers have shown the best overall
energy extraction properties.
ESCs for a III-V based HC solar cell can be realized using a QW structure in a DBS, probably
requiring either MBE or MOVPE growth. Modelling results have shown that, for a HC cell based on
an InN absorber, the optimum configuration for ESCs is constituted by an InN/InXGa1-XN DBS
[4.5.40]. This structure allows modulation of the extraction energy and shape of the transmission
probability function independently for the electrons and holes contacts. The value of extraction energy
can be modified by engineering the stoichiometry and the thickness of the QW structure. In general
the optimal extraction energy for electrons and holes is different, due to the different values of
effective mass. Thus, the hole ESC QW should be physically thinner and have a higher barrier than
the electron ESC QW to achieve similar currents at reasonable extraction energies, as illustrated in
Fig. 4.5.27.
Figure 4.5.27: Schematic of a HC solar cell based on a InN absorber and InN/InGaN ESCs.
The transmission coefficient of a double-barrier semiconductor structure has been numerically
calculated using a tight-binding method. The phonon scattering during the transmission has been
incorporated with a Greens function method. This constructs a framework for simulating a real
resonant tunneling system with ultra-thin layers. It will help in optimizing the material system for the
ESCs. Some sampling results for the DBS of AlAs/GaAs/AlAs are shown in Fig. 4.5.28.
Figure 4.5.28: The total transmission (y axis: 0~1) profile as a function of the normally-penetrating
electron energy (x axis: eV) through the zinc-blende DBS AlAs(3 unit cells)/GaAs (3 unit cells)/AlAs
(3 unit cells), with different phonon scattering properties. Left: No phonon scattering considered;
Right: with phonon scattering and the initial phonon occupation number = 0.
- 195 -
Here we mainly aim to utilize the 1st energy peak. From the transmission calculation of an
AlAs/GaAs/AlAs double-barrier structure, this peak is visible and reasonably significant if only
normally incident electrons are considered. (These constitute the majority of energy transmission in
any case.) With the consideration of phonon scattering, side-peaks appear which are not favorable for
energy-selective extraction. This indicates a preference of choosing less polar materials. Besides, the
addition of arbitrary directions off normal may broaden the transmission peak with a little loss on the
selectivity. Such a Greens function model can treat any type of atomic system in terms of the
scattering problem. It is ready for use as a tool to choose the optimal material system for the energyselective transmission.
4.5.3.3.2 Triple barrier resonant tunnelling structures for carrier selection and rectification
Researchers: Gavin Conibeer, Andy Hsieh, Santsoh Shrestha
A normal energy selective contact (ESC) using double barrier resonant tunnelling is symmetric and
does not rectify electrons from holes. An ESC which uses triple barrier resonant tunnelling through
two separately tuned quantum well or quantum dot layers can give a much greater energy selectivity
and also is a filter to separate electrons and holes in opposite directions (i.e. rectification). The miniband line-ups of the confined levels need to be tuned by the well thicknesses, effective mass of
electrons and holes and barrier heights such as illustrated schematically in Fig. 4.5.29. With
appropriate choice of these for the left contact there can be an alignment of the confined levels such as
to give a resonant channel in the conduction band for electrons but misalignment of levels in the
valence band such that there is no channel for holes. Similarly for the right contact the valence band
confined levels are aligned to allow holes to pass but misaligned in the conduction band so as to block
electrons. For non-equal electron and hole effective masses (as is the case for most materials) this
condition will in general be met [4.5.41, 4.5.42].
HotCarriersolarcellAbsorber
Electron
collector
Hole
collector
Figure 4.5.29: Triple barrier double well resonant tunnelling energy selective contacts. Wells are of
different width designed to give allignment of 1st and 2nd confined levels in the conduction band on
the left and for 1st and 2nd confined levels in the valence band on the right. In general for non-equal
effective masses allignment of first and second confined energy levels for electrons will not give the
same condition for holes thus giving electron/hole rectification and a preferred collection direction in
the external circuit.
The collection channels for electrons and holes would ideally be aligned with the peak occupancies of
energy of the hot electron and hot hole populations respectively, which in turn will depend on the
respective temperatures of these populations. This condition will give the highest efficiency, but as
discussed in Section 4.5.3.1, for high electron-electron and hole-hole carrier scattering rates, it is not
very critical.
- 196 -
Ideally there would be two of these triple barrier double QW/QD devices on either side of the
absorber, but as there is much less energy in the high effective mass hole population a simpler double
barrier or simple band contact on the hole side is also possible, without much loss in collected energy,
as investigated in Section 4.5.3.1. Various materials combinations could be used for such a design,
but specifically we are looking at III-V triple barrier QWs based on the InGaN/GaN system. Again
this is addressed to some extent in Section 4.5.3.1.
The exact thickness of QWs or size of QDs and barrier thicknesses and contact work functions will
have to be tuned for the particular confined energy levels involved in each material. Use of more than
two QW or QD layers on either side is also possible. This will increase energy selection by requiring
even more careful alignment.
This is the first application of this concept to energy selective contacts for hot carrier, Hot Lattice or
Thermoelectric cells, but the concept itself was first suggested by [4.5.43] and has been used for
energy filtering in quantum cascade lasers for several years. It is also not the first application of the
concept to solar cells as [4.5.44], at Imperial College, have used the concept for a multi-QW solar cell
design. But it is the first application to Hot Carrier or Hot Lattice devices and its ability to rectify is
likely to be a very significant advantage.
4.5.3.3 Hot Carrier Absorbers: slowing of carrier cooling
Carrier cooling in a semiconductor proceeds predominantly by carriers scattering their energy with
optical phonons. This builds up a non-equilibrium ‘hot’ population of optical phonons which, if it
remains hot, will drive a reverse reaction to re-heat the carrier population, thus slowing further carrier
cooling. Therefore the critical factor is the mechanism by which these optical phonons decay into
acoustic phonons, or heat in the lattice. The principal mechanism by which this can occur is the
Klemens mechanism, in which the optical phonon decays into two acoustic phonons of half its energy
and of equal and opposite momenta [4.5.45]. The build-up of emitted optical phonons is strongly
peaked at zone centre both for compound semiconductor due to the Frölich interaction and for
elemental semiconductors due to the deformation potential interaction. The strong coupling of the
Frölich interaction also means that high energy optical phonons are constrained to near zone centre
even if parabolicty of the bands is no longer valid [4.5.46]. This zone centre optical phonon
population determines the dominant optical phonon decay mechanism is this pure Klemens decay.
4.5.3.3.1 Modelling of electron-phonon interactions:
Researchers: Yu Feng, Hongze Xia, Santosh Shrestha, Robert Patterson, Gavin Conibeer
Another important factor is the Frohlich interactions between hot electrons and longitude optical
phonon modes. The importance of investigating the fundamental particle interactions occurring in the
device is demonstrated by the significant dependence between the carrier energy relaxation time and
the output efficiency. The underlying mechanism of phonon bottleneck effect involves two parts, the
first is the block of Frohlich interaction between electrons and phonons; while the second is the block
of phonon decay from the optical modes into acoustic modes. Also the build-up of non-equilibrium of
acoustic phonons may also contribute to reduced carrier energy relaxations. In order to understand the
mechanism of slow relaxation rates in different types of materials the three interaction processes need
to be simulated.
The Frohlich interaction occurs in polar semiconductors and creates energy exchange between
electrons and longitudinal phonons. It affects the total energy and state transition rates by an
interaction Hamiltonian. The Frohlich interaction Hamiltonian for superlattice structures involves an
overlap integral over one dimension (the direction perpendicular to planes). The derivation of the
Hamiltonian involves the interaction between electrons and polarization fields induced by
longitudinal mode vibrations. By obtaining the divergence of the polarization field an effective
phonon envelope function is developed and incorporated in the overlap integral.
- 197 -
4.5.3.3.2 Phonon decay mechanisms
Researchers: Gavin Conibeer, Santosh Shrestha, Robert Patterson, Yu Feng, Hongze Xia, Suntrana
Smyth
An optical phonon will decay into multiple lower energy phonons. The processes require the
conservation of energy and momentum. However, additionally, the principal decay path must be into
two LA phonons only, which are of equal energy and opposite momenta, via an anharmonicity in the
lattice, this was suggested by Klemens [4.5.45] and has been shown to apply to a wide range of
materials.
Wide Phononic Gaps in III-Vs and analogues
For some compounds in which there is a large difference in masses of the constituent elements, there
exists a large gap in the phonon dispersion between acoustic and optical phonon energies. If large
enough this phonon band gap can prevent Klemens’ decay of optical phonons, because no allowed
states at half the LO phonon energy exist. InN is an example of such a material with a very large
phonon gap. The prevention of the Klemens’ mechanism forces optical phonon decay via the next
most likely, Ridley mechanism, of emission of one TO and one low energy LA phonon. Such a
mechanism only has appreciable energy loss (although still much less than Klemens’ decay) if there is
a wide range of optical phonon energies at zone centre. This is only the case for lower symmetry
structures such as hexagonal. For a high symmetry cubic structure, LO and TO modes are close to
degenerate at zone centre and the Ridley mechanism is severely restricted or forbidden. Unfortunately
cubic InN is very difficult to fabricate precisely because of the large difference in masses that give it
its interesting phononic dispersion.
Slowed cooling has been observed in some III-V compounds in which there is a large difference in
atomic mass. This has been shown in slowed carrier cooling for InN [4.5.47] and for slowed carrier
cooling in InP compared to the small mass ratio GaAs [4.5.48].
Analogues of InN with abundant elements, but also with narrow Eg include II-IV-VI compounds such
as ZnSnN2, IIIA nitrides such as LaN and YN which should have large phonon gaps, and IVA nitrides
such as ZrN and HfN which are readily available. Bi and Sb compounds should also have large
phonon gaps but are not abundant. Group IV compounds have large calculated gaps and small Egs, as
well as several other advantages [4.5.49]. Fig. 4.5.30 shows calculations of phonon dispersions for
group IV compounds which have large phonon band gaps, sufficient to block Klemens’ decay.
Figure 4.5.30: Adiabatic bond charge calculations of phonon dispersions for group IV compounds.
Phonon gaps increase as the mass ratio increases with those for GeC and SnC large enough to block
Klemens’ decay.
Slowed carrier cooling in MQWs
Low dimensional multiple quantum well (MQW) systems have also been shown to have lower carrier
cooling rates. Comparison of bulk and MQW materials has shown significantly slower carrier
cooling in the latter. Fig. 4.5.31 shows data for bulk GaAs as compared to MQW GaAs/AlGaAs
- 198 -
materials as measured using time resolved transient absorption by Rosenwaks [4.5.50], recalculated to
show effective carrier temperature as a function of carrier lifetime by Guillemoles [4.5.51]. It clearly
shows that the carriers stay hotter for significantly longer times in the MQW samples, particularly at
the higher injection levels by 1½ orders of magnitude. This is due to an enhanced ‘phonon
bottleneck’ in the MQWs allowing the threshold intensity at which a certain ratio of LO phonon reabsorption to emission is reached which allows maintenance of a hot carrier population, to be reached
at a much lower illumination level. More recent work on strain balanced InGaAs/GaAsP MQWs by
Hirst [4.5.52] has also shown carrier temperatures significantly above ambient, as measured by PL.
Increase in In content to make the wells deeper and to reduce the degree of confinement is seen to
increase the effective carrier temperatures.
The mechanisms for the reduced carrier cooling rate in these MQW systems are not yet clear.
However there are three effects that are likely to contribute. The first is that in bulk material
photogenerated hot carriers are free to diffuse deeper into the material and hence to reduce the hot
carrier concentration at a given depth. This will also decrease the density of LO phonons emitted by
hot carriers as they cool and make a phonon bottleneck more difficult to achieve at a given
illumination intensity. Whereas in a MQW there are physical barriers to the diffusion of hot carriers
generated in a well and hence a much greater local concentration of carriers and therefore also of
emitted optical phonons. Thus the phonon bottleneck condition is achieved at lower intensity.
The second effect is that for the materials systems which show this slowed cooling, there is very little
or no overlap between the optical phonon energies of the well and barrier materials. For instance the
optical phonon energy ranges for the GaAs wells and AlGaAs barriers used in [4.5.50] at 210285meV and 280-350meV, respectively, exhibit very little overlap in energy, with zero overlap for
the zone centre LO phonon energies of 285 and 350meV [4.5.53]. Consequently the predominantly
zone centre LO phonons emitted by carriers cooling in the wells will be reflected from the interfaces
and will remain confined in the wells, thus enhancing the phonon bottleneck at a given illumination
intensity.
Thirdly, if there is a coherent spacing between the nano-wells (as there is for these MQW or
superlattice systems) a coherent Bragg reflection of phonon modes can be established which blocks
certain phonon energies perpendicular to the wells, opening up one dimensional phononic band gaps
(analogous to photonic band gaps in modulated refractive index structures. For specific ranges of
nano-well and barrier thickness these forbidden energies can be at just those energies required for
phonon decay. This coherent Bragg reflection should have an even stronger effect than the incoherent
scattering of the second mechanism above at preventing emission of phonons and phonon decay in the
direction perpendicular to the nano-wells.
It is likely that all three of these effects will reduce carrier cooling rates. None depend on electronic
quantum confinement and hence should be exhibited in wells that are not thin enough to be quantized
but are still quite thin (perhaps termed ‘nano-wells’). In fact it may well be that the effects are
enhanced in such nano-wells as compared to full QWs due to the former’s greater density of states
and in particular their greater ratio of density of electronic to phonon states which will enhance the
phonon bottleneck for emitted phonons. The fact that the deeper and hence less confined wells in
[4.5.52] show higher carrier temperatures is tentative evidence to support the hypothesis that nanowells without quantum confinement are all that are required. Whilst several other effects might well
be present in these MQW systems, further work on variation of nano-well and barrier width and
comparison between material systems, will distinguish which of these reduced carrier diffusion,
phonon confinement or phonon folding mechanisms might be dominant.
- 199 -
Figure 4.5.31: Effective carrier temperature as a function of carrier lifetime for bulk GaAs as
compared to GaAs/AlGaAs MQWs: time resolved transient absorption data for different injection
levels, from Rosenwaks [4.5.50], recalculated by Guillemoles [4.5.51].
4.5.3.3.3 Hot carrier cell absorber requisite properties
The above discussion allows us to estimate the major properties required for a good hot carrier
absorber material. These are listed below in approximate order of priority, although their relative
importance may well change in light of future research:
1. Large phononic band gap (EO(min) - ELA) - to suppress Klemens’ decay, requires large mass
difference between elements.
2. Narrow optical phonon energy dispersion (ELO - EO(min)) – to minimise the loss by Ridley
decay, requires a high symmetry.
3. Small Eg < 1eV – to allow broad range of photon absorption.
4. A small LO optical phonon energy (ELO) – to reduce the amount of energy lost per LO phonon
emission.
5. A small maximum acoustic phonon energy (ELA). This maximises (EO(min) - ELA) and is
important if ELO is also small.
6. Good renormalisation rates in the material, i.e. good e-e and h-h scattering. This condition is met
in most semiconductors quite easily, with e-e scattering rates of less than 100fs. But it may be
comprised in nanostructures.
7. Good carrier transport in order to allow transport of hot carriers to the contacts.
8. Ability to make good quality, ordered, low defect material.
9. Earth abundant and readily processable materials.
10. No, or low, toxicity of elements, compounds and processes.
4.5.3.3.4 Hot carrier absorber: choice of materials
InN has most of these properties, except 4, 8 & 9, and is therefore a good model material for a hot
carrier cell absorber. InN is investigated below in combination with its alloys with GaN. However the
abundance of In is very low, so it is difficult to see InN as a long term material suitable for large scale
implementation. Hence analogues of InN, which retain its interesting property of a large phonon band
gap, are also investigated.
Analogues of InN
As InN is a model material, but has the problems of abundance and bad material quality, another
approach is to use analogues of InN to attempt to emulate its near ideal properties. These analogues
can be II-IV-nitride compounds, large mass anion III-Vs, group IV compounds/alloys or
nanostructures.
- 200 -
IIA
IIIA
IB
IIB
IIIB
IVB
VB
VIB
Be
B
C
N
O
Mg
Al
Si
P
S
Ca
Sc
Cu
Zn
Ga
Ge
As
Se
Sr
Y
Ag
Cd
In
Sn
Sb
Te
Ba
La
Au
Hg
Tl
Pb
Bi
Po
Er
Figure 4.5.32: Use of the periodic table to analyse possible analogue compounds of InN based on
atomic mass combination and electro-negativity.
II-IV-Vs: ZnSnN; ZnPbN; HgSnN; HgPbN
With reference to Fig. 4.5.32, it can be seen that replacement of In on the III sub-lattice with II-IV
compounds is analogous and is now quite widely being investigated in the Cu2ZnSnS4 analogue to
CuInS2 [4.5.54].
ZnGeN can be fabricated [4.5.55] and is most directly analogous with Si and GaAs. However, its
band gap is large at 1.9eV. It also has a small calculated phononic band gap [4.5.56]. ZnSnN has a
smaller electronic gap (1 eV) and larger calculated phononic gap [4.5.56]. It is however difficult to
fabricate, and also its phononic gap is not as large as the acoustic phonon energy making it difficult to
block Klemens decay completely. HgSnN or HgPbN should both have smaller Eg and larger
phononic gaps. These materials have not yet been fabricated [4.5.57].
Large mass cation:
The Bi and Sb compounds have large predicted phononic gaps and Bi is a relatively abundant
material, with only low toxicity [4.5.57]. BiB has the largest phononic gap but AlBi, Bi2S5, Bi2O3
(bismuthine) are also attractive. Similarly SbB has a large predicted phononic gap. That for AlSb is
the same size as the acoustic phonon energy and its band gap is 1.5eV, making it marginal as an
absorber material and similar to InP.
Group IIIA III-Vs
LaN and YN both have large phononic gaps whilst that for ScN is too small.
The Lanthanides can also form III-Vs. ErN and other RE nitrides can be grown by MBE. The
phononic band gaps of the Er compounds are predicted to be large, because of the heavy Er cation,
but its discrete energy levels make it not useful as an absorber, although the combination of properties
in a nanostructure could be advantageous.
Group IV alloy/compounds:
All of the combinations Si/Sn, Ge/C or Sn/C look attractive with large gaps predicted in 1D models.
However being all group IVs they only form weak compounds. Unfortunately SiC, whether 3C, 4H
or 6H, has too narrow a phononic gap. Nonetheless GeC does form a compound and is of significant
interest [4.5.58].
- 201 -
There are also several other inherent advantages of group IV compounds/alloys all of which are
associated with the four valence electrons of the group IVs which result in predominantly covalent
bonding:
a) The elements form completely covalently bonded crystals primarily in a diamond structure
(tetragonal is also possible as in βSn). However for group IV compounds the decreasing
electronegativity down the group results in partially ionic bonding. This is not strong in SiC and
whilst it tends to give co-ordination numbers of 4, can nonetheless result in several allotropes of
decreasing symmetry: 4c, 4h, 6h. However, as the difference in period increases for the as yet
theoretical group IV compounds, so too does the difference in electronegativity and hence also the
bond ionicity and the degree of order. For a hot carrier absorber this is ideal because it is just
such a large difference in the period which is needed to give the large mass difference and hence
large phononic gaps. All of GeC, SnSi, SnC (and the Pb compounds) have computed phononic
gaps large enough to block Klemens decay, and should also tend to form ordered diamond
structure compounds.
b) Because of their covalent bonding, the group IV elements have relatively small electronic band
gaps as compared to their more ionic III-V and much more ionic II-VI analogues in the same
period: e.g. Sn 0.15eV, InSb 0.4eV, CdTe 1.5eV. In fact to achieve approximately the same
electronic band gap one must go down one period from group IV to III-V and down another
period from III-V to II-VI: e.g. Si 1.1eV, GaAs 1.45eV, CdTe 1.5eV. This means that for group
IV compounds there is greater scope for large mass difference compounds whilst still maintaining
small electronic band gaps. A small band gap of course being important for broadband absorption
in an absorber - property 3 in the desirable properties for hot carrier absorbers listed above.
c) The smaller Eg would tend to be for the larger mass compounds of Pb or Sn. Which, to give large
mass difference, would be compounded with Si or Ge. This trend towards the lower periods of
group IV also means that the maximum optical phonon and maximum acoustic phonon energies
will be smaller for a given mass ratio - the desirable properties 4 and 5.
d) Furthermore, unlike most groups, the group IV elements remain abundant for the higher mass
number elements – desirable property 9. Property 10 is also satisfied because the group IVs have
low toxicity.
Nanostructures:
As discussed in Section 4.5.3.3.9, the phonon dispersion of QD nanostructures can be calculated in
the same way as compounds. Their phononic properties can be estimated from consideration of their
combination force constants. Hence it is possible to ‘engineer’ phononic properties in a wider range
of nanostructure combinations. Of the materials discussed above the Group IVs lend themselves most
readily to formation of nanostructures instead of compounds due to their predominantly covalent
bonding, which allows variation in the coordination number. Therefore the nanostructure approaches
of Section 4.5.3.3.9 are consistent with a similar description as analogues of InN, whether it be III-V
QDs, colloidally dispersed QDs or for core shell QDs.
4.5.3.3.5 Combined phononic / mnw absorber design
Researchers: Gavin Conibeer, Yu Feng, Hongze Xia, Robert Patterson, Suntrana Smyth
The multiple nano-well structures have demonstrated reduced carrier cooling by enhancing the
phonon bottleneck, probably at a stage prior to optical phonon decay. Large phononic band gap
materials would seem to have strong potential to block Klemens’ decay of optical phonons. A
structure combining both structures should give even greater reduction in carrier cooling rates as the
mechanisms should not interfere directly with each other and hence their effects should be additive. A
combined absorber for a hot carrier cell should look similar to Fig. 4.5.33.
