Development of Ceramic – Carbon Nanotube (CNT) Nanocomposites

Transcription

Development of
Ceramic – Carbon Nanotube (CNT)
Nanocomposites
A THESIS SUBMITTED TO THE UNIVERSITY OF LONDON
FOR THE DEGREE OF DOCTOR OF PHILOSOPHY
by
Fawad Inam
School of Engineering and Materials Science
Queen Mary, University of London
London E1 4NS, United Kingdom
May 2009
Chapter 8. Mechanical properties of ceramic – CNT nanocomposites
Declaration
I certify that the present work is prepared solely by myself during the course of my
studies at the Queen Mary, University of London. It has not been submitted for a
degree at this or any other University. Any words and/ or figures from the work of
other people are fully acknowledged according to standard referencing.
This thesis fully complies with the regulations set by the University of London and
the Queen Mary, University of London.
Fawad Inam
May 2009
1
Abstract
The increasing availability of nanopowders and nanotubes combined with new
processing techniques is enabling the development of new multifunctional materials.
Carbon Nanotubes (CNTs) are one of the recently discovered allotropic forms of
carbon. They have exceptional mechanical, electrical and thermal properties. The
application of CNTs in the reinforcement of ceramic nanocomposites has not yet been
fully investigated and is the subject of this study. Alumina is the main matrix used in
this study. CNTs need to be de-agglomerated and homogeneously distributed in
ceramic nanocomposites. Dimethylformamide (DMF) produces fine and stable CNT
and alumina dispersions. All nanocomposites were sintered by Spark Plasma
Sintering (SPS). Nanocomposites prepared using DMF dispersions showed better
dispersions, higher electrical conductivity and mechanical properties as compared to
those prepared using ethanol dispersions. The addition of CNTs or Carbon Black
(CB) to alumina significantly aids its densification. The CNTs produce significant
grain growth retardation. CNTs were found to be well preserved in alumina after
being SPSed up to 1900
o
C. Structural preservation of CNTs in ceramic
nanocomposites depends on the nature of ceramic and SPS processing conditions. The
electrical conductivity of alumina – CNT nanocomposites is four times higher as
compared to alumina – CB nanocomposites due to the fibrous nature and high aspect
ratio of CNTs. Alumina coated CNTs were used for better interfacial adhesion with
the matrix. Oxidative resistance of CNTs was increased by coating them with alumina
and by decreasing the grain boundary area in alumina – CNT nanocomposites. Coated
and uncoated CNTs showed higher mechanical reinforcement in alumina
nanocomposite as compared to CB. The future for ceramic – CNT nanocomposites is
very bright, especially for applications associated with the electrical and thermal
properties. Apart from a good understanding of nanocomposites, the commercial
development of CNT based technologies heavily relies on the availability and price of
CNTs.
2
Table of Contents
Acknowledgement
8
List of Figures and Tables
9
List of Technical Abbreviations
21
Preface
22
Chapter 1.
Chapter 2.
Introduction to Carbon Nanotubes (CNTs)
24
1.1. Introduction
24
1.2. Structure of CNTs
25
1.3. History of CNTs
30
1.4. Synthesis of CNTs
31
1.5. Biocompatibility and toxicology of CNTs
33
1.6. Market overview of CNTs
34
1.7. Summary
36
Introduction to Ceramic – CNT Nanocomposites
37
2.1. Introduction
37
2.2. Fabrication of ceramic – CNT nanocomposites
37
2.2.1. Pre-consolidation processing
37
2.2.2. Consolidation techniques
38
2.3. Mechanical properties of ceramic – CNT nanocomposites
40
2.3.1. The never ending controversy
41
2.3.2. Toughening mechanisms
42
2.3.3. Effect of CNT in alumina nanocomposites
43
2.3.4. Effect of CNT in other ceramic nanocomposites
48
2.4. Electrical properties of ceramic – CNT nanocomposites
2.4.1. Percolation threshold in ceramic – CNT nanocomposites
54
55
3
Chapter 8. Table of Contents
Chapter 2 continued…
Chapter 3.
2.4.2. Effect of CNT on the electrical conductivity
56
2.5. Thermal properties of ceramic – CNT nanocomposites
61
2.6. Miscellaneous effects of CNTs on ceramics
64
2.7. Summary
66
Materials and Experimental Techniques
67
3.1. Introduction
67
3.2. Materials
67
3.2.1. Carbon nanotubes
67
3.2.2. Carbon black
69
3.2.3. Alumina
70
3.2.4. Other ceramic matrices
72
3.3. Experimental techniques
Chapter 4.
74
3.3.1. Nano-particle size analyzer
74
3.3.2. Spark Plasma Sintering (SPS)
75
3.3.3. Density measurement
79
3.3.4. Electrical conductivity measurements
80
3.3.5. Vickers indentation
82
3.3.6. Microscopic analysis
83
3.3.7. Thermo Gravimetric Analysis (TGA)
84
Homogenisation of CNTs in Ceramics
85
4.1. Introduction
85
87
4.2.1. Colloidal dispersions and characterisations
87
4.2.2. Nanocomposite powder preparation
88
89
4.2.4. Nanocomposite characterisations
90
4.3. Results and discussion
90
4.3.1. Natural drying patterns
90
4.3.2. Agglomerate size analysis
92
4
Chapter 5.
4.3.3. Re-agglomeration behaviour
93
4.3.4. Microstructure of nanocomposites
96
4.3.5. Mechanical and electrical properties
97
4.4. Conclusions
99
Sintering of Ceramic – CNT Nanocomposites
101
5.1. Introduction
101
103
5.2.1. CNTs purification
103
104
104
105
Chapter 6.
105
5.3.1. SPS of alumina and alumina – CNT nanocomposites
105
5.3.2. Sintering behaviour and possible mechanisms
107
5.3.3. Grain growth modification
112
5.3.4. Co-sintering of grain size laminate
115
5.3.5. Sintering behaviour of alumina-purified CNTs
118
5.4. Conclusions
120
Structural Stability of CNTs in Ceramics
121
(Preservation Studies)
6.1. Introduction
121
125
6.2.1. Starting materials
125
125
6.2.3. Material characterisations
127
6.2.4. Raman Spectroscopy
127
6.2.5. X-Ray Diffraction (XRD) analysis
127
128
128
5
Chapter 7.
Chapter 8.
Chapter 9.
6.3.2. SPS of other ceramics and their CNT nanocomposites
133
6.3.3. SPS of bulk CNTs
135
6.4. Conclusions
139
Electrical Properties of Ceramic – CNT Nanocomposites
140
7.1. Introduction
140
140
142
7.3.1. CNT vs. carbon black
142
7.3.2. Electrical conductivity as a function of grain size
144
7.3.3. Electrical conductivity as a function of temperature
146
7.4. Conclusions
147
Oxidative Stability of Ceramic – CNT nanocomposites
149
8.1. Introduction
150
150
150
8.3.1. Oxidation of CNTs
150
8.3.2. Oxidation of alumina coated CNTs
152
8.3.3. Oxidation of alumina – CNT nanocomposites
154
8.4. Conclusions
155
Mechanical Properties of Ceramic – CNT Nanocomposites
156
9.1. Introduction
156
157
9.2.1. Flexural bending
158
9.2.2. Thermal shock resistance during SPS (observation)
159
9.3.1. Surface finish
160
160
6
9.3.2. Vickers hardness
162
9.3.3. Indentation Fracture Toughness (IFT)
164
9.3.4. Flexural modulus
167
9.3.5. Flexural strength
168
9.3.6. Thermal shock resistance (qualitative assessment)
171
9.4. Summary of mechanical properties
173
9.5. Conclusions
174
Conclusions
175
Future Work
177
Appendix A.
Appendix B.
Properties of DMF
179
A.1. Introduction
179
A.2. Chemistry of DMF
179
Future of Ceramic – CNT Nanocomposites
181
B.1. Introduction
181
B.2. Applications of ceramic – CNT nanocomposites
182
B.2.1. Conductive ceramic – CNT nanocomposites
182
B.2.2. Electric Discharge Machining (EDM)
183
B.2.3. Special purpose electrodes
183
B.2.4. Heating elements
184
B.2.5. Porous structures
185
B.2.6. Ceramic armour
185
B.2.7. Functionally Gradient Materials (FGMs)
186
B.2.8. Nanoceramics
186
B.3. Challenges in the development of ceramic – CNT nanocomposites
186
B.3.1. CNT related issues
187
B.3.2. Understanding of nanocomposites
188
B.4. Conclusions
188
Appendix C.
Weibull statistical analysis
189
Appendix D.
Recent publications, based on carbon nanotubes
197
References
198
7
Chapter 10. Future of Ceramic – CNT nanocomposites
Acknowledgement
It is never superfluous to say thanks where acknowledgement is due. I would like to
express my sincere gratitude and appreciation to my supervisors Prof. Dr. Ton Peijs
and Dr. Michael J. Reece for their continual and valuable guidance and
encouragement. I am also very thankful to my third supervisor cum friend, Dr. Haixue
Yan, for supporting me in processing my materials. Their ideas and infectious
enthusiasm always energized me to contribute greatly in the world of nanocomposite
materials.
Grateful acknowledgement is made to Queen Mary, University of London (QMUL)
and Nanoforce Technology Limited for financially supporting my PhD studentship. I
am highly obliged to Dr. Zofia Luklinska, Dr. Monisha Phillips, Dr. Rory Wilson,
Nima Roohpur and Dr. Daniel Doni (Imperial College, UK) and for their technical
support in the characterisation part. Special thanks to Prof. Dr. Alan Weimer
(University of Colorado, USA) for supplying coated Carbon Nanotubes (CNTs). I am
indebted to my research mates, Emilano, Chris, Deng, Jamie, Rui, Peng Peng,
Manuela, Nancy, Paola, Peppe, Pavin, Edwin, Tilen, Lilly, Jean, Jianmin, Aqif Bhai,
Boon and Eric for helping me in every possible manner. I would also like to record
my gratitude to Bill Godwin, Mick Willis, Sandra Wells, Victoria Wells, Jonathan
Hills, Daniella Samos and Raj Chadha and Tara Demetriou for their time and humble
support at various stages of my work.
Words cannot express my deepest gratitude for my parents for their enduring love and
support, which means so much to me. On personal level, I wish to thank my wife,
Nageen, for her love, moral support, encouragement and tireless proof reading of this
work. Surely, this thesis is dedicated to my family.
Last but certainly not the least; I would like to acknowledge all my teachers who
showed me the wisdom. No doubt, I am at this stage because of all of them!
8
Chapter 10. Future of Ceramic – CNT nanocomposites
Chapter 1.
Figure 1.1.
Publications based on CNTs over the years. Source: Web of Science.
Figure 1.2.
Different types of CNTs based on their number of graphene cylinders:
(a) capped Single Wall CNT (SWNT); and (b) open Multi Wall CNT
(MWNT).
Figure 1.3.
Different forms of carbon fibre.
Figure 1.4.
Different types of CNTs based on their chirality: a) armchair; b)
zigzag; and c) chiral.
Figure 1.5.
Schematic diagram showing how a hexagonal sheet of graphite is
“rolled” to form a carbon nanotube [23].
Figure 1.6.
Historical forms of hollow carbon filaments. Reported in: (a) 1952
[47]; and (b) 1976. Modified from [49].
Figure 1.7.
Schematic diagrams of arc-discharge method [58].
Figure 1.8.
Vertical CVD furnace for mass production. Re-drawn from [65].
Figure 1.9.
Schematic diagram of laser ablation method [58].
Table 1.1
Properties of different engineering fibres [11-22].
Table 1.2
Different ways of producing CNTs.
Table 1.3
Some of the major CNT suppliers.
Chapter 2.
Figure 2.1.
Colloidal processing route for making ceramic – CNT nanocomposites.
9
Figure 2.2.
Common
sintering
techniques
for
making
ceramic
–
CNT
nanocomposites: (a) hot pressing; and (b) Spark Plasma Sintering
(SPS).
Figure 2.3.
Fracture toughness of alumina – carbon nanocomposites. (a) By
Vickers indentation method [140]; and (b) by SEVNB method [127].
Both papers [127, 140] were published in the same journal.
Figure 2.4.
Different toughening mechanisms in amorphous alumina – CNT (outer
diameter: ~51-56 nm) nanocomposite: a) crack deflection; b) crack
bridging; c) CNT pullout; and d) CNT collapse in shear band. Figure
modified from [192].
Figure 2.5.
Schematic of hybrid microstructure design of alumina reinforced by
MWNTs and SiC nanoparticles. Modified from [120].
Figure 2.6.
CNTs at the grain boundary: (a) schematics; and (b) high-resolution
TEM micrographs of an alumina – 10 vol% SWNT nanocomposite
showing SWNT bundles at alumina grain boundaries, with schematic
diagrams indicating corresponding SWNTs orientations. Figure
modified from [195].
Figure 2.7.
Electrical conductivity (measured at room temperature) of ceramic –
CNT nanocomposites as a function of CNT content. Percolation
threshold is: (a) 0.64 vol% for MgAl2O4. Modified from [146]; and (b)
0.79 vol% for Al2O3 [96].
Figure 2.8.
The electrical conductivity of various representative materials at room
temperature. Note the more than 13 orders of magnitude increase in
conductivity of the alumina – 15 vol% SWNT nanocomposite
compared to monolithic alumina [138].
Figure 2.9.
Thermal conductivity of different ceramic – CNT nanocomposites as a
function of temperature. (a) Barium titanate. Modified from [222]; and
(b) titanium nitride. Modified from [226].
Table 2.1.
Summary of the fracture toughness of ceramic – CNT nanocomposites.
10
Table 2.2.
Summary of the electrical conductivity of ceramic – CNT
nanocomposites.
Table 2.3.
Summary of the room temperature thermal conductivity of ceramic –
CNT nanocomposites.
Chapter 3.
Figure 3.1.
Different types of MWNTs used: (a,b) NC 7000; (c) GraphiStrength
C100; and (d,e†) alumina coated (50 ALD cycles) GraphiStrength
C100.
Figure 3.2.
Carbon black powders: (a) Vulcan XC72; and (b) Printex L6.
Figure 3.3.
Crystal structure of α – alumina. Figure modified from [256].
Figure 3.4.
Phase transitions in alumina. Redrawn from [255].
Figure 3.5.
Alumina powder used in this study.
Figure 3.6.
Reduced titanium dioxide: (a) as-received form; and (b) after ball
milling.
Figure 3.7.
Boron carbide: (a) as-received form; and (b) after ball milling.
Figure 3.8.
Boron nitride: (a) as received form; and (b) nano-sized flakes.
Figure 3.9.
Working principle of Zeta particle size analyzer. Modified from [260].
Figure 3.10.
Typical DC-pulsed current cycles (used in this study).
Figure 3.11.
Effect of DC pulse on the density of alumina [267].
Figure 3.12.
SPS facility by FCT Systeme, Germany: (a) SPS facility at Queen
Mary, University of London, UK; and (b) SPS at 1800 oC [277].
Figure 3.13.
Cross-sectional view of carbon die set.
Figure 3.14.
Electrical conductivity measurement setup: (a) room temperature jig;
and (b) high-temperature characterisation chamber in the tubular
furnace.
Figure 3.15.
A typical Vickers indent. Modified from [281].
Table 3.1.
CNTs (synthesised by CVD method) used in this study.
11
Table 3.2.
Carbon black powders used in this study.
Table 3.3.
Properties of reduced titanium dioxide as per supplier.
Table 3.4.
Properties of boron carbide as per supplier.
Table 3.5.
Properties of boron nitride as per supplier.
Chapter 4.
Figure 4.1.
Ultrasonication bath: (a) Apparatus with dispersion bottle; (b) shock
waves in distilled water, top view of the apparatus; and (c) schematic
diagram.
Figure 4.2.
Pre-sintering processing of ceramic – CNT nanocomposite powder.
Figure 4.3.
Natural drying patters formed during processing of nanocomposite
powder: (a) alumina – 2 wt% CNT; (b) alumina – alumina coated 2
wt% CNT. CNTs were supplied by NanoDynamics, USA; (c) alumina
– 2 wt% carbon black (Vulcan XC72); and (d) alumina – 2 wt% carbon
black (Printex L6).
Figure 4.4.
Alumina – 2 wt% CNT nanocomposite powder after sieving.
Individual CNT can be seen. CNTs were supplied by Nanocyl,
Belgium.
Figure 4.5.
Agglomerate size analysis with respect to ultrasonication time in
different solvents: (a) CNTs, supplied by Nanocyl, Belgium; and (b)
alumina.
Figure 4.6.
Colloidal dispersion stability comparison after 1 h bath sonication and
5 minutes hand shaking. The diameter of the bottles is 25 mm. CNTs
were supplied by Nanocyl, Belgium.
Figure 4.7.
Re-agglomeration behaviour in different solvents after 30 minutes of
ultrasonication. CNTs were supplied by NanoDynamics, USA: (a,b)
pristine CNT; (c) alumina coated CNT (25 ALD cycles); and (d)
alumina coated CNTs (50 ALD cycles).
12
Figure 4.8.
Figure 4.8. Fractured surfaces of sintered alumina – 2 wt% CNT
samples, SPSed at 1200 ◦C/100 MPa/ 3 minutes: (a) CNTs dispersed in
ethanol and (b) CNTs dispersed in DMF.
Figure 4.9.
Vickers hardness and indentation fracture toughness of alumina – 2
wt% CNT nanocomposites prepared using different solvents.
Figure 4.10.
Density and electrical conductivity of alumina – 2 wt% CNT.
Table 4.1.
Samples SPSed for this chapter#.
Chapter 5.
Figure 5.1.
Purification of CNTs by acid treatment.
Figure 5.2.
Moving punch speed as the function of temperature during SPS for
alumina and alumina – 5 wt% (~11.2 vol%) CNT (uncoated)
nanocomposite. CNTs were supplied by Nanocyl, Belgium. Inset
shows SPS program details.
Figure 5.3.
Moving punch speed as the function of temperature during SPS for
alumina – 5 wt% CNT (uncoated) and alumina – 5 wt% CNT (coated,
50 ALD cycles) nanocomposite. Heating rate: 300 oC/minute. CNTs
were supplied by Arkema, France.
Figure 5.4.
Rel. theoretical density as a function of sintering temperature for: (a)
alumina
and
nanocomposites;
and
(b)
reduced
titania
and
nanocomposites.
Figure 5.5.
Bulk density of the sintered product as the function of homogeneity of
CNTs in alumina matrix. SPS conditions: 100 MPa/ 3 minutes. CNT
content: 5 wt%.
Figure 5.6.
Schematics of graphite die sets: (a) alumina before SPS; (b) current
passing through graphite only during SPS; and (c) current passing
through graphite and alumina – CNT nanocomposite compact during
SPS.
13
Figure 5.7.
Compressibility and compactibility analysis performed at room
temperature by uni-axial pressing.
Figure 5.8.
Rel. theoretical density as a function of grain size for alumina and
nanocomposites.
Figure 5.9.
FE-SEM images of fractured surfaces of sintered nanocomposites
processed at 1800 oC/ 100 MPa/ 3 minutes: (a) alumina; (b) alumina –
2 wt% carbon black (Printex L6); (c) alumina – 2 wt% CNT (Nanocyl,
Belgium); and (d) alumina – 2 wt% alumina coated CNT (50 ALD
cycles, NanoDynamics, USA).
Figure 5.10.
Grain size refinement effect of: (a) CNTs; and (b) CNTs and CB.
Figure 5.11.
Grain size refinement in isothermal conditions (1400 oC/ 100 MPa).
Figure 5.12.
FE-SEM images of grain size laminate showing interface between
alumina and 2 wt% CNT nanocomposite: (a) high magnification
fractured surface; and (b) low magnification polished surface.
Figure 5.13.
FE-SEM images of polished grain size laminates: (a) alumina and 2
wt% CNT nanocomposite; and (b) alumina and 5 wt% CNT
nanocomposite showing cracking at interface.
Figure 5.14.
FE-SEM images of different areas of polished grain size laminates: (a)
alumina region, grain size: 43 μm; (b) 2 wt% CNT nanocomposite
region, grain size: 1.81 μm; and (c) 5 wt% CNT nanocomposite region,
grain size: 0.99 μm.
Figure 5.15.
Thermo gravimetric analysis (TGA) of CNTs, before and after acid
purification treatment. Heating rate: 5 oC/minute.
Figure 5.16.
Platinum pan (diameter: ~10 mm) used for TGA: (a) empty pan before
analysis, (b) as received-CNTs and (c) impurities left after oxidation.
Table 5.1.
Electrical conductivity measurements of graphite die sets.
Table 5.2.
Density
measurements
for
CNTs
and
alumina
–
ceramic
nanocomposites.
14
Chapter 6.
Figure 6.1.
Schematics of different Raman vibration in CNT. Figure modified
from [342, 343].
Figure 6.2.
Diamond particles formed during SPS of CNTs [344, 345].
Figure 6.3.
Raman spectra of CNTs, alumina and alumina – 5 wt% CNT
nanocomposites. Alumina and nanocomposites were SPSed at 100
MPa for 3 minutes.
Figure 6.4.
Intensity ratio (ID/IG) for CNT and alumina – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 100 MPa for 3
minutes. For 1600 oC, refer figure 6.5.
Figure 6.5.
Raman spectra of CNTs and alumina – 5 wt% CNT nanocomposites.
All nanocomposites were sintered at 1600 oC/ 100 MPa.
Figure 6.6.
Intensity ratio (ID/IG) for CNT and alumina – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 1600 oC/ 100
MPa.
Figure 6.7.
HR-TEM of alumina – 5 wt% CNT nanocomposites showing electron
diffraction patterns of different areas. SPSed at 1900 oC/ 100 MPa/ 3
minutes.
Figure 6.8.
HR-TEM of alumina – 5 wt% CNT nanocomposites showing
agglomerates of CNTs at the grain boundary.
Figure 6.9.
XRD analysis of CNTs, alumina (SPSed) and alumina – 5 wt% CNT
nanocomposite (SPSed). SPSed at 1900 oC/ 100 MPa/ 3 minutes.
Figure 6.10.
FE-SEM image of boron carbide – 5 wt% CNT nanocomposite. SPSed
at 2000 oC/ 80 MPa/ 20 minutes: a) at lower magnification; and b) at
higher magnification.
Figure 6.11.
Raman spectra of CNTs, boron carbide (SPSed) and boron carbide – 5
wt% CNT nanocomposites (SPSed). SPSed at 80 MPa for 20 minutes.
Figure 6.12.
Intensity ratio (ID/IG) for CNT and boron carbide – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 80 MPa for 20
minutes.
15
Figure 6.13.
FE-SEM analysis of boron nitride – 5 wt% CNTs (rel. theoretical
density: ~97.5%). SPSed at 2175 oC/ 80 MPa/ 20 minutes: (a) at lower
magnification; and (b) at higher magnification.
Figure 6.14.