This has narrow nano-wells of phononic band gap material, not thin enough to give quantum
confinement, separated by thin barriers. The barriers would either have an electronic band offset to
block hot carrier diffusion or an optical phonon energy offset with the nano-wells and probably a
coherent nano-well structure, or all three, depending which of the multiple nano-well (MNW)
- 202 -
mechanisms discussed above is dominant. This will enhance the phonon bottleneck either by
preventing hot carrier diffusion, by reflecting and confining optical phonons or by blocking certain
folded phonon modes, respectively, or all three. The phononic nano-wells themselves will further
enhance phonon bottleneck by preventing Klemens’ decay of optical to acoustic phonons. This
should maximise the phonon bottleneck and minimise carrier cooling for a given illumination
intensity. A lattice matched pair of large phononic band gap nano-well and barrier materials
(InN/InGaN) is suggested in Fig. 4.5.33 (b), whereas in (c) a wider range of possible InN analogue
phononic band gap materials could be used in a thin film structure with thin probably oxide barriers.
The thin barriers (a few nm) should facilitate tunnelling between wells and hence transport of the
carriers to the contacts.
Figure 4.5.33: Hot Carrier cell design combining Multiple Nano-Well (MNW) with phononic band
gap materials: (a) schematic and (b) band structure for InN/InGaN epitaxial multiple nanowell with
phononic band nano-wells; (c) band structure with nano-wels of analogues of InN with large phononic
band gap with thin film barriers.
4.5.3.3.6 Alternative quantum well structure with integrated energy selective contacts
Researchers: Dirk König, Binesh Puthen-Veettil
Use of quantum wells in the absorber region has the advantage of a continuous density of states in the
plane above the first confined energy level. This allows continuous absorption of photon energies
above this energy. With barriers of different phonon energy such that they reflect optical phonons,
phonons can be confined in the z direction and increase phonon bottleneck the consequent decrease in
carrier cooling rate. If in addition the materials and barrier are chosen such that they give two
confined energy levels below the barrier a significant distance apart, then these can select only a very
few carrier energies to be transported through the device, thus effectively playing the role of energy
elective contacts. This thus reduces the need for external energy selective contacts and allows the
contacts to be optimised for the appropriate Fermi level and work function alignment to give electron
collection and hole reflection at one contact and hole collection and electron reflection at the other,
thus giving the rectification required.
- 203 -
Figure 4.5.34: InAs/AlSb MQW on AlSb with GaP top electron contact layer to give integrated hot
carrier absorber and carrier energy selected structure without need for external ESCs.
Figure 4.5.34 shows such a structure with InAs wells and AlSb barriers, designed to give confined
energy levels such that the 2nd level aligns with the conduction band of heavily n-type GaP. The GaP
and AlSb contact materials act as hole and electron blocking layers respectively. The valence band
offset between InAs and AlSb is very small. This allows the AlSb to collect thermalised holes from
the valence band of the superlattice. InAs and AlSb have a very small lattice mismatch (1.3%) so the
whole structure could be grown epitaxially on an AlSb substrate [ 4.5.59, 4.5.60].
Such an integrated structure would give an ideal test bed for integration of absorber MQW and
internal energy selection of carriers which would not require external ESCs and which gives
rectification in the external circuit. The specific materials and their thicknesses required would not
allow optimisation for slowing of hot carrier cooling, but would be able to be optimised for
appropriately spaced energy selection and collection.
4.5.3.3.7 Phononic dispersion modelling to predict phonon lifetimes and decay rates in specific
III-V or core-shell structures
Researchers: Hongze Xia, Yu Feng, Robert Patterson, Gavin Conibeer
InN/ InxGa1−xN multiple quantum-well superlattices (MQW-SL) with wurtzite crystal structure are
studied as the absorber of the hot carrier solar cell. Such a structure will exploit both the significant
hot-phonon-bottleneck effects of the component materials and the known slowed carrier cooling in
MQWs. The former is due to the large contrast of the atomic masses, and hence a large band-gap
between high-lying and low-lying phonon modes. The phonon band-gap effectively stops Klemens
decay, i.e. one high-lying phonon decaying into two low-lying phonons, and produce hot population
on the polar modes, feeding energy back to electrons. InN and InxGa1−xN have very similar lattice
structures, with almost the same lattice constant. This benefits the solar energy conversion, in terms of
both the carrier transport and the reduction of recombination sites. The 1-dimensional superlattice
structure ensures a continuous electronic energy spectrum and hence a broad-band absorption. The
absorption is further enhanced by the small electronic band-gap of InN.
The dispersion relations of phonon modes in superlattices have been computed with a 1-dimensional
atomic-plane model. Since the epitaxial growth is intended to be along the high-symmetrical Γ − A
direction for both InN and InxGa1−xN layers, the periodicity of MQW-SL and hence the zone-folding
is along this direction. For simplicity we assumed that the mode frequency of each zone-folded miniband only depends on the vertical component of the wave-vector (i.e. along Γ − A). The reason for
this assumption partly comes from the high concentration of electron-emitted polar phonons around
the zone-center, where the dispersions are relatively flat compared to the mini-gaps. In wurtzite
- 204 -
structure periodic atomic planes with alternating elements align perpendicular to the Γ−A direction. If
only considering the vertical component of wave-vectors, all atoms in one atomic plane vibrate in
phase and can be treated as one uniform displacement. The phononic model adopted here treats a 1dimensional chain with an equal length of the superlattice periodicity, taking into account the planeto-plane force constants. The plane-to-plane force constants are calculated by adopting the
conventional Keating potentials for all bonds. A sample phononic dispersion relation is shown in Fig.
4.5.35. The electronic structures, involving the wavefuntions and energies of all the states, can be
calculated by using the Kronig-Penney model for superlattices. In equilibrium conditions space
charges occur in the well layers (negative) and in the barrier layers (positive) due to the difference of
their electron affinities. The electron affinity of InGaN is significantly lower than that of InN, giving a
larger off-set of the conduction band edge rather than that of the valence band edge. A sample
electronic dispersion relation is shown in Fig. 4.5.36.
Figure 4.5.35: The left figure indicates the phonon decay paths and the sample dispersion for a SL
structure with 3 layers of unit cells inside each layer (barrier or well). The right figure shows the
phonon modulation function of a representative mode from each category of modes, computed from
the eigenfunctions of the lattice dynamic equation. The transverse modes have similar modulation
functions and hence only that part of the longitude modes is shown.
Figure 4.5.36: The electronic dispersion relation and the zone-center wave functions for a SL
structure with 3 layers of unit cells inside each layer (barrier or well). The wave vector is along the
Gamma-A crystal direction.
The calculation of the rate of polar interaction between hot electrons (here emission by holes is not of
concern) and polar phonons are based on the Frohlich-type Hamiltonian and the 1-st order
- 205 -
perturbation theory. A hot reservoir of electrons at 1000K is assumed, with heat transferred to the cold
reservoir of lattice modes at 300K. The energy relaxation times referring to the polar phonon emission
are illustrated in Fig. 4.5.37, for all the MQW-SL configurations, i.e. for different well and barrier
thicknesses and different barrier materials. 18×18 combinations of the thicknesses are sampled for
representing all the possible structures from 2nm to 9nm. The number of combinations comes from
the fact that a complete well/barrier layer should include an integer number of unit cells. Figure 4.5.37
is generated by interpolating the relaxation time data of the sampling combinations.
Figure 4.5.37: Energy relaxation times of the hot electron system (1000K), in the reservoir of
phonons (300K), with superlattice structures of different well/barrier thicknesses and InxGa1−xN
compositions (Left: x=0 Right: x=0.2)
The energy relaxation times of high-lying longitude optical phonons are demonstrated in Fig. 4.5.38,
for different combinations of well and barrier thicknesses. For Indium mole fraction x=0 (left figure),
the contrast of relaxation times are larger than that for the case for x=0.2 (right figure); this is
reasonable for as the higher Indium content in the barrier layer would make the structure closer to
bulk InN. Both figures show the same trend with different barrier/well thicknesses: the relaxation time
increases with thicker well layers and thinner barrier layers, in spite of some irregular local variations.
The regular variation mainly results from the change of numbers of InN-like modes and InGaN-like
modes. Here a hot reservoir of high-lying modes at 1000K is assumed, with heat transferred to the
cold reservoir of low-lying lattice modes at 300K. The non-equilibrium between the two systems is
physical due to their significant frequency separation.
Figure 4.5.38: Energy relaxation times of the high-lying LO phonon system (1000K), in the reservoir
of 1000K high-lying phonons and 300K low-lying phonons, with superlattice structures of different
well/barrier thicknesses and InxGa1−xN compositions (Left: x=0 Right: x=0.2)
To explain the variation, we need to first examine the phonon dispersions of the MQW-SL. Taking
x=0 as an example, the left figure in Fig. 4.5.35 shows the dispersions of InN/GaN SL structure, with
- 206 -
6 layers of nitrogen atoms inside each layer (the same for both the InN layer and the GaN layer).
From the number of atomic layers involved we could expect the same number of high-lying optical
modes which are InN-like and GaN-like respectively. The GaN-like optical modes (blue color: solid
line for LO, dash line for TO) have much higher frequencies than the InN-like optical modes (green
color), for the bond force constants of GaN is larger than for InN and the atomic mass of Ga is smaller
than In. There are also four branches (orange colour: two of LO and two of TO) corresponding to the
interfacial modes (IF), with frequencies between GaN-like modes and InN-like modes. From the right
figure of Fig. (4), the vibrations can be seen to be almost completely confined in the respective layers,
for InN-like and GaN-like optical modes. For IF most of the vibrational energy is within the 1st and
2nd nitrogen atoms from the interface, and hence has little overlap with InN-like or GaN-like modes.
Considering all three-phonon processes, only the Ridley channel and the Shrivistava-Barman channel
(decaying into a high- lying optical phonon and a low-lying optical phonon) are allowed due to the
energy conservation law. As in wurtzite-structured SL, the difference between the acoustic branches
and the low-lying optical branches becomes almost indistinguishable with both having partial opticallike character. These low-lying branches can be separated into two categories: GaN-confined modes
(red lines) and GaN-InN mixed modes (grey lines). In fact the InN-like low-lying modes sit within the
allowed frequency band of GaN, hence vibrations in GaN layers are excited too. Therefore no InNconfined modes exist in the low-lying branches. Among all allowed three-phonon processes, GaN-like
optical modes can decay into both types of low- lying modes (See the solid arrow and the dashed
arrow in Fig. 4.5.35), while InN-like optical modes can only decay into the mixed modes as they only
overlap with the mixed modes. Besides, the energy gap between GaN-like LO modes and GaN-like
TO modes is relatively large; hence the resulting low-lying phonons have relatively high energies,
compared to the decayed low-lying phonons from the InN-like modes. Since the low-lying branches
with high energies are generally flatter, leading to a larger joint density of states of transition, the
decay rates of GaN-like LO modes could be enhanced further. Due to the two reasons explained
above the GaN-like LO modes decay faster than the InN-like LO modes. Therefore with a thicker
well layer (or a thinner barrier layer), the energy relaxation time of the high-lying LO phonon system
becomes longer, for it introduces more InN-like modes (or fewer InGaN-like modes).
According to Fig. 4.5.38, the relaxation time of high-lying phonons could go up to 300ps,
significantly longer than the bulk value of around 1ps. And according to Fig. 4.5.37 the relaxation
time of hot electrons (if phonons are completely thermalized) can go up to more than 1ps, this is also
much longer than the bulk value. Combining these two, an optimized MQW-SL structure should
involve a thin barrier layer and a thick well layer. The contrast in phonon energies between the well
and the barrier also needs to be large, indicating a small Indium content in the barrier. In addition,
such a structures could potentially prevent hot phonons diffusing out, making it even more attractive
than bulk materials.
4.5.3.3.8 Time resolved photoluminescence measurements of bulk phononic band gap materials
The potential efficiency boost, which can be achieved by hot carrier solar cells, is directly related to
the possibility of extracting high energy carriers from the absorber layer before thermalisation,
increasing the voltage and hence the conversion efficiency. The poor conversion efficiency of photons
with energies above the band gap of the absorber is the main loss mechanism in conventional single
junction solar cells. The investigation of thermalisation time constants of hot carriers is a crucial step
towards the engineering of hot carrier cells. The efficiency of an InN based hot carrier solar cell has
been calculated using the theoretical model in Sections 4.5.3.3.7 and 4.5.3.1. It was found that the
limiting efficiency is strongly related to hot carriers relaxation velocity in the absorber [4.5.61].
A comparison of femtosecond time resolved photoluminescence (tr-PL) spectroscopy between InP
and GaAs was reported in the annual report two years ago and now published in [4.5.48]. This
showed a distinctly longer carrier cooling time constant for the wide phononic gap InP as compared to
almost zero phononic gap of GaAs. It also showed further evidence for excitation into higher side
valleys for both GaAs and InP for appropriate excitation wavelengths. Hot carrier cooling in InN has
- 207 -
been investigated using tr-PL. The wide gap between optical and acoustic branches in the InN
phononic dispersion relation (wider than that for InP) prevents the Klemens decay of optical phonon
into acoustic phonons. This can lead to slower carrier cooling due to the “Hot Phonon Effect”
[4.5.62]. The decay of hot carriers for different excitation wavelengths InN has been investigated.
Figure 4.5.39: Schematic representation of time resolved photoluminescence arrangement.
Tr-PL experiments have been performed on an InN sample grown on sapphire substrate, obtained
from collaborators in Taiwan, using the measurement configuration shown in Fig. 4.5.39. In this
technique a laser pulse acts as a switching gate relating the photoluminescence signal to the time
domain. The PL signal is collected from the sample, after a femtosecond laser excitation, and focused
in a non linear crystal. The gate signal is generated from the same laser and is focused on the same
crystal after passing through an optical delay stage. The signal is detected using a monochromator and
a PMT. Our system configuration provides 150 fs pulses with tunable wavelength over a range of 256
nm (4.84 eV) to 2.6 µm (0.48 eV).
Figure 4.5.40: 3D representation of InN time resolved PL data.
Figure 4.5.40 is a three dimensional representation of PL as a function of time for all the probed
wavelengths. It can be observed that the PL sharply rises when the carriers are photo-excited by the
laser pulse. The fast decay of the PL shows the thermalisation of carriers towards respective band
- 208 -
edges. The decay is faster for highly energetic carriers compared to carriers closer to the bandgap.
Thus the carrier population quickly degenerates towards the band edges during the thermalisation
process. In InN the thermalisation is most probably due to interaction of highly energetic electrons
and holes with LO phonons.
To investigate the rate of the carrier cooling process, the effective temperature of the carrier
population has been calculated fitting the high energy tail of the PL spectrum for every single time
during the cooling transient. The PL has been fitted assuming that carriers form a Boltzmann-like
distribution in a femtosecond time scale using the following equation.
E  
L( )   ( ) 2 exp G

 k BTC 
L(ε) represents the measured PL intensity at energy ε, α(ε) is the measured sample absorption
coefficient, EG is the InN energy gap, 0.7 eV in this case, and kB is the Boltzmann constant. TC is the
fitted parameter and represents the hot carrier temperature. Figure 4.5.41 shows the tr-PL data for two
sets of samples (a) In N grown on Sapphire substrate and (b) InN grown by MBE at Saitama
University for zinc blende structure on sapphire and as wurtzite structure on MgO substrate. For the
data in (a) the carrier temperature transient is seen to follow an exponential behaviour quite well (blue
line – fit). The data for the higher quality samples in (b) show significantly slower carrier cooling
rates than for (a).
Figure 4.5.41: Carrier relaxation curves for InN using tr-PL: (a) Lower quality InN with
soPhotoexcited carrier density is 1.5 x 1019 cm-3. The blue dashed curve is a single exponential fit. (b)
Zinc blende and wurtzite InN samples grown by MBE on MgO substrates by Shuei Yagi, Saitama
University (Cubic=crosses, Wurtzite structure = circles). This high quality material shows much
slower carrier cooling rates.
The fitting of the calculated temperature data has been performed using a single exponential.
 t
T (t )  C  exp
  TH

  K

Here τTH represents the carriers thermalisation time constant, whereas C and K are two constant
parameters. The fitted value for τTH is 7 ps. This long cooling constant can be attributed to the hot
phonon effect due to the long lifetime of the A1(LO) phonon [4.5.48,4.5.63]. The variation in carrier
cooling rates measured is attributed to the difficulty of growing good quality InN on sapphire
substrate, such that the carrier cooling velocity is strictly related to the quality of the material. Other
slow carrier cooling constants have been reported for InN in the literature [4.5.64].
- 209 -
4.5.3.3.9 Nanostructures for absorbers
Nanostructures offer the possibility of modification of the phonon dispersion of a composite material.
III-V compounds or indeed most of the cubic and hexagonal compounds can be considered as very
fine nanostructures consisting of ‘quantum dots’ of only one atom (say In) in a matrix (say N) with
only one atom separating each ‘QD’ and arranged in two interpenetrating fcc lattices. Modelling of
the 1D phonon dispersion in this way gives a close agreement with the phonon dispersion for zincblende InN extracted from real measured data for wurtzite material.
Similar ‘phonon band gaps’ should appear in good quality nanostructure superlattices, through
coherent Bragg reflection of modes such that gaps in the superlattice dispersion open up [4.5.65].
There is a close analogy with photonic structures in which modulation of the refractive index in a
periodic system opens up gaps of disallowed photon energies. Here modulation of the ease with
which phonons are transmitted (the acoustic impedance) opens up gaps of disallowed phonon
energies.
Force constant modelling of III-V QD materials by SK growth
3D force constant modelling, using the reasonable assumption of simple harmonic motion of atoms in
a matrix around their rest or lowest energy position, reveals such phononic gaps [4.5.66]. The model
calculates longitudinal and transverse modes and can be used to calculate dispersions in a variety of
symmetry directions and for different combinations of QD sub-lattice structure and super-lattice
structure.
III-V Stranski-Krastinov grown QD arrays of InAs in InGaAs and InGaAlAs matrices are fabricated
at the University of Tokyo using MBE. We are investigating these for evidence of phonon dispersion
modulation and potential slowed carrier cooling. In order to understand the expected phonon
dispersions these are being modelled using the 3D force constant technique.
Lattice matched and strain compensated material pairs that may produce large phonon bandgaps are
of interest. Previous iterations in the design of these structures indicated the importance of separating
“light” and “heavy” atoms to different parts of the nanostructure. Initially, the lightest atom in the
system, As, was present in both the QD and the matrix. This meant that the reduced mass of both
regions (proportional to the sum of the inverses of each atomic mass) was very similar as the light
element dominates in this case.
On this iteration structures with an In0.5Ga0.5-xAlxAs matrix (with x=0.4) and InAs QDs were grown.
Significant Al content was introduced into these structures with the expectation that this light element,
segregated to the matrix material, might produce appreciable phonon bandgaps. Some images derived
from characterisation of the structure are presented in Fig. 4.5.42. The superlattice of QDs has a
simple hexagonal structure. Extraordinary periodic out-of-plane stacking is achievable and largely
defect free structures can be grown on the order of microns.
- 210 -
Figure 4.5.42: Some images of the structure
as grown and modelled. (Above) the nearperfect stacking of the structure is shown.
(Above right) in-plane InAs QD arrangement.
(Right) simple hexagonal superlattice
structure showing stacking. The QDs are
similar to flattened disks with in plane
dimensions of 40-50 nm and heights of 7 nm.
ES
TA
Ef
e- selective
energy contact
Hot carrier distribution
µA
= qV
TH
Force constant modelling of this structure predicts an appreciable phonon bandgap, as shown in Fig.
4.5.43 and Fig. 4.5.44. This bandgap is due almost entirely to the presence of the Al in the system. A
small bandgap is present due to the mass difference between In, Ga and As, but it is less than a quarter
of the size shown in Fig. 4.5.43 and Fig. 4.5.44. Due to computational constraints the size of the QDs
is quite small, only about a nanometre in diameter. While actual sizes for these structures are too
computationally intensive to model without extreme effort, recent modelling with gradually
increasing size suggests that the dispersion relations scale linearly with size. That is, once the discrete
distances (bond lengths) are small relative to the superlattice unit cell dimension, the dispersion
should look exactly the same when scaled such that the relative dimensions are preserved. This will
be investigated in greater detail in further work.
- 211 -
Figure 4.5.43: Dispersion curves for InAs
QDs in a simple hexagonal SL. The matrix is
In0.5Ga0.5-xAlxAs with (x = 0.4). There is a
large phonon bandgap present in the system
due to the presence of Al. The gap seems
very tolerant to Al location. (Top): <100>,
(Middle):
<110>,
(Bottom):
<111>
directions.
Figure 4.5.44: Densities of states (DOS) for
the dispersion curve immediately to the left.
Note small changes in the DOS could be due
to pseudo-random locations of the Ga and Al
particles or to differences in crystalline
direction. (Top): <100>, (Middle): <110>,
(Bottom): <111> directions respectively.
4.5.3.3.10 Fabrication and characterisation of highly ordered nanoparticle arrays for hot
carrier absorber
As demonstrated by the work on modelling phonon modulation, periodic core-shell QD arrays offer a
way to significantly change the phonon modulation in a superlattice because the core and shell can be
of materials of very different force constant, directly leading to a strong phonon confinement.
Deposition methods are being investigated to fabricate such highly ordered QD arrays.