FE-SEM images of bulk CNTs SPS processed at different
temperatures. a) Pressed at room temperature; b) SPS processed at
1000 oC; c) at 1500 oC; and c) at 2000 oC.
Figure 6.15.
HR-TEM images of bulk CNTs: a) as received; b) SPS processed at
1000 oC/ 100 MPa/ 3 minutes.
Figure 6.16.
HR-TEM of CNT SPS processed at 2000 oC/ 100 MPa/ 20 minutes:
inset a) nano-onion; and inset b) CNT after SPS, maintaining its aspect
ratio.
Figure 6.17.
Raman spectra of CNTs SPS processed at different conditions. All
samples were processed at 80 MPa for 20 minutes.
Table 6.1.
Samples SPS processed for this chapter.
Chapter 7.
Figure 7.1.
Electrical conductivities of alumina – carbon nanocomposites.
Figure 7.2.
HR-TEM image of alumina – 5 wt% CNT nanocomposite: (a) CNTs
around grains; and (b) percolating network highlighted.
Figure 7.3.
Voltage vs. current measured at room temperature for alumina – CNT
nanocomposites. Both nanocomposites were SPSed at 1800 oC/ 100
MPa/ 3 minutes.
Figure 7.4.
Electrical conductivities of alumina – carbon nanocomposites as the
function of grain size. The grain sizes were coarsened by using higher
sintering temperatures.
Figure 7.5.
Electrical conductivities of alumina – 5 wt% nanocomposites as the
function of SPS processing durations. CNTs were supplied by Nanocyl,
Belgium.
16
Figure 7.6.
Electrical conductivities of alumina – 5 wt% nanocomposites as the
function of temperature. The grain sizes were coarsened by using
longer processing durations. Heating rate: 2 oC/ minute.
Table 7.1.
Samples SPSed for this chapter#.
Chapter 8.
Figure 8.1.
TGA of raw CNTs from different suppliers. Heating rate: 5 oC/ minute.
Figure 8.2.
TGA of uncoated CNTs and alumina coated CNTs. CNTs were
supplied by NanoDynamics, USA. Heating rate: 5 oC/ minute.
Figure 8.3.
Platinum pan used for TGA: a) alumina coated (50 ALD cycles) CNTs
before oxidation; and b) alumina nanotubes left after oxidation of
coated CNTs. CNTs were supplied by NanoDynamics, USA.
Figure 8.4.
Alumina nanotube left after the oxidation of coated CNTs. CNTs were
supplied by NanoDynamics, USA: a) at lower magnification; and b) at
higher magnification.
Figure 8.5.
TGA of raw CNTs and SPSed alumina – 5 wt% uncoated CNT
nanocomposites.
CNTs were supplied
by
Nanocyl,
Belgium.
Heating rate: 5 oC/ minute.
Table 8.1.
Alumina – 5 wt% CNT nanocomposites SPSed for this chapter.
Chapter 9.
Figure 9.1.
Special jig for improved 3 – point flexural testing. Sample thickness:
1.5 mm.
Figure 9.2.
Chamfered edges to avoid stress concentration points on corners.
Sample: Fractured alumina surface.
17
Figure 9.3.
Polished
surfaces
of
alumina
–
CNT
(Nanocyl,
Belgium)
nanocomposites containing: (a) no CNTs, thermally etched at 1500 oC
for 10 minutes; (b) 2 wt% (~4.7 vol%) CNTs; (c) 5 wt% (~11.2 vol%)
CNTs; and (d) 10 wt% (~21 vol%) CNTs.
Figure 9.4.
Polished surface of alumina – 10 wt% CNT nanocomposite. Arrows
showing alumina grain ex-sites. CNTs were supplied by Nanocyl,
Belgium.
Figure 9.5.
Alumina coated CNT (50 ALD cycles) encapsulated in an alumina
grain of nanocomposite. CNTs were supplied by NanoDynamics, USA.
Figure 9.6.
Different types of fracture modes in alumina – 2 wt% CNT
nanocomposites; (a) intragranular fracture mode in coated CNT
nanocomposite (50 ALD cycles, NanoDynamics, USA); and (b)
intergranular fracture mode in CNT nanocomposite (Nanocyl,
Belgium).
Figure 9.7.
Vickers indent after applying 2.5 kg load in: (a) alumina; and (b)
alumina – 5 wt% CNT nanocomposite. CNTs were supplied by
NanoDynamics, USA.
Figure 9.8.
Sub-surface damage analysis after Vickers indentation. Optical
micrograph (dark field image) showing sub-surface cracking in: (a)
alumina; and (b) alumina – 5 wt% uncoated CNT (Nanocyl, Belgium)
nanocomposite.
Figure 9.9.
Sub-surface damage analysis after Vickers indentation. Cross-section
of: (a) alumina, showing major crack; and (b) alumina – 5 wt%
uncoated CNT (Nanocyl, Belgium) nanocomposite, showing no major
sub-surface damage.
Figure 9.10.
Low energy failure in 3 – point flexural testing: (a) Optical micrograph
(side view) of alumina – 5 wt% uncoated CNT nanocomposites; and
(b) schematics as per ASTM C1161-02c [368].
Figure 9.11.
Fractured surfaces of alumina – uncoated CNT nanocomposites. Two
tensile surfaces are mounted together. Fracture origin can be identified.
CNT concentration: (a) 2 wt%; and (b) 5 wt%.
18
Figure 9.12.
Effect of rapid cooling (300 oC/ minute) from 1800 oC after sintering.
Disk diameter: 20 mm; (a) alumina; (b) alumina – 2 wt% CB
nanocomposite; and (c) alumina – 2 wt% uncoated CNTs (Nanocyl,
Belgium) nanocomposite.
Table 9.1.
Vickers hardness of alumina and alumina – carbon nanocomposites.
Table 9.2.
Indentation fracture toughness of different materials.
Table 9.3.
Flexural modulus of different materials.
Table 9.4.
Flexural strength of different materials.
Table 9.5.
Comparison of thermal shock resistance for different materials.
Table 9.6.
Summary of mechanical properties of alumina and alumina – carbon
nanocomposites#.
Appendix A.
Figure A.
Chemical formula of DMF, C3H7ON.
Table A.
Properties of DMF [377, 381, 382].
Appendix B.
Figure B.1.
Different sports goods based on non-ceramic – CNT nanocomposites
[386-389].
Figure B.2.
Heating element based on alumina – CNT nanocomposite. Sample
diameter is 20 mm: (a) alumina – 5 wt% CNT (Nanocyl, Belgium); (b)
demonstration; and (c) after oxidation, white area shows oxidation of
CNTs.
Figure B.3.
CNT network on the grain boundaries of alumina grains. (a) Before
oxidation; and (b) After oxidation of CNTs, percolating porosity.
19
Appendix C.
Figure C1.
Weibull plot for data shown in table D1. Weibull modulus: 9.05.
Figure C2.
Figure C3.
Figure C4.
Figure C5.
Figure C6.
Table C.
Weibull modulus (fracture strength) of different materials.
Table C1.
Weibull statistical analysis (fracture strength) for alumina.
Table C2.
Weibull statistical analysis (fracture strength) for alumina – 5 wt% CB
nanocomposite. CNTs were supplied by NanoDynamics, USA.
Table C3.
Weibull statistical analysis (fracture strength) for alumina – 2 wt%
uncoated
CNT
nanocomposite.
CNTs
were
supplied
by
NanoDynamics, USA.
Table C4.
coated CNT nanocomposite. CNTs were supplied by NanoDynamics,
USA.
Table C5.
uncoated
CNT
nanocomposite.
CNTs
were
supplied
by
NanoDynamics, USA.
Table C6.
coated CNT nanocomposite. CNTs were supplied by NanoDynamics,
USA.
20
List of figures and tables
List of Technical Abbreviations
AFM
Atomic Force Microscopy
ALD
Atomic Layer Deposition
CB
Carbon Black
CCE
Carbon Ceramic Electrode
CNF
Carbon Nanofibre
CNT
Carbon Nanotube
CVD
Chemical Vapour Deposition
CoMoCAT
Co-Mo CATalyst
DMSO
DiMethyl SulfOxide
DMF
N,N-DiMethylFormamide
DWNT
Double Wall Carbon Nanotube
DDS
Drug Delivery System
DLS
Dynamic Light Scattering
EDM
Electric Discharge Machining
FE-SEM
Field Emission Scanning Electron Microscopy
FGM
Functionally Gradient Material
GPS
Gas Pressure Sintering
HR-TEM
High Resolution Transmission Electron Microscopy
HiPCO
HIP
High-Pressure CO
Hot Isostatic Pressing
IET
Impulse Excitation Technique
IFT
Indentation Fracture Toughness
JCPDS
Joint Committee on Powder Diffraction Standards
MWNT/ MWCN
Multi Wall Carbon Nanotube
NMP
N-MethylPyrrolidone
NAFB
Nano Agglomerate Fluidised Bed
PCS
Photon Correlation Spectroscopy
PVD
Physical Vapour Deposition
RBM
Radial Breathing Mode
SHS
Self-propagating High-temperature Synthesis
SWNT/ SWCN
Single Wall Carbon Nanotube
SEVNB
Single-Edge V-Notched Beam
SPS
Spark Plasma Sintering
TZP
Tetragonal Zirconia Polycrystals
TBC
Thermal Barrier Coating
TGA
Thermo Gravimetric Analysis
XRD
X-Ray Diffraction
21
List of figures and tables
Preface
The discovery of new material structures is shifting research interest from developing
traditional materials like metals, ceramics, polymers and composites, to the more
novel nanocomposites. “Nanocomposites” are not new. Nature has evolved ways to
make many nanocomposites, such as seashell [1], spider silk [2, 3], ivory [4] and bone
[5]. Bone shows ceramic-like properties and it is a natural hybrid nanocomposite of
plate-shaped hydroxyapatite mineral particles in an oriented collagen polymer matrix,
with an attractive balance of stiffness, toughness, and vibrational damping properties
[5]. In the context of scientific knowledge, the term “nanocomposite” can be found in
literature as early as 1986, when Roy et al. [6] prepared hybrid ceramic – metal
nanocomposite by the sol-gel method. Nanocomposite materials can be defined as
composites of more than one Gibbsian solid phase where at least one of the phases
shows dimensions in the nanometre range [6]. The solid phases can exist either in
amorphous, semi-crystalline or crystalline states [7].
CNT is one of the recently discovered allotropes of carbon. A comprehensive
introduction to CNTs is given in chapter 1. Since 1991 [8], CNT reinforced
nanocomposites have been the focus of intense global research. The main emphasis of
these global research efforts has been on the synthesis of CNTs, developments in
electronics and development of CNT reinforced polymer nanocomposites.
The application of CNTs in the reinforcement of ceramic nanocomposites has not yet
been fully investigated and is the subject of this thesis. A literature review on the
development of ceramic – CNT nanocomposites is presented in chapter 2. Alumina,
which is the most commonly used structural [9] and bio- ceramic [10], is used as a
model ceramic for this study.
22
Preface
The main topics of the current study are:
1. Homogenisation and de-agglomeration of CNTs in ceramic matrices (chapter 4).
2. Effect of CNTs on the sintering and the grain growth behaviour of different
ceramics (chapter 5).
3. Survivability of CNTs sintered at high temperatures and their suitability as a
ceramic reinforcement (chapter 6).
4. Electrical (chapter 7), chemical (chapter 8) and mechanical (chapter 9) properties of
ceramic – CNT nanocomposites.
23
Preface
Chapter 1.
Introduction to Carbon Nanotubes (CNTs)
1.1. Introduction
Carbon Nanotubes (CNTs) are the subject of intense global research. In 2008, CNTs
research produced more than 20 publications per day (figure 1.1). Among all other
nanotubes (BN, BC3, BC2N, C3N4, CN), CNTs appear to have the highest commercial
interest. Medical science, electronics and composite technology are the main sectors
that might benefit most from the properties of CNTs (table 1.1). CNTs have tensile
strengths significantly higher than steel and carbon fibre, electrical conductivity
similar to silver and platinum, an ability to carry higher current densities than copper,
thermal conductivity better than diamond and a density much lower than aluminium.
Figure 1.1. Publications based on CNTs over the years. Source: Web of Science.
24
Chapter 1. Introduction to Carbon Nanotubes (CNTs)
Table 1.1. Properties of different engineering fibres [11-22].
Young's
Tensile
Thermal
Electrical
Modulus
Strength
Conductivity
Conductivity
(GPa)
(GPa)
(W/mK)
(S/m)
0.01-0.04
1300
20-63
1800-6000
10
3
15-150
400
2.4
10
10-7-10
Boron
2.6
2.3-2.5
400
4
38
10 -10
HM Carbon
1.95
1.7-8
380
2-7
105
10 -10
HS Carbon
1.75
1.7-8
230
3.4
24
10 -10
Glass
2.56
11
76
2
0.05-13
10 -10
Aramid
1.4
12
70-180
3.6
0.3
Alumina
3.4
3-20
300
2
5
10
Cellulose
1
1-5
80
2
0.035-0.06
10
Fibres/
Density
Diameter
Properties
(g/cm3)
(microns)
CNTs
1.3-2.1
SiC
7
-3
-4
-2
4
5
4
5
-6
-12
-2
10 -10
-13
-3
CNTs appear in many different forms, i.e. short, long, thick, thin, single wall, multi
wall, functionalised, open, capped, stacked, containing different structural defects,
spirality and rolling structures. Each type has its own advantages and disadvantages
for different applications. This chapter comprehensively covers the structure, history,
synthesis, biocompatibility, toxicology and market overview of CNTs.
1.2. Structure of CNTs
CNTs can be conceptually visualized as rolled graphene. The main cylindrical part of
CNTs is like chicken wire, based on hexagons, whereas the end caps correspond to
25
half a fullerene molecule, made of hexagons and pentagons. CNTs can be classified in
many ways:
1. Based on the number of graphene cylinders (x), CNTs (figure 1.2) can be termed as
Single Wall CNTs (SWNT, x = 1), Double Wall CNTs (DWNT, x = 2) or Multi Wall
CNTs (MWNT, x > 2). For MWNTs, the intershell spacing between two successive
tubes is in the range from 0.344 nm to 0.36 nm and the carbon-carbon bond length is
0.144 nm [23]. To date, Shanov and Schulz [24] have synthesised the longest CNT of
18 mm, which is 900,000 times larger than its diameter. The aspect ratio of CNTs
makes them very appealing for many applications. On the other end, Peng et al. [25]
produced the thinnest SWNT of 0.33 nm. Another variation of MWNT is Carbon
Nanofibre (CNF). CNF is relatively thick (diameter exceeding 100 nm) and long (4-5
μm) nanotube [26]. To date, most of the work has been focused on SWNTs and
MWNTs. For diameter comparison, different forms of carbon fibre are shown in
figure 1.3.
(b)
(a)
(b)
Figure 1.2. Different types of CNTs based on their number of graphene cylinders: (a)
capped Single Wall CNT (SWNT); and (b) open Multi Wall CNT (MWNT).
26
Carbon Nanofibre
(CNF ~ 100 nm)
Single Wall CNT
(SWNT ~1.5 nm)
Multi Wall CNT
(MWNT ~10 nm)
Carbon fibre/
human hair (~ 5 μm)
100 nm
Figure 1.3. Different forms of carbon fibre.
2. Based on the symmetry of the carbon bonds, they can be chiral or achiral [27]. In an
achiral CNT, the cylindrical structure follows mirror symmetry in both axes,
longitudinal and transverse. In chiral CNT (also called helical CNT), the mirror
symmetry is not obeyed. There are two types of achiral CNTs, i.e. arm-chair and
zigzag (figure 1.4). Another interesting structure is cup-stacked CNT that comprises
several truncated conical graphene layers [28]. The various ways to roll graphene into

tubes are mathematically defined by the vector of helicity C (figure 1.5), and the
angle of helicity θ, as follows [23, 27, 29]:



C = na 1 + ma 2
Equation 1.1
 m 3 
And   tan 1 

 m  2n 
Equation 1.2
27
(a)
(b)
(c)
Figure 1.4. Different types of CNTs based on their chirality: a) armchair; b) zigzag;
and c) chiral.
Where n and m are integers, and can be grouped together to make lattice translational


indices (n,m). a1 and a 2 are the vectors of the hexagonal lattice that corresponds to a
section of the nanotube perpendicular to the nanotube axis (figure 1.5). The angle of
helicity θ is the tilt angle of the hexagons with respect to the rolling axis and
determines the spiral symmetry. Due to the six fold symmetry of the honeycomb
lattice, the value of angle of helicity θ falls in the range 0o – 30o. When n = m, the
nanotube is armchair type (θ = 0o); when m = 0, then it is of the “zigzag” type (θ =
0o); and when n= m, it is a “chiral”.
The value of (n,m) determines the chirality of CNT and affects the optical,
mechanical and electronic properties. CNT with |n − m| = 3q are metallic and those
with |n − m| = 3q ± 1 are semiconducting (where q is an integer) [23].
28
Figure 1.5. Schematic diagram showing how a hexagonal sheet of graphite is “rolled”
to form a carbon nanotube [23].
3. Based on the presence of defects or an attached foreign molecule, CNTs can be
functionalised for better performance. It has been widely accepted among polymer
researchers that the functionalisation of CNTs significantly increases the interfacial
interactions [30-32], dispersion [33-38] and the mechanical properties [33, 39-43] of
polymer – CNT nanocomposites. Apart from the attachment of chemical groups,
functionalisation can also be carried out by introducing different types of defects such
as: point defects like vacancies, topological defects caused by forming pentagons and
heptagons, hybridisation defects due to functionalisation [44]. However, the effect of
functionalised CNTs on ceramic matrices has not yet been considered.
29
1.3. History of CNTs
There is a lot of controversy about the discovery of CNTs [45]. The first patent on the
forming of carbon filaments was issued in 1889 [46]. The patent proposed the use of
such filaments in light bulbs [45]. Editors of the journal “Carbon” [45] believe that
Radushkevich and Lukyanovich [47] should be credited for the discovery of carbon
filaments or CNTs. These filaments were hollow and had a nanometre-size diameter
(figure 1.6a). The paper [47] was published in the Russian language in the Journal of
Physical Chemistry of Russia (1952). Due to the cold war, access to Russian scientific
publications by Western scientists was not easy at that time, and the use of the
Russian language made comprehension difficult [45]. In 1959, Walker et al. [48]
synthesised carbon filaments for the first time by thermal decomposition (now called
chemical vapour decomposition) of hydrocarbons. SWNTs were first reported in 1976
by Oberlin et al. [49] (figure 1.6b), but due to the low resolution of their TEM, they
could not be discovered. The magnification used was too low and graphene fringes
could not be resolved so the number of walls could not be determined [45]. However,
considering the diameter of this tube (~5 nm), it may have been SWNT or DWNT.
The discovery of C60, Buckminsterfullerene, in 1985 by a group led by R.E. Smalley
[50] motivated a number of researchers to work on the nano structures of carbon. In
1990, Kratschmer et al. [51] found that the soot produced by graphite electrodes
contained C60 and other fullerenes. The discovery of CNTs was an accidental event.
An electron microscopist, Iijima from NEC, Japan reported the observation of
“Helical microtubules of graphitic carbon” as a by-product of fullerene research [8].
Initially, CNTs were called “graphitic microtubules” [8, 52, 53]. In 1992, these
microtubules were referred to as “carbon nanotubes” for the first time by Ajayan and
Iijima [54]. Later, SWNTs were reported in 1993 in two papers that were published
in the same issue of NATURE, one by Iijima and Ichihashi [55], the other by Bethune
et al. [56]. In 1995, Bandow and Saito [57] made the first ceramic – CNT
nanocomposite by surrounding ultrafine particles of ZrC and V4C3 in graphene layers
by the arc burning of metal carbide graphite composites.
30
(a)
(b)
Figure 1.6. Historical forms of hollow carbon filaments. Reported in: (a) 1952 [47];
and (b) 1976. Modified from [49].
1.4. Synthesis of CNTs
The quality of CNTs is dependent on their method of production. As-synthesised,
CNTs contain carbonaceous impurities, typically amorphous carbon and graphite
nanoparticles, as well as particles of the transition-metal catalyst [58, 59]. To
eliminate these unwanted impurities, several physical and chemical methods are used,
e.g. filtration, centrifugation or microfiltration of ultrasonic-treated solutions,
chromatography, oxidation, selective reduction, and sublimation in vacuum at high
temperature [60].
CNTs are currently synthesised by different techniques (figures 1.7-1.8). The arcdischarge method (figure 1.7) is the one by which CNTs were first produced and
recognised [58]. The final properties of CNTs and their market value are much
dependant on their method of synthesis. The basic and most common methods are
summarised in table 1.2.
31
Table 1.2. Different ways of producing CNTs.
Inventor
Summary
Yields
(%)
Spacetime yield
[61]
Amorphous
carbon
[61]
Carbon
purity
[61]
Arcdischarge
Iijima, 1991 [8]
Graphitic anode
evaporates to form CNT
by a plasma via high
current
(figure 1.7).
> 70
[62]
Low
High
High
CVD
Yacaman et al.,
1993 [53]
CNTs grow from a carbon
source by the catalysis
action of metal particles
(figure 1.8).
100
[63]
High
Low
Medium
Laser
ablation
Guo et al., 1995
[64]
Direct laser vaporisation
of graphite based target
(figure 1.9).
70 - 90
[44]
Low
High
High
Method
Furnace
Furnace
Catalyst and
carbon source
CNTs
collection
Figure 1.7. Schematic diagrams of
Figure 1.8. Vertical CVD furnace for
arc-discharge method [58].
mass production. Re-drawn from [65].
32
Figure 1.9. Schematic diagram of laser ablation method [58].
The most crucial factor for the success of CNT based technologies is the production
cost, and unless cost is not reduced to a competitive level, large-scale use of
nanotubes is unlikely [66]. Some of the other techniques for the synthesis of CNTs
are HiPCoTM [67, 68], CoMoCATTM [69], electrolytic method [70], solar furnace
method [71], arc-plasma jet method [72], floating catalyst method [73], templating
technique [74], diffusion flame synthesis [75], Nano Agglomerate Fluidised Bed
(NAFB) reactor [76] and pyrolysis for the production of aligned CNTs [77]. The
synthesis of CNTs and their actual and potential industrial uses have attracted the
attention of many scientists worldwide, but relatively little attention has been paid, so
far, to their potential detrimental effects on human health and the environment [78].
1.5. Biocompatibility and toxicology of CNTs
CNTs could find application as local Drug Delivery Systems (DDSs) and scaffolds to
promote and guide bone tissue regeneration. Usage of CNTs as biomaterials is still at
33
a preliminary stage [79]. Many reports [80, 81] have confirmed that functionalised
CNTs have greater biocompatibility as compared to pristine CNTs. Kam et al. [82]
reported that the selective cancer cell destruction can be achieved by functionalisation
of SWNT without harming normal cells.
In contrast, studies [83] examining the toxicity of engineered nanomaterials in cell
cultures and animals have shown that size, surface area, surface chemistry, solubility
and possibly shape all play a role in determining the potential for nanomaterials to
cause harm [84]. CNTs fall in the category of nanoparticles and all nanoparticles are
considered toxic [85]. SWNTs have been shown to be acutely toxic [86-88] in cells,
and direct evidence was presented by Porter et al. [86]. CNTs have the ability to cause
oxidative stress and inflammation just like asbestos and quartz [87]. Some researchers
[89, 90] have argued over the toxicity of CNTs. Pulskamp et al. [90] reported that,
instead of CNTs, the metallic impurities associated with the CNTs are solely toxic
[90]. These metallic impurities can be reduced by: optimizing the synthesis of CNTs
[91]; and, purifying CNTs after production [59].
Until a clear toxicity appraisal is available, CNTs should be treated as a toxic material
[92] and strict preventive and protective measures should be taken to limit inhalation
exposure of CNTs in occupational settings [78, 87]. It should be noted that the
harmful effects of CNTs are significantly reduced, once they are encased into
composite materials.
1.6. Market overview of CNTs
Nanotubes and fullerenes already represent a significant niche market with global
revenue exceeding US$300 million in 2008. By 2015 it is predicted that the market
will exceed US$4.6 billion [93]. It is expected that CNT based composites will have
largest share of the market by a wide margin.
34
At present, the major producer of CNTs is the USA [22]. Recently, there has been a
shift of production from the USA to Asia Pacific. It is predicted that in the near future,
the major producer and supplier of all types of CNTs will be China and Korea
respectively [22]. Some of the leading suppliers of CNTs with their selling prices are
shown in table 1.3. The prices are much dependant on the quality and production
methods of CNTs.
Table 1.3. Some of the major CNT suppliers.*
Selling price (£/gram)
Supplier
Country
SES Research
SWNT
DWNT
MWNT
USA
50 - 492
52 - 125
27 - 340
Nano Lab
USA
112 - 1000
250 - 257
1.2 - 80
MER
USA
17.5 - 30
175 - 200
2.5 - 30
Helix
USA
3 - 105
62 - 105
2 - 41
NanoAmor
USA
15 - 175
27 - 53
1.3 - 38
Cheap Tubes
USA
5 - 150
NA
0.1 – 17.5
Bucky
USA
62 - 75
NA
27 - 45
Nanocs
USA
125 - 1250
NA
40 - 1000
SWeNT
USA
25 - 250
NA
NA
Carbolex
USA
30 - 50
NA
NA
Carbon Solutions
USA
25 - 200
NA
NA
Apex
USA
1.4 – 2.8
NA
NA
Nanocyl
Belgium
154 - 347
116 - 231
2.3 - 77
Arry
Germany
25 - 199
42 - 199
0.27 -26
Bayer
Germany
NA
NA
From 0.05
Nano Thinx
Greece
47.8 – 96.3
NA
5.4 – 25.4
Rosseter
Cyprus
NA
NA
10 - 12
Nano Carblab
Russia
30 - 800
125 - 750
NA
Chengdu Org. Chem.
China
22 - 35
10 - 28
0.75 - 2
Alpha Nano powder
China
20 - 75
25 - 50
0.075 - 10
NanoTechnologies
China
40 - 50
40 - 125
NA
* Surveyed on March 16, 2009, NA = Not available
35
Other major suppliers are Thomas Swan (UK), NanoDynamics (USA), Arkema
(France), Sun Nanotech (China), Hanwha (Korea) and Mitsubishi (Japan).
1.7. Summary
CNTs were comprehensively introduced in this chapter. Their unique and outstanding
properties and high aspect ratio make them an ideal reinforcement for composite
materials as compared to other engineering fibres. SWNTs are seamless rolled
graphene sheets. MWNTs are concentric graphene sheets or in other words, Russian
dolls made out of SWNTs. MWNTs were first reported in 1952. However, the time of
discovery of CNTs is still quite controversial. CNTs are currently synthesised by
different techniques and their quality relies on the method of production. With the
increase in global production, the price of CNTs is decreasing. The development of
CNT based technologies heavily relies on the availability and price of CNTs. CNTs
are potentially toxic because of their size and shape. Health hazards associated with
CNTs are greatly reduced once they are encapsulated into ceramics to create
nanocomposites.
36
Chapter 2.
Introduction to Ceramic – CNT Nanocomposites
2.1. Introduction
Since 1991, CNT reinforced nanocomposites have been the focus of intense global
research. The excitement for CNTs originates from their unique and unrivalled
properties, which were discussed in the previous chapter. A majority of the work and
reviews has been presented on the development of CNT reinforced polymer
nanocomposites. However, the application of CNTs in the reinforcement of ceramic
nanocomposites has not yet been fully exploited and is subject of major on-going
research efforts.
This chapter reviews the development of the fabrication methods and properties of
ceramic – CNT nanocomposites. Important key results and mechanisms relating to the
fracture toughness, electrical and thermal conductivity of the ceramic – CNT
nanocomposites are discussed.
2.2. Fabrication of ceramic – CNT nanocomposites
2.2.1. Pre-consolidation processing
The homogenisation and de-agglomeration of CNTs is a wide topic that falls outside
of the context of this chapter. Briefly speaking, when dispersing nano particles in a
37
Chapter 2. Introduction to Ceramic – CNT Nanocomposites
suspension, there are three stabilisation mechanisms: electrostatic stabilisation, steric
stabilisation, and electrosteric stabilisation [94, 95]. For ceramic – CNT
nanocomposites, the most common dispersion method before consolidation is the
colloidal processing route [26, 96-144]. The sol-gel method is a variation of colloidal
processing and has been used in some studies [145, 146]. The common steps in
colloidal processing are summarized in figure 2.1.
Purification of CNTs in acids and/or solvents
De-agglomeration and homogenisation of CNTs
in solvents with/without surfactants
Mixing homogenised solution of CNTs with ceramic powder
Thorough extraction of CNTs
Sieving of the powder to break agglomerates
Nanocomposite powder ready for consolidation
Figure 2.1. Colloidal processing route for making ceramic – CNT nanocomposites.
2.2.2. Consolidation techniques
Hot-pressing [104, 109-112, 114, 116, 117-119, 128, 130, 131, 133, 137, 143, 146150] and SPS [26, 96-98, 101, 103, 107, 113, 122, 136, 139-142, 145, 151-159] are
the most common consolidation techniques (figure 2.2) used for making ceramic –
CNT nanocomposites.
38
In hot pressing (figure 2.2a) the simultaneous application of pressure and heat to a
‘green’ component is responsible for the sintering of ceramic powders. Heat can be
applied directly (induction or resistance heating) or indirectly (convection or
radiation). Pressure is applied statically or dynamically to the heated component
[160]. A vacuum or controlled atmosphere can also be applied to prevent degradation
of the ceramic powder during consolidation.
(a)
(b)
Optical
pyrometer
Optical pyrometer
DC pulse
generator/
Heat
source
Heat
source
Vacuum
chamber
Vacuum
chamber
Figure 2.2. Common sintering techniques for making ceramic – CNT
nanocomposites: (a) hot pressing; and (b) Spark Plasma Sintering (SPS).
SPS (figure 2.2b) is a variation of hot-pressing that involves a different mechanism of
heat transfer. A detailed introduction of the technique is presented in section 3.2.2.
Compared to SPS, hot-pressing methods, involving longer durations and higher
temperatures, which can damage CNTs in the nanocomposite, leading to a decrease or
total loss of reinforcing effects without producing fully dense nanocomposites [97,
138, 140, 141, 147]. Other techniques used for consolidation are pressureless sintering
[123], hot isostatic pressing (HIP) [107, 124], hot extrusion [161], tape casting [102],
plasma spraying [162-165], high pressure reactive sintering [144], microwave
39
sintering [166], melt-infiltration reaction bonding [167], thermolysis [168],
solvothermal synthesis [169], laser surface alloying [100] and spray pyrolysis [170].
The effect of CNTs on the sintering behaviour of ceramics is given in chapter 5.
2.3. Mechanical properties of ceramic – CNT nanocomposites
CNTs are one of the strongest and stiffest fibres known so far, due to the intrinsic
strength of the carbon–carbon sp2 bond. The experimental evidence of such strength
can be found in a short communication by Hung et al. [171], where a single CNT
failed after stretching 280% before breaking at high temperature. Lourie et al.
estimated that the stress required for producing buckling or collapse of a CNT is
approximately 100–150 GPa [172]. Earlier work by Treacy et al. [173] showed
individual CNTs with Young’s moduli of more than 3 TPa (highest: 3.8 TPa). It has
been predicted that CNTs have the highest modulus of all the different types of
nanotubes (BN, BC3, BC2N, C3N4, CN etc.) [174].
To translate the superior properties of CNTs to ceramics, the processing route is
critical. CNTs can provide multi-axial damage tolerance to ceramic nanocomposites
[175] if they are homogeneously dispersed without agglomeration in ceramic
matrices. Mechanical properties are very dependent on the agglomeration of CNTs in
ceramic matrices [109, 128]. Another important factor is the interfacial compatibility.
Good interfacial bonding is required to achieve load transfer across the CNT – matrix
interface, a condition necessary for improving the mechanical properties of ceramic
nanocomposites [176]. CNTs and the matrix are bonded by a combination of residual
thermal stress and a diffusion layer, which makes the interface shear strength as high
as ~1 GPa [118]. CNTs can be added to ceramic matrices up to a limit. High volume
contents of CNTs in ceramics produces inhomogeneous dispersion of CNTs and more
porosity because of trapped gases in agglomerates, which reduce the mechanical and
electrical properties of the nanocomposite. This has been widely reported in the
literature [112, 176, 177].
40
2.3.1 The never ending controversy
Zhan et al. [138, 140] SPSed alumina – SWNT nanocomposites and reported a
threefold improvement in the fracture toughness at the expense of a 20% decrease in
hardness (figure 2.3a). The group used the Vickers indentation method [178] to
evaluate KIC. However, in the last 5 years, the effect of CNTs on the toughness of
ceramic has become a very controversial topic [127, 179-183]. It is due to the fact,
that, to date, no research group has been able to reproduce Zhan et al.’s [138, 140]
results. Several authors [111, 120, 128, 156, 162, 180, 182-189] argued over the
validity of the techniques used for determining the fracture toughness of ceramics and
their CNT nanocomposites (figure 2.3b). In 2007, Quinn and Bradt [185] suggested
abandoning the indentation fracture method for the evaluation of the fracture
toughness, but the research community [117, 130, 131, 136, 155, 156, 165, 190, 191]
still use the method. This is primarily due to the unavailability of an easy, effective,
quick and economical alternative method. The indentation fracture toughness method
is only useful for comparison purposes. The technique should not be used to
determine KIC as it tends to overestimate the value [185].
(a)
(b)
Figure 2.3. Fracture toughness of alumina – carbon nanocomposites. (a) By Vickers
indentation method [140]; and (b) by SEVNB method [127]. Both papers [127, 140]
were published in the same journal.
41
2.3.2 Toughening mechanisms
Fracture toughness evaluates the ability of a material to resist the crack propagation
until fracture. Different toughening mechanisms (figure 2.4) are reported for ceramic
– CNT nanocomposites: crack deflection at the CNT – matrix interface [130, 140,
141, 152, 192]; crack bridging by CNTs [110, 112, 137, 141, 153, 163, 168, 192];
CNT pullout on the fracture surfaces [104, 110, 112, 141, 152, 153, 168, 192] and
CNT shear band collapse [192].
(a)
(b)
(c)
(d)
Figure 2.4. Different toughening mechanisms in amorphous alumina – CNT (outer
diameter: ~51-56 nm) nanocomposite: a) crack deflection; b) crack bridging; c) CNT
pullout; and d) CNT collapse in shear band. Figure modified from [192].
42
Residual stresses are also responsible for enhancing the toughness of ceramic – CNT
nanocomposites [192], as they draw the transverse cracks into the compression
stresses produced at the CNT – matrix interface, at which point the crack must either
deflect or penetrate the CNT [192].
2.3.3 Effect of CNT in alumina nanocomposites
After Zhan et al.’s [140] paper on the fracture toughness of alumina – CNT
nanocomposites (section 2.3.1), Wang et al. [127] reported that alumina – SWNT and
the alumina – graphite prepared by SPS are not tough, but their contact-damage
resistance properties are attractive for applications where contact loading is prevalent.
CNTs induced anelasticity/ viscoelasticity in the alumina matrix that enhanced the
contact-damage resistance of the nanocomposite material [127]. In the same report,
they showed that alumina – SWNT has inferior fracture toughness as compared to
alumina – graphite. They used the single-edge V-notch beam method and showed
only a 3% improvement in the fracture toughness of alumina – SWNT nanocomposite
over monolithic alumina. However, the dependence of the notch radius on the
toughness was not considered, especially when the notch radius is hardly well
controlled using a razor blade. One can argue on the reliability of such results. Only a
really sharp notch can be regarded as identical to a pre-existing crack for the purpose
of meaningful long crack toughness testing [179].
Mo et al. [145] dispersed CNTs in alumina by sol-gel process followed by SPS. They
reported an enhancement in Vicker’s hardness (~7%) as well as fracture toughness
(~10%). Fan et al. [118] prepared SWNTs in alumina by heterocoagulation of
SWNTs into matrix grains. The addition of 1 wt % SWNTs in the alumina
nanocomposite increased the fracture toughness by 103% and flexural strength by
20% higher compared to unreinforced alumina ceramics. In another report, Fan et al.
[109] hot-pressed alumina – MWNT nanocomposites and reported an 80%
improvement in the fracture toughness at the cost of a 4% decrease in bending
43
strength. Sun and Gao [128] reported that 1 wt% addition of MWNTs to an alumina
matrix produced a 32% and 10% improvements in fracture toughness and bending
strength respectively compared to monolithic alumina. It was observed that hotpressing for longer duration at high temperature (1500 oC) sharply decreased the
mechanical properties as compared to lower temperatures (< 1400 oC) [128]. In
another report, Sun et al. [101] found that the addition of only 0.1 wt % CNTs in an
alumina nanocomposite increased the fracture toughness from 3.7 to 4.9 MPa√m
(~32%). Chang et al. [149] reported the fracture toughness for a nanocomposite
containing 10 vol% of the MWNT was 4.2 MPa√m; an improvement of 24 % when
compared with that of the monolithic alumina.
Siegel et al. [143] used MWNTs to toughen an alumina matrix and achieved a 24%
increase in fracture toughness. Cha et al. [142] fabricated alumina – MWNT
nanocomposites by molecular level mixing and SPS and reported a ~15% and 30%
improvements in the hardness and fracture toughness of the nanocomposite over
pristine alumina. Recently, Wei et al. [110] reported a fracture toughness increase of
79% and bending strength increase of 13% for 3 vol% CNTs – alumina
nanocomposite as compared to that of pure nanocrystalline alumina.
Zhu et al. [111] aligned CNTs in alumina using an electric field and reported an
improvement in the mechanical properties along the normal direction. The mechanical
properties of the hot pressed nanocomposites were characterised in two directions,
with fracture toughness of 4.66 MPa√m and 3.65 MPa√m, and flexural strengths of
390 MPa and 191 MPa, respectively in the normal and the perpendicular directions.
These results showed significant improvements when compared with the respective
fracture toughness and flexural strengths of 3.78 MPa√m and 302 MPa for pristine
alumina, and 4.09 MPa√m and 334 MPa for alumina nanocomposites filled with 2
wt% MWNTs prepared without the effect of an electric field [111].
Balani et al. [162, 165] used plasma spraying to fabricate alumina – CNT
nanocomposite coating. Better densification was achieved due to strong interfacial
44
bonding between MWNTs and the alumina matrix, which enhanced the hardness
[165] and the fracture toughness [162, 165] as compared to pristine alumina.
Maensiri et al. [137] fabricated alumina – CNF for the first time. An improvement in
the fracture toughness of 13% over pure alumina was reported in the alumina
reinforced with 2.5 vol% CNF. However, hardness and bending strength of the
nanocomposites were reduced with increasing volume fraction of CNF. Hirota et al.
[26] SPSed alumina – CNF nanocomposite and reported a ~25% and ~5%
improvement in the bending strength and the fracture toughness, respectively,
compared with those of monolithic alumina.
Peigney et al. [148] used a catalytic method that produced alumina – Fe
nanocomposite powders that contained in situ grown MWNTs and SWNTs. They
found that the fracture strength and the fracture toughness of alumina – CNT
nanocomposites were higher than that of alumina, but still lower than those of the
carbon free alumina – Fe nanocomposites. They reported big agglomerates of CNTs,
which resulted in poor mechanical properties. In other reports [147, 193], the same
group reported the same observations when they processed these hybrid
nanocomposites by hot pressing.
Thomson et al. [136] added 5 vol% SWNTs to 10 vol% Nb–alumina and reported
inferior indentation fracture toughness (~12% less) and hardness (~ 16% less) as
compared to 10 vol% Nb–alumina due to poor dispersion of SWNTs in the hybrid
nanocomposite. Yoo et al. [159] SPSed alumina – iron – MWNT nanocomposite and
found that the fracture strength was less than that of the alumina – iron
nanocomposite. However, Ahmad and Pan [120] SPSed alumina – SiC – MWNT
hybrid nanocomposites (figure 2.5) and reported an improvement of ~117% in
fracture toughness and ~44% in bending strength, while the hardness remained
unaffected.
45
Intra/ Inter type
Mutually redundant mechanisms
CNT
Grain
boundary
strengthening
SiC
Alumina
Nano-fibre type
SiC
CNT
Alumina
Fibre
toughening
Crack bridging by
CNTs and SiC
Figure 2.5. Schematic of hybrid microstructure design of alumina reinforced by
MWNTs and SiC nanoparticles. Modified from [120].
Other mechanical properties were also characterized to study the effect of CNT
addition in ceramics. Solvas et al. [194] SPSed alumina – SWNT nanocomposite and
found an improvement of two orders of magnitude in the creep-resistance compared
to a pure alumina, which was attributed to partial blocking of grain-boundary sliding
by SWNTs in the nanocomposites. CNTs prevent grain-boundary sliding because of
the entangled network of agglomerates (figure 2.6). This may not be true for MWNTs
because of their telescoping effect in tensile stress.
46
(a)
(b)
Figure 2.6. CNTs at the grain boundary: (a) schematics; and (b) high-resolution TEM
micrographs of an alumina – 10 vol% SWNT nanocomposite showing SWNT bundles
at alumina grain boundaries, with schematic diagrams indicating corresponding
SWNTs orientations. Figure modified from [195].
An et al. [132] studied the influence of MWNT content on the tribological properties
of alumina – MWNT nanocomposites. A 56% reduction in the wear loss and 30%
improvement in the microhardness of the nanocomposite were observed as compared
to pure alumina. Xia et al. [196] reported an ~80% reduction in the coefficient of
friction for well-aligned alumina – MWNT nanocomposites as compared to
monolithic alumina. For high wear resistance and low friction coefficients of the
nanocomposites, highly aligned CNTs and thick CNTs are useful [196]. The higher
content of CNTs in wear debris resulted in a lower coefficient of friction for an
alumina – CNTs nanocomposite [102, 197]. On the other hand, Wasche and Klaffke
[198] reported that graphite powder had no effect on the tribological properties of the
ceramics. Lim et al. [102] prepared alumina – SWNT nanocomposite by tape casting
followed by hot pressing and reported 71% less wear loss in a 12 wt% alumina –
SWNT nanocomposite as compared to monolithic alumina. For the same composition
47
of nanocomposite, they also reported a 68% improvement in the fracture strength as
compared to monolithic alumina. Lu [199] observed that during processing, the green
strength of the alumina – CNT nanocomposites increased with the CNT content when
CNTs are well separated.
2.3.4. Effect of CNT in other ceramic nanocomposites
Apart from alumina, CNTs were used to reinforce other ceramic matrices as well
[103, 107, 112, 115-117, 119, 129, 134, 141, 200, 201]. Manocha [202] reported
improved mechanical properties for C-C nanocomposite by the inclusion of CNTs.
Huang et al. [151] commented on the phase transformation and stresses in BaTiO3 –
CNT nanocomposites. Due to the presence of a high compressive stress at the grain
boundary and the CNT – BaTiO3 interface, an increase of 143% in the fracture
toughness over CNT-free BaTiO3 was reported [151].
Balazsi et al. [107, 200] compared the mechanical properties of silicon nitride –
MWNT nanocomposite, prepared by SPS and hot-pressing. Only a 2% improvement
in the fracture toughness was reported for the nanocomposite processed by SPS. On
the other hand, 13% and 17% decrements were reported in modulus and hardness
respectively for the nanocomposite as compared to pristine silicon nitride. Hot
pressing damaged the MWNTs and produced coarser grains as compared to SPS. In
other reports by Balazsi et al. [115, 133], the bending strength and elastic modulus of
MWNT – silicon nitride nanocomposites showed a considerable improvement
compared to matrices with carbon fibre, carbon black or graphite additions. But the
mechanical properties of silicon nitride – MWNT were inferior as compared to pure
silicon nitride. The decrease of modulus and strength can be related to the lower
densification rate. [133].
Hwang and Hwang [99] compared the hardness of SiO2 – CNT glass micro-rods SiO2 nanocomposites and SiO2 glass micro-rods – SiO2 nanocomposites. They found
48
that the hardness of nanocomposite discs containing ~80 wt% SiO2 – CNT glass
micro-rods increased by ~210% and ~65%, as compared to pure SiO2 and SiO2 glass
micro-rods – SiO2 nanocomposite respectively. CNTs improved the wetting and
interface between glass micro-rods and the matrix [99]. Ning et al. [119] hot-pressed
SiO2 – MWNT nanocomposite and reported that the bending strength and fracture
toughness of SiO2 – MWNTs nanocomposite, compared with the monolithic SiO2,
were enhanced by 65 and 100%, respectively. Gou et al. [141] incorporated MWNTs
in SiO2 and found a 158% and 38% improvement in the fracture toughness and the
Young’s modulus of 10 vol% MWNT – SiO2 as compared to pure SiO2. Ye et al.
[112] dispersed MWNTs in a glassy-ceramic (barium aluminosilicate) and reported an
enhancement in flexural strength (~192%) and fracture toughness (~143%) with up to
10 vol% of MWNT content in the nanocomposite. Boccaccini et al. [116] dispersed
10 wt% MWNTs in a glassy-ceramic (borosilicate) and reported a 11% improvement
in the fracture toughness of the nanocomposite. However, the presence of the CNT
aggregates and porosity weakened the material rather than reinforcing it, which
resulted in a 41%, 21% and 14% reduction in fracture strength, hardness and elastic
modulus respectively.