The Langmuir-Blodgett (LB) deposition technique is conventionally used for the fabrication and
characterisation of single and multilayer films with precise control of thickness, molecular orientation
and packing density. In recent years, it has been applied to metallic nanoparticle depositions for
formation of highly ordered array.
We have employed an LB system to fabricate gold (Au) and silicon (Si) nanoparticle layers. A LB
system has been set up at UNSW. We have designed the trough top for nanoparticle deposition to
optimise the dispersability. The system is shown in Fig. 4.5.45, which consists of a trough, a pair of
barriers, a surface pressure sensor, a dipping unit for film transferring and an interface controller.
- 212 -
Surfacepressure
sensor
barrier
Dippingunit
Interfaceunit
LBtrough
Figure 4.5.45: A customised LB system for nanoparticle deposition, installed in Room 1013
Chemical Engineering Building UNSW.
LB deposition and formation of monolayer with Si nanoparticles:
Figure 4.5.46(a) is a schematic diagram showing the LB system, consisting of an LB trough, a solid
substrate, a dipping unit, a Wilhelmy plate and two barriers. To make a Si NP array, a small amount
of 0.5 mg/mL chloroform solution with dodecane-capped Si NPs dissolved with diameters of 2.7 nm
or 6.9 nm, was spread onto the water surface drop wise over a time interval of 30 seconds. After the
complete evaporation of the solvent (chloroform), the Langmuir monolayer of Si NPs on the water
surface was dynamically manipulated by compression with the barriers. During the compression, the
surface pressure (SP) was measured using the Wilhelmy plate, and the isotherm curve (SP versus film
area) was recorded. Fig. 4.5.46(b) shows the isotherm curve for a 30 mL chloroform solution of 2.7nm-diameter Si NPs with a concentration of 0.5 mg/ml. Upon application of an appropriate level of
compression, suitable packing within the monolayer is achieved on the water surface. Subsequently,
the monolayer was transferred onto a hydrophilic substrate (Si or Quartz wafer washed in piranha
solution) through vertical-dipping controlled by the dipping unit. From the AFM image (Fig.
4.5.46(c)), a film is observed. The cross-sectional view clearly shows the thickness of the film is ~2.7
nm, equal to the size of Si NPs, which indicates a monolayer.
Figure 4.5.46: (a) a schematic of the LB system; (b) measured isotherm curve for Si NPs of 2.7 nm
and the chosen depositing surface pressure (SP); (c) AFM image and cross-sectional profile of the
fabricated 2.7 monolayer on a Si wafer.
- 213 -
Amorphous and short-range ordered close-packed arrays:
Unlike other semi-conductor NPs such as CdSe or metallic NPs, due to their irregular shapes and
lighter weight, Si NPs are not mobile enough for self-assembly to happen in a short timescale; this is
due to the weaker attractive capillary forces between the particles. Therefore, the reported results
always show a random array of Si NPs or aggregated clusters. In our work, we found that by
increasing the stabilizing time for monolayer formation on the water surface or slowing down the
moving rate of the barriers before performing the vertical dipping, we can achieve decently uniform
ordered close-packed arrays as shown in Fig. 4.5.47 (d).
Figure 4.5.47: TEM images of (a) amorphous (b) short-range-ordered (c) more uniform short-rangeordered close packed arrays deposited at a SP of 50 mN/m with different LB processes (d) an
enlarged area in (c); Process 1 is to prolong the stabilizing time of the monolayer on the water surface,
2 is to slow down the moving rate of the barriers.
In summary, we have established a process for deposition of nanoparticle monolayers on a solid
substrate, with reasonably close packed ordered arrays of nanoparticles. Future work will concentrate
on improving periodicity of the monolayer for both HC absorbers and energy selective contacts.
4.5.3.4 Implementation of nanostructures for hot carrier cells
In order to use a nanoparticle array as a hot carrier absorber, and hence utilise slowed carrier cooling
due to modified phonon dispersion, then a way to fabricate a complete structure with ESCs must be
established. A structure in which nanoparticles are arranged in a uniform 3D array should give the
required phononic band gap. The degree to which this needs to be ordered is still under investigation.
Certainly a large difference between masses is beneficial between the nanoparticles and the matrix
- 214 -
and this can be best achieved in an array of ‘core shell QDs’ in which the core and shell have very
different acoustic impedance in order to promote coherent reflection and hence confinement of
phonons [4.5.65]. The matrix in which such an array is embedded needs to allow transport of carriers
to contacts and also electron-electron and preferably hole-hole scattering to renormalise carrier
energies. Such a structure should also have a narrow electronic band gap so as to absorb a wide range
of photon energies. This combination of properties is challenging but not mutually exclusive because
phononic properties are largely independent of electronic properties and such a structure would not
need to have electronic confinement of carrier energies just vibronic modulation of phonon energies.
Energy selective contacts are also required for such a structure. These would most ideally be
arranged at the top and bottom of the absorber. However, in a realistic device a selective hole contact
might not be required as the hole population only contains a small fraction of the hot carrier energy
and hence thermalisation of holes is less important, whereas a structure with only one contact
(presumably the top contact) as an ESC would be much easier to fabricate. These ESCs would be QD
or QW double barrier structures. (Addition of extra layers to give double QD/QW triple barrier
structures (or even triple QD/QW quadruple barriers) would give greater resonance and enebale the
contact to be rectifying – important if it is to collect carriers of only one. Such a QD array / double
barrier QD ESC structure is shown schematically in Fig. 4.5.48.
External contact
Selective energy contact: double barrier QD resonant
tunnelling contact. Or triple barrier to also give rectification.
Slowed carrier cooling absorber layer:
 Bulk phononic band gap material – e.g. III-nitrides
 and/or Periodic QW superlattice array;
 and/or periodic QD array tuned to block optical phonon
decay by Klemens mechanism;
 Must also allow renormalization of carrier energy by e-e
& h-h scattering.
Proof of concept with MOVPE/MBE III-V QW or QW array.
Selective energy contact: selecting opposite carrier type.
Or for rear hole contact a simple band contact may suffice.
External contact
Figure 4.5.48: Conceptual integrated hot carrier device, incorporating double (or triple) barrier QD
resonant tunnelling contacts and absorber probably based on either a bulk or nanostructure phononic
band gap material and/or MQW nanostructure, for slowed carrier cooling.
4.5.3.5 Summary of hot carrier solar cell research
There has been significant development in most areas of hot carrier solar cell work. The modelling of
efficiencies now includes combinations of absorber and energy selective contacts with reasonable
values chosen using real material parameters based on III-nitrides. This has been carried out both for
bulk absorber InN and for multiple quantum well materials again based on nitrides. Such devices are
capable of being fabricated and then tested against these predicted results and work with collaborators
is moving towards growing such devices. In addition the theory of the slowed carrier cooling
behaviour in MQWs is being investigated with plans to test the competing mechanisms again using
material from a range of collaborators. Measurement of InN with time-resolved PL has indicated
some evidence for slowed carrier cooling, further corroborating the importance of a large phonon
band gap to block optical phonon decay, but also highlights the importance of material quality.
Progress on the Langmuir-Blodgett approach to ordered nanoparticle arrays has seen development of
highly ordered arrays of nanoparticles. The potential application of nanostructures to fully integrated
devices has been investigated conceptually, with various designs considered. Similarly the possibility
of absorber materials which are analogous to InN is also being investigated. These many aspects of
Hot carrier cells will see further development and consolidation in 2013 with working devices
planned for early 2014.
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4.5.4 Up-Conversion
Researchers: Sanghun Woo, Craig Johnson, Supriya Pillai, Shujuan Huang, Richard Corkish, Gavin
Conibeer
Collaboration with: Peter Reece (Physics, UNSW), John Stride (Chemistry, UNSW)
Up-conversion in novel silicon-based materials
Up-conversion (UC) in erbium-doped phosphor compounds (particularly NaYF4:Er) has been shown
to be a promising means of enhancing the sub-band-gap spectral response of conventional Si solar
cells without modification of the electrical properties of the cell [4.5.67]. In this scheme, a layer
containing the phosphor is applied to the rear of a high-efficiency bifacial cell. After absorbing two
long-wavelength (~1500nm) photons - which are transmitted by the cell - the excited Er ions can relax
by emitting a photon with an energy greater than the Si band gap, thereby increasing the current that
can be extracted from the cell.
Synthesis of Er doped NaYF4 phospors for UC
Fabrication of in-house phosphors allows control over the growth process and the ability to control Er
doping levels and the possibility to include co-dopants.
NaY(1-x)F4:Erx nanocrystals were synthesized by thermal decomposition of metal trifluoroacetate
(TFA) precursors [4.5.68, 4.5.69]. After the reaction at 80°C for two hours, oleic acid (OA) and 1octadecene (1-ODE) were added to the reaction and slowly heated for 1 hour at 280°C. At this stage,
NaYF4:Er NCs formed in the alpha phase. To convert NCs from alpha to the beta phase, the
temperature was slowly increased to 320°C for 3 hours. The resulting solution was treated to remove
unreacted precursors, excess OA and 1-ODE. NaYF4:Er NC films were fabricated by spin casting
followed by nucleophilic ligand exchange. NC films were coated on the glass by multiple spin
castings.
Fig. 4.5.49(A) shows HRTEM images of hexagonal beta phase NaYF4:Er NCs. The dominant and
most distinguishable lattice (100) of the hexagonal beta phase NaYF4 NCs had a fringe distance of
0.54 nm. The histogram in Fig. 4.5.49 (B)indicates that particles had sizes ranging from 25 to 35 nm
with a narrow size distribution peaking around 30 nm.
Figure 4.5.49: Electron microscope data for NaYF4: 10% Er. (A) low and high resolution TEM. Both
scale bars correspond to 50nm. (B) Histogram of diameters of Er dope NCs.
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The absorption spectrum of UC NCs is shown in Fig. 4.5.50 and corresponds to all energy levels up to
3.1 eV of NaYF4:Er NCs. Fig. 4.5.50(B) shows five red-shifted peaks compared to the absorbance
which were resulted from radiative relaxations from 2H11/2, 4S3/2, 4F9/2, 4I9/2 and 4I11/2. As shown by UC
luminescence images in insets of Fig. 4.5.50(B) increased Er concentrations in UC NCs accelerated
multi-phonon relaxation processes between 4S3/2 and 4F9/2, which emphasizes the importance of
optimization between slowing non-radiative decay rates and increasing absorption sites in NCs.
Figure 4.5.50: Absorption spectra for NaYF^4:Er NCs (A). (B) detil 300-1000nm, with insets
showing the luminescence.
The SEM images of spin cast films before and after the soaking process in a MPA solution are shown
in Fig. 4.5.51(A) and (B) respectively. Without the post soaking process, NCs were coated separately
from neighbouring crystals, and packed closely. By treating with a 3-MPA solution, NCs were
agglomerated and linked to each other resulting in a bulk-like feature. The cross-linked NC film made
with 3-MPA brings about easy handling, light trapping in NC layers, and a reduced film thickness.
Figure 4.5.51: (A) SEM images of morphology of upconverting layer (A) before and (B) after
soaking in a MPA solution. (C) Up-convesrion spectroscopy showing the intensities of emitted
photons in terms of laser power densities and Er doping concentration. (D) Film thickness dependent
upconversion under three different excitation powers.
Figure 4.5.51 (C) shows the total intensity of upconverted photons from energy levels above 4I11/2,
which are all able to contribute to the photoconversion process in c-Si solar cells, depending on the
excitation power between 30 μW and 215 mW and the Er concentration between 2 and 100%
substituting Y sites of the NaYF4 host structure. To minimize the effect of the total number of Er ions
- 217 -
in UC samples, all samples were prepared thicker than the required thickness for saturated
upconversion, as described in the thickness dependent analysis below. Considering the result of Er
concentration effects, 15% Er doped UC layers always gave the best upconversion regardless of the
excitation power. This means that increasing Er doping concentration did not affect the number of
photons that could meet with each trivalent Er ion. The ratio between the Er doping concentration and
the photon flux was not a predominant factor because the number of photons that could be absorbed
by each Er ion does not change when the films were thick enough. Thus, increased or decreased Er
content under fixed laser density did not have the same effect as changing the laser power at the fixed
Er concentrations. By increasing Er densities in UC NCs, the number of neighbouring ions is
increased, and the distance between Er ions is reduced. Both can boost non-radiative cross relaxation
processes as well as energy transfer upconversion mechanisms. From the results, it can be concluded
that 15% Er doped NaYF4 can maximize the rate of the energy transfer upconversion minus cross
relaxation decay. This optimisation is practical because it is independent of the solar concentrations.
Using the brightest UC NCs with a 15% doping concentration of Er, the film thickness was optimized
under various excitation fluxes, as illustrated in Fig. 4.5.51 (D). From the excitation coefficient
information, the UC films had to be thicker than 20 μm, so that all photons within energies of 4I13/2
could be absorbed by the UC phosphor. However, the required film thicknesses to perform the
saturated emission were only about 2.5 μm under a 1 mW laser, 3 μm under a 5 mW laser and 5 μm
under a 10 mW laser. This is because upper UC NC layers re-absorbed upconverted photons emitted
from the lower layers. Also, the stable bulk-like layers may have a higher absorption coefficient than
simply packed NCs, resulting in relatively thinner thicknesses required for the saturated UC
efficiency.
Light trapping in slowed light modes to enhance UC efficiency
While phosphors have demonstrated high-efficiency UC behaviour, their use presents particular
challenges with regard to fabrication and cost. Last year we presented data on use of Er-doped porous
Si (PSi:Er) as an alternative UC material. PSi is unique in that its porosity - and hence the material
refractive index - can be varied as a function of depth, allowing for the elaboration of high-quality
monolithic Si optical structures such as distributed Bragg reflectors (DBRs).
A cross-sectional electron micrograph of such a structure is shown in Fig. 4.5.52. Its porous
substructure also allows for deep infusion of dopant atoms via techniques such as electroplating.
Figure 4.5.52: Cross-sectional electron micrograph of a 30-bilayer porous silicon distributed Bragg
reflector.
Figure 4.5.53 shows the calculated field intensity profile for these DBRs in the spectral region
between 1500 and 1600nm. The results have been normalized to the incident field intensity and each
- 218 -
plot is independently colour-scaled to show maximum detail. They clearly demonstrate that
considerable field enhancement is possible throughout the depth of such a multilayer structure, with
peak relative intensities of more than 11 for the 30-bilayer structure and more than 18 for the 40bilayer structure. However, it is also clear that the regions of enhancement are increasingly narrowlyconcentrated as the number of bilayers increases. This is due to the additional interference fringes in
the reflectivity characteristic of the DBR that arise with additional layers, as is apparent from Fig.
4.5.53. Calculations for more than 40 layers show that enhancements of up to 80 are possible for a
large number of bilayers though the peak narrows to less than 1nm wide for 100 bilayers.
Figure 4.5.53: Intensity-valued mode profile in DBRs with varying numbers of high/low-porosity
bilayers for a leading band edge fixed at ~1550nm. All intensities are normalized to the incident field
intensity and each plot is independently color-scaled to show maximum detail. The air and substrate
interfaces with the DBR are indicated by the dashed black (top) and red (bottom) lines, respectively.
Sensitisation of Er phosphors to a wider wavelength range
Er has an absorption window at the I13/2 level from 1480nm to 1580nm. This narrow range means that
it is inefficient at absorbing below Si band gap photons. Sensitizing to the 1100-1500nm range would
significantly enhance the absorption and hence have a dramatic effect on the up-conversion
efficiency, because of its non-linear dependence.
PbS nanoparticles have a wide absorption range below 1100nm. They can also be tuned in size to
emit at 1500nm, with absorbed photons being downshifted to 1500nm at some quantum efficiency
less than 100%.
Such PbS NCs have been synthesized based on the procedure developed by Cademartiri et al. [4.5.70]
by decomposition of PbCl2 with sulphur solution. PbS QDs and Er-doped NaYF4 NC films were
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fabricated by the simple spin casting method followed by nucleophilic cross-linking of NCs.
The sizes, the size distribution and the fringe distance of resulting UC NCs were determined using
HRTEM and glancing incident x-ray diffraction (GIXRD) was carried out to investigate the
crystalline structures of PbS QDs.
Figure 4.5.54: Absorbance and emission data from UV-Vis spectroscopy. (a) for PBs quantm dots
showing the desired slight mismatch between absorbance and emission such that re-absorption is
unlikely. (b) absorbnce for Er doped NCs and emission from PbS QDs, showing the good match
between PbS emission to pump Er absorption at the.
Figure 4.5.54 shows absorbance and emission data for PbS QDs and Er doped NaYF4 NCs. The range
of absorbance for PbS QDs is close to ideal for absorbance of light below the Si bandgap but stopping
short of the emission. This emission, tuned to 1500nm, is ideal for pumping the I13/2 level at 1520nm
in Er, without being re-absorbed by the PbS. Thus the sensitization of Er to a significantly wider
wavelength range should be very valuable in increasing the flux of photons which can be absorbed
and up-converted by the Er doped phosphor.
Conclusion
Up-conversion in Er doped phosphors shows very promising results. Fabrication of NaYF4 NCs
allows greater control over doping levels and geometries for efficient up-conversion. Light trapping in
DBR structures based on P-Si should enable greater absorbance of 1500nm light through the greater
density of photonic states. Progress towards sensitisation of Er to a wider wavelength range using
down shifting PbS quantum dots tuned to emit at 1500nm has been made. Combination of all these
elements in an up-conversion structure is expected to significantly boost UC efficiencies for Si cells.
4.5.5 Concluding Remarks for the Third Generation Section
Work has proceeded significantly in all the areas of Third Generation research, with improved
fabrication and characterisation of materials and complexity of modelling which together give an
overall better understanding and optimisation of devices.
Group IV based nanostructure materials have seen significant improvement in understanding of the
parameters required for growth and the influence of these on material and device properties. The
modelling of equivalent circuits for the devices is giving greater understanding and now allowing
prediction of growth conditions to give improvement. The realisation that the effective bandgap and
hence the VOC in these materials is dominated by defects has led to a greater effort on characterising
and reducing defects. Light trapping through back reflectors and the use of plasmonics has led to
more than double the current in nanostructure devices. Further optimisation in this area is very likely.
The use of conductive substrates to reduce the lateral resistivity has made promising progress, with
SiC and Mo substrates showing some improvements. Optimisation of these will continue to give
better lateral transport and should lead to better photovoltaic properties. Higher VOCs over 500mV
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have been measured at lower temperatures. The justification for this is that due to the high
resistivities, devices heat up considerably and cooling is valid to bring them back to room
temperature. A temperature of 50degress below ambient is calculated to achieve this condition and
gives VOC of 503mV, although significantly high VOCs occur at even lower temperatures as would be
expected. Work on other related materials has seen lower resistivities for Si3N4 interlayers in Si QDs
in SiO2; high quality Ge QW growth; initial data for growth of Ge:C materials with the aim of
engineering band gaps through alloying and the potential for transfer of growth processes to PECVD
with it lower levels of damage and thermal budget.
Hot carrier cells have seen a great deal of advance in the last year. Modeling of overall devices now
includes a much larger range of real material properties in a more integrated structure. In addition
modelling of combined multiple quantum well an phononic bandgap structures is leading to new
designs for combined structures which offer promise for slowed carrier cooling. Modelling of energy
selective contacts has progressed again with real material properties and integration into full
conceptual devices. Such devices are entirely feasible and plans to grow these with collaborators are
underway. New designs for rectifying ESCs or incorporation of ESCs within the absorbed material
have been suggested and these also can be incorporated into real structures. Increased understanding
of the mechanisms involved in both phononic band gap materials and in multiple quantum well
structures is leading to new experiments to determine the exact relationship between reduced carrier
cooling mechanisms. The results of these will lead to the optimum design requirements for combined
structures which should allow significantly slower carrier cooling rates. Deposition of nanoparticle
arrays has made progress with very ordered arrays now possible using Langmuir-Bodgett deposition.
These arrays have been modelled to give significant phonon bandgaps and characterisation of these is
now underway. New materials suitable for some of these phononic and nanostructure properties are
being identified and growth and characterisation of these materials is planned in-house and with a
range of collaborators.
Up-conversion has seen a further significant development with synthesis of NaYF4 doped with Er
now allowing greater flexibility in doping and geometry of devices. Light trapping work is continuing
to increase the number of photonic modes available for absorption in Er. Sensitisation of Er to a wider
wavelength range is being enabled using PbS QDs which are now being synthesised and characterised
and which are showing close to ideal properties for such sensitisation. Combination of all these
approaches is expected to give significantly improved up-conversion.
The development of all the 3rd generation projects now allows much greater understanding of the
materials and devices. Work in 2013 will see consolidation of this into improved devices. Several
new areas of funding will contribute to this and also allow development of new project areas.
References:
4.5.1
4.5.2
4.5.3
4.5.4
4.5.5
4.5.6
4.5.7
M. A. Green, “Third Generation Photovoltaics: Ultra-High Efficiency at Low Cost”
(Springer-Verlag, 2003).
J. Nelson, “The Physics of Solar Cells”, Imperial College press, 2003.
G. Conibeer, “Third Generation Photovoltaics”, Materials Today, 10 (11) (2007) 42-50.
E-C. Cho, S. Park, X. Hao, D. Song, G. Conibeer, S-C. Park, M.A. Green, “Silicon quantum
dot/crystalline silicon solar cells”, Nanotechnology, 19 (2008) 245201
Hao X J, Podhorodecki A, Shen Y S, Zatryb G, Misiewicz J, Green M A, “Effects of Si-rich
oxide layer stoichiometry on the structural and optical properties of Si QDs/SiO2 multilayer
film”, Nanotechnology 20, 485703 (2009).
M.A. Green, G. Conibeer, D. König, I. Perez-Wurfl, A. Gentle, S. Huang, D. Song, X.J.
Hao, F. Gao, S. Park, C. Flynn, Y. So, B. Zhang, D. Di, P. Campbell, Y. Huang and T.
Puzzer, “Progress with All-Silicon Tandem Cells On Glass Using Silicon Quantum Dots in
Silicon-Based Dielectric Matrices”, 24th European Photovoltaic Solar Energy Conference,
Hamburg, Germany, 21-25 September 2009.
M. Zacharias, J. Heitmann, R. Scholz, U. Kahler, M. Schmidt, Appl. Phys. Lett. 80(2002)
- 221 -
4.5.8
4.5.9
4.5.10
4.5.11
4.5.12
4.5.13
4.5.14
4.5.15
4.5.16
4.5.17
4.5.18
4.5.19
4.5.20
4.5.21
4.5.22
4.5.23
4.5.24
4.5.25
4.5.26
4.5.27
4.5.28
4.5.29
4.5.30
661.
X.J. Hao, I. Perez-Wurfl, G. Conibeer, M.A. Green, 19th PVSEC, Korea, Nov 2009.
I. Perez-Wurfl, X. Hao, A. Gentle, D. Kim, G. Conibeer, M.A. Green, “Si nanocrystal p-i-n
diodes fabricated on quartz substrates for third generation solar cell applications,”Applied
Physics Letters, 95 (2009) 153506.
X.J. Hao, E-C. Cho, G. Scardera, Y.S. Shen, E. Bellet-Amalric, D. Bellet, G. Conibeer,
M.A. Green, “Phosphorus-doped silicon quantum dots for all-silicon quantum dot tandem
solar cells”, Solar Energy Materials and Solar Cells, 93 (2009) 1524.
G. Conibeer, M.A. Green, R. Corkish, Y.-H. Cho, E.-C. Cho, C.-W. Jiang, T.
Fangsuwannarak, E. Pink, Y. Huang, T. Puzzer, T. Trupke, B. Richards, A. Shalav, K.-L.
Lin, Thin Solid Films 511/512 (2006) 654.
E.-C. Cho, Y.H. Cho, T. Trupke, R. Corkish, G. Conibeer and M.A. Green, Proc. 19th
European Photovoltaic Solar Energy Conference, Paris, (June 2004), p. 235.
F. Meillaud, A. Shah, C. Droz, E. Vallat-Sauvain, C. Miazza, Sol. Energy Mater. Sol. Cells
90, 29522 (2006).
T. Arguirov, T. Mchedlidze, M. Kittler, R. Rölver, B. Berghoff, M. Först, B. Spangenberg,
Appl. Phys. Lett. 89 (2006) 053111.
T.-W. Kim, C.-H. Cho, B.-H. Kim, S.-J. Park, Appl. Phys. Lett. 88 (2006) 123102.
Y.-H. Cho, E.-C. Cho, Y. Huang, C.-W. Jiang, G. Conibeer, M.A. Green, Proceeding of
20th European Photovoltaic Solar Energy Conference, Barcelona, Spain, June 6-10, 2005,
p.47.
Y. Kurokawa, S. Miyajima, A. Yamada and M. Konagai, Jpn. J. Appl. Phys. 45, L1064
(2006).
C.-W. Jiang, M.A. Green, J. Appl. Phys. 99 (2006) 114902.
X.J. Hao, E-C. Cho, C. Flynn, Y.S. Shen, S.C. Park, G. Conibeer, M.A. Green, Solar Energy
Materials & Solar Cells, “Synthesis and characterization of boron-doped Si quantum dots
for all-Si quantum dot tandem solar cells”, 93 (2009) 273–279.
I. Perez-Wurfl, A. Gentle, X. Hao, M.A. Green, G. Conibeer and D-H. Kim, Proceedings of
24th EPVSEC, Hamburg, Germany, (2009).
J.M. Shah, Y.-L. Li, Th. Gessmann, E. F. Schubert., Journal of applied physics. 94 (2003)
2627.
K. Boer, “Survey of Semiconductor Physics”, van Nostrand Reinhold, 1990.
Di D, Perez-Wurfl I, Green MA, Conibeer G, Formation and photoluminescence of Si
quantum dots in SiO2/Si3N4 hybrid matrix for all-Si tandem solar cells, Solar Energy
Materials and Solar Cells, 2010; 94: 2238-43.
Zhang B, Shrestha S, Green MA, Conibeer G, Surface states induced high p-type
conductivity in nanostructured thin film composed of Ge nanocrystals in SiO2 matrix,
Applied Physics Letters, 2010; 97: 132109.
Huang S, Cho EC, Conibeer G, Green MA, Bellet D, Bellet-Amalric E, Cheng S,
Fabrication and Characterisation of Tin-based Nanocrystals, Journal of Applied Physics,
2007; 102: 114304.
Huang S, So YH, Conibeer G, Green MA, In-situ formation of tin nanocrystals embedded in
silicon nitride matrix, Journal of Applied Physics, 2005; 105: 124303.
QI F, Conibeer G, Shrestha S, Huang S, The study of potential application of Group IV
alloys in photovoltaic field, Australian Solar Energy Society conference, 2010.
Wan Z, Huang S, Green MA, Conibeer G, Rapid thermal annealing and crystallization
mechanisms study of silicon nanocrystal in silicon carbide matrix, Nanoscale Research
Letters, 7541289134788882, accepted January 2011.
Vnuk F, DeMonte A, et al. The composition dependence of the grey tin --> white tin
transition in dilute tin-germanium alloys. Materials Letters, 1983; 2: 67-70; Goodman CHL,
Direct-gap group IV semiconductors based on tin. IEE Proceedings I: Solid State and
Electron Devices, 1982; 129: 189-92.
Bauer M, Ritter C, et al. Synthesis of ternary SiGeSn semiconductors on Si(100) via
SnxGe1-x buffer layers., Applied Physics Letters, 2003; 83: 2163-5.
- 222 -
4.5.31
4.5.32
4.5.33
4.5.34
4.5.35
4.5.36
4.5.37
4.5.38
4.5.39
4.5.40
4.5.41
4.5.42
4.5.43
4.5.44
4.5.45
4.5.46
4.5.47
4.5.48
4.5.49
4.5.50
4.5.51
4.5.52
4.5.53
4.5.54
4.5.55
4.5.56
4.5.57
4.5.58
D'Costa VR, Fang YY, et al., Tunable optical gap at a fixed lattice constant in group-IV
semiconductor alloys. Physical Review Letters, 2009; 102: 107403.
Sekkal W, Zaoui A, Predictive study of thermodynamic properties of GeC, New Journal of
Physics, 2002; 4: 9.1–9.9.
Liu ZT, Zhu JZ, et al. Structure and properties of germanium carbide films prepared by RF
reactive sputtering in Ar/CH4. Japanese Journal of Applied Physics, 1997; 36: 3625.
Hu CQ, Xu L, et al. Effects of radio frequency power on the chemical bonding, optical and
mechanical properties for radio frequency reactive sputtered germanium carbide films.
Journal of Physics D: Applied Physics, 2006; 39: 5074-9.
G. Scardera, T. Puzzer, G. Conibeer, M. A. Green, Journal of Applied Physics, “Fourier
transform infra-red spectroscopy study of silicon nanocrystals embedded in silicon nitride”,
104 (2008) 104310.
Y.-H. Cho, D. Song, E.-C. Cho, G. Conibeer and M.A. Green, “Silicon quantum dot
superlattices in SiC matrix for all silicon tandem cells”, Technical Digest of 17th
International Photovoltaic Science and Engineering Conference (PVSEC-17), Fukuoka,
Japan, 3-7 December 2007.
Santosh K. Shrestha, Pasquale Aliberti, Gavin J. Conibeer, Energy selective contacts for hot
carrier solar cells, Solar Energy Materials and Solar Cells, 94 (Sept 2010) 1546-1550.
P. Würfel, Sol. Energy Mats. and Sol. Cells. 46 (1997) 43.
R. Ross, A.J. Nozik, J. Appl Phys (1982) 53, 3318.
Y. Feng, P. Aliberti, B.P. Veettil, R. Patterson, S. Shrestha, M.A. Green, G. Conibeer, Nonideal energy selective contacts and their effect on the performance of a hot carrier solar cell
with an indium nitride absorber, Applied Physics Letters, 100, 053502 (2012).
G. Conibeer, C.W. Jiang, D. König, S.K. Shrestha, T. Walsh, M.A. Green, “Selective energy
contacts for hot carrier solar cells”, Thin Solid Films, 216 (2008) 6968-6973.
A. Hsieh, PhD Thesis UNSW, 2013, “Investigation of Si/SiO2 and Si/SiOx Quantum Well
Structures for Applications as Energy Selective Contacts and All-Silicon Tandem Cells”.
Summers and Brennan, Appl. Phys. Lett. 48 (1986) 806.
Barnham et al J. Appl. Phys. 67 (1990) 3490.
P.G. Klemens, Phys. Rev. 148 (1966) 845.
G. Conibeer, R. Patterson, P. Aliberti, L. Huang, J-F. Guillemoles, D. König, S. Shrestha, R.
Clady, M. Tayebjee, T. Schmidt, M.A. Green “Hot Carrier solar cell Absorbers”, 24th
European PVSEC, Hamburg (Sept. 2009).
F. Chen, A.N. Cartwright, Appl Phys Lett, 83 (2003) 4984.
R. Clady, M.J.Y. Tayebjee, P. Aliberti, D. Konig, N.J. Ekins-Daukes, G.J. Conibeer, T.W.
Schmidt, M.A. Green, submitted PIP, 20 (Jan 2012) 82–92. DOI: 10.1002/pip.1121.
G. Conibeer, R. Patterson, P. Aliberti, H. Xia, S. Huang, K. König, B. Puthen-Veettil, S.
Shrestha, M.A. Green, “Hot carrier cell absorber materials,” Proc. 26th European
Photovoltaic Solar Energy Conference Hamburg, (2011).
Y. Rosenwaks, M. Hanna, D. Levi, D. Szmyd, R. Ahrenkiel, A. Nozik, “Hot-carrier cooling
in GaAs: quantum wells versus bulk,” Phys Rev B, 48, 14675-14678 (1993).
J-F. Guillemoles, G. Conibeer, M.A. Green, “On achievable efficiencies of Hot carrier solar
cell absorbers based on reported experimental thermalisation properties,” Proc. 21st
European Photovoltaic Solar Energy Conference, Dresden Germany, 234–237 (2006).
L. Hirst, M. Führer, A. LeBris, J-F. Guillemoles, M. Tayebjee, R. Clady, T. Schmidt, Y.
Wang, M. Sugiyama, N. Ekins-Daukes, Proc. 37th IEEE PV Specialists Conference, Seattle,
talk 927 (2011).
C. Colvard, T.A. Gant, M.V. Klein, “Folded acoustic and quantized optic phonons in
(GaAl)As superlattices,” Physical Review B, 31, 2080-2091 (1985).
T.K. Todorov, K.B. Reuter, D.B. Mitzi, Advanced Materials, 22 (21010) 1–4.
W.R. L. Lambrecht, “Structure and phonons of ZnGeN2”, Phys. Rev. B 72 (2005) 155202
T.R. Paudel, W.R.L. Lambrecht, Phys. Rev. B, 79 (2009) 245205.
S.C. Erwin, I. Zutic, Nature Materials, 3 (2004) 410-414.
Z.T. Liu, J.Z. Zhu, N.K. Xu, X.L. Zheng, Jap. J. Appl. Phys., 36 (1997) 3625.
- 223 -
4.5.59
4.5.60
4.5.61
4.5.62
4.5.63
4.5.64
4.5.65
4.5.66
4.5.67
4.5.68
4.5.69
4.5.70
D. Konig, Y. Takeda, B. Puthen-Veettil, G. Conibeer, Japanese Journal of Applied Physics
51 (2012) 10ND02, Lattice-Matched Hot Carrier Solar Cell with Energy Selectivity
Integrated into Hot Carrier Absorber.
D. Konig, Y. Takeda, B. Puthen-Veettil, Appl. Phys. Lett., 101, 153901 (2012),
Technology-compatible hot carrier solar cell with energy selective hot carrier absorber and
carrier-selective contacts.
P. Aliberti, Y. Feng, Y. Takeda, S.K. Shrestha, M.A. Green, G.J. Conibeer, “Investigation of
theoretical efficiency limit of hot carrier solar cells with bulk InN absorber”, Journal of
Applied Physics, 108 (2010) 094507-10.
K.T. Tsen, J.G. Kiang, D.K. Ferry, H. Lu, W.J. Schaff, H.W. Lin, S. Gwo, “Direct
measurements of the lifetimes of longitudinal optical phonon modes and their dynamics in
InN”, Applied Physics Letters, 90 (2007) 3.
Y.C. Wen, Y.C., C.Y. Chen, C.H. Shen, S. Gwo, C.K. Sun, “Ultrafast carrier thermalization
in InN”, Applied Physics Letters, 89 (2006) 3; M.D. Yang, et al., “Hot carrier
photoluminescence in InN epilayers”, Applied Physics A - Materials Science & Processing,
90 (2008) 123-127.
F. Chen, A.N. Cartwright, H. Lu, W.J. Schaff, “Ultrafast carrier dynamics in InN epilayers”,
Journal of Crystal Growth, 269 (2004) 10-14.
G. Conibeer, N.J. Ekins-Daukes, J-F. Guillemoles, D. König, E-C. Cho, C-W. Jiang, S.
Shrestha and M.A. Green, “Progress on Hot Carrier Cells”, Solar Energy Materials and
Solar Cells, 2009, 93, 713.
R. Patterson, M. Kirkengen, B. Puthen Veettil, D. Konig, M.A. Green, G. Conibeer, Sol.
Ener. Mats. and Sol. Cells, 94 (2010) 1931-1935.
A. Shalav, B.S. Richards, M.A. Green, “Luminescent layers for enhanced silicon solar cell
performance: Up-conversion”, Sol. Energy Mats. and Solar Cells, 91 (2007) 829-842.
J.-C. Boyer, L. A. Cuccia, and J. A. Capobianco, "Synthesis of Colloidal Upconverting
NaYF4: Er3+/Yb3+ and Tm3+/Yb3+ Monodisperse Nanocrystals," Nano Letters, vol. 7,
pp. 847-852, 2007/03/01 2007.
C. Rüssel, "A pyrolytic route to fluoride glasses. I. Preparation and thermal decomposition
of metal trifluoroacetates," Journal of Non-Crystalline Solids, vol. 152, pp. 161-166, 1993.
Cademartiri, L., et al., Multigram Scale, Solventless, and Diffusion-Controlled Route to
Highly Monodisperse PbS Nanocrystals. The Journal of Physical Chemistry B, 2005.
110(2): p. 671-673; Cademartiri, L., et al., Size-Dependent Extinction Coefficients of PbS
Quantum Dots. Journal of the American Chemical Society, 2006. 128(31): p. 10337-10346.
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4.6
Silicon Photonics and Device Characterisation
4.6.1
Photoluminescence based characterisation of silicon
University Staff
A/Prof. Thorsten Trupke
Project scientists and technicians
Allen Yee
Postgraduate Students
Yael Augarten
Mattias Juhl
Bernhard Mitchell
External Collaborators
Ron Sinton, Sinton Consulting
BT Imaging Pty Ltd
Daniel Macdonald, ANU
Johannes Greulich, Fraunhofer ISE, Germany
4.6.1
4.6.1.1
Photoluminescence based characterisation of silicon
Background
In photoluminescence (PL) imaging an intense light source is used to illuminate large area samples
such as silicon bricks, wafers or complete solar cells homogeneously with typically one-sun equivalent
illumination intensity. A CCD camera captures a high resolution picture of the luminescent light that
the sample emits. This technology, which was first demonstrated on silicon wafers and solar cells by
our group in 2005 [4.6.1.1] is now widely used in research laboratories worldwide and is increasingly
adopted also in industrial manufacturing, for example for quality inspection of as-cut wafers for
outgoing (wafer manufacturers) or incoming (cell manufacturers) quality control.
Main benefits of PL imaging, apart from its contactless and non-destructive nature include typically
very high spatial resolution and short measurement times. Megapixel luminescence images can be
captured within seconds or even fractions of a second, allowing very fast and detailed studies of
specific material and device properties. Over the last few years PL imaging has seen the development
of a range of specific applications that are applicable across the PV value chain. Examples include
series resistance imaging [4.6.1.2] and diffusion length imaging [4.6.1.3] on fully processed cells,
which were both first demonstrated at UNSW. The PL group at UNSW continues to further develop PL
imaging and find new applications and analysis methodologies. This year the research by our group
continued on the application of the photoluminescence intensity ratio method for measurements of the
bulk lifetime on silicon bricks (Section 4.6.1.2) with emphasis on quantification of limitations of that
method and experimental and analytical methods to mitigate the impact of these limitations.
UNSW spin-off company BT Imaging Pty Ltd commercialises the UNSW developed and patented PL
imaging technology since 2007 and has since then not only significantly extended the associated patent
portfolio but also made significant progress with the further improvement of the hardware and analysis
software, with specific emphasis on the development of fast systems that enable as-cut wafer
inspection on the fly at line speed. Line scanning PL imaging systems could be shown to allow not
only measurements on as-cut wafers at 3,600 wafers per hour throughput, but also enables fascinating
new possibilities, for example megapixel images of the series resistance of fully processed cells to be
measured without contact to the cell and on-the fly with the cell moving on a belt at 250mm/s (section
4.6.1.5).
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4.6.1.2
PL imaging on Si bricks
Photoluminescence imaging is an ideal tool for fast characterisation of silicon bricks, giving instant
information about the position of low lifetime regions that are commonly observed near the bottom
and top of a brick and also exposing areas of high structural defect density. This information can be
used for example as process feedback during the production of silicon wafers or as a cutting guide for
the removal and recycling of low lifetime regions, the processing of which would result in underperforming cells.
Fig.4.6.1.1 Bulk lifetime cross section from
bottom to top measured on a mulsticrystalline
silicon brick obtained from the PL intensity
ratio method (solid line) and from a single PL
image. Strong artefacts are observed in the
PLIR results in the so-called “red-zones” near
the bottom and the top of the brick.
Our previous work has shown that spectral
photoluminescence imaging on side facets of
silicon bricks can provide quantitative bulk
lifetime and doping images if applied on silicon bricks or thick silicon wafers [4.6.1.4]. A
photoluminescence intensity ratio (PLIR) is measured in that technique using two different spectral
filters in front of the camera. We carried out a more comprehensive study of this method that
addresses previously reported artefacts in low lifetime regions and provides a more complete
understanding of the technique [4.6.1.5].
Initial studies showed good quantitative agreement between bulk lifetime data obtained from PLIR
with results obtained from a number of other measurement techniques over a wide range of bulk
lifetimes and across the central region of cast multicrystalline bricks. However, artefacts were
observed at the top and bottom in the so-called “red-zone”, where the presence of interstitial Fe
reduces the bulk lifetime to values on the order of one microsecond or less. The PLIR was found to
over-report lifetime in these regions, see Fig.4.6.1.1.
PL normalised
10
0
10
-1
10
-2
10
-3
10
-4
10
-5
10
-6
10
-7
25 mm
17 mm scaled
8 mm scaled
1 mm scaled
0.1 mm scaled
0
200
400
600
800
1000
Pixel
(a)
(b)
Fig. 4.6.1.2 (a) PL images of luminescence light penetrating through a 100 µm diameter pinhole,
detected by a Si CCD camera. The luminescence light is filtered with a 1063 nm long pass filter. (b)
Experimental point spread function. For accurate deconvolution the PSF needed to be measured
with a dynamic range equivalent to seven orders of magnitude in PL intensity.
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Fig. 4.6.1.3 Bulk lifetime
image of a directionally
solidified multicrystalline
silicon brick side face with
(left) and without (right)
optical deconvolution of the
individual PL images. Both
images are shown on the
same colour scale (in µs)
for direct comparison. A
significant reduction in
image
aretefacts
and
improvement in contrast is
obtained.
Defect luminescence from the highly impurity contaminated regions was suspected to be the cause for
this experimental artefact. Spectrally resolved photoluminescence measurements on several silicon
bricks were performed, which revealed that luminescence originating from sub band gap defects do
not cause those artefacts. Rather, we find that optical light spreading within the silicon CCD is
responsible for most of the distortion in image contrast. Fig.4.6.1.2a shows a PL image of a silicon
sample that is covered with a 100 mm diameter pinhole. The image is measured with a Si CCD
camera, using a 1063nm long pass filter. The detected image intensity shows that light spreading
causes a smearing of the image contrast, i.e. a point source appears blurred in the image. Fig.4.6.1.2b
shows the experimentally measured point spread function, which quantifies this effect of light
spreading. We investigated a method to mitigate the impact of this light spreading effect via image
deconvolution. Fig.4.6.1.3 shows bulk lifetime images obtained using the PLIR method with (left) and
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without (right) image deconvolution. A significant improvement in image contrast and much
improved correlation of resulting bulk lifetime with established methods was demonstrated.
Point spread deconvolution is not only useful for the analysis of silicon bricks, but more broadly for
PL imaging on silicon wafers and cells. Further investigation are being conducted in collaboration
with the Australian National University [4.6.1.6].
Further potential sources of experimental error inherent in the PLIR are related to lateral carrier
diffusion and to free carrier absorption. These effects were analysed in more detail [4.6.1.7].