Wang et al. [105, 144] prepared SiC – CNT and SiC – diamond – CNT
nanocomposites by high pressure reactive sintering and reported superior mechanical
properties as compared to CNT free matrices. Thostenson et al. [167] made SiC – Si –
MWNT nanocomposite and reported that a very small amount of CNT content (<2.1
vol%) has no influence on the mechanical properties of the nanocomposite. Morisada
et al. [103] mixed nanometre-sized SiC powders with 1–5 vol% SiC – coated
MWNTs and SPSed them. An improvement of ~20% in microhardness and ~13% in
the fracture toughness was reported for the nanocomposites as compared to
monolithic SiC, due to the improved adhesion between the MWNTs and the SiC
matrix by the SiC coating. Ma et al. [134] hot pressed nano SiC – MWNT
nanocomposite and reported a 10% enhancement in the three-point bending strength
and fracture toughness over monolithic SiC ceramic. Hirota et al. [26] SPSed SiC –
CNF nanocomposite and reported a ~23% improvement in the bending strength
49
compared to monolithic SiC. An et al. [201] employed MWNTs to reinforce a
polymer-derived ceramic (SiCN) matrix and reported a 60% and 52% improvement in
the modulus and hardness respectively. Katsuda et al. [168] showed marginal
improvement in the modulus and 70% improvement in the fracture toughness of a
ternary Si-C-N matrix reinforced with MWNTs. Burghard et al. [203] showed
marginal improvement in the hardness and ~36% improvement in the modulus of SiC-N matrix reinforced with SWNTs. Wang et al. [104, 204] hot pressed a mulliteMWNT nanocomposite. The addition of 5 vol% MWNTs led to a 10% increase in
bending strength and 78% increase in fracture toughness, respectively, compared with
the monolithic Mullite [104].
Many researchers have reported the effect of CNTs on zirconia matrices [108, 124,
125, 130, 131, 205-207]. Ionascu and Schaller [125] and Daraktchiev et al. [108]
showed that introducing MWNTs into zirconia reduces the grain boundary sliding and
consequently the mechanical loss at high temperatures, leading to better creep
resistance. Recently, this was also observed for alumina – SWNT nanocomposite
[194]. For comparison, Daraktchiev et al. [108] compared yttria-stabilized zirconia –
MWNT nanocomposite with silica doped yttria-stabilized zirconia, and found that
MWNT based nanocomposite possessed better creep resistance at high temperatures.
Ionascu and Schaller [125] compared yttria-stabilized zirconia – MWNT
nanocomposite with yttria-stabilized zirconia – SiC whiskers, and found that MWNT
based nanocomposite possessed better creep resistance at high temperatures. Duszova
et al. prepared yttria-stabilized zirconia – CNF nanocomposite [131] and yttriastabilized zirconia – CNT nanocomposite [130] by hot pressing. In both
nanocomposites, due to significant porosity, the hardness and fracture toughness of
the nanocomposites were lower than that of pure zirconia. Sun et al. [139] reported
that the existence of the agglomerated CNTs in the grain boundary of ceramic grains
and the weak bonding between CNT and zirconia matrix led to the failure in
reinforcement. Ukai et al. [124] prepared yttria-stabilized zirconia – 1 wt% MWNT
nanocomposite and reported a ~60% loss in fracture strength as compared to pristine
yttria-stabilized zirconia.
50
Iron aluminide (Fe3Al) nanocomposites reinforced with CNTs were fabricated by hot
press consolidation for the first time by Pang et al. [208]. As compared to iron
aluminide, the hardness, compressive yield strength, and bending strength hardness of
nanocomposites were higher by 63%, 1%, and 5%, respectively. Zhang et al. [150]
hot pressed titanium diboride – 5 wt% nickel – 0.5 wt% MWNT nanocomposite and
showed a clear increase in the hardness (17%), bending strength (15%) and fracture
toughness (60%) of the nanocomposite as compared to that of titanium diboride – 5
wt% nickel nanocomposite. Zhang et al. [153] reported that the WC – Co – 0.5 wt%
CNTs nanocomposites possess superior hardness and toughness. The hardness was
about 15% and fracture toughness was about 40% higher than that of the pure nanoWC – Co cermets consolidated processed under the same conditions.
Singh et al. [135] reported remarkable improvements of 1069% and 1101% in elastic
modulus and hardness, when 0.1 wt% MWNTs were added to a PMMA modified HA
matrix for biomedical bone cement and implant applications. Meng et al. [117] hotpressed MWNT – HA nanocomposite and reported (50% and 28%) improvements in
the fracture toughness and flexural strength respectively when the volume percentage
of MWNTs reached 7%. Balani et al. [163] used plasma spraying to fabricate HA –
CNT nanocomposite coatings and reported a fracture toughness improvement of
~56%. Chen et al. [100] prepared MWNT reinforced hydroxyapatite nanocomposite
coating by laser surface alloying and reported a 41% and 21% improvement in the
hardness and the modulus respectively as compared to a MWNT-free nanocomposite
coating. The higher the value of the modulus for the coating, the stronger the
mismatch between the coating and the living bone tissues. This slight increase in the
modulus is very beneficial for bio applications of such coatings [100]. Kealley et al.
[106] reported that the inclusion of increasing amounts of CNT material has a slight
effect on the strain in the HA, suggesting that there may be a small amount of bonding
between the two materials. Wang et al. [174] showed that by addition of as-received
MWNTs and bio-mineralized MWNTs, the compressive strength of the calcium
phosphate cement increased by 24% and 120%, respectively.
51
The literature discussing the mechanical performance of CNT reinforced ceramic
nanocomposites is at a relatively early stage. The reports of modest improvements in
mechanical properties do not provide clear evidence linking the quantitative
performance data to the actual mechanisms involved. In short, a brief summary of the
CNT toughness effect on ceramics is shown in table 2.1.
Table 2.1. Summary of the fracture toughness of ceramic – CNT nanocomposites.
Research group
†
Year
Ceramic matrix
Processing method
Reinforcement
type
Fracture
toughness
improvement (%)
Zhan et al. [140]
2003
Alumina
SPS
SWNT
194†
Guo et al. [141]
2007
Silica
SPS
MWNT
158†
Ye et al. [112]
2006
Hot pressing
MWNT
143
Huang et al. [151]
2005
Barium titanate
SPS
MWNT
143†
Fan et al. [118]
2006
Alumina
Hot pressing
SWNT
103
Ning et al. [119]
2003
Silica
Hot pressing
MWNT
100†
Fan et al. [109]
2006
Alumina
Hot pressing
MWNT
80
Wei et al. [110]
2008
Alumina
Hot pressing
MWNT
79
Wang et al. [104]
2007
Mullite
Hot pressing
MWNT
78
Wang et al. [144]
2006
Silicon carbide
MWNT
75†
Lei et al. [152]
2008
Alumina
SPS
MWNT
70†
Katsuda et al. [168]
2006
Ternary Si-C-N
Themolysis
MWNT
70
Barium aluminosilicate
glass
High pressure
reactive sintering
Measured by Vickers indentation method
52
Table 2.1. continued
Research group
Zhan et al. [150]
2006
Meng et al. [117]
2008
Balani et al. [162]
2007
Zhang et al. [153]
2005
Sun and Gao [128]
2005
Sun et al. [101]
Ceramic matrix
Titanium diboride +
Processing method
Reinforcement
type
Fracture
toughness
improvement (%)
Hot pressing
MWNT
60
Hydroxyapetite
Hot pressing
MWNT
50†
Alumina
Plasma spraying
MWNT
43†
SPS
MWNT
40†
Alumina
Hot pressing
MWNT
33†
2002
Alumina
SPS
MWNT
32†
Cha et al. [142]
2005
Alumina
SPS
MWNT
30†
Chang et al. [149]
2000
Alumina
Hot pressing
MWNT
24†
Siegel et al. [143]
2001
Alumina
Hot pressing
MWNT
24†
Zhu et al. [111]
2007
Alumina
Hot pressing
MWNT
23
Maensiri et al. [137]
2007
Alumina
Hot pressing
CNF
13†
Morisada et al. [103]
2007
Silicon carbide
SPS
Ma et al. [134]
1998
Boccaccini et al.
[116]
Ukai et al. [124]
Mo et al. [145]
†
Year
2007
2006
2005
nickel
Tungsten carbide +
cobalt
Silicon carbide +
boron carbide
Borosilicate glass
Yttria-stabilized
zirconia
Alumina
SiC coated
MWNT
13†
Hot pressing
MWNT
13
Hot pressing
MWNT
11†
MWNT
10†
MWNT
10†
Pressureless
sintering + hot
isostatic pressing
SPS
53
Table 2.1. continued
Research group
†
Year
Ceramic matrix
Processing method
Fracture
Reinforcement
type
toughness
improvement (%)
Peigney et al. [148]
2000
Alumina + iron
Hot pressing
MWNT + SWNT
9
Thomson et al. [136]
2008
Alumina + niobium
SPS
SWNT
7
Hirota et al. [26]
2007
Alumina
SPS
CNF
5†
Sun et al. [139]
2005
SPS
MWNT
5†
Wang et al. [127]
2004
Alumina
SPS
SWNT
3
Hirota et al. [157]
2007
Silicon carbide
SPS
CNF
2†
Balazsi et al. [107]
2006
Silicon nitride
SPS
MWNT
2†
Duszova et al. [131]
2008
Hot pressing
CNF
-10†
Yttria-stabilized
zirconia
Yttria-stabilized
zirconia
2.4. Electrical properties of ceramic – CNT nanocomposites
The effect of the addition of CNT on the electrical conductivity of ceramics is much
more significant than their effect on mechanical properties. Adding very little amount
of CNT to a ceramic that is inherently an insulator can make it a good conductor.
Carbon, a group IV element like Si, has a lot of potential in electronics. Prior to the
discovery of nano fullerene structures of carbon, diamond (bandgap ~6 ev) and
graphite (semi-metal) were the most commonly known forms [209]. The axial
electrical conductivity of CNTs was found to be extremely high, reaching 2 x107 S/m
[210], comparable to that of platinum i.e. 6.8 x107 S/m [211]. The electrical properties
of CNTs vary with the type, diameters, chirality and defects in the structure.
54
SWNTs have much better electrical properties as compared to MWNTs due to their
perfect structure as compared to MWNTs [97]. MWNTs can be metallic or semiconducting [158]. In semi-conducting category, MWNTs have a p-type conducting
semiconducting behaviour, which can be converted to n-type semiconductor by
doping alkali earth ions or annealing in a reducing atmosphere [158]. Based on the
current-voltage characteristics, several authors [121, 122, 212] reported that the
conductivity mechanisms involved in CNT are fluctuation-assisted tunnelling [213],
whereas variable range hopping could be the conductivity mechanism as indicated in
other reports [212].
2.4.1. Percolation threshold in ceramic – CNT nanocomposites
A very small volume content of CNTs, as low as 0.3 volume %, resulted in a 75%
(measured by two-point method) reduction in the electrical resistivity of SiC –
MWNTs nanocomposite [167]. Rul et al. [146] uniaxially hot pressed SWNT +
MWNT – MgAl2O4 and reported a percolation threshold of 0.64 vol% (figure 2.7a).
To date, this is the lowest percolation threshold reported. However, the waviness of
CNTs is not considered in the scaling law of percolation theory. The effective
(theoretical) percolation threshold may be lower than 0.64 vol% and the probable
existence of paths with a low conductivity due to damage of some of the CNTs during
hot pressing [146]. Ahmad et al. [96] showed that the electrical conductivity increased
sharply as the content of MWNTs in alumina was close to percolation threshold of
0.79 vol% (figure 2.7b). This percolation threshold is 20 times smaller than that of
micron scale two-phase random composites, and this low value is attributed to the
enormous aspect ratio of MWNTs [96].
55
0.64%
0.79%
(a)
(b)
Figure 2.7. Electrical conductivity (measured at room temperature) of ceramic – CNT
nanocomposites as a function of CNT content. Percolation threshold is: (a) 0.64 vol%
for MgAl2O4. Modified from [146]; and (b) 0.79 vol% for Al2O3 [96].
Hirota et al. [26] SPSed alumina – CNF nanocomposite and reported a percolation
threshold of 1.5 vol%. Shi and Liang [122, 126] SPSed 3 mol% yttria (Y) stabilized
zirconia – MWNT and experimentally found a percolation threshold of 1.7 wt% (4.7
vol%), four times smaller than that of Ti3SiC2 – 3Y – TZP (Tetragonal Zirconia
Polycrystals) nanocomposites [214].
2.4.2. Effect of CNT on the electrical conductivity
Zhan et al. [97, 138] SPSed alumina – SWNT nanocomposite and reported an
increase of 13 orders of magnitude in the electrical conductivity over pure alumina.
To date, this group has reported the highest electrical conductivity of 3345 S/m for 15
vol% SWNT – alumina nanocomposite (figure 2.8).
56
Figure 2.8. The electrical conductivity of various representative materials at room
temperature. Note the more than 13 orders of magnitude increase in conductivity of
the alumina – 15 vol% SWNT nanocomposite compared to monolithic alumina [138].
Modelling shows that the effective electrical conductivity of the MWNT – alumina
nanocomposite in the aligned case is about three times of that in random distribution
[215]. Peigney et al. [161] aligned CNTs in a Fe – Co – MgAl2O4 matrix by hightemperature extrusion and reported a 20 times greater conductivity in the direction of
extrusion as compared to the transverse one. However, for MgO – Co – CNT
nanocomposites, high-temperature extrusion damaged the structure of the CNT,
which reduced the electrical conductivity from 20 S/m (un-extruded) to 10-6 S/m
[161]. Recently, Zhu et al. [111] prepared alumina – 2 wt% MWNT nanocomposite
by aligning MWNTs with AC electrical field. The electrical conductivities of these
nanocomposites in both parallel and perpendicular directions to the MWNTs
alignment were 6.2 x 10−2 Sm−1 and 6.8 x 10−9 Sm−1, respectively, compared with that
of 4.5 x 10−15 Sm−1 for pristine alumina ceramics. Rul et al. [146] uniaxially hot
pressed 24.5 vol% SWNT + MWNT – MgAl2O4 and reported an electrical
57
conductivity of 853 S/m. Flahaut et al. [147] uniaxially hot pressed 10 vol% CNT –
4.3 vol% Fe – Al2O3 and reported an electrical conductivity of 280-400 S/m.
Thostenson et al. [167] fabricated SiC – Si – MWNT nanocomposites by melt
infiltration of silicon and reported an electrical conductivity of 1538 S/m, a 96%
decrease in electrical resistivity was observed for the ceramics with the highest CNT
volume fraction of 2.1%. Balazsi et al. [133] prepared silicon nitride nanocomposites
with MWNTs, carbon black and graphite by hot isostatic pressing. They reported a
DC conductivity for a 5 wt% MWNT – silicon nitride nanocomposite of 85% less
than that of 5 wt% carbon black – silicon nitride nanocomposite due to porosity and a
poor interface at CNT/ matrix [133]. Tatami et al. [129] showed that hot pressing after
Gas Pressure Sintering (GPS) of silicon nitride – MWNT nanocomposite improves
the electrical conductivity from 30 to 79 S/m. Densification occurred at a much lower
temperature using Y2O3 – Al2O3 – TiO2 – AlN as sintering aids during hot pressing
[129].
Jian and Gao [123] reported a 44.7% and 11.5% enhancement in electrical
conductivity when they added ~12 vol% CNT to TiN and Fe2N systems respectively.
In another report, Jiang and Gao [169] reported that the addition of CNTs in a
magnetite nanocomposite increased the electrical conductivity by about 32% from 1.9
to 2.5 S cm-1, compared with magnetite.
Boccaccini et al. [116] dispersed 10 wt% MWNTs in glassy–ceramic (borosilicate)
and reported an electrical conductivity of 7.7 S/m for the nanocomposite as compared
to 10-3 S/m for monolithic Duran glass. Guo et al. [113] SPSed silica – MWNT
nanocomposite and reported that the electrical conductivity of the system increased
linearly with temperature from 5 to 300 K, showing a negative temperature coefficient
of resistivity.
Huang and Gao [114, 158] observed that the conductivity of BaTiO3 – MWNT
decreased 47% when MWNTs are located within the grains. When CNTs are located
58
in the grain, they are not connected to the conductive network. The space between
CNT and the grain is mainly responsible for the decreased electrical conductivity of
the nanocomposite [114, 158]. Whereas in other case, MWNTs have been found to be
effective in improving the electrical conductivity at the grain boundaries [96, 216] in
ceramics.
Zhan et al. [98] SPSed yttria stabilized zirconia – SWNT and alumina – yttria
stabilized zirconia – SWNT nanocomposite and reported that the thermoelectric
power of these nanocomposites increases with increasing temperature. But as
compared to other thermoelectric materials, the electrical conductivity of the CNT/
ceramic nanocomposites is still low. Hirota et al. [26] SPSed alumina – 10 vol% CNF
nanocomposite and reported an electrical conductivity of ~588 S/m. Duszova et al.
[131] prepared yttria–stabilized zirconia – CNF nanocomposite and reported an
electrical conductivity of 952 S/m, measured by the four point method. A brief
summary of the room temperature electrical conductivity of ceramic – CNT
nanocomposite is given in table 2.2.
Table 2.2. Summary of the electrical conductivity of ceramic – CNT nanocomposites.
Group
Hirota et al.
[26]
Zhan et al.
[97]
Balazsi et
al. [133]
Shi and
Liang [122]
Guo et al.
[113]
Zhan et al.
[98]
Year
Ceramic matrix
Processing method
Reinforcement
type
Electrical
Percentage
conductivity
difference
(S/m)
(%)
2007
Alumina
SPS
CNF
588
1020
2003
Alumina
SPS
SWNT
3345
1017
2006
Silicon nitride
Hot pressing
MWNT
130
1016
SPS
MWNT
65
1016
SPS
MWNT
64.5
1016
SPS
SWNT
55
1016
2006
2007
2006
Yttria stabilized
zirconia
Silica
Yttria stabilized
zirconia
59
Table 2.2. continued.
Group
Tatami et al.
[129]
Zhu et al.
[111]
Duszova et
al. [130]
Rul et al.
[146]
Ukai et al.
[124]
Ahmad et al.
[96]
Boccaccini et
al. [116]
Thostenson
et al. [167]
Jian and Gao
[123]
Jian and Gao
[169]
Jian and Gao
[123]
Huang and
Gao [114]
Huang and
Gao [158]
Peigney et
al. [161]
Flahaut et al.
[147]
Flahaut et al.
[147]
Year
Ceramic matrix
2005
Silicon nitride
2007
Alumina
2008
2004
2006
2006
2007
2005
Yttria stabilized
zirconia
Magnesium
aluminate
Electrical
Percentage
conductivity
difference
(S/m)
(%)
MWNT
79
1015
Hot pressing
MWNT
0.062
1015
Hot pressing
MWNT
952
1014
Hot pressing
SWNT + MWNT
853
1014
MWNT
13.2
1013
Processing method
Gas pressure sintering
+ hot pressing
Reinforcement
type
Yttria stabilized
Pressureless sintering
zirconia
+ hot isostatic pressing
Alumina
SPS
MWNT
7
105
Hot pressing
MWNT
7.7
105
Melt infiltration
MWNT
1538
1588
MWNT
73500
44.7
MWNT
2
32
Borosilicate
glass
Silicon carbide
+ silicon
2005
Titanium nitride
2003
Magnetite
2005
Iron nitride
MWNT
88500
11.5
2004
Barium titanate
Hot pressing
MWNT
363
-47
2005
Barium titanate
SPS
MWNT
790
-64
Hot extrusion
SWNT + MWNT
2000
-
Hot pressing
MWNT
400
-
Hot pressing
MWNT
180
-
Solvothermal
synthesis
Magnesium
2002
aluminate + iron
+ cobalt
2000
Alumina + iron
2000
aluminate + iron
Magnesium
+ cobalt
60
Table 2.2. continued.
Group
Peigney et
al. [161]
Zhan et al.
[98]
Flahaut et
al. [147]
Yoo et al.
[159]
Peigney et
al. [161]
Year
2002
Ceramic matrix
Alumina + iron
Processing method
Reinforcement type
Electrical
Percentage
conductivity
difference
(S/m)
(%)
Hot extrusion
SWNT + MWNT
158
-
SPS
SWNT
34
-
Hot pressing
MWNT
20
-
SPS
MWNT
8.47 x 10-4
-
Hot extrusion
SWNT + MWNT
1.8 x 10-6
-
Alumina + yttria
2006
stabilized
zirconia
2000
2006
2002
Magnesium
oxide + cobalt
Alumina + iron
Magnesium
oxide + cobalt
2.5. Thermal properties of ceramic – CNT nanocomposites
Due to a large phonon mean free path in the strong carbon sp2 bond network of CNT
walls, a very high theoretical value of thermal conductivity (~6600 W/mk) was
predicted [217]. This high thermal conductivity is very much dependant on the
structure, vacancies and defects associated with the individual CNTs [218].
Earlier results by Zhan et al. [138] showed that SWNTs decreased the thermal
conductivity of the alumina nanocomposite as compared to pure alumina.
Agglomeration and high thermal resistivity at the CNT – ceramic interface were
mainly responsible for the poor behaviour. Calculations by Bakshi et al. [219] based
on Xue’s model [220] for alumina – MWNT nanocomposite shows that thermal
conductivity increases with the quality of the MWNTs dispersion. The strong tubetube coupling can decrease the thermal conductivity of MWNT bundles by an order of
magnitude relative to individual tubes [138, 217]. Zhang et al. [221] reported that
CNTs exhibited thermal conductivity as low as 4.2 W/m/K when SPSed into bulk
61
samples. The tube–tube interaction, entangled tubes, tube kinks and tube crosssectional heat diffusion were responsible for such a lower thermal conductivity.
Huang et al. [222] reported a decrease in the thermal conductivity after adding
MWNTs due to an interfacial thermal barrier between CNTs and the BaTiO3 matrix
(figure 2.9a). The lattice mismatch between CNTs and BaTiO3 may account for this
thermal barrier [222]. For other ceramics, the reduced thermal conductivity is useful
for the applications of Thermal Barrier Coatings (TBCs) where they are used to
protect critical components in the hot sections of gas turbine engines [223]. In
ceramic – CNT nanocomposites, the thermal conductivity is affected by many factors
like the crystallite boundaries, porosity, CNT content, interphase boundaries, and
phase content [219].
Sivakumar et al. [224] reported that the thermal conductivity of the silica – MWNT
nanocomposites was improved due to incorporation of MWCNT compared with a
pure silica matrix. The improvement in thermal conductivity was enhanced with
increasing amounts of MWNTs [224]. Chin et al. [225] fabricated a ceramic – CNT
catalyst and reported that the novel structure possessed superior thermal conductivity.
The CNTs allowed efficient heat removal from catalytic active sites during
exothermic reaction. Jiang and Gao [226] found that the thermal conductivity
increased with increasing MWNT content and temperature for a TiN – MWNT
nanocomposite. In the presence of 5 wt% MWNTs, there was a 97% and 11%
enhancement in the thermal conductivity at 430 oC and 27 oC, respectively, compared
with that of TiN [226]. Apart from CNTs, Nakamatsu et al. [227] found that carbon
coating on aluminium nitride powder increased the thermal conductivity of the final
sintered product by 22%.