1.0
b normalised
0.8
1.5625
3.125
6.25
12.5
25
50
100
200
400
800
0.6
0.4
0.2
0.0
0.0
0.5
1.0
1.5
2.0
Fig.4.6.1.5. PLIR model prediction of
measured bulk minority carrier lifetime as a
function of distance to the grain boundary and
of grain bulk lifetime [µs]. A line indicates the
distance for which the respective calculated
bulk lifetime varies more than 10% from the
actual grain lifetime.
2.5
Distance to GB [mm]
The basic theoretical model underlying the PLIR is one dimensional, i.e. lateral variations in bulk
lifetime are ignored. Especially in multicrystalline silicon bricks, this assumption breaks down due to
the presence of structural defects and grain boundaries. Fig.4.6.1.4 shows the excess carrier profiles of
the grain/ grain boundary system, which was modelled two dimensionally using Sentaurus Device in
collaboration with Johannes Greulich from Fraunhofer Institute for Solar Energy Systems. Here, the
impact of a local grain boundary with infinite surface recombination velocity on the local carrier
density profile was calculated as a function of the bulk lifetime in the surrounding area. Employing
the optical photoluminescence intensity ratio model, a criterion can then be given by which the impact
of minority carrier diffusion on the bulk lifetime image is rated depending on both the distance to the
grain boundary and the local grain bulk lifetime (see Fig. 4.6.1.5).
4.6.1.3
Raw wafer inspection
Multicrystalline wafers PL images taken on as-cut multicrystalline wafers show different defects that
impact solar cell efficiency. These defects can broadly be classified as structural defects (grain
boundaries and dislocation networks) and impurities (often strongly dominated by iron). PL imaging
based quality control extracts metrics from PL images that are related to these defects using
automated image processing methods. The metrics can be correlated with cell data such as the
efficiency, VOC or ISC. The correlation of these metrics with cell performance data can then be used for
a range of actions, all aiming at higher yield, and higher average efficiency in production: (i) wafer
rejection, (ii) wafer quality based pricing, (iii) wafer quality specific processing, or (iv) assignment of
specific types of wafers to specially tuned production lines. Separating specific wafers for high
efficiency processing lines such as a selective emitter line, which requires higher quality wafers to
realise the full efficiency potential and may not be as robust against specific types of defects as the
common screen printed process, is one specific example for that latter application.
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Wafer to cell correlation trials typically consist of
wafer inspection, grading and sorting into a
number of quality bins, followed by cell processing
and reporting of IV data per bin by the cell
manufacturer. Results from such trials between
UNSW spin off company BT Imaging and two
industrial partners are shown in Fig. 4.6.1.6. They
represent trials using 2,000 and 1,150 wafers,
respectively. In each case wafers were binned by
the impurity area fraction (AE), and dislocation
area fraction (14) and each figure plots the
absolute difference in efficiency for each bin
relative to the lowest efficiency bin in each trial.
Both manufacturing lines exhibit a reduction in the
average efficiency as the dislocation density
increases from Bin 1 to Bin 4, irrespective of the
impurities. These dependencies are moderate, and
stronger dependencies have been observed in other
trials. It is interesting to note however that the two
manufacturing lines differ substantially in their
sensitivity to impurities. In Fig. 4.6.1.6a, there is a
significant reduction in efficiency as the impurity
metric increases from Bin A to Bin E. In this case,
the cell process is not effective in terms of
reducing the negative impact of impurities by
gettering during the phosphorus diffusion or by
hydrogenation. By contrast, Fig.4.6.1.6b does not
contain any obvious dependence on the impurity
metric; the manufacturing process is effectively
gettering all impurities in grades AD.
Fig.4.6.1.6 Average efficiency from two trials
for bins defined by impurities (AE) and
dislocations (14).See text for details. Data
from [4.6.1.12].
Cast mono wafers In early 2012 the development of cast mono wafers was pursued by many
companies as an effective path towards further cost reduction in silicon PV, with the promise of
combining the high efficiency potential of Czochralsky (Cz) monocrystalline wafers with the cost
benefits of casted multicrystalline material, with the added benefit of lower oxygen content and thus
reduced BO-complex related degradation. Not least the availability of PL imaging has resulted in a
significant reduction in activity in this area. PL inspection reveals a strong variation in structural
defect concentration in these wafers, which in turn causes strong variations in solar cell performance,
as will be shown below.
a)
b)
c)
d)
e)
Fig.4.6.1.7 PL images of five wafers from different equidistant heights of a cast mono center brick. a)
wafer near the bottom, just above the impurity rich so-called “red zone”, e) wafer from near the top.
High solar cell efficiencies obtained from cast mono wafers were presented e.g. by Clark et al., who
demonstrated a tight efficiency histogram with efficiencies of 18.8±0.15% for cells made from 100
wafers from the near bottom region of a cast ingot [4.6.1.8]. However, the average cell efficiency of
cells from 300 wafers representing the entire brick was reported to be 17.83%, thus suggesting a total
- 229 -
efficiency spread significantly exceeding 2% absolute efficiency. The likely cause for this large
efficiency spread can be identified in PL images taken on as-cut cast mono wafers.
Efficiency
Figure 4.6.1.7 shows PL images taken on 100% monocrystalline wafers from different heights in a
cast mono center brick (note, these wafers are not related to the study by Clark). Substantial variations
in the presence of strongly recombination active dislocation networks are observed from bottom to
top. In the PL images these dislocation networks are visible as dark lines and can be quantified using
automated image processing algorithms. These
wafers (plus 20 additional wafers from the same
0.174
brick) were processed into industrial screen printed
0.172
solar cells. Fig.4.6.1.8 shows the cell efficiencies as
0.170
a function of the defect density obtained from the
0.168
PL images. The structural defects caused an
0.166
absolute efficiency spread of 1.6%. The dislocation
0.164
density can therefore clearly be identified as the
0.162
main cause for the efficiency spread from wafers in
0.160
a cast mono ingot. In addition to these electronic
0.158
material quality variations cast mono wafers also
exhibit a variation in the monocrystalline area
2
3
4
5
6
7
fraction. Given these variations and the large cell
Defect metric
Fig.4.6.1.8 Efficiency of 25 solar cells made efficiency variations it remains to be seen whether
or not cast mono can provide a real cost benefit in
from 100% monocrystalline wafers across the
$/W. It appears that by the end of 2012 most wafer
height of a casted mono centre brick as a and cell manufacturers started to reduce R&D in
function of the defect metric obtained from PL this area.
images taken on the as-cut wafers.
4.6.1.4
Commercialisation of PL imaging by BT Imaging
A significant technological advancement was made by BT Imaging in
the area of inline PL imaging on as-cut wafers. The previous inline tool,
the iLS-W1 used similar illumination and detection optics as employed
in laboratory tools, i.e. full area illumination and detection.
Measurements are performed in steady state and require stopping the
wafer for the measurement. That constraint turned out to be
incompatible with a number of specific applications, for example for
outgoing quality control in automated wafer lines, where the wafers
normally don’t stop. The iLS-W2 (Fig.4.6.1.9) is a line scanning PL
system, which achieves 3,600 wafer per hour throughput at line speeds
of up to 300mm/s. Importantly measurements are carried out “on the
fly”, i.e. without the need to stop the wafer. The iLS-W2 was introduced
in the PV market in 2012 and is now offered by BT Imaging in a variety
of automation configurations.
The flagship laboratory tool, the LIS-R1 was replaced by a technically
improved LIS-R2 in 2012. The R2 incorporates improved optics, which
were adapted from the iLS-W1 inline systems, enabling further
improvements in image quality even beyond the already industry
leading standard of the R1 and substantial reduction in measurement
time for wafers and bricks.
4.6.1.5
Fig.4.6.1.9. BT Imaging
iLS-W2, line scanning PL
system with a throughpout
of up to 3,600 wafers per
hour.
Contactless on the fly series resistance imaging
Experimental data suggesting the analysis of luminescence images in terms of series resistance
images on fully processed silicon solar cell were first presented by our group in 2006 [4.6.1.9]. A first
quantitative analysis of PL images taken under different operating conditions, resulting in series
resistance images in cm2 was demonstrated shortly after that [4.6.1.2].
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In principle resistance information is contained in any luminescence image taken under operating
conditions that cause current flow, including PL images taken with simultaneous current extraction or
electroluminescence (EL) images. Generally, the PL based methods are seen to be more accurate,
since they allow a more precise distinction of series resistance related features from lifetime related
features and since they represent the series resistance that matters for the cell operation near the
maximum power point.
Previous RS imaging techniques require electrical contact to the cell for current injection or current
extraction. A contactless method to obtain qualitative series resistance data was suggested by
Kasemann [4.6.1.10]. The method is based on PL images taken with non-uniform illumination,
resulting in lateral current flow within the cell from illuminated to non-illuminated areas. Kasemann
used shading masks for the creation of the non-uniform illumination. His approach requires a
minimum of three images (two partially shaded images and one non-shaded image). These images are
taken while the cell is not in motion and require mechanical or manual insertion of the shading masks.
As discussed above, the BTi iLS-W2 is a line scanning PL imaging system. Only a small fraction of
the wafer is illuminated at any time during the measurement, which causes lateral current flow within
Fig.4.6.1.10. Left: Series resistance image of a multicrystalline silicon solar cell. Right: Single PL
image taken on the same cell using a BT ImagingiLS-W2 line scanning system. For the latter image
the cell was not contacted and was moving at a constant speed of 250mm/s. Data from [4.6.1.12].
the cell from the illuminated to the non-illuminated areas (as in the above shadow mask technique).
Within the typical time frames of such measurements the cell is in quasi steady state with regard to
these lateral current flows at any time. Importantly the illumination intensity within the illuminated
line can be used to manipulate the fraction of photocurrent that is extracted. At low illumination
intensity (e.g. 1 Sun equivalent or less) efficient lateral photocurrent extraction takes place. In
contrast, with increasing illumination intensities significantly above 1 Sun only a reduced fraction of
the photocurrent is extracted and the resulting PL image approaches operating conditions near VOC.
In analogy to the various previous approaches for Rs-imaging we can thus use iLS-W2 line-scan PL
images taken with different illumination intensities to create a number of images with different
operating conditions over a wide range from near Voc (at high intensity) to below the maximum
power point (which is in contrast to the shading method, which is limited to less than half the
photocurrent being extracted and therefore unable to approach MPP conditions). Even non-contact
and non-stop electroluminescence images can be obtained in this fashion if the imaged area is chosen
to be outside the illuminated area, i.e. we rely on lateral current injection via the emitter.
In order to detect local features with enhanced series resistance which are not unambiguously
measureable from IV data in full scale production, a rather qualitative analysis may be sufficient.
Initial proof of concept studies were performed, which clearly demonstrate the feasibility of obtaining
qualitative series resistance data from a single PL image. Fig.4.6.1.10 (left) shows a series resistance
image taken on an industrial six-inch multicrystalline silicon solar cell using a BT Imaging R2
- 231 -
luminescence imaging system and using the RS-imaging method proposed by Kampwerth [4.6.1.11].
Fig. 4.6.1.10 (right) shows a single PL image taken on an iLS-W2 inline system using 2 Suns
equivalent illumination intensity. That measurement was taken without making contact to the cell and
with the cell in constant motion at 250mm/s belt speed. Good qualitative agreement is observed
between strong series resistance features in both images. As expected, some local lifetime related
features, which are removed in the quantitative method, remain visible in the line scan image.
This preliminary data supports the idea that spatially resolved RS information, particularly on strong
local features of enhanced series resistance, as caused by broken grid fingers or by firing problems,
can be gained from luminescence images with no need to contact or stop the cell. A fully quantitative
series resistance analysis appears feasible and will be investigated in further work.
4.6.1.6
Summary
An exceptional variety of material and solar cell parameters can already be measured on silicon
bricks, silicon wafers and silicon solar cells with high lateral spatial resolution and short measurement
time using luminescence imaging techniques. The range of applications continues to grow, with new
applications for PL imaging being developed at UNSW, and increasingly also in other research
institutes and by R&D groups in PV companies. PL imaging, introduced at UNSW only a few years
ago, has been rapidly adopted as a standard characterisation method, with PL tools now in use at
virtually all leading research institutes worldwide and also by leading wafer and solar cell
manufacturers. A range of opportunities for in-line quality control, process monitoring and process
control have been demonstrated and are enabled by the high resolution and speed of PL imaging. The
adoption of these techniques continues and will contribute to the further cost reduction in silicon solar
cell manufacturing and to the widespread adoption of PV as a cost competitive energy source.
References
4.6.1.1 Trupke, T. et al. Photoluminescence imaging of silicon wafers. Applied Physics Letters 89,
44107 (2006).
4.6.1.2 Trupke, T. et al. Spatially resolved series resistance of silicon solar cells obtained from
luminescence imaging. Applied Physics Letters 90, 93506 (2007).
4.6.1.3 Würfel, P. et al., Diffusion lengths of silicon solar cells obtained from luminescence images.
Journal of Applied Physics 101, 123110 (2007).
4.6.1.4 Mitchell B. et al. Bulk minority carrier lifetimes and doping of silicon bricks from
photoluminescence intensity ratios, Journal of Applied Physics 109, 083111 (2011).
4.6.1.5 Mitchell B., On the method of photoluminescence spectral intensity ratio imaging of silicon
bricks : advances and limitations, Journal of Applied Physics 112, 063116 (2012).
4.6.1.6 Walter D. et al., Contrast Enhancement of Luminescence Images via Point-Spread
Deconvolution, in: IEEE 38th Photovoltaic Specialists Conference, 2012.
4.6.1.7 Mitchell B. et al., Quantifying the effect of minority carrier diffusion and free carrier
absorption on photoluminescence bulk lifetime imaging of silicon bricks, Solar Energy
Materials and Solar Cells 107, 75–80 (2012).
4.6.1.8 Clark, R. Fet al.. Status of Mono 2 TM Gen5 Casting at AMG IdealCast Solar. 38th IEEE
Photovoltaic Specialists Conference, Austin, USA, (2012).
4.6.1.9 Trupke, T. et al. (2006). Fast photoluminescence imaging of silicon wafers. 4th World
Conference on Photovoltaic Energy Conversion, WCPEC-4, Waikoloa, USA (2006).
4.6.1.10 Kasemann, M. et al. Contactless qualitative series resistance imaging on solar cells. IEEE
Journal of Photovoltaics 2(2), 181 (2012).
4.6.1.11 Kampwerth, H. et al., Advanced luminescence based effective series resistance imaging of
silicon solar cells. Applied Physics Letters 93, 202102 (2008).
4.6.1.12 Trupke T. et al. Inline Photoluminescence Imaging for Industrial Wafer- and Cell
Manufacturing, 22nd Workshop on Crystalline Silicon Solar Cells & Modules: Materials and
Processes, Vail, USA, July 2012.
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4.6.2
Other Optical Characterisation Techniques
University Staff
Dr Henner Kampwerth (Group Leader)
Prof Martin A Green
A/Prof Throsten Trupke
Dr Yang Yang
Postgraduate Students
Chao Shen (PhD student)
Kai Wang (PhD student)
External Collaborators
Prof Ottwin Breitenstein, Max Planck Institute of Microstructure Physics, Germany
Dr Jürgen Carstensen, Christian-Albrechts-Universität zu Kiel, Germany
A/Prof Daniel Macdonald, ANU
Dr Bill McLean, UNSW
Dr Martin Schubert, FHG-ISE, Germany
Andreas Schütt, Christian-Albrechts-Universität zu Kiel, Germany
4.6.2.1 Background
The photovoltaic industry is constantly trying to decrease the cost-per-watt of solar cells. This is
either achieved by (a) lowering the manufacturing costs, which involves the use of new
manufacturing techniques and lower-quality materials without lowering the established cell efficiency
or (b) the use of higher-efficiency cell designs without increasing the cost. Both approaches require
the development of new, more accurate and easier-to-interpret measurement techniques. Lowerquality materials with their higher levels of impurities and defects need to be better understood. Also,
higher-efficiency cell designs with their stricter process tolerances need to be monitored more
precisely.
Our Centre, with its leading role in research and development for industry and academia, is
particularly interested in measurement techniques that can produce accurate and meaningful data. The
improvement of measurement techniques has a direct impact on both research activities and the
optimisation of procedures used in manufacturing. The time and number of experiments needed to
understand a certain phenomenon can be greatly reduced. Correct interpretation of these results is also
a central concern.
This research group focuses primarily on the significant improvement of existing techniques and the
development of completely new measurement concepts that would be of value for multiple research
projects. Since mid 2012, our research has focused especially on increasing the applicability of
advanced time-resolved photoluminescence spectroscopy (Section 4.6.2.2). We have also continued to
develop algorithms for the extraction of local diode characteristics from photoluminescence images
(Section 4.6.2.3), an activity connected to the work of A/Prof. Trupke’s group (Section 4.6.1). The
third research activity in this group is focused on the optical reflection characteristics of solar cell
back reflectors. Dr Yang Yang developed in this group the experimental setup, now being used for
research in the High Efficiency Cells group under Dr Anita Ho-Baillie. For details, please be referred
to the report of that group (Section 4.3.1.1).
4.6.2.2 Fast Time-Resolved Photoluminescence Spectroscopy Measurements
The minority carrier lifetime is a key parameter in determining the performance of a solar cell. The
detailed study of this parameter and its dependency on various types of defects are the focus of an
- 233 -
international research project under the lead of Dr Kampwerth and supported by ASI/ARENA
[4.6.2.1]. Project partners are Dr Macdonald (ANU, Canberra) and Dr Schubert (Fraunhofer ISE
Freiburg, Germany).
Highly contaminated or defect-rich silicon, which is often found in thin film solar cells or is found in
grain boundaries, has very short minority carrier lifetimes of often less than 1 μs. Measurement of this
parameter with commonly-used techniques (such as quasi-steady-state photoluminescence decay or
standard photoconductance decay techniques) is impossible as they are limited to larger lifetimes and
probing areas. An additional difficulty is that short lifetimes, thin test-samples or small areas create
only very weak luminescence and photoconductance signals. This is often beyond the speed and
sensitivity of typically-used analogue detectors and electronics.
Experimental Setup
In this project, we at UNSW are building a unique and state of the art system for very detailed
minority carrier lifetime measurements. It will close the vital gap between the ultrafast optically-gated
PL spectrometer—which measures lifetimes in the fs to ps range—used by the Third Generation
group (Section 4.5.3.3.2), and the more commonly-used optical lifetime techniques used in
photovoltaics which work in the range of µs to ms.
The system is designed to measure:
 (photo-) luminescence with 600 ps time resolution. This allows the measurement of a
minimal minority carrier lifetime of 1 ns. The longest lifetime to be measured will be beyond
1 µs. (Overcoming technical barriers to exceed 1 µs are the focus of current work.)
 with a very high usable light sensitivity of 10 fA equivalent, or 1x104 photons/s and a wide
dynamic range. This is equivalent to highly amplified cooled analogue detectors
 with detectors sensitive in a wavelength range between 500 nm and 1,600 nm with sub-nm
spectral resolution.
 with an illumination wavelength range variable between 500 nm and 2,200 nm with several
nm spectral resolution.
 on small spots on the test sample of less than 40 µm in diameter.
 while the test sample is heated or cooled between 105 K and 370 K (-168°C and +97°C).
This system combines several characteristics that are usually distributed over multiple individual
systems. For example, common spectroscopes measure the luminescence spectrum without any time
resolution and time resolved setups often do not feature optical spectral resolution, measuring at one
or a few wavelengths. By combining both features it becomes possible to obtain luminescence spectra
within the first 1 ns after excitation. In this time scale charge carriers will not have moved and spatial
scans will be blur free.
The functionality of the system, which is depicted with its components [Fig. 4.6.2.1], can be
explained as follows: a very short laser pulse illuminates the test sample and generates minority
carriers. These carriers recombine after a statistical lifetime, a material parameter of upmost
importance to photovoltaic device operation. A certain percentage of these charge carriers will decay
via a radiative recombination process that creates photons. These photons are then measured with an
extremely sensitive and fast detector. Spectral filters for the excitation and detection beams allow the
further investigation of energy levels, concentrations and lifetimes of various contaminants and
defects in detail. To investigate very small and often difficult-to-measure sections such as grain
boundaries, laser doped areas or Si-metal contact regions, a confocal microscope is used to pump and
probe the sample on very small spots. Attenuators are used to change the excitation light intensity on
the sample and to avoid overexposure of the detectors. For practical reasons, a camera with its own
illumination source can be temporarily used to find spots of interest on the test sample before the
actual measurement. Two optical fibres mechanically separate the main sections of excitation light
conditioning, probing and analyzing. This allows not only a plug-and-play interface with other setups,
but also limits re-alignment and optimisation procedures to the section that is modified.
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Figure 4.6.2.1: The basic components of the time resolved PL spectroscopy setup. It is
separated into three major operational blocks of the conditioning of the excitation beam
(upper section), probing of the test sample (lower right hand side) and the analysing of the
probe beam (lower left hand side). top: A pulsed white light (hyperspectrum) laser is first
spectrally filtered via automated or individual fixed band-pass filters, attenuated in its
intensity and further pulse-picked to the required repetition rate. bottom right: A confocal
microscope focuses the excitation beam onto the test sample, which is controlled in its
temperature and x,y,z position. For spot finding on the sample, a camera with its own
illumination source can be used temporarily. bottom left: The probe beam is first spectrally
filtered via a monochromator, intensity adjusted and then spectrally split to feed two fast APD
detectors. Not shown are the timing electronics.
Theoretical Work
Luminescence decay measurements with high time resolution provide more details than what has been
commonly used for lifetime measurements. To make use of all details of the measurement signal, we
have begun to develop more sophisticated mathematical models. These will then be used for iterative
curve fitting to measured data to extract material parameters.
PhD candidate Kai Wang is focusing on the equations for time- and spectral-resolved luminescence.