62
0 wt% CNT
0.1 wt% CNT
1 wt% CNT
5 wt% CNT
0 wt% CNT
0.5 wt% CNT
1 wt% CNT
3 wt% CNT
(a)
(b)
Figure 2.9. Thermal conductivity of different ceramic – CNT nanocomposites as a
function of temperature. (a) Barium titanate. Modified from [222]; and (b) titanium
nitride. Modified from [226].
The thermal conductivity of the alumina – CNT nanocomposites [138, 219] decreased
with increasing temperature due to the dominant effect of Umklapp scattering
(phonon-phonon scattering) in reducing phonon mean-free path length [138]. This
trend has also been reported for the thermal conductivity of CNTs [217, 221],
diamond and graphite [217]. However, in titanium nitride – MWNT nanocomposite
[226] phonons dominate thermal transport at all temperatures, which resulted in
higher thermal conductivity at high temperatures (figure 2.9b). Ning et al. [228]
showed that at 650 oC, the thermal conductivity of a 10 vol% CNT – silica
nanocomposites was ~21% more as compared to that of silica. The exact mechanism
for increased thermal conductivity at high temperatures was not reported. A brief
summary of the room temperature thermal conductivity of ceramic – CNT
nanocomposite is given in table 3.3.
63
Table 2.3. Summary of the room temperature thermal conductivity of ceramic – CNT
nanocomposites.
Processing
Reinforcement
Percentage
method
type
difference (%)
Silica
SPS
MWNT
65
2008
Alumina
Plasma spraying
MWNT
15
Jiang and Gao [226]
2008
Titanium nitride
SPS
MWNT
11
Huang et al. [222]
2005
Barium titanate
SPS
MWNT
-3
Zhan et al. [219]
2004
Alumina
SPS
SWNT
-73
Thostenson et al. [167]
2005
Melt infiltration
MWNT
-9
Group
Year
Ceramic matrix
Sivakumar et al. [224]
2007
Bakshi et al. [219]
Silicon carbide +
silicon
2.6. Miscellaneous effects of CNTs on ceramics
Luo et al. [205] revealed that the addition of CNTs during the synthesis of ZrO2
assisted the transition from monoclinic ZrO2 to cubic ZrO2. CNTs stabilized cubic
ZrO2 and prevented agglomeration of ZrO2 nanoparticles as well [205]. Balani et al.
[163] found that CNTs improved the crystalline content of HA by 27%, which is
beneficial for the bio applications of hydroxyapatite coatings [229].
Shi and Liang [122] reported that the dielectric constant was greatly increased when
the MWCNT concentration in yttria stabilized zirconia was close to the percolation
64
threshold, which was attributed to dielectric relaxation, the space charge polarisation
effect, and the percolation effect. Dou et al. [230, 231] and Kovac et al. [232] sintered
magnesium boride – CNT nanocomposites and reported superior magnetic properties
of the superconducting wires. Sun et al. [233] made iron oxide – CNT
nanocomposites that showed ferromagnetic and super magnetic behaviours, when
processed using different processing conditions. Huang and Gao [114] observed that
BaTiO3 semiconductor transformed from n-type to p-type after doping with 0.1 wt%
of CNTs due to the formation of a Schottky barrier constructed at the CNT–matrix
contact. Huang and Gao [158] fabricated bi-layer ceramics, stacked by one layer of
MWNT free BaTiO3 and another layer of 1 wt % MWNT – BaTiO3 nanocomposite
that showed excellent rectification properties.
For improved photovoltaic properties, Cao et al. [234] coated CdS and Lee et al.
[235] coated TiO2 on MWNTs. Different authors found that ceramic coatings such as
TiC [236], BN [237, 238], NiO [239] and BaO/ SrO [240] on the surfaces of CNTs
improve the field emission properties of CNTs as compared to uncoated CNTs.
However, a dielectric MgO coating on CNT decreased the strength of electric field on
the CNT surface and increased the tunnel barrier for field emission [241]. Other
ceramics like SiO2 [242], SnO2 [243], TiO2 [244] were coated on CNTs for the
development of functional nanocomposites.
Espinosa et al. [245] showed that the addition of a small quantity of MWNTs to
different ceramics can significantly improve the detection capability of metal oxidebased sensors at low operating temperatures.
Ma et al. [246] and Li et al. [247] showed that tungsten carbide – CNT
nanocomposite has improved electro catalytic activity compared to tungsten carbide.
Addition of MWNT improves the electrochemical activity of LiNi0.7Co0.3O2 [248] and
LiCoO2 [249], which are widely used in batteries.
65
2.7. Summary
Application of CNTs as a reinforcement in ceramic nanocomposites has not yet been
fully exploited and is the subject of major on-going research efforts. This chapter
reviewed recent studies conducted on the development of ceramic – CNT
nanocomposites. CNTs have been demonstrated to increase the mechanical, electrical
and thermal properties of the ceramic matrices. However, a significant degree of
discrepancy still exists, especially with regard to the mechanical properties of these
materials. The effect of CNTs on the electrical properties of the ceramics is the most
promising area. An improved understanding of ceramic nanocomposites and
breakthroughs in materials processing are need to be achieved for the successful
placement of ceramic – CNT nanocomposites in high technology applications.
66
Chapter 3.
Materials and Experimental Techniques
3.1. Introduction
This chapter provides detailed information about the materials, equipments and
experimental procedures used in this study.
3.2. Materials
3.2.1. Carbon Nanotubes
A detailed introduction to CNTs was covered in Chapter 1. Pristine and alumina
coated MWNTs were used in this study (table 3.1, figure 3.1).
Table 3.1. CNTs (synthesized by CVD method) used in this study.
Primary source /
Nanocyl
Arkema
NanoDynamics
Properties
(Belgium)
(France)
(USA)
Commercial grade
NC 7000
GraphiStrength C100
NDCNT (E-1005-03)
Density (g/cm3)
1.66
1.82
3.04
1.72
2.70
2.96
Average no. of graphitic shells
10 - 15
5 - 15
5 - 15
10 - 30
10 - 30
10 - 30
% wt carbon
90
> 90
~46
> 92
~ 56
~ 45
Average diameter (nm)
9.5
10 - 15
20 - 45
10 - 30
15 - 35
20 - 40
Average length (µm)
1.5
0.1 - 10
0.1 - 10
1-2
1-2
1-2
ALD* coating cycles
0
0
50
0
25
50
* ALD was performed in Prof. A. Weimer’s lab, University of Colorado, USA.
67
Chapter 3. Materials and Experimental Techniques
(a)
(b)
(c)
(d)
(e)
Figure 3.1. Different types of MWNTs used: (a,b) NC 7000; (c) GraphiStrength
C100; and (d,e†) alumina coated (50 ALD cycles) GraphiStrength C100.
Atomic Layer Deposition (ALD) is an ideal method for depositing thin films on high
aspect ratio materials as it is independent of line of sight and self-limiting [250].
Sequential surface chemical reactions deposit highly conformal films with precise
control at the atomic scale [250]. The method has been shown to be a viable technique
to deposit a coating on a single CNT without adversely affecting its inherent
† Figure 3.1e was provided by Prof. A. Weimer’s group
68
properties [251, 252]. Alumina coating on CNTs was done by elsewhere by ALD
method [250].
3.2.2. Carbon black
Carbon-derived powders and particles comprise a family of synthetic materials,
known under the generic term of carbon black, made by burning hydrocarbons in air.
Carbon black (CB) are aggregates of graphite micro crystals, each only a few unit
cells in size [253]. Some other names for CB are acetylene black, channel black,
furnace black, lamp black, lampblack and thermal black [254]. For comparative
studies, we used different types of carbon black powders (figure 3.2) summarised in
table 3.2.
(a)
(b)
Figure 3.2 Carbon black powders: (a) Vulcan XC72; and (b) Printex L6.
Table 3.2. Carbon black powders used in this study.
Primary source / Properties
Cabot (USA)
Degussa (Germany)
Commercial grade
Vulcan XC72
Printex L6
Density (g/cm3)
2.01
1.80
Average particle size (nm)
13
18
Particle shape
Very spherical
Roughly spherical
69
3.2.3. Alumina
Alumina was used as the main ceramic matrix because of its industrial significance.
Alumina is the most commonly used structural [9] and bio- ceramic [10]. It exists in a
number of crystalline phases (polymorphs).
The most important, and common,
polymorphs are denoted α, γ, θ, and κ. These phases of alumina are unique for
different applications. For example, the α and κ phases are widely used as wear
resistant coatings due to their high hardness and thermal stability, while γ- and θalumina are more suited for catalytic applications due to their high surface energies,
leading to larger active surface areas for catalytic reactions [255]. In addition to these,
there are more than twenty other crystalline phases of alumina [255].
In this study α – alumina (figure 3.3) is the main phase that existed in the final
sintered product. The α – alumina is also known as corundum (the name comes from
the naturally occurring mineral corundum. It is used not only in materials science, but
exists also as gemstones. Ruby is α – alumina doped with chromium, whereas
sapphire is α – alumina doped with iron and titanium [255]. α – alumina has a
variation of rhombohedral structure, whereas γ – alumina has a defected cubic spinel
structure [255-256]. The thermodynamic stability of α-alumina makes it the most
suited phase for use in many industrial applications. The corundum structure can be
visualized as layers of hexagonal close-packed oxygen atoms with small Al atoms in
two-thirds of the octahedrally coordinated holes between the oxygen atoms [255,
257]. The atomic positions consist of 12 aluminium atoms and 18 oxygen atoms. The
unit cell dimensions are: a = b = 4.7588 A° and c = 12.992 A° [255, 257]. The α
structure is thermodynamically stable at all temperatures up to its melting point at
2051 °C (figure 3.4). However, the metastable phases (e.g., γ and θ) still appear
frequently in alumina growth studies [255].
The alumina powder used in this study was commercially available “544833
aluminium oxide” nanopowder from Sigma-Aldrich, UK (figure 3.5). As supplied by
70
the supplier, the main features of this product are: γ phase; particle size: < 50 nm;
surface area 35-43 m2/g; melting point 2040 oC; and density 3.97 g/cm3.
    0.4759nm
c  1.2992nm
    90o
  120o
Aluminium
Oxygen
Stacking order
Unit cell (variation of rhombohedral crystal structure).
Rhombohedras combine to form hexagon.
Figure 3.3. Crystal structure of α – alumina. Figure modified from [256].
Figure 3.4. Phase transitions in alumina.
Figure 3.5. Alumina powder used in
Redrawn from [255].
this study.
71
3.2.4. Other ceramic matrices
Some other ceramics studied for comparison in Chapter 4 and Chapter 6 are:
3.2.4a. Reduced titanium dioxide (Red. titania)
Reduced titanium dioxide (Red. Titania/ TiO2) was supplied by Atraverda Inc., UK.
This powder contained coalesced particles, which formed strong big agglomerates
(figure 3.6a). After 24 hours of ball milling in ethanol, these micron-sized particles
were broken up (figure 3.6b). As per supplier, the main features of this powder are
given in table 3.3. This material is used as electrode because of its unique
combination of metallic-like electrical conductivity along with the characteristic high
corrosion resistance of ceramics [258].
(a)
(b)
Figure 3.6. Reduced titanium dioxide: (a) as-received form; and (b) after ball milling.
Table 3.3. Properties of reduced titanium dioxide as per supplier.
Surface area (m2/g)
Grade
Density
(g/cm3)
Melting point
(oC)
Ebonex
4.29
1800
† Measured by BET surface analyzer
Mean particle size (µm)
Before ball
milling
After ball
milling†
Before ball
milling
After ball
milling*
0.32 – 0.34
1.6 – 2.1
52
24.5
* Measured by Nano-particle size analyzer (section 3.3.1)
72
3.2.4b. Boron carbide (B4C)
Boron carbide was provided by H.C. Starck, Germany. The powder was ball milled
for 24 hours (figure 3.7) in ethanol before mixing with CNTs. As per supplier, the
main features of this powder are given in table 3.4. Typical applications of boron
carbide are abrasive grit, polishing, lapping, light weight armour, wear resistant
engineering components, sintering additives for different ceramics, neutron shielding
[256].
Table 3.4. Properties of boron carbide as per supplier.
Surface area (m2/g)
Grade
Density
(g/cm3)
B:C ratio
HT 03
2.52
3.8 – 4.2
† Measured by BET surface analyzer
Mean particle size (µm)
Before ball
milling
After ball
milling†
Before ball
milling
After ball
milling*
2.5 – 4
4 – 4.6
4.5
3.8
* Measured by Nano-particle size analyzer (section 3.3.1)
(a)
(b)
Figure 3.7. Boron carbide: (a) as-received form; and (b) after ball milling.
3.2.4c. Boron nitride (BN)
Boron nitride was supplied by H.C. Starck, Germany. As per supplier, the main
crystalline phase of the raw powder is hexagonal. Details are given in table 3.5. The
73
raw powder consists of soft nano-sized flakes (figure 3.8). Typical applications of this
material are solid lubricant for high temperature applications, mould release, raw
material for cubic-BN (second hardest material known), evaporation boats, thermally
conductive for polymers, refractories etc [259].
Table 3.5. Properties of boron nitride as per supplier.
Grade
C
Density
3
(g/cm )
2.1
% B2O3
Surface area (m2/g)
5-8
10 - 20
(a)
(b)
Figure 3.8. Boron nitride: (a) as received form; and (b) nano-sized flakes.
3.3.1. Nano-particle size analyzer
Zetasizer nano-particle analyzer (series Malvern nano ZS) was used to study the
colloidal stability and re-aggregation trend of the powder in Chapter 4. The results
were analyzed using standard software “Dispersion Technology software, ver. 4.00”.
The instrument (figure 3.9) performs size measurements using a process called
“Dynamic Light Scattering (DLS)”, also known as “Photon Correlation Spectroscopy
(PCS)”. The technique analyzes the Brownian motion of the suspended particles/
agglomerates and relates it to their size [260]. It does this by illuminating the
74
particles/ agglomerates with a laser and analyzing the intensity fluctuations in the
scattered light. DLS is very sensitive to the intensity of light scattered by the particles/
agglomerates. Large particles/ agglomerates scatter more light than smaller ones.
Hence, this technique is very good for studying particle/ agglomerate sizes and
colloidal dispersions. Both, Zetasizer nano-particle analyzer and software were
Attenuator
Laser source
supplied by Malven Instruments Ltd., UK.
Digital signal
processor
Detector
Cuvette
Computer
Detector
Zetasizer Nano ZS
Figure 3.9. Working principle of Zeta particle size analyzer. Modified from [260].
Spark Plasma Sintering (SPS) is a variation of hot-pressing. It involves the rapid
heating of graphitic dies by pulsed DC electric currents (figure 3.10). This rapid
heating rate (up to 600 oC/ minute) combined with high pressure (up to 1 GPa) [261]
is the main feature of SPS. During SPS, the detailed mechanism of enhanced
densification is unclear. This is due to the particular electrical, thermal and
75
mechanical processes that are associated with the SPS process. Modelling suggests
that high heating rates reduce the duration of densification-noncontributing surface
diffusion, that favours powder systems’ sinterability and the densification is
intensified by grain-boundary diffusion [262]. It has also been suggested that the
direct current pulse could generate several effects such as spark plasma, spark impact,
Joule heating, and electric field assisted diffusion [263-265]. Three factors that
contribute to the rapid densification process can be discerned: (i) the application of a
mechanical pressure; (ii) the use of rapid heating rates; and (iii) the use of pulsed
direct current (figure 3.11), implying that the samples are also exposed to an electrical
15 ms
5 ms
Transformer voltage (V)
Electric current (kA)
field [266].
Time (milliseconds)
Figure 3.10. Typical DC-pulsed current
Figure 3.11. Effect of DC pulse on the
cycles (used in this study).
density of alumina [267].
The commonly used name for this technique is very misleading and quite debatable.
To date, there is no experimental observation of a “Spark Plasma” during SPS. Other
names for SPS are pressure-assisted resistance sintering, electric-discharge sintering,
discharge powder compaction, electro-consolidation, plasma activated sintering, fieldassisted sintering, electric pulse sintering, pulse electric current sintering and
electromagnetic-field-assisted powder consolidation [262, 264].
SPS is also a cost-effective sintering technique. The entire processing time to sinter
dense ceramic composites is about 1/10 of that required by conventional sintering
processes [268]. It is also shown that SPS typically achieves maximum density at
76
temperatures of about 150-200 oC [269] and 250-300 oC [270] lower than hot
pressing. SPS also limit the grain growth. Lee et al. [271] reported grain size of ~200
nm in 99% densified TiO2 by SPS. Compared to this, the microwave sintering
resulted in grain size of ~300 nm, as against 1–2 μm grain size in conventionally
sintered TiO2 [272].
Other techniques that also involve rapid heating and sintering are Self-propagating
High-temperature Synthesis (SHS) [273, 274] and microwave sintering [275, 276].
However, temperature and heating rate cannot be practically controlled in these
techniques.
For ceramic – CNT nanocomposites, the use of the SPS technique allows sintering in
very short times, limiting the matrix grain growth and damage to the CNTs [197].
Compared to SPS, hot-pressing methods, involving longer durations and high
temperature, damage carbon nanotubes in the nanocomposite, leading to a decrease or
total loss of reinforcing effects without producing fully dense nanocomposites [97,
138, 140, 141, 147].
In this study, all the samples were Spark Plasma Sintered (SPSed) in a HPD 25/1
furnace by FCT Systeme, Germany (figure 3.12). The current set up at Queen Mary,
University of London allows samples of up to about 80 mm to be produced with
sintering temperatures up to 2200°C. Details of the graphite parts and their assembly
are shown in figure 3.13. The furnace has an optical pyrometer above the furnace and
focussed inside a hole in the top graphite punch (figure 3.13). Typically, the voltage
applied between the upper and lower punches is in the order for few voltages and the
current can be as high as few thousand amperes. Graphite dies were lined with
graphite paper (Le Carbone, UK, thickness: 0.38 mm) before pouring powder in them.
It was used to prevent direct contact between graphite parts and the ceramic powder
and to guarantee electrical contacts between all parts. Graphite dies were covered with
carbon insulation (SGL, UK, thickness: 7mm) to avoid heat loss during SPS
operation.
77
(a)
(b)
Figure 3.12. SPS facility by FCT Systeme, Germany: (a) SPS facility at Queen Mary,
University of London, UK; and (b) SPS at 1800 oC [277].
Pyrometer measuring temperature
through this channel
Top alloy piston (moving)
Carbon reducer
Carbon die
Carbon insulation jacket
Powder (before sintering)
Carbon punch
Bottom alloy piston (fixed)
Figure 3.13. Cross-sectional view of carbon die set.
78
3.3.3. Density measurements
For the evaluation of bulk density, all of the sintered samples were ground using SiC
paper to remove the carbon paper and diffused carbon layer. The bulk density was
measured by water buoyancy method and the density of powders was measured by
Helium pycnometery.
3.3.3a. Water buoyancy method
The actual density (  A ) of the sintered samples was measured by weighing them in
air ( mA ). The material was then submerged in distilled water and put on heating plate
to boil the water. In this way, distilled water penetrated into the open porosities. After
10 minutes of boiling, the distilled water was left to cool down to room temperature,
as the density of distilled water (  W ) changes with the temperature. Using a
Archimedes density kit, the sintered material was submerged in the distilled water,
and the submerged mass ( mw ) was recorded. The actual density (  A ) was then
calculated using:


mA
 
 
m m 
A
A
W
Equation 3.1
W
XRD analysis confirmed that there was no reaction between the CNTs and the
ceramics used in this project. Hence, the theoretical density (  T ) of the composites
was calculated according to the rule of mixtures. Rel. theoretical density (  R ) was
calculated by:

 

R

  100
T 
A
Equation 3.2
79
3.3.3b. Helium pycnometery
The density of the CNTs was measured by He pycnometer. The Micromeritics
AccuPyc 1330 pycnometer is a fully automatic gas displacement pycnometer. The
instrument determines the skeletal density and volume of a sample by measuring the
pressure change of helium in a calibrated volume chamber. The instrument is very
accurate as the Helium molecule has a diameter of less than 0.1 nm. The instrument
has a density resolution of 0.001 g/cm3 [278]. The technique is good for powders
only. It should not be used for sintered products, as it cannot evaluate the amount of
closed porosity.
3.3.4. Electrical conductivity measurements
The electrical conductivity of the sintered materials was measured using the twoprobe method [167] for the temperature range 30 – 500 oC. Silver electroded
specimens (3 × 3 × 3 mm) were characterised (equation 3.3) with a high sensitivity
digital micro-ohmmetre (Keithley 580).

l
R A
Equation 3.3
Where,  = electrical conductivity or specific conductance, l = sample thickness (3
mm), R = electrical resistance and A = cross-sectional area (9 mm2). The samples
were held in a copper jig (figure 3.14a). A power supply (Keithley 2602) and digital
multimeter (Keithley 6517A) were used to measure the current-voltage (I-V)
characteristics of the samples. To measure electrical conductivity at high temperature,
a specially designed alumina chamber in a tubular furnace was employed (figure
3.14b). Platinum wires and electrodes were selected as contacts. A heating rate of 2
o
C/ minute was selected and the temperature was measured using a K-type
thermocouple positioned next to the sample.
80
Connecting to
micro ohm-metre
Copper connects
(a)
Alligator gold
connects
Spring for uniform
contact pressure
Silver electroded
sample
Polymer (insulating)
base
(b)
Vertical tube furnace
Silver electroded
ceramic plate
Support for hollow
alumina tube
chamber
Thermocouple
measurement
Silver electroded
sample (placed next to
thermocouple)
Platinum wires
(protected in
alumina tube)
Micro ohm-metre
Figure 3.14. Electrical conductivity measurement setup: (a) room temperature jig; and
(b) high-temperature characterisation chamber in the tubular furnace.
81
3.3.5. Vickers indentation
Due to its simplicity, its non-destructive nature, and the fact that minimal machining
is required to prepare the sample, the use of the Vickers indentation method to
quantify toughness has become quite popular [279]. A diamond indenter was applied
to the surface of the specimens. Upon removal, the impression of indent was used for
the quantification of hardness. The length of the radial cracks (figure 3.15) reflects of
the crack toughness of the material which can be used to calculate the toughness of
the material by semi-empirical formulation. Load of 2.5 kg was used for a duration
time of 5 seconds. Vickers hardness was evaluated in accordance with ASTM C132703 [280].
H  0.0018544
P
d2
Equation 3.4
Figure 3.15. A typical Vickers indent. Modified from [281].
82
Where, H = hardness, P = indentation load and d = average length of indentation
diagonals (figure 3.15). For indentation fracture toughness Anstis’s equation [178]
was used:
 P  E
IFT  0.016  3/ 2 
c  H
Equation 3.5
Where, IFT = indentation fracture toughness and c = half of the mean radial crack
length (figure 3.15).
3.3.6. Microscopic analysis
3.3.6a. Optical Microscopy
For microstructural characterisations in Chapter 4, 5 and 9, optical microscopy was
performed using an Olympus BX60F fitted with a live camera assembly.
3.3.6b. Field-Emission Scanning Electron Microscopy (FE-SEM)
FE-SEM was used extensively (chapter 3-10) in this study. Two different FE-SEMs
were used, a JEOL (JSM-6300, 20 kV) and FEI (Inspect F, 20 kV). All of the
powders, fractured and polished surfaces were gold coated prior to SEM examination.
3.3.6c. High Resolution transmission Electron Microscopy HR-TEM
In Chapter 6 and chapter 7, HR-TEM (JEOL 2010, 200 kV) is used to study the
survivability of CNTs after SPS. The electron transparent nanocomposite films
(thickness <100 nm) were prepared by mechanical grinding, polishing, dimpling and
focused ion milling.
83
3.3.7. Thermo gravimetric analysis (TGA)
TGA was performed in chapter 4 and 8 using TA Instruments SDT Q600 TGA
thermo gravimetric analyzer. All specimens were examined on platinum pans in the
range 30 – 1000 oC. A heating rate of 5 oC/ minute in flowing air (at 180 ml/ minute)
was used. Powder sample masses ranged from 30 – 40 mg, whereas sintered sample
masses ranged from 30 – 50 mg.
84
Chapter 3. Homogenisation of CNTs in Ceramics
Chapter 4.
Homogenisation of CNTs in Ceramics
4.1. Introduction
A critical step in the processing of ceramic – CNT nanocomposites is the preparation
of a suspension of homogeneously isolated CNTs that can be added to different
ceramic powders to make nanocomposites. Ultrasonication (figure 4.1) in solvents is a
common primary step, and high power bath ultrasonication has been shown to be one
of the best methods for producing homogeneous and relatively aggregate-free
dispersions [95]. The dispersant used needs to overcome the strong van der Waals
force between CNTs and then resist their re-agglomeration [95].
(a)
(b)
(c)
Metallic
chamber
Wave
generators
Dispersion
bottle
De-ionised
water bath
Figure 4.1. Ultrasonication bath: (a) Apparatus with dispersion bottle; (b) shock
waves in distilled water, top view of the apparatus; and (c) schematic diagram.
The best solvents reported for generating CNT dispersions are amides, particularly
N,N-Dimethylformamide (DMF) and N-methylpyrrolidone (NMP) [282-284]. All of
85
these solvents are characterized by high values for β (electron pair donicity),
negligible values for α (the hydrogen bond donation parameter of Taft and Kamlet),
and high values for π (solvochromic parameter) [284, 285]. Thus, Lewis basicity (i.e.,
the availability of a free electron pair) without hydrogen donors is key to the good
dispersion of CNTs [284, 286]. However, this seems to be a necessary but not
complete set of conditions, as Dimethyl Sulfoxide (DMSO), a mediocre solvent,
meets these criteria [284, 287]. Ham et al. [288] illustrated that solvents with high
values of dispersion component (δd) of the Hildebrand solubility parameter (δt) are
the best for making homogeneous and agglomerate-free dispersions of CNTs. In this
regard, DMF (δd = 17.4 MPa1/2) proves to be better than ethanol (δd = 15.8 MPa1/2),
water (δd = 15.6 MPa1/2), acetone (δd = 15.5 MPa1/2) and methanol (δd = 15.1 MPa1/2)
for making CNT dispersions [288]. Other approaches to make stable dispersions are
the use of surfactants [289-291], acid treatments [123, 286, 290], and chemical
functionalisation [32, 43, 292, 293], which change the surface energy of CNTs,
improving their adhesion/wetting characteristics and reducing their tendency to
agglomerate in the solvents [95]. However, in all these approaches, the selection of
the solvent still remains a very important factor.
Despite the significant differences in the chemical properties of various solvents [284,
285, 288], many authors [111, 112, 114, 119, 129, 137, 140, 153, 181] have
repeatedly employed ethanol for dispersing CNTs in different ceramics. This seems
primarily due to the fact that alcohols are a common media for ball milling of
ceramics. Wang et al. [127] compared the use of methanol and DMF for the
dispersion of CNTs in an alumina matrix, and reported that the choice of dispersant
made no difference to the nanocomposites in terms of their densities and
microstructures. They did not describe the processing details for their ultrasonic
agitation, so there is insufficient evidence to assess their observations. The
effectiveness of a dispersion route depends upon various factors like solvent
properties, bath properties, energy applied, solution concentration, geometry of the
vessel, and the vessel position [95].
86
A significant amount of work has been done on the dispersion of CNTs in solvents.
DMF (appendix A) is well recognized among polymer researchers [31, 92, 294-297]
as a good dispersant for processing CNT – based nanocomposites. Lau et al. [294]
reported solvent effects in the order of DMF > ethanol > acetone for making polymer
– CNT nanocomposites. After synthesizing, Moniruzzaman et al. [295] stored CNTs
in DMF to avoid agglomeration and later also employed DMF to disperse their CNTs
in epoxy nanocomposites. Ciselli [92] showed smaller agglomerate sizes for CNTs in
DMF solution as compared to other solvents. The use of DMF as a dispersant by the
ceramic community making CNT nanocomposites is very rare [127] and not fully
realized.
In this chapter, we study the agglomeration and re-aggregation behaviour of coated
and uncoated CNTs, and carbon black in ethanol and DMF. This chapter also
compares the use of ethanol and DMF for making alumina – CNT nanocomposites by
analyzing the pre-sintering (colloidal stability and agglomerate size analysis) and
post-sintering (dispersion profile and electrical conductivity measurements) stages.
4.2. Experimental procedure
4.2.1. Colloidal dispersions and characterisations
To monitor the colloidal stability, a 77 mg/l concentration of CNTs in DMF was hand
mixed for 15 seconds and high power bath ultrasonicated (Engisonic plus, Engis Ltd.,
UK) for 1 hour. It was then hand shaken for another 5 minutes to remove any
gradients produced by non-uniform ultrasonication. The dispersion was then placed in
front of a luminescent light box to observe its re-aggregation behaviour. For
agglomerate size analysis, CNTs were hand mixed for 15 seconds in DMF solution
(100 mg/l) and high power bath ultrasonicated for different durations. The solutions
were then transferred to standard glass cuvettes (10 mm × 10 mm × 45 mm) and
placed in a Malvern Zetasizer nanoparticle size analyser (Nano ZS). The software was
programmed to record the average of at least 30 readings for the quantification of the
87
agglomerates’ size and distribution. Using the same procedure as above, alumina
dispersion in DMF was also characterized (150 mg/l). All of the above procedures
were then repeated with ethanol to allow a comparison between the behaviour of
DMF and ethanol.
Re-agglomeration behaviour of CNTs was studied systematically using CNTs
supplied by NanoDynamics, USA. To study the re-agglomeration behaviour, a 1.75
g/l concentration of CNTs in DMF was hand mixed for 15 seconds and high power
bath ultrasonicated in a glass cuvette for 30 minutes. The solution was then placed in
a nanoparticle size analyzer. The software was programmed to record the average of
at least 30 readings. The re-agglomeration behaviour was studied for 20 minutes. The
same experiment was conducted for alumina coated CNTs (25 cycles and 50 cycles).
All of the above procedures were then repeated with ethanol to allow a comparison
between the behaviour of DMF and ethanol.
Alumina – 2 wt% (~4.7 vol%) CNT nanocomposites were prepared. The CNTs were
dispersed in DMF via high power bath sonication for 2 h and then hand mixed with
the alumina nanopowder for another 5 minutes. The liquid mixture was transferred to
another jar filled with zirconia balls (milling media) of two different sizes (10 and 5
mm, mass ratio: 3:2). The jar was sealed and rotation ball milled for 8 h at ~200 rpm.
The milled powder was then shifted through a steel pan. The milled slurry mixture
was dried at 75 oC for 12 hours on a heating plate and then transferred to a vacuum
oven (100 oC) for 3 days for complete removal of the dispersant. A solvent trap (filled
with ice) was connected between the vacuum pump and the oven.
The same
procedure was followed for making alumina – alumina coated CNTs (50 cycles)
nanocomposite powder and alumina – carbon black nanocomposite powder.
88
The dried mixture (alumina – CNT nanocomposite powder) was ground and sieved
using a 250 mesh and then returned to the vacuum oven for another 4 days at the same
temperature for thorough extraction of the solvent. This lengthy drying procedure was
followed because any residual solvent has a detrimental effect on the properties of
CNT-reinforced nanocomposites [287, 288]. The same method was employed to make
nanocomposite powder using ethanol. A brief summary of the dispersion process is
given in figure 4.2.
Mixing solvent with CNTs
Ultrasonicating mixture for 2 hours
Mixing CNT dispersed mixture with ceramic powder
Ball milling mixture for 8 hours
Drying milled mixture on hot-plate for 12 hours
Vacuum drying powder mixture for 7 days
Nanocomposite powder ready for Spark Plasma Sintering
Figure 4.2. Pre-sintering processing of ceramic – CNT nanocomposite powder.
Dried nanocomposite powder (~2 g) was poured into a carbon die and cold pressed at
0.62 MPa for 5 s before sintering. Nanocomposite discs (thickness 2 mm and diameter
89
20 mm) were prepared by SPS. A pressure of 100 MPa was applied concurrently with
the heating (rate 300 oC/ minute) and released at the end of the sintering time, which
was 3 minutes for all of the samples. The sintering temperatures were in the range
1200 – 1950 oC. All of the samples were slowly cooled (~50 oC/ minute) to avoid
fracture due to thermal shocks and differential contractions. The same sintering
procedure was repeated for the dried nanocomposite powder dispersed using ethanol.
The SPSed samples were ground using SiC paper and diamond polished down to 1
μm. Density measurements were conducted using helium pycnometer (AccuPyc 1330,
Micrometics) and water buoyancy methods. There were no significant differences in
the results from both techniques, so the mean was used to characterize the density of
the SPSed samples. SPSed samples were fractured in order to observe the
agglomeration and dispersion of CNTs. Nanocomposite powder (alumina – CNT) and
the fractured surfaces were gold coated and observed in a field emission scanning
electron microscope (FE-SEM). The electrical conductivities of the samples were
measured (section 3.3.4) with a high sensitivity digital micro-ohmmetre (Keithley
580) using the two-point method on silver electroded specimens (3 mm × 3 mm× 3
mm) prepared using a diamond cutting machine.
4.3.1. Natural drying patterns
Before sieving, the natural drying patterns (figure 4.3) of the nanocomposite powder
hinted at the strength of the secondary bonding of nano carbon fillers. Because of high
entanglement and aspect ratio of CNTs, they formed centimetre-sized agglomerates
with alumina (figure 4.3a). It appeared that coating CNTs with alumina decreased the
van der forces between CNTs, as they did not form entangled networks (figure 3b).
90
The same was also observed for alumina – carbon black nanocomposite powders
because of the non-fibrous nature of carbon black (figure 3c and 3d). This is another
qualitative assessment that highlights the strong entanglement in CNTs because of
their fibrous nature. A good dispersion was observed after sieving alumina – 2 wt%
CNT (figure 4.4).
(a)
(b)
(c)
(d)
Figure 4.3. Natural drying patters formed during processing of nanocomposite
powder: (a) alumina – 2 wt% CNT; (b) alumina – alumina coated 2 wt% CNT. CNTs
were supplied by NanoDynamics, USA; (c) alumina – 2 wt% carbon black (Vulcan
XC72); and (d) alumina – 2 wt% carbon black (Printex L6).
91
Figure 4.4. Alumina – 2 wt% CNT nanocomposite powder after sieving.
Individual CNT can be seen. CNTs were supplied by Nanocyl, Belgium.
4.3.2. Agglomerate size analysis
The presence of agglomerates in powders with very fine grain size contributes to
grain coarsening during sintering and produces non-uniformity in the resulting
microstructure [261]. Agglomeration is particularly significant in CVD-grown
nanotubes because substantial entanglement of the tubes occurs during nanotube
synthesis [167]. CNTs and alumina powder were separately ultrasonicated for various
durations and their agglomerate sizes were then measured immediately (figure 4.5).
For all of the ultrasonication durations, DMF disperses CNTs more efficiently as
compared to ethanol, by reducing the agglomerate size (figure 4.5a). It should be
noted that these results are presented for comparison and do not represent the
optimum conditions for dispersing the CNTs. DMF also showed better dispersion of
alumina compared to ethanol (figure 4.5b). All these observations related to colloidal
dispersions can be explained by the higher values of Hildebrand solubility parameter
(δt) and Lewis basicity of DMF as compared to ethanol [284, 285, 288]. Looking at
the lengthy processing procedures followed by many researchers [38, 110, 111, 114,
92
153, 294] before sintering (for ceramics) or curing (for polymers), the fact cannot be
ignored that the efficiency of DMF for de-bundling and making stable CNT
dispersion is much better as compared to that of ethanol.
(b)
(a)
Figure 4.5. Agglomerate size analysis with respect to ultrasonication time in different
solvents: (a) CNTs, supplied by Nanocyl, Belgium; and (b) alumina.
4.3.3. Re-agglomeration behaviour
Successful fabrication of nanocomposites depends crucially on maintaining stable
colloidal mixtures of the nanotubes and matrix phase [298] before ceramic sintering
or polymer curing. Figure 4.6 shows the colloidal dispersion for CNTs in DMF and
ethanol after ultrasonication and at different time intervals. The DMF dispersion is
very stable, showing no signs of agglomeration even after several months (figure 4.6).
These observations are consistent with previous work that showed that CNT – DMF
dispersions aggregate on a timescale of days [283] and weeks [299]. The CNT –
ethanol dispersion
re-agglomerated
significantly
within
half
an
hour
of
ultrasonication. This qualitative analysis shows that CNTs are much more stable in
DMF as compared to ethanol.
93
DMF
Ethanol
Soon after ultrasonication
After 24 hours
After 15 months
Figure 4.6. Colloidal dispersion stability comparison after 1 h bath sonication and 5
minutes hand shaking. The diameter of the bottles is 25 mm. CNTs were supplied by
Nanocyl, Belgium.
To quantify such observation, re-agglomeration with the passage of time was
observed for all types of CNTs in a nano particle size analyzer. The slopes (figure 4.7)
indicated the rate of re-agglomeration of the CNTs in ethanol and DMF. The higher
the magnitude of slope, the faster the re-agglomeration and vice versa [300]. At time
= 0 minutes, it appears that the dispersion process employed was not appropriate for
all samples because of the presence of large agglomerates. As mentioned before, these
results are for comparison and do not represent the optimum conditions for dispersing
the CNTs.
94
(a)
(b)
(c)
(d)
Figure 4.7. Re-agglomeration behaviour in different solvents after 30 minutes of
ultrasonication. CNTs were supplied by NanoDynamics, USA: (a,b) pristine CNT; (c)
alumina coated CNT (25 ALD cycles); and (d) alumina coated CNTs (50 ALD
cycles).
Comparing the slopes, faster re-agglomeration was observed in ethanol as compared
to DMF for all types of CNTs. Higher colloidal stability is only possible if the CNTs
have a charge on their surface preventing aggregation after dispersion by repulsive
electrostatic forces [301]. The lower charge density of the nanotubes dispersed by
ethanol is responsible for their lower stability in ethanol compared to DMF. DMF (pH
= 9) is also more basic as compared to ethanol (pH = 6), which is another important
factor for its better dispersion properties [286]. Other important factors are the high
values of Hildebrand solubility parameter (δt) and Lewis basicity of DMF as
compared to ethanol [284, 285, 288]. It was also observed that coating CNTs with
alumina reduced the re-agglomeration rate. The larger the thickness of alumina
coating, the slower the re-agglomeration. This is attributed to weaker van der Waals
attraction between coated CNTs.
95
4.3.4. Microstructure of nanocomposites
Samples SPSed for this chapter are summarised in table 4.1.
Table 4.1. Samples SPSed for this chapter#.
#
Rel. theoretical
Matrix
Weight %
Dispersant
SPS conditions
Alumina
2
Ethanol
1200 oC/ 100 MPa/ 3 minutes
~77
Alumina
2
Ethanol
~80
Alumina
2
Ethanol
~87
Alumina
2
Ethanol
~96
Alumina
2
Ethanol
~100
Alumina
2
Ethanol
-
Alumina
2
DMF
~78
Alumina
2
DMF
~87
Alumina
2
DMF
~97
Alumina
2
DMF
~100
Alumina
2
DMF
-
density (%)
All CNTs were supplied by Nanocyl, Belgium. For CNTs, 2 wt% = ~4.7 vol%
96
Representative images of the fractured surfaces of the SPSed nanocomposites were
selected for studying the distribution of the CNTs (figure 4.8). Bright micron-sized
agglomerates of CNTs on alumina grains (100-300 nm) are visible in the sample
prepared using ethanol as the dispersant (figure 4.8a). These agglomerates were
produced because of the inability of ethanol to make an agglomerate-free dispersion
of CNTs before sintering. The sample prepared using DMF as the dispersant has a
homogeneous distribution of individual CNTs (figure 4.8b). It is interesting to note
the different grain sizes of the nanocomposites produced with ethanol and DMF.
Those prepared with ethanol have a noticeably larger grain size (figure 4.8a)
compared to the equivalent nanocomposite prepared with DMF (figure 4.8b). The
effect of CNTs on the grain size refinement of nanocomposites is the subject of
Chapter 5.
(a)
(b)
Figure 4.8. Fractured surfaces of sintered alumina – 2 wt% CNT samples, SPSed at
1200 ◦C/100 MPa/ 3 minutes: (a) CNTs dispersed in ethanol and (b) CNTs dispersed
in DMF.
4.3.5. Mechanical and electrical properties
The presence of agglomerates in nanocomposites is property limiting. They reduce the
mechanical properties [302, 303] and electrical properties [303, 304] of the
nanocomposites. Figure 4.9 shows the Vickers hardness and indentation fracture
97
toughness for alumina – 2 wt% nanocomposites prepared using different solvents. The
densities and electrical conductivities of the nanocomposites are shown in figure 4.10.
The SPS processing of the colloidally dispersed starting powder mixtures produced
nanocomposites of high density with well-distributed CNTs. There was no major
difference between the densities of the SPSed samples prepared from the different
dried nanocomposite powders. However, nanocomposite prepared using ethanol
showed inferior Vickers hardness and indentation fracture toughness as compared to
nanocomposite prepared using DMF. Agglomeration favours an inhomogeneous
densification and thus inhomogeneous grain size distribution [177]. Agglomerates
contain very fine porosity that is responsible for easy crushing under load resulting
poor hardness. The toughness improving mechanisms like fibre pull-out, crack
deflection and crack bridging are not possible in the presence of agglomerates of
CNTs. This resulted in poor indentation fracture toughness of nanocomposite
prepared using ethanol as compared to the nanocomposite prepared using DMF
(figure 4.9).
Figure 4.9. Vickers hardness and indentation fracture toughness of alumina – 2 wt%
CNT nanocomposites prepared using different solvents.
98
Samples
melted
Figure 4.10. Density and electrical conductivity of alumina – 2 wt% CNT.
Error bars are not marked in the electrical conductivity measurements as they are very
small (figure 4.10). There is a significant difference in the electrical conductivities of
the SPSed samples prepared from ethanol and DMF dispersions, particularly as the
density of the nanocomposites increases. This is due to a better homogeneous
dispersion of the conductive CNTs in DMF solution as compared to ethanol. Alumina
is inherently an insulator (electrical conductivity: 10−13 S/m [211]), so a uniform
distribution of highly conductive CNTs is critical for making the nanocomposites
good electrical conductors and reducing the percolation threshold.
4.4. Conclusions
A prerequisite for the ceramic nanocomposites with good electro-mechanical
properties is the homogeneous dispersion and distribution of the CNTs in the ceramic
99
matrices. The extraordinary high specific surface area of CNTs results in very high
van der Waals forces between them, inducing a strong tendency to agglomerate. The
selection of the ultrasonication medium is very important for the final properties of
the nanocomposite. Non-hydrogen bonding Lewis bases are the best solvents for CNT
dispersions.
From
dispersion
stability
observations
and
agglomerate
size
measurements, it is clear that DMF produces fine and stable CNT and alumina
dispersions. Faster re-agglomeration was observed in ethanol as compared to DMF for
pristine and coated CNTs. Coating CNTs with alumina reduced the re-agglomeration
rate. The larger the thickness of alumina coating, the slower the re-agglomeration. No
evidence of agglomeration and a good distribution of the CNTs was observed in FESEM micrographs of the SPSed samples when they were mixed with alumina in
DMF. Nanocomposites prepared using DMF dispersions showed better dispersions
and higher electrical conductivity as compared to those prepared using ethanol
dispersions. Therefore, it is concluded that DMF is a good dispersant for making
homogeneous and agglomerate-free slurries by any type of colloidal processing.
100
Chapter 5. Sintering of Ceramic – CNT nanocomposites
Chapter 5.
Sintering of Ceramic – CNT Nanocomposites
5.1. Introduction
Control of microstructure and improvement in densification is one of the objectives of
using dopants/ reinforcements in ceramics. Improved mechanical [127, 140, 200],
electrical [129] and thermal properties [226] have also been reported for ceramic –
CNT nanocomposites produced by rapid processing using SPS. However, the role of
CNTs in the sintering of ceramics is not clear in the literature, possibly because of
differences in the materials and, particularly, dispersion and mixing of the CNTs.
An and Lim [305] suggested that CNTs in alumina decreased mass transportation
during sintering, which inhibited the densification process. Tatami et al. [129] and
Jiang and Gao [226] reported inhibitation of densification of silicon nitride – CNT
and titanium nitride – CNT nanocomposites. For glass-ceramics, Boccaccini et al.
[116, 306] and Ning et al. [119] observed that the presence of MWNTs in a glass
matrix hindered the densification of the material; the CNTs served as nucleation
points, and the crystallized phases acted as a rigid body, which hindered densification.
However, Morisada et al. [103] and Wei et al. [110] reported that the addition of
CNTs to silicon carbide and alumina had no effect on the densification behaviour of
the nanocomposites. These negative and nil effects are possibly because of differences
in the materials studied, dispersion and mixing of the CNTs, and the presence of
agglomerates. Guo et al. [141] reported that CNTs improved the densification and
101
Chapter 5. Sintering of Ceramic – CNT Nanocomposites
mechanical properties of silica nanocomposite. Huang et al. [222] attributed this
effect to the good combination of SPS sintering and the high electrical and thermal
conductivity of CNTs. In some studies [110, 114, 140, 147, 158, 161, 222] it was
reported that the addition of CNTs modified the grain size. Zhan et al. [140] prepared,
using SPS, 100% dense alumina and alumina – 10 vol% CNTs at the same processing
temperature (1150 oC); and the grain size of the alumina in the nanocomposite was
~39% smaller as compared to in the alumina. Even carbon nanofibres (diameter 100200 nm) were found to retard alumina grain growth [137]. However, there has been
no systematic investigation into the effect of CNTs on the sintering behaviour and
grain growth of ceramics.