Under the lead of Prof Green, he has been able to find an analytical solution [4.6.2.3] of Luke and
Cheng’s work [4.6.2.2] for thick wafers under transient light pulse excitation, see fig [Fig. 4.6.2.2].
The result provides the researcher with one relative simple equation which combines surface- and
bulk- recombination, as well as the carrier diffusion involved. One of the main results of this work is
shown in equations [Eq. 4.6.2.1]. It describes the time-resolved photoluminescence spectrum after the
illuminating light of the chosen wavelength is switched off. BB is the black body photon flux emission
at given wavelength and temperature, Na doping level, ni intrinsic carrier concentration, F* photon
flux entering the semiconductor, D diffusitivity, t time and S surface recombination velocity. Function
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f() can be approximated with a remaining inaccuracy of only 0.08%. If required, a better but slightly
longer approximation is provided in the article. All variables with a bar are normalised values. Figure
[Fig. 4.6.2.3] shows the graphic representation of this equation for selected parameters in normalised
values. Figure [Fig. 4.6.2.4] visualises the shortcomings of the commonly used simple exponential
decay equation to calculate the effective lifetime τeff. Fraction τeff/τ should correctly be unity, as τ is
the real bulk lifetime. By using equation [Eq. 4.6.2.1], the bulk lifetime τ and surface recombination
velocity S can now be correctly separated.
(4.6.2.1)
with
Ongoing work looks very promising to extend this to wafers of all thicknesses and different types of
illuminating light pulses.
Figure 4.6.2.2: A silicon block with an
absorption coefficient α1 is illuminated by
monochromatic light at wavelength λ1. The
photoluminescence is measured at a wavelength
λ2 where its absorption coefficient is α2. Rf(λ2)
indicates the internal reflection of the front
surface.
Figure 4.6.2.3: Computed ‘‘on-off’’ transient
photoluminescence (PL) from a brick with
normalised parameters. = 3,
= 0.3 and
= 2, with τ representing the bulk minority
carrier lifetime. The photoluminescence is
normalised to the steady-state value and shows
the symmetry between ‘‘on’’ and ‘‘off’’
transients.
Figure 4.6.2.4: Normalised τeff, the effective
lifetime corresponding to the negative inverse
of the slope of the logarithmic off-transient
for normalised ranging from S = 0 (upper
curve) to S  ∞ (lower curve). This
demonstrates the shortcomings of using a
simple exponential decay equation for
transient light sources.
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4.6.2.3 Local I-V Characteristics from PL Images
In recent years, several publications have focused on the extraction of spatially resolved maps of
electrical parameters from photoluminescence (PL) images, since Trupke et al. [4.6.2.4] demonstrated
the feasibility of this idea in 2006. These maps show a local electrical parameter for each point of the
solar cell. Most of these parameters correspond to a relatively simple equivalent circuit, the one-diode
model with series resistance.
PhD candidate Chao Shen successfully extended these theories to the more commonly used two-diode
equation and derived several valuable parameter maps from this [4.6.2.5].
Even though the newly developed technique requires a longer measurement time than other PL
imaging techniques, it offers the advantages of producing maps of the following local parameters
(denoted with index xy, referring to spatial coordinates):
 series resistance RS,xy
 dark saturation current density J01,xy (diode 1)
 dark saturation current density J02,xy (diode 2)
 net generated / sunk current density Jxy(Vterm) (at arbitrary terminal voltages Vterm)
 voltage Vxy(Vterm) (at arbitrary terminal voltages Vterm)
 power density Pxy(Vterm) (at arbitrary terminal voltages Vterm)
 efficiency ηxy
The maps of parameters J02,xy, Jxy(Vterm), Pxy(Vterm), ηxy are new to PL-based imaging techniques and
promise to be a useful addition. The different parameter maps are demonstrated on a single multicrystalline Si solar cell in Figures [Fig. 4.6.2.5] to [Fig. 4.6.2.9]. The series resistance RS,xy highlights
all ohmic resistance problems in Figure [Fig. 4.6.2.5]. In this image a typical pattern of a conveyer
belt is visible. It belongs to a belt furnace that was used during co-firing. It is likely that the belt was
at a lower than targeted temperature, causing an insufficient formation of a low metal-Si contact. Also
visible are interruptions of the metal grid, such as broken fingers or even cracks. The advantage of an
image displaying electrical parameters becomes here obvious as the impact of a crack becomes
visible, which might be otherwise not detectable via normal optical inspection due to its small size.
In the images of dark saturation current of the first and second diode a two diode equivalent model
J01,xy and J02,xy indicate recombination activity in the often associated bulk and junction areas, see
Figure [Fig. 4.6.2.6]. A serial number is visible in the J02,xy map, indicating damage to the junction
and a reduction of cell performance, see Figure [Fig. 4.6.2.8].
In Figure [Fig. 4.6.2.7] are maps of the local voltage and effective current density at one-sun
equivalent illumination and maximum operating voltage at the cell terminals. The cracks in the metal
grid mostly impact the voltage in this particular cell. Only the highest series resistance area from the
belt furnace step and the damaged junction had a noticeable impact on the current.
Figure [Fig. 4.6.2.8] shows the local values of the fill factor FF and the efficiency ηxy. Thanks to the
quantitative values in the maps, underperforming areas can be easily identified and their overall
impact estimated.
Figure 4.6.2.9 is a comparison between the efficiency calculated by this PL method and the power
loss calculated via CELLO [4.6.2.6].
Our current research is focused on improving these algorithms. For example, the homogeneous series
resistance patterns in Figure [Fig. 4.6.2.5] are also visible in the dark saturation current density maps
J01,xy and J02,xy. This is likely due to an insufficient separation of these parameters in the calculation.
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Series Resistance RS,xy
[Ω·cm2]
Figure 4.6.2.5: The local effective series
resistance RS,xy shows a pattern belonging to
the conveyer belt of a belt-furnace. The
impact of broken fingers and cracks are also
visible.
Dark Saturation Current Density J01,xy
[pA/cm2]
Dark saturation current density J02,xy
[0.1μA/cm2]
Figure 4.6.2.6: left: the dark saturation current density of the first diode in a two diode model
J01,xy and right: the second diode J02,xy. This image also shows a serial number, which has
obviously damaged the junction.
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Voltage at MPP Jxy(Vterm,MPP)
[V]
Current Density at MPP Jxy(Vterm,MPP)
[A/cm2]
Figure 4.6.2.7: left: The voltage map at 1 sun illumination and terminal at maximum power
point. The crack, section A, the pattern of the belt furnace, section B and C, and the serial
number, section D, all reduced the local voltage. right: The current density map under the
same operation conditions is more homogeneous. Only the most severe belt furnace marks
and the serial number, section B and D, reduced the current.
Fill Factor FF
[-]
Power Density Pxy(Vterm,MPP)
Efficiency ηxy
[mW/cm2]
[%]
Figure 4.6.2.8: left: The fill factor FF and right: the efficiency ηxy are good maps that
summarise the performance of a cell. Areas that underperform are clearly visible, and the
magnitude of the parameter is presented quantitatively.
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Efficiency ηxy by PL
Power Loss by CELLO
Figure 4.6.2.9: bottom: A comparison
between the efficiency calculated from PL
images and power loss via CELLO method.
(Please note that one colour scheme was
inverted for compatibility reasons.) The
agreement is a good preliminary result.
4.6.2.4
Summary
The results of the ongoing and new research activities in this growing group have been excellent. New
ideas for photoluminescence imaging algorithms created noticeable advances in meaningful
quantitative imaging of solar cells. The new project for fast and microscopic photoluminescence
spectroscopy has also seen first results in the theoretical work. The Future measurement capabilities,
together with advances in its related data processing will create an additional world leading
competence in our centre.
References
4.6.2.1
4.6.2.2
4.6.2.3
4.6.2.4
4.6.2.5
4.6.2.6
ASI / ARENA 1-GER010 Time- and spectrally- resolved Photoluminescence for Silicon
Solar Cell Characterisation, 2012–2015.
K.L. Luke, and L. Cheng, “Analysis of the interaction of a laser pulse with a silicon
wafer: determination of bulk lifetime and surface recombination velocity”, J. Appl. Phys.
61 (1987), p. 2282–2293.
K. Wang, M.A. Green, H. Kampwerth, “Transient photoconductance and
photoluminescence from thick silicon wafers and bricks: Analytical solutions”, Solar
Energy Materials and Solar Cells, 111 (2013), p. 189–192.
Trupke, T., R. A. Bardos, et al. "Photoluminescence imaging of silicon wafers." Applied
Physics Letters. vol. 89 (2006) 044107.
C. Shen, H. Kampwerth, M.A. Green, T. Trupke, J. Carstensen, A. Schütt, “Spatially
Resolved Photoluminescence Imaging of Essential Silicon Solar Cell Parameters and
comparison with CELLO measurements”, Solar Energy Materials and Solar Cells, vol.
109 (2013), p. 77–81.
J. Carstensen, A. Schütt, and H. Föll, "CELLO local solar cell resistance maps: Modeling
of data and correlation to solar cell efficiency", 22nd European Photovoltaic Solar
Energy Conference, (2007) Milan.
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5.
ORGANISATIONAL STRUCTURE
When the Key Centre for Photovoltaic Engineering was awarded in 1999, it was established as an
autonomous Centre within the Faculty of Engineering of UNSW, becoming independent from the
School of Electrical Engineering where UNSW photovoltaic activities had been previously located.
At that time, the Key Centre was given the same operational independence and rights as the other
Engineering Schools, with the Director Stuart Wenham being given the same status and authority as
other Heads of Schools. With the awarding of the ARC Centre of Excellence in 2003, the University
committed to the Key Centre for Photovoltaic Engineering becoming the School of Photovoltaic
Engineering, within which the ARC Photovoltaics Centre of Excellence would be located. This
transition occurred in January 2006. The corresponding organisational relationships are shown in the
upper part of Fig. 5.1.
Figure 5.1: Centre organisational structure.
Within the Centre of Excellence, Deputy Directors have been appointed for the major strands of
research as shown. The Management Committee of the Centre comprises the six Directors and
Deputy Directors, along with the Head of School and the Business and Operations Manager. This
committee meets fortnightly on the 1st and 3rd Fridays of each month, with each Deputy Director
giving a report on the activities in his or her area over the preceding fortnight. This committee takes
responsibility for decision making within the Centre that affects the Centre as a whole, while the
individual Deputy Directors are responsible for their own annual budgets and make and implement
decisions that impact only their own laboratory areas and research activities.
The Advisory Committee for the Centre comprises the Centre Directors, the Head of School for
Photovoltaic Engineering, leading academics from other institutions, industry leaders such as CEO’s
of various companies involved in the field, and research leaders. This committee provides high level
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advice, feedback and recommendations in relation to the Centre’s activities and their relevance. The
Advisory committee meets annually, taking into account the geographical separation of its members
(Australia, USA, China, Germany, Italy, Spain), although more frequent correspondence takes place
when necessary with individual committee members either by email or telephone. The membership of
the Advisory Committee has included:














Professor Stuart Wenham, Centre Director
Professor Martin Green, Centre Executive Research Director
Dr Gavin Conibeer, Deputy Director (3rd Generation PV Research)
Associate Professor Thorsten Trupke, Deputy Director (Silicon Photonics)
Dr Sergey Varlamov, Deputy Director (2nd Generation PV research)
Professor Allen Barnett: world leading academic in the field; recipient of many prestigious
international prizes such as the William Cherry Award; former President of AstroPower, one
of the world’s largest solar cell manufacturers, before being purchased by GE.
Prof. Andres Cuevas, ANU;
Dr Francesca Ferrazza, Chief Scientific Officer, Enitechnologie;
David Hogg: until recently COO of Suntech Power; formerly CEO of CSG Solar AC,
Germany and former CEO of Pacific Solar, Australia;
Mr. David Jordan, until recently, Manager of New Technology, BP Solar International;
Dr Larry Kazmerski, until recently Head, Photovoltaic Division, US National Renewable
Energy Laboratory;
Prof. Antonio Luque, Head, ESTI, Polytechnical University of Madrid; Leading European
academic in the photovoltaic field; founder of Isofoton, one of Europe’s most established
manufacturers;
Dr Zhengrong Shi, former CEO and Chair of Suntech-Power, the world’s largest silicon cell
manufacturer; recipient of various industry prizes; former Deputy Research Director of
Pacific Solar; major collaborator of the Centre;.
Prof. Peter Würfel, Institut fur Angewandte Physik der Universität Karlsruhe.
Other distinguished researchers and industrialists have been co-opted to the Committee as required.
Also within the Management structure and a member of the Centre’s Management Committee is the
Centre’s Business, Technology and Operations Manager Mark Silver, given the extensive
involvement of Centre staff with industry, large number of collaborative research projects, and the
high level of success of the Centre in generating, marketing and commercialising technology. The
Head of Administrative Support for the Centre during 2012 was Joyce Ho. Not shown in the
Management Structure is Elena Estrina who has looked after the Centre Accounts, interfacing with
the University Financial System, helping the Centre Directors and Deputy Directors with budgeting
and financial matters.
A primary responsibility of the Director in the Centre is to ensure the Management structure as shown
functions efficiently and effectively, with appropriate support and resources to facilitate the Centre’s
achievement of its milestones and performance targets. An important aspect of this is managing the
Centre’s finances to ensure suitable levels of funding are made available for supporting the major
research strands, while simultaneously ensuring adequate funding is available for repair and
maintenance of existing facilities and the purchasing of appropriate new infrastructure and equipment.
The Director works closely with the Centre’s Executive Research Director to assist in setting high
level research direction and priorities for the Centre’s programs. In the teaching area, with the Key
Centre for Photovoltaic Engineering and its activities being incorporated into the Centre of
Excellence, the Director also takes significant responsibility in ensuring the successful development
and implementation of the Centre’s educational programs, both at undergraduate and postgraduate
levels. In this area, the Director works closely with Dr Richard Corkish, the Head of School of
Photovoltaic and Renewable Energy Engineering.
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6.
6.1
PUBLICATIONS
BOOKS
Stuart R. Wenham, Martin A. Green, Muriel E. Watt, Richard P. Corkish, Alistair Sproul, Applied
Photovoltaics, Earthscan from Routledge, London, 3rd Edition (ISBN 9781849711425 (paperback),
9781849711418 (hardback), 9781849776981 (e-book)) 2012.
http://www.routledge.com/books/details/9781849711425/
6.2
BOOK CHAPTERS
Allen Barnett and Xiaoting Wang, “High Efficiency Photovoltaics Leading to Low Energy Cost”,
Secure and Green Energy Economies, February 2013 (in press).
M.A. Green, “Limits to Photovoltaic Energy Conversion”, in Imperial College Series on
Photoconversion, " Clean Electricity from Photovoltaics", 2nd Edition, M.D. Archer and M.A.Green
(Eds.) (Imperial College Press, London, 2013).
M.A. Green, “Crystalline Silicon Solar Cells”, in Imperial College Series on Photoconversion,
“Photovoltaics for Clean Electricity”, 2nd Edition, M.D. Archer and M.A.Green (Eds.) (Imperial
College Press, London, 2013).
Martin A Green, “Photovoltaic Material Resources”, Chapter 6, in “Semiconductors and Semimetals
(Vol. 87): Advances in Photovoltaics (Vol. 1)”, G.P. Willeke and E.R. Weber (Eds.), pp. 143-183,
Academic Press, Elsevier, 2012. (ISBN: 978-0-12-388419-0; ISSN 0080-8784) (invited).
S. Huang and G. Conibeer, “Silicon Quantum Dots for Next Generation of Silicon Based Solar Cells”,
in “Handbook on silicon Photonics”, Lorenzo Pavesi and Laurent Vivien (Eds.), in press, accepted
June 2012 by Taylor & Francis / CRC Press (invited).
D. König, “Introduction of Majority Carriers into Group IV Nanodots, Chapter 4 (pp. 1 - 30)” in J.
Valenta (Ed.), “Nanotechnology and Photovoltaic Devices: Light Energy Harvesting with Group-IV
Nanostructures”, Pan Stanford Press, in progress, scheduled for Aug. 2013 (invited).
S. Pillai, M.A. Green, “Plasmonics for Photovoltaics”, In: “Comprehensive Renewable Energy”, A
Sayigh (Ed.), Vol. 1, pp. 641-656, Oxford: Elsevier Ltd., 2012. (Comprehensive Renewable Energy
was awarded the PROSE Award 2012, Reference Work: Best Multivolume Reference/Science, by the
American Association of Publishers.)
Steven A. Ringel, Timothy J. Anderson, Martin Green, Rajendra Singh and Robert J. Walters,
"Photovoltaic Devices", Chapter 16, IEEE Electron Devices Society, 2013.
A. Slaoui, J. Poortmans, M. Topic and G. Conibeer, EOLSS chapter for UNESCO Encyclopedia,
“Evolution of Photovoltaic Materials for Renewable Energy Development”, article 6.106.56, accepted
May 2012.
M. Tayebjee, T. Schmidt and G. Conibeer, Chapter 31: “Downconversion” in Volume 1:
“Photovoltaics” W. van Sark (Ed.), in book “Comprehensive Renewable Energy” A. Sayigh (Ed.),
Vol. 1, Chapter 31, Oxford: Elsevier Ltd., 2012 (Comprehensive Renewable Energy was awarded the
PROSE Award 2012, Reference Work: Best Multivolume Reference/Science, by the American
Association of Publishers.)
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A. Uddin, “Photovoltaic Devices” in “Handbook of Research on Solar Energy Systems and
Nanotechnology”, S. Anwar, S. Qazi, and H. Efstathiadis (Eds.), Chapter 5, page 126 – 162,
published by IGI Global publisher, USA, 2012.
Ashraf Uddin and Martin A Green, “Donor-Acceptor Interface and Conversion Efficiency”, in
“Organic Solar Cells: New Research”, Nova Science Publishers, Inc. Hauppauge, NY, 2013 (in
press).
A. Uddin and M. A. Green, “Donor-Acceptor Interface and Conversion Efficiency”, in “Organic Solar
Cells: New Research”, Nova Science Publishers, Inc. Hauppauge, NY, 2012 (in press).
A. Uddin, “Doping of Organic Electronic Materials” in “Doping: Properties, Mechanism and
Applications”, L. Yu (Ed.), Nova Science Publishers, Inc. Hauppauge, NY, 2012 (in press).
6.3
PATENT APPLICATIONS
C. Chan, A. Wenham, B. Hallam and S. Wenham, “A monolithically integrated solar cell system”,
Australian provisional application 2013900704, filed 1 March 2013.
G. Conibeer, S. Shrestha, D. König and M.A. Green, “Hot carrier energy conversion structure and
method of fabricating the same”, CN201080029571.8, DE112010002821.4, JP2012-516441 and
US13/378580, filed 3 January 2012.
M.A. Green, H. Kampwerth and C. Shen, “A method for determining properties of a semiconductor
element”, Australian provisional application 2012900401, filed 3 February 2012.
M. Green, “Si cells with III-V multijunction cells”, US13/577617, filed 12 August 2012.
A. Lennon, S.R. Wenham, “A method for selective delivery of material to a substrate”, US13/389735,
filed 11 February 2012.
A. Lennon, Z. Li, S.R. Wenham and D. Lu, “Metallisation Method for rear Contact Cells idea #1 ‘lift
off’”, PCT/AU2012/000763 filed 28 June 2012; and TW101123599, filed 29 June 2012.
A. Lennon, Z. Li, S.R. Wenham and D. Lu, “Metallisation Method for rear Contact Cells idea #2
‘Aluminium Oxide’”, PCT/AU2012/000764 filed 28 June 2012 and TW101123600, filed 29 June
2012.
B. Puthen-Veettil and D. König, “Front-end hot carrier absorption boosted solar cell” Australian
provisional application 2012903645, filed 23 Aug. 2012.
L. Mai, S.R. Wenham and A. Wenham, “Photoplating of metal electrodes for solar cells”, US
13/505518, AU 2010314804, CN 201080049916.6 and KR 10-2012-7014531, filed 3 May 2012.
S.R. Wenham, L. Mai, “Hybrid solar cell contact”, PCT/AU2012/000367, filed 10 April 2012.
S.R. Wenham, B. Tjahjono, N. Kuepper and A. Lennon, “Improved metallization for silicon solar
cells”, US13/504253 and CN201080048559.1, filed 26 April 2012.
S.R. Wenham, et al., “Advanced hydrogenation of silicon solar cells”, Australian provisional
application 2012902090, filed 21 May 2012.
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S.R. Wenham, et al., “Formation of localised molten regions in silicon containing multiple impurity
types”, Australian provisional application 2012901267 filed 29 March 2012; PCT/AU2012/001302
filed 25 October 2012 and TW101139464, filed 25 October 2012.
S.R. Wenham, V. Vais, A. Lennon, Y. Zhao, B. Tjahjono, A. Wenham, “Field induced plating
method”, PCT/AU2012/001394 filed 12 November 2012 and TW101142651, filed 15 November
2012.
N. Western, S. Bremner, “A method of forming a contact for a photovoltaic cell”, Australian
provisional application 2012903616, filed 22 August 2012.
Y. Yao, A. Lennon, S.R. Wenham, et al., “Formation of metal contacts”, Australian provisional
application 2012904893, filed 9 November 2012.
6.4
PAPERS IN REFEREED SCIENTIFIC AND TECHNICAL JOURNALS
S. Bambrook and A. Sproul, “Maximising the energy output of a PVT air system”, Solar Energy, Vol.
86, pp.1857-1871, 2012.
K.Y. Ban, D. Kuciauskas, C.B. Honsberg and S. Bremner, “Observation of band alignment transition
in InAs/GaAsSb quantum dots by photoluminescence”, Journal of Applied Physics, Vol. 111(10), p.
104302 (4 pages), 2012.
A. Basch, F. J. Beck, T. Soderstrom, S. Varlamov and K. R. Catchpole, “Enhanced light trapping in
solar cells using Snow Globe Coating”, Progress in Photovoltaics: Research and Application (A),
2012, DOI: 10.1002/pip.2240.
A. Basch, F.J. Beck, T. Soderstrom, S. Varlamov, K Catchpole, “Combined plasmonic and dielectric
reflectors for enhanced photocurrent in solar cells”, Applied Physics Letters (A*), Vol. 100,
p. 243903, 2012.
T. Bray, Y. Zhao, P. Reece and S. P. Bremner, “Photoluminescence of antimony sprayed indium
arsenide quantum dots for novel photovoltaic devices”, Journal of Applied Physics, Vol. 113,
p. 093102, 2013.
R. Clady, M. Tayebjee, P. Aliberti, D. Konig, N.J. Ekins-Daukes, G. Conibeer, T. Schmidt and M.A.
Green, “Interplay between the hot phonon effect and intervalley scattering on the cooling rate of hot
carrier in GaAs and InP”, Progress in Photovoltaics, Research and Applications, Vol. 20 (1), pp. 8292, January 2012.
F. Comez-Gil, X. Wang and A. Barnett, “Energy production of photovoltaic systems: Fixed, tracking,
and concentrating”, Renewable and Sustainable Energy Reviews, Vol. 16, pp. 306-313, 2012.
F. Comez-Gil, X. Wang and A. Barnett, “Analysis and Prediction of Energy Production in
Concentrating Photovoltaic (CPV) Installations”, Energies, Vol. 5, pp. 770-789, 2012.
G. Conibeer, I. Perez-Wurfl, X. Hao, D. Di and D. Lin, “Si solid-state quantum dot-based materials
for tandem solar cells”, Nanoscale Research Letters, Vol. 7, p. 193, 2012.
R.P. Corkish, A. Barnett, S. Bremner, A.G. Bruce, M.J. Kay, A.J. Lennon, I. Perez-Wurfl, A.B.
Sproul, G.J. Stapleton, S.K. Shrestha, E.D. Spooner, R. Taylor, A. Uddin and D. Wentworth,
“Renewable Energy Education at UNSW”, Journal of Materials Education, Vol. 34, pp. 117-132,
2012.
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H. Cui and P. Campbell, “A glass thinning and texturing method for light incoupling in thin film
polycrystalline silicon solar cells application”, Physica Status Solidi (RRL), Rapid Research Letters 6,
p. 334, 2012.
H. Cui, P. Campbell and M.A. Green, “Compatibility of glass texture with e-beam evaporated silicon
thin-film solar cells”, Applied Physics A, DOI:10.1007/s00339-012-7318-3, 2012.
H. Cui, P. Campbell and M.A. Green, “Optimisation of back surface reflectors for polycrystalline
silicon thin film solar cells application”, Energy Procedia, accepted January 2013.
H. Cui, P.J. Gress, P.R. Campbell and M.A. Green, “Developments in the aluminium induced
texturing (AIT) glass process”, Glass Technology – European Journal of Glass Science & Technology
Part A, Vol. 53(4), pp. 158-165, August 2012.
H. Cui and M.A. Green, P. Campbell, O. Kunz, S. Varlamov, “A Photovoltaic light-trapping
estimation method for textured glass based on surface decoupling calculation”, Solar Energy
Materials and Solar Cells, Vol. 109, pp. 82-90, 2013.
H. Cui, S. Pillai, P. Campbell and M.A. Green, “A novel silver nanoparticle assisted texture as
broadband antireflection coating for solar cell applications”, Solar Energy Materials and Solar Cells,
Vol. 109, pp. 233-239, 2013.
J. Dore, S. Varlamov, B R. Evans, Eggleston, D. Ong, O. Kunz, J. Huang, U. Schubert, K. Kim, R.
Egan, M. Green, “Performance potential of low-defect density silicon thin-film solar cells obtained by
electron beam evaporation and laser crystallisation”, EPJ Photovoltaic, 2013, v.4, article no. 40301,
http://dx.doi.org/10.1051/epjpv/2012012.
J. Dore, R. Evans, U. Schubert, B. Eggleston, D. Ong, K. Kim, J. Huang, O. Kunz, M. Keevers, R.
Egan, S. Varlamov, M. Green, “Thin-film polycrystalline silicon solar cells formed by diode laser
crystallisation”, Progress in Photovoltaics: Research and Application (A), 2012, DOI:
10.1002/pip.2282.
B. Eggleston, S. Varlamov and M.A. Green, “Large area diode laser defect annealing of
polycrystalline silicon solar cells”, IEEE Transactions on Electron Devices, Vol. 59(10), pp. 28382841, October 2012.
Y. Feng, P. Aliberti, B.P. Veettil, R. Patterson, S. Shrestha, M.A. Green and G. Conibeer, “Non-ideal
energy selective contacts and their effect on the performance of a hot carrier solar cell with an indium
nitride absorber”, Applied Physics Letters, Vol. 100(5), p. 053502, January 2012.
Y. Feng, S. Lin, M. Green and G. Conibeer, “Effect of static carrier screening on the energy
relaxation of electrons in polar-semiconductor multiple-quantum-well superlattices”, Journal of
Applied Physics, Vol. 113(2), p. 024317, 2013.
F. J. Gómez-Gil, X. Wang and A. Barnett, “Analysis and prediction of energy production in
concentrating photovoltaic (cpv) installations”, Energies, Vol. 5, pp. 770-789, 2012.
F. J. Gomez-Gil, X. Wang and A. Barnett, “Energy production of photovoltaic systems: fixed,
tracking, and concentrating”, Renewable and Sustainable Energy Reviews, Vol. 16(1), pp. 306-313,
2012.
M.A. Green, K. Emery, Y. Hishikawa, W. Warta and E.D. Dunlop, “Solar cell efficiency tables
(Version 39), Progress in Photovoltaics, Vol. 20, pp. 12-20, January 2012.
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M.A. Green and S. Pillai, “Harnessing Plasmonics for Solar Cells”, Nature Photonics, Vol. 6(3), pp.
130-132, March 2012.
M.A. Green, "Analytical Treatment of Trivich-Flinn and Shockley-Queisser Photovoltaic Efficiency
Limits Using Polylogarithms" Progress in Photovoltaics, Research and Applications, Vol. 20(2),
pp.127-134, March 2012.
M.A. Green, “Radiative efficiency of state-of-the-art photovoltaic cells”, Short Communication,
Progress in Photovoltaics, Vol. 20(4), pp. 472-476, June 2012.
(Version 40), Progress in Photovoltaics, Vol. 20, pp. 606-614, August 2012.
M.A. Green, “Time-Asymmetric Photovoltaics”, Nano Letters, Vol. 12(11), pp. 5985-5988,
November 2012.
M.A. Green, “Limiting photovoltaic efficiency under new ASTM G173-based reference spectra”,
Progress in Photovoltaics, Vol. 20(8), pp. 954-959, December 2012.
(Version 41), Progress in Photovoltaics, Vol. 21, pp. 1-11, 2013.
M.A. Green, “Silicon solar cells: State-of-the-art”, A Philosophical Transactions of the Royal Society
of London Series A: Mathematical Physical and Engineering Sciences, 2012.
P. Gress and S. Varlamov, “Quantification of power losses of the interdigitated metallisation of the
crystalline silicon solar cells on glass”, Int. J. Photoenergy, Vol. 2012, ID 814697,
doi:10.1155/2012/814697, 2012.
B. Hallam, C. Chan, A. Sugianto and S. Wenham, “Deep junction laser doping for contacting buried
layers in silicon solar cells” Solar Energy Materials and Solar Cells, Vol. 113, pp. 124-134, 2013.
R. Hao, C. P. Murcia, C. Leitz, A. Gerger, A. Lochtefeld, K. Shreve, R. Opila and A. Barnett,
“Investigating voltage as a function of the reduced junction area for thin silicon solar cells that utilise
epitaxial lateral overgrowth”, IEEE Journal of Photovoltaics, Vol. 3(1), pp 119-124, January 2013.
L. Hua and S. Wenham, “Influence of a-Si:H deposition power on surface passivation property and
thermal stability of the a-Si:H/SiNx stack”, American Institute of Physics Advances, Vol. 2(2),
pp. 022106, 2012.
S. Huang and G. Conibeer, “Sputter-grown Si quantum dot nanostructures for tandem solar cells”,
Journal of Physics D: Applied Physics Vol. 46, p. 024003, 2013. (invited)
S. Huang, Y.H. So, G. Conibeer and M.A. Green, “Doping of silicon quantum dots embedded in
nitride matrix for all-silicon tandem cells”, Japanese Journal of Applied Physics, Vol. 51(10), Special
Issue Si, October 2012.
M.E. Ikpi, P. Atkinson, S.P. Bremner and D.A. Ritchie, “Fabrication of a self-aligned cross-wire
quantum-dot chain light emitting diode by molecular beam epitaxial regrowth”, Nanotechnology, Vol.
23(22), p. 225304, 2012.
X. Jia, D. Di, H. Xia, L. Wu; Z. Lin and G. Conibeer, “Size-dependent optical absorption of silicon
nanocrystals embedded in SiO2/Si3N4 hybrid matrix”, Journal of Non-crystalline Solids, accepted
November 2012.
- 247 -
C.M. Johnson and G.J. Conibeer, “Limiting efficiency of generalized realistic c-Si solar cells coupled
to ideal up-converters”, Journal of Applied Physics, Vol. 112, p. 103108, 2012.
C.M. Johnson, P.J. Reece, and G.J. Conibeer, “Theoretical and experimental evaluation of multilayer
porous silicon structures for enhanced erbium up-conversion luminescence”, Solar Energy Materials
and Solar Cells, Vol. 11, pp. 168-181, 2013.
C. Kerestes, Y. Wang, Yi, K. Shreve, J. Mutitu, T. Creazzo, P. Murcia and A. Barnett, “Design,
Fabrication, and Analysis of Transparent Silicon Solar Cells for Multi-Junction Assemblies”, Progress
in Photovoltaics: Research and Applications (Online: DOI: 10.1002/pip.1232), 23 February 2012.
T. Kim, P. Gress, S. Varlamov, “Metallisation and interconnection of e-beam evaporated
polycrystalline silicon thin-film solar cells on glass”, Int. J. Photoenergy, v. 2012, article ID 271738,
doi:10.1155/2012/271738, 2012.
D. König, Y. Takeda, B. Puthen-Veettil and G. Conibeer, “Lattice-matched hot carrier solar cell with
energy selectivity integrated into hot carrier absorber”, Japanese Journal of Applied Physics, Vol. 51,
p. 10ND02, 2012.
D. König, Y. Takeda and B. Puthen-Veettil, “Technology-compatible hot carrier solar cell with
energy selective hot carrier absorber and carrier-selective contacts”, Applied Physics Letters, Vol.
101, p. 153901, 2012.
D. König and J. Rudd, “Formation of Si or Ge nanodots in Si3N4 with in-situ donor modulation
doping of adjacent barrier material”, AIP Advances, Vol. 3, p. 012109, 2013.
D. König, D. Hiller, S. Michard, C. Flynn and M. Zacharias, “Static hot carrier populations as a
function of optical excitation energy detected through energy selective contacts by optically assisted
IV, Progress in Photovoltaics, available online (Early View® open access), Jan. 2013;
http://onlinelibrary.wiley.com/doi/10.1002/pip.2367/abstract.
M. Lenio, J. Howard, D. Lu, F. J. Jentschke, Y. Augarten, A. Lennon and S. Wenham, “Series
resistance analysis of passivated emitter rear contact cells patterned using inkjet printing”, Advances
in Materials Science and Engineering, p. 965418, 2012.
A. Lennon, Y. Yao and S. Wenham, “Evolution of metal plating techniques for silicon solar cell
metallisation”, Progress in Photovoltaics, accepted for publication, March 2012.
H. Li and S.R. Wenham, “Influence of a-Si:H deposition power on surface passivation property and
thermal stability of a-Si:H/SiNx:H stacks”, AIP Advances, Vol. 2, p. 022106, 2012.
H. Li and S. Wenham, “Influence of a-Si:H deposition temperature on surface passivation property
and thermal stability of a-Si:H/SiNx:H stack”, IEEE Journal of Photovoltaics, Vol. 2, pp. 405-410,
2012.
Z. Li, D. Cordiner, N. Borojevic, J. Rodriguez and A. Lennon, “Lift-off contact separation method for
rear contact solar cells,” Journal of Imaging Science and Technology, Vol. 56(4), p. 40505, 2012
Y Li, Z. Li, Y. Zhao and A. Lennon, “Modelling of light trapping in acidic-textured multicrystalline
silicon wafers”, International Journal of Photoenergy, Vol. 2012, Article ID 369101, pp. 1-8
(doi:10.1155/2012/369101), 2012.
P. H. Lu, K. Wang, Z. Lu, A. J. Lennon and S. R. Wenham, “Anodic aluminum oxide passivation for
silicon solar cells,” IEEE Journal of Photovoltaics, Vol. 3(1), pp. 143-151, January 2013.
- 248 -
X. Lu, S.R. Huang, M. Diaz, N. Kotulak, R. Hao, R.L. Opila and A. Barnett, “Wide band gap gallium
phosphide solar cells”, IEEE Journal of Photovoltaics, Vol. 2(2), pp. 214-220, 2012.
L. Mai, E.J. Mitchell, K.S. Wang, D. Lin and S.R. Wenham, “The development of the advanced
semiconductor finger solar cell”, Progress in Photovoltaics, accepted for publication, 2012.
A. Michael, O. Kazuo, Y. W. Xu, C. Y. Kwok, T. Puzzer and S. Varlamov, “Characterization of UHV
e-beam evaporated low-stress thick silicon film for MEMS application”, Procedia Engineering, 2012,
v. 47, pp. 690-693.
B. Mitchell, J. Greulich and T. Trupke, “Quantifying the effect of minority carrier diffusion and free
carrier absorption on photoluminescence bulk lifetime imaging of silicon bricks”, Solar Energy
Materials and Solar Cells, Vol. 107, pp. 75–80, 2012.
B. Mitchell, J.W. Weber, D. Walter, D. Macdonald and T. Trupke, “On the method of
photoluminescence spectral intensity ratio imaging of silicon bricks : advances and limitations”,
Journal of Applied Physics, Vol. 112, p. 063116, 2012.
K. Ohdaira, K. Sawada, N. Usami, S. Varlamov and H. Matsumura, “Large-grain polycrystalline
silicon films formed through flash-lamp-induced explosive crystallization”, Japanese Journal of
Applied Physics, Vol. 51, p. 10NB15, 2012.
I Perez-Wurfl, L. Ma, D. Lin, X. Hao, M.A. Green, G. Conibeer, “Silicon nanocrystals in an oxide
matrix for thin film solar cells with 492 mV open circuit voltage”, Solar Energy Materials and Solar
Cells, Vol. 100, pp. 65-68, May 2012.
J. Rao, S. Varlamov, J. Park, S. Dligatch, A. Chtanov, “Optimization of dielectric-coated silver
nanoparticle films for plasmonic-enhanced light trapping in Thin Film Silicon Solar Cells”, Journal of
Plasmonics, DOI 10.1007/s11468-012-9473-y, 2013.
J. Rodriguez, A. J. Lennon, M. Luo, Z. Li, Y. Yao, P. H. Lu, C. Chan, S. R. Wenham, “Dielectric
Patterning using Aerosol Jet Printing,” Science and Technology, Vol. 56(4), pp. 40502 (7 pages),
2012.
C. Shen, H. Kampwerth, M.A. Green, T. Trupke, J. Garstensen and A. Schütt, “Spatially resolved
photoluminescence imaging of essential silicon solar cell parameters and comparison with CELLO
measurements”, Solar Energy Materials and Solar Cells, Vol. 109, pp. 77-81, 2013.
L. Song and A. Uddin, “Design of high efficiency organic solar cell with light trapping”, Optics
Express, Vol. 20, pp. A606-621, 2012.
Y. Tao, S. Varlamov, O. Kunz, Z. Ouyang, J. Wong, T. Soderstrom, M. Wolf, R. Egan, “Effects of
annealing temperature on crystallisation kinetics, film properties and cell performance of silicon thinfilm solar cells on glass”, Sol Mat (A), DOI: 10.1016/j.solmat.2012.01.039, 2012.
T. Trupke and R. Sinton, “Limitations on dynamic excess carrier lifetime calibration methods”,
Progress in Photovoltaics: Research and Applications, Vol. 20, pp. 246 – 249, 2012.
T. Trupke and D. Macdonald, “Imaging crystal orientations in multicrystalline silicon wafers via
photoluminescence”, Applied Physics Letters, Vol. 101, pp. 082102, 2012.
T. Trupke, B Mitchell, J. Weber, R. Bardos, W. McMillan and R. Kroeze, “Photoluminescence
imaging for photovoltaic applications”, Energy Procedia, pp. 135-146, 2012.
- 249 -
A. Uddin, H.S. Wong and C.C. Teo,, “CdSe/ZnS quantum dot size dependent carrier relaxation in
hybrid organic/inorganic system”, Journal of Nanoscience and Nanotechnology, Vol. 12, pp. 1-7,
2012.
S. Varlamov, J. Rao, T. Soderstrom, “Polycrystalline silicon thin-film solar cells with plasmonicenhanced light-trapping”, J. Vis. Exp (JoVE), 65, pii: 4092, DOI: 10.3791/4092, 2012.
E.-C. Wang, S. Mokkapati, T. P. White, T. Soderstrom, S. Varlamov, K. Catchpole, “Light trapping
with titanium dioxide diffraction gratings fabricated by nanoimprinting”, Progress in Photovoltaics:
Research and Application (A), DOI: 10.1002/pip.2294, 2012.
E.-C. Wang, S. Mokkapati, T. Soderstrom, S. Varlamov, K. Catchpole, “Effect of nanoparticle size
distribution on the performance of plasmonic thin-film solar cells: monodisperse versus multidisperse
arrays”, IEEE Journal of Photovoltaics, DOI: 10.1109/JPHOTOV.2012.2210195, 2012.
K. Wang, M.A. Green and H. Kampwerth, “Transient photoconductance and photoluminescence from
thick silicon wafers and bricks: analytical solutions”, Solar Energy Materials and Solar Cells, Vol.
111, pp. 189-192, 2013.
Q. Wang, T. Soderstrom, K. Omaki, A. Lennon, S. Varlamov, “Etch-back silicon texturing for lighttrapping in electron-beam evaporated thin-film polycrystalline silicon solar cells”, Energy Procedia,
Vol. 15, pp. 220-228, 2012.
S. Wang, A. Lennon, B. Tjahjono, L. Mai, B. Vogl and S.R. Wenham, “Overcoming over-plating
problems for PECVD SiNx passivated laser doped p-type multi-crystalline silicon solar cells”, Solar
Energy Materials and Solar Cells, Vol. 99, pp. 226 - 234, 2012.
X. Wang, J. Byrne, L. Kurdgelashvili and A. Barnett, “High efficiency photovoltaics: on the way to
becoming a major electricity source”, WIREs Energy and Environment, Vol. 1(2), pp. 132–151, 2012.
X. Wang and A. Barnett, “The Effect of spectrum variation on the energy production of triplejunction solar cells”, IEEE Journal of Photovoltaics, Vol. 2(4), pp. 417-423, 2012.
J. Wong and M.A. Green, “From junction to terminal: Extended reciprocity relations in solar cell
operation”, Physical Review B00, 005200 (pp1-8), June 2012.
J. Wong, J. Huang, S. Varlamov M.A. Green and M. Keevers, “The roles of shallow and deep levels
in the recombination behavior of polycrystalline silicon on glass solar cells” Progress in
Photovoltalics, Vol. 20(8), pp. 915-922, December 2012.
M. Wright and A. Uddin, “Organic - Inorganic Hybrid Solar Cells: A Comparative Review”, Solar
Energy Materials and Solar Cells, Vol. 107, 87-111, 2012.
X. Yang, A. Uddin, and M. Wright, “Plasmon enhanced light absorption in bulk-heterojunction
organic solar cells”, Physical Status Solidi, RRL 6, pp. 199-201, 2012.
X. Yang, A. Uddin and M. Wright, “Effects of annealing on bulk-heterojunction organic solar cells”,
Advanced Science, Engineering & Medicine, Vol. 4, pp. 1-7, 2012.
Y. Yang, S. Pillai, H. Mehrvarz, H. Kampwerth, A. Ho-Baillie, M.A. Green, “Enhanced light trapping
for high efficiency crystalline solar cells by the application of rear surface plasmons”, Solar Energy
Materials and Solar Cells, Vol. 101, pp. 217-226, June 2012.
- 250 -
Y. Yang, M. A. Green, A. Ho-Baillie, H. Kampwerth, H. Mehrvarz and S. Pillai, “Characterisation of
2-D reflection pattern from textured front surfaces of silicon solar cells”, Solar Energy Materials and
Solar Cells, accepted March 2013.
Y. Yao, D. König and M.A. Green, “Investigation of boron antimonide as hot carrier absorber
material, Solar Energy Materials and Solar Cells”, Vol. 111, pp. 123-126, 2013.
Y. Yao, A. Sugianto, A. Lennon, B. Tjahjono and S. Wenham, “Use of inductively coupled plasma
measurements to characterise light induced plating for silicon solar cells”, Solar Energy Materials and
Solar Cells, VOl. 96, pp. 257-265, 2012.