It is well known that carbon black is one of the best sintering aids for various
ceramics [307-310], and very little addition is required as compared to other additives
[311, 312]. Erkalfa et al. [312] reported that due to the self-lubricating nature of
carbon, it enhances compactibility and compressibility, which favours densification.
The increased densification of non-oxide ceramics is possibly due to the removal of
surface oxide layer on the powder by the addition of carbon, which increases surface
diffusion during sintering [227, 313].
In this chapter, the sintering and grain growth behaviour of alumina – CNTs and
alumina – carbon black nanocomposites, alumina prepared by Spark Plasma Sintering
(SPS) were studied. The influence of CNT addition on the sintering and grain growth
of reduced titanium dioxide was also studied. In the last section of this chapter, the
effect of residual impurities (left during the synthesis of CNTs) on the sintering
behaviour of alumina – CNT nanocomposites is analysed.
102
5.2.1. CNTs purification
As per supplier (Nanocyl, Belgium), CNTs were >90% pure. An acid treatment was
performed using a mixture of concentrated nitric (HNO3, 90%) and sulfuric (H2SO4,
90%) acids. Distilled water (~20 vol%) was used to dilute the acids. In order to
produce pure CNTs, the as-received CNTs (400 mg) were mixed with 200 ml dilute
acidic solution. Both acids were equally mixed in the solution. The acid-CNT mixture
was homogenized by stirring with a glass rod on heating plate (~85 oC) for 30 minutes
and then bath ultrasonicated for 2 hours. The resulting CNT dispersion was
thoroughly washed with distilled water until the filtrate was colourless and neutral
(pH ~7) after filtration. A Whatman filter paper of 1 μm was used. The purified CNTs
were then dried for 48 hours at 100 oC in an oven. The quality of CNTs was
quantified by thermo gravimetric analysis (TGA). A short summary of the purification
process is shown in figure 5.1.
Mixing CNTs with dilute acid mixture
Mixing on hot-plate (~110 oC) for 30 minutes
Bath Ultrasonicating for 2 hours
Filtering with ultra fine filter paper
Cleaning with distilled water until pH ~ neutral
Drying CNTs at 100 oC for 48 hours
Purified CNTs
Figure 5.1. Purification of CNTs by acid treatment.
103
In this section, composite powders were prepared using different dispersion routes.
Those prepared by a colloidal dispersion method (using DMF, section 4.2.1) gave the
best results. This method will be referred to as “optimum method of dispersion” in
this study. To compare the role of dispersion on the sintering behaviour, different
routes were adopted to prepare in-homogenous dispersions of CNTs. The properties
of the ceramic nanocomposites prepared by the different routes were then compared.
1. Viscous solution mixing: The same process was followed as in section 4.2.1, with
the exception of using 1 gram of CNTs in 150 ml of DMF. In “optimum method of
dispersion”, 400 mg of CNTs in 150 ml of DMF was used.
2. Hand mixing: The same process was followed as in section 4.2.1, with the
exception of doing hand mixing of CNTs in DMF in place of ultrasonication and ball
milling. The hand mixing was done for 30 minutes.
Ceramic and nanocomposite pellets (diameter 20 mm and thickness 2 mm) were
prepared by SPS. A pressure of 100 MPa was applied concurrently with the heating
(rate 300 oC/ minute) and released at the end of the sintering period (3 minutes) for all
samples. An alumina / alumina – CNTs / alumina laminate sample was prepared by
subsequent compacting of alumina nanopowder, nanocomposite powder and alumina
nanopowder in a graphite die. The powder compact was cold pressed at 0.64 MPa for
a few seconds and co-sintered at 1800 oC for 3 minutes under a pressure of 85 MPa.
104
All of the sintered samples were ground using SiC paper down to 4000 grit. The
density of the ground samples was measured using the water buoyancy method.
Selected samples were thermally etched at 250 oC less than the SPS temperature for
10 minutes. Field emission scanning electron microscopy (FE-SEM) was used to
observe fractured surfaces in order to determine the grain sizes. The polished laminate
and fractured surfaces were coated with a very thin layer of gold and transferred to an
FE-SEM for examination. Grain sizes were measured with the aid of software (Image
tool for Windows, version 3.00, developed by UTSHCSA, USA). A minimum of 200
readings was taken to measure the grain sizes of each material.
With the aid of user-friendly, in-built touch screen, SPS can be monitored very
intelligently during sintering. Whenever a material shrinks, sinters or melts the speed
of the moving punch increases noticeably at that instant. This observation is
reproducible and a very clever way of knowing the important events during sintering
of unknown and new materials.
The bulk initial phase of alumina was γ. During SPS, the phase transformation (γ to α)
in alumina and alumina – CNT nanocomposite was observed distinctly (figure 5.2) at
~1200 oC. Such phase transition is reported to result in the formation of interlocking
vermicular structure, which is detrimental to densification [272].
105
Figure 5.2. Moving punch speed as the function of temperature during SPS for
alumina and alumina – 5 wt% (~11.2 vol%) CNT (uncoated) nanocomposite. CNTs
were supplied by Nanocyl, Belgium. Inset shows SPS program details.
At 1600-1700 oC, there is another peak in the speed of the moving punch. This relates
to the bulk sintering or liquid phase sintering (figure 5.2). In both peaks, CNTs
enhanced the compressibility and compactibility of the nanocomposite, which is
evident from the higher speed of the moving punch as compared to the lower speed in
alumina (figure 5.2). However, when CNTs are encased within the alumina coating,
the speed at these temperatures (~1200 oC and 1650 oC) was reduced and the events
were delayed as well (figure 5.3). This illustrates that uncoated CNTs have a very
clear effect on the sintering of alumina – CNT nanocomposites.
106
Figure 5.3. Moving punch speed as the function of temperature during SPS for
alumina – 5 wt% CNT (uncoated) and alumina – 5 wt% CNT (coated, 50 ALD
cycles) nanocomposite. Heating rate: 300 oC/minute. CNTs were supplied by Arkema,
France.
5.3.2. Sintering behaviour and possible mechanisms
Figure 5.4 illustrates the rel. theoretical density versus sintering temperature of
ceramic – CNT nanocomposites SPS processed for 3 minutes at their sintering
temperature at a pressure of 100 MPa. The addition of CNTs significantly reduces the
sintering temperature required to achieve full densification of the nanocomposites as
compared to pristine ceramics (figure 5.4). For example, at 1200 oC, the alumina – 5
wt% CNT has 100% rel. theoretical density, while the alumina has 59% rel.
theoretical density.
107
(a)
(b)
Figure 5.4. Rel. theoretical density as a function of sintering temperature for: (a)
alumina and nanocomposites; and (b) reduced titania and nanocomposites.
108
For comparison, the sintering behaviour of alumina – 2 wt% carbon black
nanocomposites prepared by the same processing route were also investigated. The
effect of CNTs and carbon black, for the same carbon content, on densification
behaviour is very similar (figure 5.4a).
Homogenous dispersion of CNTs is crucial for good densification of CNT
nanocomposites. The increased density of the nanocomposite with increasing CNT
content (figure 5.4) suggests that the CNTs were well dispersed, which is not true for
nanocomposites prepared using in-homogenous powder mixtures (figure 5.5).
Figure 5.5. Bulk density of the sintered product as the function of homogeneity of
CNTs in alumina matrix. SPS conditions: 100 MPa/ 3 minutes. CNT content: 5 wt%.
In electrically insulating powders, like alumina, the current only flows through the
graphite dies and punches during SPSing (figure 5.6a and figure 5.6b). Huang et al.
[222] speculated that CNTs promote densification of ceramic powders due to the flow
of electrical current through the conductive powder (figure 5.6c) as well as graphite
dies and punches during SPS. The high pulse current in SPS is preferentially
transported through CNTs due to their low resistance, which locally increases the
109
temperature near CNTs. In this manner, CNTs become the dominant heat generator in
the composites [222]. However, no evidence of such phenomenon was presented.
Table 5.1 shows the electrical resistance of graphite die set and its contents. It is
confirmed that as opposed to alumina compact, current flows through alumina – CNT
nanocomposite compact during SPS. Electrical current flows through CNTs that
promotes local Joules heating in the powder compact.
(a)
(b)
(c)
Figure 5.6. Schematics of graphite die sets: (a) alumina before SPS; (b) current
passing through graphite only during SPS; and (c) current passing through graphite
and alumina – CNT nanocomposite compact during SPS.
Table 5.1. Electrical conductivity measurements of graphite die sets.
Graphite die set
Electrical resistance† (mΩ)
Comments
Empty die
2.85 ± 0.3
-
Die + CNTs
2.97 ± 0.4
-
Die + air cavity
3.81 ± 0.1
-
Die + alumina powder
3.72 ± 0.8
Before/ after sintering
Die + alumina – 5 wt% CNT
3.31 ± 0.3
Before sintering
Die + alumina – 5 wt% CNT
3.02 ± 0.4
After sintering
† All die sets contain same volume of powder/ cavity (0.25 cm3) and pressed to 1 ton load.
110
Figure 5.7 shows the displacement of the rams during cold pressing (loading and
unloading cycle) for alumina and alumina – 5 wt% CNT powders. Subtracting the
effect of the compliance of the SPS loading train and the die set, the alumina – 5 wt%
CNT powder has better compactibility as compared to alumina powder. The green
relative density for the alumina – 5 wt% CNT compact is 47.6%. It was not possible
to measure the same for alumina, as the compact did not stay intact after removal
from the die. The improved sinterability of alumina with carbon additions is partly
due to improved self-lubricating properties, which promotes compactibility and
compressibility of the nanocomposite powder.
Figure 5.7. Compressibility and compactibility analysis performed at room
temperature by uni-axial pressing.
111
5.3.3. Grain growth modification
The density versus grain size dependence of alumina, alumina – 2 wt% carbon black
and alumina – 2 wt% CNT nanocomposites are shown in Figure 5.8. This shows that
it is possible to produce completely dense ceramics with sub-micrometer grain
structures by the addition of CNTs. The grain growth during densification is significantly
less in the alumina – 2 wt% CNT and alumina – 2 wt% carbon black nanocomposites,
compared to the alumina even though they are dense and the alumina is not.
Figure 5.8. Rel. theoretical density as a function of grain size for alumina and
nanocomposites.
Figure 5.9 shows the microstructure of alumina (figure 5.9a) and the nanocomposites
(figure 5.9b and c) sintered under the same conditions of 1800 C for 3 minutes. The
CNTs are located at the grain boundary (figure 5.9c), which has been previously
112
reported [195]. Comparing the microstructures of the monolithic alumina and the
nanocomposites, the addition of carbon black and CNTs retards grain growth. The
CNTs have a greater grain size retardation effect than the carbon black, which may be
attributed to the different geometric contributions of CNTs and carbon black powder.
(a)
(b)
(c)
(d)
Figure 5.9. FE-SEM images of fractured surfaces of sintered nanocomposites
processed at 1800 oC/ 100 MPa/ 3 minutes: (a) alumina; (b) alumina – 2 wt% carbon
black (Printex L6); (c) alumina – 2 wt% CNT (Nanocyl, Belgium); and (d) alumina –
2 wt% alumina coated CNT (50 ALD cycles, NanoDynamics, USA).
The grain sizes versus sintering temperature are shown in figure 5.10. There is a very
large difference in the grain size of the alumina and the alumina – 5 wt% CNT
nanocomposite (figure 5.10a). Figure 5.10b shows the comparison of the effect of
CNTs and carbon black, and CNT content on grain size. The carbon black, like the
alumina, shows an exponential dependence of grain size with temperature (same
sintering time of 3 minutes). An parabolic dependence of the grain size with
113
temperature is reported for monolithic ceramics [279], particulate composites [314317] and doped ceramics [266, 318-320] The alumina – CNT nanocomposites show a
decreasing rate of grain growth with increasing CNT content (figure 5.10a). The 5
wt% CNT nanocomposites show a nearly linear dependence (figure 5.10b), which
suggests a different type of mechanism for grain growth. The CNTs form a strong
entangled network around the grains, which appears to constrain the grain growth.
This effect was not observed when CNTs were coated with alumina. There was no
web of CNTs, therefore large grains are very clear in figure 5.9d.
Figure 5.10. Grain size refinement effect of: (a) CNTs; and (b) CNTs and CB.
In order to observe the grain growth retardation after full densification, lengthy dwell
times were used (figure 5.11). The grain size strongly depends on the density of
sintered samples [321], therefore, the comparison of grain growth between the two
materials should be determined at the same density. In general, grain growth law for
polycrystalline materials during isothermal holding can be described using equation
5.1 and 5.2 [321].
D
G n  Gon   t
T 
Equation 5.1
 E 
D  Do exp 

 RT 
Equation 5.2
114
where G and Go are the grain sizes at holding time t=t and t=0, n is a constant related
to grain growth mechanism, D is diffusion coefficient related to the grains, and Do is
diffusion activation energy. The equations are applicable for the solid state sintering
of alumina [321]. Because of the decrease in atomic diffusion coefficient caused by
the presence of CNTs at the grain boundary, the presence of CNTs reduced the grain
growth by a factor of ~5.3 (figure 5.11).
Grain growth:
80 nm/minute
15 nm/minute
Figure 5.11. Grain size refinement in isothermal conditions (1400 oC/ 100 MPa).
5.3.4. Co-sintering of grain size laminate
In monolithic ceramics, a coarse-grained microstructure is desired for applications
requiring higher modulus [322], higher creep resistance [279], higher thermoelectric
properties (Hall mobility and figure of merit) [323], higher thermal conductivity [323,
324] higher electrical conductivity [323, 325] and higher optical transparency [279].
115
On the other hand, a very fine-grained microstructure is desired for applications
requiring higher strength [279], higher wear resistance [326], higher thermal shock
resistance [327], higher cyclic fatigue resistance [279], low dielectric loss [279] and
higher optical transparency [279]. Grain-size FGMs are advantageous for bio-medical
applications [328]. Morsi et al. [328] hot-pressed alumina with different particulate
sizes and produced grain-size FGMs with a difference of only ~2 times. Moreover,
there was cracking at the interface due to poor bonding. An alumina / 2 wt% CNT
nanocomposite laminate structure was fabricated in the current work to demonstrate
the grain refinement size effect of CNTs and their ability to produce grain size
laminated materials (figure 5.12).
(a)
(b)
Figure 5.12. FE-SEM images of grain size laminate showing interface between
alumina and 2 wt% CNT nanocomposite: (a) high magnification fractured surface;
and (b) low magnification polished surface.
It should be noted that the monolithic alumina layers and the nanocomposite layers
were co-sintered at the same temperature (1800 oC). Figure 5.12a shows a fractured
surface and figure 5.11b a polished and thermally etched surface. Both layers were
successfully co-sintered without any cracking at the interface (figure 5.12a). The
upper half part of figure 5.11b shows large grains of alumina (~ 20 μm), whereas the
other lower part of the image shows very fine grains (~ 2 μm) in the nanocomposite
layer. This image illustrates the potential of CNTs to control the microstructure of
ceramics.
116
An alumina/ alumina – 5 wt% CNT nanocomposites/ alumina laminate was also
SPSed (figure 5.13). Due to large difference in the grain size, there were cracks at the
interface (figure 5.13b). These cracks were developed during the slow cooling of the
laminate. Further optimisation of sintering process may avoid this cracking. However,
clear grain growth retardation effect was observed in these materials (figure 5.14).
(a)
(b)
Figure 5.13. FE-SEM images of polished grain size laminates: (a) alumina and 2 wt%
CNT nanocomposite; and (b) alumina and 5 wt% CNT nanocomposite showing
cracking at interface.
(a)
(b)
(c)
Figure 5.14. FE-SEM images of different areas of polished grain size laminates: (a)
alumina region, grain size: 43 μm; (b) 2 wt% CNT nanocomposite region, grain size:
1.81 μm; and (c) 5 wt% CNT nanocomposite region, grain size: 0.99 μm.
117
5.3.5. Sintering behaviour of alumina – purified CNTs
CNTs supplied by Nanocyl, Belgium were studied in this section. To the best of
authors’ knowledge, there is no comparative study for analyzing the effects of
residual impurities (left during the synthesis of CNTs) on the sintering behaviour of
ceramic – CNT nanocomposites. Figure 5.15 shows the oxidation behaviour of as
received-CNT and pure-CNTs.
94.7% loss
98.6% loss
Figure 5.15. Thermo gravimetric analysis (TGA) of CNTs, before and after acid
purification treatment. Heating rate: 5 oC/minute.
Thermo gravimetric analysis (TGA) is a good tool for quantifying the noncarbonaceous species e.g. catalytic metals and oxides (figure 5.16). The as receivedCNT were ~95 % pure, whereas purified CNTs were ~98.5% pure, at the cost of
decreased oxidation resistance.
118
(a)
(b)
(c)
Figure 5.16. Platinum pan (diameter: ~10 mm) used for TGA: (a) empty pan before
analysis, (b) as received-CNTs and (c) impurities left after oxidation.
No significant effect of residual impurities was observed on the densification of the
alumina – CNT nanocomposite (table 5.2) as appeared in the literature as well [155].
The CNT content used in this chapter was up to 5 wt%. Therefore, the catalytic
impurities and oxides have no influence on the densification because of the lower
content (<0.25 wt%) in the final product.
Table 5.2. Density measurements for CNTs and alumina – ceramic nanocomposites.
Rel. theoretical density (%) of alumina – 5 wt%
Material
Powder density
CNT nanocomposite (g/cm3)†
(g/cm3)*
SPSed at 1400 oC
SPSed at 1600 oC
As received CNTs
1.66 ± 0.4
99.7 ± 0.8
100 ± 0.6
Purified CNTs
1.59 ± 1.1
99.5 ± 0.5
100 ± 0.4
* Measured by Helium pycnometery.
† Measured by water buoyancy method.
119
5.4. Conclusions
The addition of CNTs or carbon black to alumina significantly increases its sintering
rate. The sintering temperature required to achieve full densification of alumina –
CNT nanocomposites was reduced by 500
o
C as compared to alumina. An
improvement in the densification was also observed in reduced titanium dioxide –
CNT nanocomposite by the addition of CNTs. Clear evidence is presented of the
effect of CNTs on grain growth. The CNTs, which form entangled networks at the
grain boundaries, produce significant grain growth retardation. Using this effect, an
alumina / nanocomposite laminate structure with a grain size difference of about ten
times was successfully co-sintered. However, the grain growth retardation was not
evident when alumina coated CNTs were employed in the alumina matrix. The effect
of residual impurities (left during synthesis) was not observed on the densification of
alumina – CNT nanocomposites.
CNTs should not only be considered as an additive for improving the properties of
ceramics because of their excellent intrinsic physical properties, but also as a means
of controlling their sintering behaviours and microstructures. This will allow materials
with improved and novel microstructures to be fabricated. This includes
nanoceramics, co-fired ceramic multilayered structures, and functionally gradient
materials.
120
Chapter 6.
Structural Stability of CNTs in Ceramics
(Preservation Studies)
6.1. Introduction
Significant improvements in the electrical and mechanical properties of polymer [31,
329] and ceramic [98, 111, 127, 140, 146, 179, 330] based CNT nanocomposites were
reported. CNTs are well preserved in polymers due to low processing temperatures
and pressures [31, 329]. However, to date, there is a lot of controversy about the
stability of CNTs during high temperature processing in ceramic composites [111,
112, 134, 146, 164, 179, 181, 330-333].
It is difficult to examine the exact amount of structural defects in CNTs. Raman
spectroscopy is a well-known tool to characterise graphitic carbons [334, 335].
Different types of Raman vibrations of CNT are illustrated in figure 6.1. The peak at
~100-400 cm–1 corresponds to a Radial Breathing Modes (RBMs). RBMs can provide
information about CNT chirality (i.e., (n,m) indices) and diameter distributions since
the RBM frequency is inversely proportional to nanotube diameter [181, 336].
However, RBMs signals (~100-400 cm–1) are very weak for thicker MWNTs. RBMs
cannot be used for quality assessment of MWNTs [335, 337]. The peak at ~13201350 cm–1 corresponds to a disorder–induced phonon mode (D band) of MWNTs, and
121
Chapter 6. Structural Stability of CNTs in Ceramics (Preservation Studies)
the peak at ~1500-1600 cm–1 can be assigned to the G band of MWNTs or tangential
stretching of sp2 bonded carbon atoms in a two-dimensional hexagonal lattice [179,
226, 335]. The D band is the signature of defects. The D band is a double–resonance
Raman mode, which provides a measure of the structural disorder produced by
amorphous carbon and any defects [335]. Local defects in the walls of the CNTs lead
to a reduction in the activation energy and lower oxidation temperatures [335].
According to Tuinstra and Koenig [338], an increase in ID/IG corresponds to an
increase in the amount of “unorganized” carbon and/or decrease in the mean crystal
size [335]. ID/IG > 2 indicates a highly disordered form of carbon [338, 339]. There is
a lot of experimental evidence that any interfacial interaction with a matrix [226] and
sidewall derivatisation [340] of CNTs significantly increase ID/IG. Moreover, during
graphitisation, the ID becomes smaller than IG which indicates a more improved
graphene structure [341]. Therefore, the authors used ID/IG ratios supported with highresolution transmission electron microscopic analysis to study the degradation of
CNTs.
Radial Breathing Modes (RBM)
-1
Range: 100-400 cm
Tangential stretching (G band)
-1
Range: 1500-1600 cm
Defects/ disorders (D band)
-1
Range: 1320-1350 cm
SWNT/ DWNT
SWNT/ DWNT/ MWNT
SWNT/ DWNT/ MWNT
Figure 6.1. Schematics of different Raman vibration in CNT. Figure modified from
[342, 343].
122
During high temperature processing of CNTs and ceramic – CNT nanocomposites,
many different outcomes have been reported. This is mainly attributed to the
difference in processing and characterisation methods. Menchavez et al. [330, 331]
reported the formation of some refractory phase, aluminium carbide, during the
sintering of alumina – carbon nanocomposites at 1700 oC. Rul et al. [146] reported
damage of the SWNT structure during hot–pressing in vacuum, which caused a
decrement in the electrical conductivity of magnesium aluminate – CNT
nanocomposite. Jiang et al. [179] reported conversion of SWNT to graphite at higher
sintering temperatures (>1150 oC) due to the disappearance of a shoulder on the G
band peak in the Raman spectra. This resulted in a decrease in the fracture toughness
of the alumina – CNT nanocomposite [179]. Poyato et al. [286] reported that the SPS
process (1550 oC/ 40 MPa/ 3 minutes) is responsible for selective destruction of
SWNTs, and the conversion of some SWNTs into disordered graphite, diamond, and
carbon ‘nano-onions’. In another study, diamond particles were found with diameters
close to 10 μm after SPS of CNTs at 1500 oC for 20 minutes [344]. Diamond particles
(figure 6.2) were formed from CNTs by nucleation from the cores of carbon nanoonions, which are formed from CNTs [344]; subsequently these crystals grow as the
sintering time increases [345]. Many other techniques have also been successful in
synthesising diamond from CNTs. These include laser irradiation [346], shock waves
[347] and radio-frequency hydrogen plasma [348].