Y. Wang, A. Gerger, A. Lochtefeld, L. Wang, C. Kerestes, R. Opila and A. Barnett, “Design,
fabrication and analysis of germanium: silicon solar cell in a multi-junction concentrator system”,
Solar Energy Materials & Solar Cells, Vol. 108, pp. 146-155, 2013.
B. Zhang, W. Truong, S. Shrestha, M.A. Green and G. Conibeer, “Structural, mechanical and optical
properties of Ge nanocrystals embedded in superlattices fabricated by in situ low temperature
annealing”, Physica E – Low Dimensional Systems & Nanostructures, Vol. 45, pp. 207-213, August
2012.
H. Zhang, S. Huang and G. Conibeer, “Study of photo-cathode materials for tandem
photoelectrochemical cell for direct water splitting”, Energy Procedia, Vol. 22, pp. 10-14, 2012.
6.5
CONFERENCE PAPERS AND OTHER PRESENTATIONS
J.I. Bilbao and A.B. Sproul, “Experimental results of a PVT-Water module design for developing
countries”, Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7 December 2012.
J.I. Bilbao and A.B. Sproul, “Analysis of flat plate photovoltaic-thermal (PVT) models”, World
Renewable Energy Forum (WREF 2012), Denver, 2012.
O. Breitenstein, C. Shen, H. Kampwerth and M.A. Green, “Comparison of DLIT- and PL-based local
solar cell efficiency analysis”, SiliconPV 2013, Hamelin, Germany, 25-27 March 2013 (abstract
submitted).
S. Bremner, C. Hongsberg, “Intermediate band solar cell with non-ideal band structure under AM1.5
spectrum”, 38th IEEE Photovoltaic Specialists Conference, Austin, 3-8 June 2012, pp. 21-24.
G. Conibeer, “Hot carrier cells: an example of third generation photovoltaics,” Proceedings, SPIE
Conference, San Diego, 12-16 August 2012, 8256, 82560Z.
G. Conibeer, S. Shrestha, H. Huang and M.A. Green, “Hot carrier solar cell absorbers:
superstructures, materials and mechanisms for slowed carrier cooling”, 38th IEEE Photovoltaic
Specialists Conference (PVSC), Austin, TX, 3-8 June 2012.
G. Conibeer, S. Shrestha, S. Huang, R. Patterson, P. Aliberti, H. Xia, Y. Feng, N. Gupta, S. Smyth, Y.
Liao and M. A. Green, “Hot carrier solar cell absorbers: superstructures, materials and mechanisms
for slowed carrier cooling”, Proceedings, SPIE Conference, San Diego, 12-16 August 2012.
G. Conibeer, S. Shrestha, S. Huang, R. Patterson, P. Aliberti, H. Xia, Y. Feng, N. Gupta, S. Smyth, Y.
Liao, Y. Kamikawa and M.A. Green, “Hot carrier solar cell absorbers: superstructures, materials and
mechanisms for slowed carrier cooling”, Proceedings, 27th European Photovoltaic Solar Energy
Conference, Frankfurt, Germany, 24-28 September 2012.
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J. Fan and R.P. Corkish, “Performance of PV systems with back-contact cells with and without
positive grounding”, International Conference on Clean Energy (ICCE 2012), Quebec City, Canada,
10-12 September 2012.
H. Cui, P. Campbell and M.A. Green, “Nano-imprinting for polycrystalline silicon thin-film solar
cells on glass superstrate application”, 27th European Photovoltaic Solar Energy Conference,
Frankfurt, Germany, 24-28 September 2012, pp. 2408 - 2412.
H. Cui, S. Pillai, P.R. Campbell and M.A. Green, “A novel broadband antireflection coating using Ag
nano-particles as etch mask for poly-si thin-film solar cells applications”, Proceedings, 27th European
Photovoltaic Solar Energy Conference, Frankfurt, Germany, 24-28 September 2012.
H. Cui, P. Campbell and M.A. Green, “Optimisation of back surface reflectors for polycrystalline
silicon thin film solar cells application”, PV Asia Pacific Conference, Singapore, 23-25 October 2012
(oral).
H. Cui, C. Xue, W. Li, O. Kunz, P. Campbell and M.A. Green, “A novel glass thinning and texturing
method for Si thin-film solar cells application”, 22nd International Photovoltaic Science and
Engineering Conference, Hangzhou, China, 5-9 November 2012 (oral).
H. Cui, H., K. Omaki, O. Kunz, P. Campbell and M.A. Green, “Compatibility of glass texture with ebeam evaporated silicon thin-film solar cells”, Proc. 22nd International Photovoltaic Science and
Engineering Conference and Exhibition, Hangzhou, China, 5-9 November 2012. (oral)
J. Cui, A. To, Z. Li, J. Rodriguez and A. Lennon, “Inkjet patterned anonic aluminum oxide for rear
metal contacts”, 3rd International Conference on Silicon Photovoltaics (Silicon PV 2013), Hamelin,
Germany, 25-27 March 2013. (accepted)
J. Cui, J. Colwell, Z. Li and A. Lennon, “Local back surface field formation via different patterning
approaches”, Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7 December
2012. (oral)
J. Dore, R. Evans, U. Schubert, B. D. Eggleston, D. Ong, K. Kim, J. Huang, O. Kunz, M. Keevers, R.
Egan, S. Varlamov, M. A. Green, “Thin-film polycrystalline silicon solar cells formed by diode laser
crystallisation”, (oral), Frankfurt, Germany, 24-28 September 2012.
J. Dore, S. Varlamov, R. Evans, B. Eggleston, D. Ong, O. Kunz, J. Huang, U. Schubert, K. H. Kim,
R. Egan, M. Green, “Performance potential of low-defect density thin-film Si solar cells obtained by
electron beam evaporation and diode laser crystallisation”, PVTC-2012, 6-8 June, Aix-en-Provence,
France. (oral)
J. Dore, R. Evans, B. Eggleston, S. Varlamov, M. Green, “Intermediate layers for thin-film pc-Si solar
cells formed by diode laser crystallisation”, MRS Spring Meeting, San-Francisco, USA, 9-13 April
2012. (oral)
B. Eggleston, S. Varlamov, J. Huang, R. Evans, J. Dore, M. Green, “Large grained, low defect density
polycrystalline silicon on glass substrates by large-area diode laser crystallisation”, MRS Spring
Meeting, San-Francisco, USA, 9-13 April 2012. (oral)
Y. Feng, P. Aliberti, R. Patterson, B. Puthen-Veettil, S. Shrestha, M.A. Green and G. Conibeer,
“Analysis on carrier energy relaxation in superlattices and its implications on the design of hot carrier
solar cell absorbers”, Proceedings, SPIE Optics and Photonics, San Diego, 12-16 August 2012, p.
847103.
- 252 -
M.A. Green, “Advanced photovoltaic conversion: beyond the Shockley-Queisser limit”, Innovative
PV Symposium, University of Tokyo, 5-6 March 2012.
M.A. Green, “Advanced PV concepts”, Tutorial Lectures, Research Center for Advanced Science and
technology (RCAST), University of Tokyo, 7 March 2012.
M.A. Green, “Silicon photovoltaics: where is it heading?”, 1st China Photovoltaic Technology
International Conference (CPTIC) 2012, Shanghai, 19-21 March 2012.
M.A. Green, “Solar cells: power source for the future?”, 2012 ISEEE Distinguished Speaker Series,
University of Calgary, Canada, 5 April 2012. (invited)
M.A. Green, “Towards 20% efficient, 50 cents/watt silicon solar cells”, MRS Spring 2012
Symposium, San Francisco, 9-12 April 2012.
M.A. Green, “Future trends, opportunities and challenges in solar photovoltaics”, Solar Future Forum,
Australian Embassy Berlin, 5 June 2012. (invited)
M.A. Green, “Advanced silicon solutions: from high efficiency crystalline solar cells to thin films”,
Photovoltaic Technical Conference 2012, Aix en Provence, S. France, 6-8 June 2012 (invited).
M.A. Green, “Inorganic nanophotovoltaics”, Nobel Symposium on Nanoscale Energy Converters,
Örenäs Castle, Sweden, 12-16 August 2012. (invited)
M.A. Green, “Future of Si wafer technology”, 7th International Conference on Advanced Materials
(ROCAM 2012), Brasov, Romania, 28-31 August 2012. (invited)
M.A. Green, “Harnessing plasmonics for photovoltaics”, MNE Conference, Toulouse, France, 16-20
September 2012. (invited)
M.A. Green and S. Pillai, “Harnessing plasmonics for photovoltaics”, 1AP.1.3, 27th European
Photovoltaic Solar Energy Conference, Frankfurt, Germany, 24-28 September 2012 (Opening
Plenary).
M.A. Green, “Recent development in photovoltaics”, ANSTO seminar, Lucas Heights, Sydney, 9
October 2012 (Distinguished Lecture Series).
M.A. Green, “Cost reduction in PV technology”, PV Asia Pacific 2012, Singapore, 23-25 October
2012. (Opening Plenary)
M.A. Green, “Silicon solar cells: recent progress and future directions”, 22nd International
Photovoltaic Science and Engineering Conference (PVSEC-22), Hangzhou, China, 5-9 November
2012. (Opening Plenary)
M.A. Green, “Nanomaterials for photovoltaics”, Advances in Functional Nanomaterials – Energy and
Environment Conference, University of New South Wales, Sydney, 15-16 November 201. (Opening
Plenary)
X. J. Hao, N. Song, S. Huang and M.A. Green, “Promise of epitaxial growth of CZTS for tandem
cells”, 3rd European Kesterite Workshop, Luxemburg, 2012. (poster)
X. J. Hao, N. Song, X. Liu, S. Huang and M.A. Green, “Fabrication and properties of CZTS thin film
epitaxially grown on silicon substrate”, Proceedings, 22nd International Photovoltaic Science and
Engineering Conference, Hangzhou, China, November 2012. (oral)
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J. Huang, J. Dore, B. Eggleston, R. Evans, M. Keevers, R. Egan, S. Varlamov, M. A. Green, “Lowdefect polycrystalline silicon solar cells on glass formed by diode laser crystallisation”, (oral)
(PVSEC-22), Hangzhou, China, 5-9 November 2012.
S. Huang, L. Treiber, P. Zhang and G. Conibeer, “Highly periodic nanoparticle arrays for Hot Carrier
solar cell applications”, Proceedings, 22nd International Photovoltaic Science and Engineering
Conference (PVSEC-22), Hangzhou, China, 5-9 November 2012.
H. Kampwerth, Y. Yang and M.A. Green, “Measurement of internal optical reflection characteristics
of solar cell back reflectors”, 38th IEEE Photovoltaic Specialists Conference (PVSC), Austin, TX, 3-8
June 2012.
H. Kampwerth, Y. Yang and M. A. Green, “Experimental setup to measure the optical scattering
properties of solar cell back reflectors”, 27th European Photovoltaic Solar Energy Conference,
Frankfurt, Germany, 24-28 September 2012, pp. 436-438.
D. König, Y. Yao, R. Patterson, B. Puthen-Veettil and G. Conibeer, “Hot carrier solar cells:
principles, structures and materials”, 3rd International Symposium on Frontier Photonics, Saitama
University, Tokyo, 6-7 March 2012. (invited)
D. König, “The intrinsic nanocrystal junction – replacing doping by interface energetics”, IHT,
RWTH Aachen University, 13 June 2012. (invited)
D. Lazos and A.G. Bruce, “The value of commercial BIPV systems in the urban environment”,
Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7 December, 2012.
S. Lee, S. Huang, G. Conibeer and M.A. Green, “Application of Ge quantum wells fabricated by laser
annealing as energy selective contacts for hot-carrier solar cells”, 38th IEEE Photovoltaic Specialists
Conference (PVSC), Austin, TX, 3-8 June 2012, pp. 801-805.
M. Lenio, N. Borojevic, J. Durrant and S. Wenham, “Towards the production of a fully inkjet
patterned PERC cell”, 27th European Photovoltaic Solar Energy Conference, Frankfurt, Germany, 2428 September 2012.
H. Li, S. Wenham, W. Zhenjiao and H. Peiyu, “Thermally stable A-Si:H/SiNx stack passivating
system and the application on rear-localized contact solar cells on CZ p-type crystalline silicon”, 38th
IEEE Photovoltaic Specialists Conference (PVSC), Austin, TX, 3-8 June 2012, pp. 1069-1072.
H. Li and S.R. Wenham, “Industrially feasible rear passivation and contact scheme for high efficiency
crystalline silicon solar cells on p-type CZ substrate”, 27th European Photovoltaic Solar Energy
Conference, Frankfurt, Germany, 24-28 September 2012, pp. 883-887.
Y. Li, Z. Lu and A. Lennon, “Optical modeling of anodic aluminium oxide for light-trapping in
silicon solar cells”, Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7
December, 2012.
Y. Liao, S. Shrestha, S. Huang and G. Conibeer, “Fabrication of γ-Al2O3/Si/γ-Al2O3 quantum well
structures for energy selective contact of hot carrier solar cell”, Australian Solar Energy Society
Conference (Solar2012), Melbourne, 6-7 December 2012.
D. Lin, K. Wang, X. An, L. Mai, E. Mitchell and S. Wenham, “Further development of the advanced
semiconductor finger solar cell”, 27th European Photovoltaic Solar Energy Conference, Frankfurt,
Germany, 24-28 September 2012.
- 254 -
X. Liu, X. Hao, M. Green, S. Huang and G. Conibeer, “RF magnetron sputtered molybdenum back
contact for Cu2ZnSnS4 thin film solar cells”, Conference on Optoelectronic and Microelectronic
Materials and Devices (COMMAD 2012), Melbourne, 12-14 December 2012, pp. 61-62.
Z. Liu, X. Hao and M.A. Green, “Influence of substrate temperature on epitaxial growth of Ge on Si
by sputtering”, 27th European Photovoltaic Solar Energy Conference, Frankfurt, Germany, 24-28
September 2012.
Z. Liu, X. Hao and M.A. Green, “Investigation of thin epitaxial Ge films on Si substrates”, 22nd
International Photovoltaic Science and Engineering Conference (PVSEC-22), Hangzhou, China, 5-9
November 2012.
S. Pillai, J.J. Hohn, C.M. Johnson and G. Conibeer, “Effect of surface plasmon resonance on the
photoluminescence from Si quantum dot structures for third generation solar cell applications”, MRS
Fall Meeting, Boston, USA, 28 November - 2 December, 2011; MRS Online Proceedings Library, p.
1391, 2012. (refereed)
S. Pillai, “Plasmonic photovoltaics – the design evolution”, LIMA project Workshop, Valencia,
Spain, 6-7 November, 2012.
F. Qi, R. Chen and A. Uddin, “Study of anodic aluminium oxide (AAO) mold for the ordered organic
PV device morphology control”, 27th European Photovoltaic Solar Energy Conference and
Exhibition, Frankfurt, Germany, 24-28 September 2012, pp. 2871-2874.
K. J. Schmieder, A. Gerger, M. Diaz, Z. Pulwin, C. Ebert, A. Lochtefeld, R. Opila and A. Barnett,
“Analysis of tandem III-V/SiGe devices grown on Si”, 38th IEEE Photovoltaic Specialists
Conference, Austin, 3-8 June 2012, pp. 968-973.
K. Schmieder, M. Diaz, R. Opila, A. Gerger, A. Lochtefeld, Z. Pulwin, C. Ebert and A. Barnett,
“Progression of III-V/SiGe tandem solar cell growth and fabrication on Si substrates”, 27th European
Photovoltaic Solar Energy Conference and Exhibition, Frankfurt, Germany, 24-28 September 2012.
C. Shen, H. Kampwerth, M.A. Green, “Spatially resolved photoluminescence imaging of essential
silicon solar cell parameters”, 38th IEEE Photovoltaic Specialists Conference (PVSC), Austin, TX, 3-8
June 2012, pp. 1855-1859.
C. Shen, H. Kampwerth, M.A. Green, T. Trupke, J. Carstensen and A. Schütt, “Luminescence based
efficiency and other important parameters imaging of silicon solar cells”, 27th European Photovoltaic
Solar Energy Conference, Frankfurt, Germany, 24-28 September 2012.
S.K. Shrestha, N. Gupta, P. Aliberti and G. Conibeer, “Growth and characterization of germanium
carbide films for Hot Carrier solar cell absorber”, 38th Photovoltaic Specialists Conference, Austin,
TX, 3-8 June 2012, pp. 002601:002604.
S.K. Shrestha and J. Liu, “Structural properties of germanium nanocrystals deposited with RF
sputtering”, Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7 December
2012.
N. Song, X. Hao, J. Huang, X. Liu, S. Huang and M. Green “Film thickness and substrate temperature
effects on rf mag-netron sputtered Al:ZnO window layer for Cu2ZnSnS4 thin film solar cells”,
Conference on Optoelectronic and Microelectronic Materials and Devices, Melbourne, 12-14
December 2012, pp. 153-154.
- 255 -
A.B. Sproul, J.I. Bilbao and S.M. Bambrook, “A novel thermal circuit analysis of solar thermal
collectors”, Australian Solar Energy Society Conference (Solar2012), Melbourne, 6-7 December
2012.
L.S. To, A. Zahnd and R. Riek, “Building capacity for solar photovoltaics in rural Nepal, in: World
Renewable Energy Forum”, World Renewable Energy Forum, Colorado, USA, 13-17 May 2012.
T. Trupke, K.R. Macintosh, J.W. Weber, W. McMillan, L. Ryves, J. Haunschild, C. Shen and H.
Kampwerth, “Inline photoluminescence imaging for industrial wafer- and cell manufacturing”, 22nd
Workshop on Crystalline Silicon Solar Cells & Modules: Materials and Processes, Colorado, USA,
22-25 July 2012.
A. Uddin and C.C. Teo, “Fabrication of high efficient organic/CdSe quantum dots hybrid OLEDs by
spin-coating method”, SPIE International Symposium on SPIE OPTO, San Francisco, California,
USA, 2-7 February 2013.
S. Varlamov, B. Eggleston, J. Dore, R. Evans, D. Ong, O. Kunz, J. Huang, U. Schubert, K. H. Kim,
R. Egan, M. Green, “Diode laser processed crystalline silicon thin-film solar cells”, (invited-oral)
SPIE Photonics West, San Francisco, USA, 2-7 February 2013.
K.S. Wang, D. Lin, X.R. An, L. Mai, E. Mitchell and S.R. Wenham, “18.8% efficient laser doped
semiconductor fingers screen printed silicon solar cell with light induced plating”, 38th IEEE
Photovoltaic Specialists Conference (PVSC), Austin, TX, 3-8 June 2012.
Y. Wang, R. Opila, A Gerger, A. Lochtefeld and A. Barnett, “High performance light trapping for Ge:
Si solar cell in a multi junction solar cell”, 27th European Photovoltaic Solar Energy Conference,
Frankfurt, Germany, 24-28 September 2012.
Y. Wang, A. Gerger, A. Lochtefeld, R. Opila and A. Barnett, “Design and demonstration of light
trapping for Ge:Si solar cell below Si solar cell in a multi-junction solar cell system”, 38th IEEE
Photovoltaic Specialists Conference, 2012, pp 3328-3332.
M.E. Watt, A.G. Bruce, I.F. MacGill, S. Lewis and M. Hancock, “Understanding high PV penetration
effects”, Clean Energy Week 2012 - ATRAA, Sydney, 25-27 July 2012.
M. Wright, X. Yang, F. Caragay and A. Uddin, “Effect of solution concentration ratio on
PCPDTBT:PC70BM bulk heterojunction organic solar cells”, 27th European Photovoltaic Solar
Energy Conference, Frankfurt, Germany, 24-28 September 2012.
M. Wright, X. Yang, F. Caragay and A. Uddin, “Effect of solution concentration ratio on
PCPDTBT:PC70BM bulk heterojunction organic solar cells”, 27th European Photovoltaic Solar
Energy Conference and Exhibition, Frankfurt, Germany, 24-28 September 2012, pp. 2893-2896.
X. Yang, A. Uddin and M. Wright, “Plasmon enhanced light absorption in PCPDTBT:PC70BM bulkheterojunction organic solar cells”, 27th European Photovoltaic Solar Energy Conference and
Exhibition, Frankfurt, Germany, 24-28 September 2012, pp. 2871-2874.
Y. Yang, H. Mehrvarz, S. Pillai and M.A. Green, “The effect of rear surface passivation layer
thickness on high efficiency solar cells with planar and scattering metal reflectors”, Proceedings, 38th
IEEE Photovoltaic Specialists Conference (PVSC), Austin, TX, 3-8 June 2012, pp. 1172-1176.
P. Zhang, S. Huang and G, Conibeer, “Fabrication of two dimensional close-packed arrays of
colloidal silicon nanoparticles via Langmuir-Blodgett deposition”, Conference on Optoelectronic and
Microelectronic Materials and Devices, Melbourne, 12-14 December 2012.
- 256 -
6.6
OTHER REPORT
G. Francisco, W. Harcourt, B.C. Farhar, B. Osnes, P. Karve, L.S. To, N. Speer and C. Sweetman,
Report in “Views, events, and debates”, Journal of Gender & Development, Vol. 20, pp. 607–621 (a
report of a workshop on “Engaging Women in Clean Energy Solutions”, World Renewable Energy
Forum, Denver, 13 May 2012.
6.7
MEDIA
ChemViews/M.A. Green, “Solar Cell Efficiency Tables”, ChemistryViews, Progress in Photovoltaics,
6 February 2012.
http://www.chemistryviews.org/details/ezine/1444365/Solar_Cell_Efficiency_Tables.html
- 257 -

UNSW Photovoltaics Annual Report 2012

Transcription

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