(a)
(b)
Figure 6.2. Diamond particles formed during SPS of CNTs [344, 345].
123
In contrast, Thomson et al. [181] found that the SWNT structure was preserved in
alumina nanocomposites at SPS temperatures up to 1250 oC, even when a very high
pressure (105 MPa) was used. Wang et al. [127] reported SWNT in an alumina
composite were undamaged when processed by SPS at 1550 oC. Recently, Wei et al.
[332] reported a 200% improvement in the fracture toughness (indentation fracture
toughness) of alumina – 1 wt% SWNT as compared to monolithic alumina. The
nanocomposite was hot–pressed at 1600 oC for one hour. MWNTs are less sensitive to
high temperature degradation as compared to SWNTs because of their concentric
shells [349]. MWNT bulk samples were prepared by SPS at 1600 oC/ 60 MPa/ 1
minute and were found to be well-preserved at this temperature [350]. Ye at al. [112]
reported an increase in fracture toughness and fracture strength of bariumaluminosilicate glass – 10 vol% MWNT after hot–pressing the composite at 1600 oC/
20 MPa/ 1 hour. In another report, MWNTs were found well preserved at 1500 oC /
40 MPa/ 1 hour based on Raman spectroscopy data [111]. Raman spectrum of
sintered MWNTs shows the intensity of the G peak was about 1.3 times stronger than
that of the D peak, which demonstrated that the graphite sheet structure of MWNTs
was markedly improved after hot press sintering [111]. Balaszi et al. [333] sintered
silicon nitride – MWNT nanocomposite by hot isostatic pressing (HIP) at 1700 oC/ 3
hours and reported preserved MWNTs as observed by showing HR-TEM images. Ma
et al. [134] hot-pressed silicon carbide – 10 wt% MWNTs at 2000 oC for 1 hour (25
MPa in Ar environment). The MWNTs were not damaged in the composite and 10%
improvements in bending strength and fracture toughness over monolithic SiC were
reported.
In some instance CNTs were subjected to very high temperatures. MWNTs were
clearly observed by Laha et al. [164] after plasma spraying blended powder (Al – S –
MWNTs) to a rotating metallic mandrel. In the same report, MWNTs were sprayed at
very high temperatures (9,700 – 14,700 oC), but for very short durations. Also, the
graphitisation of CNTs starts from 2000 oC in an inert atmosphere or vacuum. In the
graphitisation process, defects are removed, leading towards a more perfect graphene
cylinder, which can have a large radius curvature [341]. The graphitisation
124
temperature of carbon can be reduced to 1900 oC by the addition of alumina [351]
and to 1000 oC by the addition of catalytic metals (Fe, Co and Ni) [352].
Alumina is the most common structural ceramic [9] and very popular ceramic matrix
for CNTs [98, 111, 127, 140, 179, 181, 286, 330-332]. So far there has been no
systematic study of the stability of CNTs during high temperature processing in
alumina. To study the stability of the CNTs in ceramics that require very high
temperature (>2000 oC) and long sintering times, the authors SPSed boron carbide –
CNT nanocomposites and boron nitride – CNT nanocomposites. To date, these
nanocomposites have not been sintered to full density. These ceramics would be
difficult to sinter by conventional techniques because of their high sintering
temperatures. This chapter studies the structure of CNTs after SPS of alumina – CNT,
boron carbide – CNT and boron nitride – CNT nanocomposites.
6.2.1. Starting materials
CNTs were supplied by Nanocyl, Belgium. Details of alumina, boron carbide and
boron nitride are provided elsewhere (section 3.2.4).
Composite powders were prepared by colloidal dispersion method (using DMF,
section 4.2.1). Bulk CNTs, ceramic and nanocomposite pellets (diameter 20 mm and
thickness 2 mm) were prepared by SPS. A pressure of 80-100 MPa was applied
concurrently with the heating (rate 300 oC/ minute) and released at the end of the
sintering period for all samples. All powder compacts were cold pressed at 0.64 MPa
125
for a few seconds before SPS. Samples SPSed for this chapter are summarised in table
6.1.
Table 6.1. Samples SPS processed for this chapter.
Matrix
CNT weight %
SPS conditions
Rel. theoretical
density (%)
Alumina
0
~100
Alumina
5
~98.5
Alumina
5
~100
Alumina
5
~100
Alumina
5
~100
Alumina
5
~100
Alumina
5
~100
Alumina
5
~100
Alumina
5
~100
Boron carbide
0
~99
Boron carbide
5
~85
Boron carbide
5
~93
Boron carbide
5
~100
Boron carbide
5
~100
Boron nitride
5
~97.5
Bulk CNT
-
-
Bulk CNT
-
-
Bulk CNT
-
-
Bulk CNT
-
-
126
6.2.3. Material characterisations
All of the sintered samples were ground using SiC paper down to 4000 grit. The
density of the ground samples was measured using the Archimedes’ method. Field
emission scanning electron microscopy (FE-SEM) was used to observe fractured
surfaces. The fractured surfaces were coated with a very thin layer of gold and
transferred to an FE-SEM for examination. High-resolution transmission electron
microscopy (HR-TEM) was carried to study the structure of CNTs after SPS. The
electron transparent nanocomposite films (thickness less than 100 nm) were prepared
by mechanical grinding, polishing, dimpling and focus ion milling. Bulk CNT SPS
processed disks were cut and specimens were scratched with a razor blade from the
centre. The scratched specimens were ultrasonicated in ethanol. A drop of suspension
was left to evaporate on a carbon coated copper grid for HR-TEM analysis.
6.2.4. Raman spectroscopy
Structural characterisation of CNTs in ceramic nanocomposites was performed by
Raman spectroscopy. A Nicolet Almega Dispersive Raman Spectrometer was used.
Raman spectra were excited with a 488 nm Ar+ laser line at a power of 35 mW.
Spectra were detected with an imaging photomultiplier (1024 x 1024) with 5 cm-1
resolution. All samples were cut and Raman laser was focused on the different areas
of the cross section. Typical collection times were 4 minutes. At least 10 locations
were examined to determine the ratio of the intensities of the D and G band.
6.2.5. X-Ray Diffraction (XRD) analysis
The X-Ray Diffraction (XRD) patterns of different phases were obtained by using Cu
Kα radiation, generated with X’PERT PRO (Phillips) at 45 kV, 40 mA. The data was
compared with JCPDS standards using the in-built software.
127
Raman spectra of alumina – 5 wt% CNTs processed at different conditions is shown
in figure 6.3. Alumina has no Raman signals for the given range (1100-1800 cm–1)
that may interfere with the Raman signals of CNTs. Slight peak shifts indicate the
presence of residual stresses [18, 20, 24, 39]. Figure 6.4 shows the intensity ratio of
the D and G bands. The CNTs were well preserved in the composite as ID/IG is < 2 for
all samples. It should be noted that sample prepared at 1400 oC and above were 100%
dense.
o
Composite (1900 C)
o
Relative Raman Intensity
Composite (1700 C)
o
Composite (1400 C)
o
Composite (1100 C)
o
Alumina (1800 C)
D
G
CNT
-1
Wavelength (cm )
Figure 6.3. Raman spectra of CNTs, alumina and alumina – 5 wt% CNT
nanocomposites. Alumina and nanocomposites were SPSed at 100 MPa for 3 minutes.
128
Figure 6.4. Intensity ratio (ID/IG) for CNT and alumina – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 100 MPa for 3 minutes.
For 1600 oC, refer figure 6.5.
Figure 6.5 shows the stability of CNTs with dwell time when sintered at 1600 oC. The
CNTs were well preserved in the nanocomposite sintered for 3 minutes (figure 6.6).
The ratio ID/IG was observed to be greater than 2 for dwell times of more than 13
minutes. The structural stability of CNTs is dependent on the sinter dwell time.
129
30 minutes
20 minutes
13 minutes
3 minutes
D
G
CNT
Wavelength (cm-1)
Figure 6.5. Raman spectra of CNTs and alumina – 5 wt% CNT nanocomposites. All
nanocomposites were sintered at 1600 oC/ 100 MPa.
Figure 6.6. Intensity ratio (ID/IG) for CNT and alumina – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 1600 oC/ 100 MPa.
130
Figure 6.7 shows the HR-TEM image of the composite SPSed at 1900 oC/ 100 MPa/ 3
minutes. Electron diffraction patterns from the different areas confirmed the structural
preservation of CNTs after SPS. It was difficult to visualize individual CNT as they
were overlapping with each other (figure 6.8).
Alumina
CNT
Figure 6.7. HR-TEM of alumina – 5 wt% CNT nanocomposites showing electron
diffraction patterns of different areas. SPSed at 1900 oC/ 100 MPa/ 3 minutes.
The grain boundaries and ceramic – CNT interfaces play an important role in the
resulting properties of the ceramic composites. From HR-TEM analysis (figure 6.7
and figure 6.8) and XRD studies (figure 6.9), no grain boundary or other phases were
observed. All XRD peaks observed (figure 6.9) were representing α alumina
(corundum).
131
CNTs
CNTs
Alumina
Alumina
Figure 6.8. HR-TEM of alumina – 5 wt% CNT nanocomposites showing
agglomerates of CNTs at the grain boundary.
Alumina – 5 wt% CNT
CNTs
x 10 times
Alumina
Figure 6.9. XRD analysis of CNTs, alumina (SPSed) and alumina – 5 wt% CNT
nanocomposite (SPSed). SPSed at 1900 oC/ 100 MPa/ 3 minutes.
132
6.3.2. SPS of other ceramics and their CNT nanocomposites
Using SPS, we prepared 100% dense boron carbide – CNT nanocomposite without
adding any sintering aid. FE-SEM inspection of the fracture surface showed fibrous
CNTs (figure 6.10). Figure 6.10 shows a FE-SEM image of a sample that was SPSed
at 2000 oC/ 80 MPa/ 20 minutes.
(a)
(b)
Figure 6.10. FE-SEM image of boron carbide – 5 wt% CNT nanocomposite. SPSed at
2000 oC/ 80 MPa/ 20 minutes: a) at lower magnification; and b) at higher
magnification.
Raman spectroscopy (figure 6.11) revealed that when a holding time of 20 minutes
was used, the CNTs were structurally degraded (ID/IG increase) with the increase in
processing temperature. ID/IG > 2 indicates a highly disordered form of carbon [35],
which was observed at SPS temperatures >1400 oC (figure 6.12). In figure 6.11, at
temperature > 1800 oC, no shoulder (~1060 cm-1) was observed in boron carbide
peak. This may be an indication of CNTs reaction with boron carbide that is not the
subject of this chapter. At 2000 oC, no Raman peak for CNT was detected, which
indicates that there was severe structural degradation of the CNTs during SPS.
133
o
Composite (2000 C)
o
Composite (1800 C)
o
Composite (1600 C)
o
Composite (1400 C)
o
Boron carbide (2000 C)
D
G
CNT
Wavelength (cm-1)
Figure 6.11. Raman spectra of CNTs, boron carbide (SPSed) and boron carbide – 5
wt% CNT nanocomposites (SPSed). SPSed at 80 MPa for 20 minutes.
Figure 6.12. Intensity ratio (ID/IG) for CNT and boron carbide – 5 wt% CNT
nanocomposites. All nanocomposites were SPSed at 80 MPa for 20 minutes.
134
To observe the severity of degradation at high temperatures, boron nitride – 5 wt%
CNTs was SPSed using the highest furnace operating temperature of 2175 oC. To
produce dense boron nitride, a high sintering temperature because of its inherent
strong covalent bonding [51]. No Raman peaks of CNTs were observed in boron
nitride – CNT nanocomposite. However, like boron nitride – CNT nanocomposite, the
CNTs were found to be fibrous and with a high aspect ratio (figure 6.13).
(a)
(b)
Figure 6.13. FE-SEM analysis of boron nitride – 5 wt% CNTs (rel. theoretical
density: ~97.5%). SPSed at 2175 oC/ 80 MPa/ 20 minutes: (a) at lower magnification;
and (b) at higher magnification.
6.3.3. SPS of bulk CNTs
To analyze the degree of structural deterioration of the CNTs during high temperature
processing, monolithic bulk CNTs disks were SPS processed. Fabricating bulk CNT
for various applications is not new [350]. However, none of the previous works
discussed the structure of the CNTs after the hot pressing. FE-SEM analysis of bulk
CNTs prepared at different temperatures are shown in figure 6.14.
135
(a)
(b)
(c)
(d)
Figure 6.14. FE-SEM images of bulk CNTs SPS processed at different temperatures.
a) Pressed at room temperature; b) SPS processed at 1000 oC; c) at 1500 oC; and c) at
2000 oC.
CNTs appear thick as the function of SPS temperature (figure 6.14). The CNTs used
in this study were >90% pure. It can be concluded that the impurities (catalytic metals
and oxides) might have reduced the graphitisation temperature of the CNTs and
graphitisation caused the formation of additional graphene layers on CNTs. This is
previously reported in the literature [351, 352]. It was proposed that during hot press
sintering, the graphite sheets will grow along their original orientations (the axial and
circumferential directions) with diffusions of the carbon atoms, and the growing
graphite sheets may join together to form bigger sheets [111]. However, HR-TEM
analysis (figure 6.15) shows that CNTs are well deformed, de-shaped and wavy at
higher SPS temperatures and pressures (1000 oC/ 100 MPa). Because of the limited
FE-SEM resolution, the deformed CNTs appeared thick in figure 6.14. This is more
136
clear in figure 6.16, where CNTs were SPS processed at 2000 oC/ 80 MPa/ 20
minutes.
(b)
(a)
Figure 6.15. HR-TEM images of bulk CNTs: a) as received; b) SPS processed at
1000 oC/ 100 MPa/ 3 minutes.
Amorphous carbon
(a)
CNT shells
Maintaining high
aspect ratio
(b)
Figure 6.16. HR-TEM of CNT SPS processed at 2000 oC/ 100 MPa/ 20 minutes: inset
a) nano-onion; and inset b) CNT after SPS, maintaining its aspect ratio.
137
During high-pressure SPS, the ends of the CNTs are more likely to deform. The stress
concentrates in defective areas damaging the end caps and the defective regions along
the nanotubes axis. Electron diffraction pattern of different areas of bulk CNTs SPS
processed at 2000 oC revealed partly damaged CNT (figure 6.16). These observations
are well supported with Raman spectroscopy (figure 6.17). A very slight peak shift
was observed in the Raman spectroscopy, which is an indication of structural
transformations. Intensity of D band line is lower than that of G band line, which
indicates high temperature graphitisation, leading towards large CNTs [341]. This is
well supported with the FE-SEM result (figure 6.14). At these conditions, CNTs are
going through a series of transformation as reported in the literature. Here, such
transformations are not fully complete to reveal a clear microscopic evidence. It can
be graphitisation and/or transformation into nano-onions and diamond. Some HRTEM observations revealed that the layers of the outer shells of the nanotubes break
and transform into curled graphitic structures, termed as nano-onions (figure 6.16,
inset a).
Amorphous carbon
1332 Cubic diamond
o
2000 C
o
1500 C
o
1000 C
Raw form
Wavelength (cm-1)
Figure 6.17. Raman spectra of CNTs SPS processed at different conditions. All
samples were processed at 80 MPa for 20 minutes.
138
In figure 6.17, the Raman signal for bulk CNTs SPS processed at 2000 oC shows a
Raman peak at 1332 cm-1, which is a Raman signal for the cubic diamond phase (C–C
sp3 bond) [353]. It should be noted that the motive of the work in this chapter was not
to synthesize diamond from CNTs. Also, no diamond particle or electron diffraction
pattern of diamond was observed during HR-TEM analysis. This is because of low
processing durations and pressures. The Raman signal also shows amorphous carbon,
which was observed in the HR-TEM analysis (figure 6.16).
6.4. Conclusions
The structure of CNTs should be preserved in the ceramic matrix if they are to
provide an effective reinforcement. CNTs were found to be preserved in alumina after
being SPSed up to 1900 oC/ 100 MPa/ 3 minutes. In alumina and boron carbide
matrices, structural degradation of CNTs started from 1600 oC. This can be avoided
by the use of additional sintering aids that lowers the sintering temperature. CNTs
maintained their high aspect ratio and fibrous nature after being SPSed in boron
nitride at 2175 oC for 20 minutes. However, no Raman vibrations of CNTs were
observed for nanocomposites processed at temperatures > 2000 oC. Structural
preservation of CNTs in ceramic nanocomposites depends on the nature of ceramic,
SPS temperature and dwell times. CNTs are not suitable for matrices that require higher
temperatures (> 1600 oC) and longer processing times (> 13 minutes).
CNTs went through a series of incomplete transformations during high temperature
SPS processing of bulk CNTs. Partly amorphous CNT and nano-onions were very
distinctly observed. Even in bulk form, CNTs maintained their high aspect ratio and
fibrous nature after being SPSed at 2000
o
C/ 80 MPa/ 20 minutes. Raman
spectroscopy provided peaks for cubic diamond phase. However, in contrast to the
literature, no diamond phase was seen in the electron microscopic analysis. This is
because of low processing durations and pressures.
139
Chapter 6. Structural Stability CNTs in Ceramics (Preservation Studies)
Chapter 7.
Electrical Properties of Ceramic – CNT
Nanocomposites
7.1. Introduction
The addition of CNTs for improving electrical conductivity of ceramics is widely
appreciated (table 2.2). Alumina is inherently an insulator (electrical conductivity:
10-13 S/m). Adding a small amount of CNT (~0.79 vol%) [96] to alumina can make it
electrically conductive (10-4 S/m). A detailed literature review is presented elsewhere
(section 2.4). In this chapter, the electrical conductivity of alumina – CNT and
alumina – carbon black nanocomposites were compared. The effect of grain size and
sintering conditions on the electrical conductivity of the alumina – carbon
nanocomposites was also studied. For alumina – CNT nanocomposites, electrical
conductivities were also measured as a function of temperature.
MWNTs (supplied by Nanocyl, Belgium) and CB (supplied by Degussa, Germany)
were used to prepare alumina nanocomposites. The preparation method is described in
chapter 4 and 5. All nanocomposites were fully dense (rel. theoretical density:
~100%). Samples prepared for this chapter are described in table 7.1.
140
Chapter 7. Electrical Properties of Ceramic – CNT Nanocomposites
Table 7.1. Samples SPSed for this chapter#.
#
Matrix
Filler (wt %)
SPS conditions
Rel. theoretical
density (%)
Alumina
-
~100
Alumina
CB (2)
~84
Alumina
CB (2)
~98
Alumina
CB (2)
~99
Alumina
CNT (2)
~78
Alumina
CNT (2)
~88
Alumina
CNT (2)
~97
Alumina
CNT (2)
~100
Alumina
CNT (3.5)
~67
Alumina
CNT (3.5)
~84
Alumina
CNT (3.5)
~98
Alumina
CNT (3.5)
~100
Alumina
CNT (3.5)
~100
Alumina
CNT (3.5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
Alumina
CNT (5)
~100
For CNTs, 2 wt% = ~4.7 vol%, 3.5 wt% = ~8 vol% and 5 wt% = ~11.2 vol%
141
The electrical conductivities of the samples were measured (section 3.3.4) with a high
sensitivity digital micro-ohmmetre (Keithley 580) using the two-point method on
silver electroded specimens (3 mm × 3 mm× 3 mm) prepared using a diamond cutting
machine. Details are given elsewhere (section 3.3.4). A power supply (Keithley 2602)
and digital multimeter (Keithley 6517A) were used to measure the current-voltage (IV) characteristics of the samples.
7.3. Results and Discussion
7.3.1. CNT vs. carbon black
The DC electrical conductivities of alumina – CNT and alumina – carbon black
nanocomposites are shown in figure 7.1. The conductivity mechanisms involved in
CNT nanocomposites could be variable range hopping [212] or fluctuation-assisted
tunnelling [213].
Figure 7.1. Electrical conductivities of alumina – carbon nanocomposites.
142
The higher electrical conductivity of the ceramic – CNT nanocomposite is the result
of the large aspect ratio of CNTs (~150) as compared to that of carbon black (~1).
This is well supported in figure 7.1. The large aspect ratio of CNTs resulted in
entangled network of conductive pathways (figure 7.2) on the grain boundaries of
alumina, which did not occur in alumina – carbon black nanocomposites (figure 5.9b).
As compared to alumina – 2 wt% carbon black (30 S/m), the electrical conductivity of
alumina – 2 wt% CNT nanocomposites is 125 S/m, i.e. four times higher than alumina
– carbon black nanocomposite. A larger CNT content resulted in higher electrical
conductivity. The increased electrical conductivity is attributed to the presence of
undamaged CNTs, due to the utilisation of SPS technique that allowed lower sintering
temperatures and shorter sintering times. Figure 7.3 shows the voltage – current
relations measured at room temperature for alumina – CNT nanocomposite. The
current increased linearly with voltage for any instance, indicating a good ohmic
behaviour. The slope of the current – voltage curve corresponds to the CNT content
(figure 7.3). Higher CNT content in the nanocomposite resulted in a lower slope/
electrical resistance and vice versa.
(a)
(b)
Figure 7.2. HR-TEM image of alumina – 5 wt% CNT nanocomposite: (a) CNTs
around grains; and (b) percolating network highlighted.
143
Figure 7.3. Voltage vs. current measured at room temperature for alumina – CNT
nanocomposites. Both nanocomposites were SPSed at 1800 oC/ 100 MPa/ 3 minutes.
7.3.2. Electrical conductivity as a function of grain size
Figure 7.4 shows the electrical conductivity of alumina – carbon nanocomposites as a
function of grain size. The grain sizes were coarsened by using higher sintering
temperatures. For alumina – CNT nanocomposites, the electrical conductivity
increased significantly with increasing grain size. In large grained nanocomposites,
higher electrical conductivity is due to the availability of fewer paths for the current to
flow. Large grains of alumina concentrate CNTs in fewer conductive paths. However,
due to the particulate nature of carbon black, this was not observed for the alumina –
carbon black nanocomposites (figure 7.4). With the growth of alumina grains, the
particulates or agglomerates of carbon black may isolate, resulting in higher
percolation threshold and lower electrical conductivity. Due to the fibrous nature and
144
aspect ratio of CNT, such isolation is not possible that resulted higher electrical
conductivity for alumina – CNT nanocomposite.
o
SPSed at 1800 C
o
SPSed at 1600 C
Figure 7.4. Electrical conductivities of alumina – carbon nanocomposites as the
function of grain size. The grain sizes were coarsened by using higher sintering
temperatures.
Apart from addition of CNTs and using higher SPS temperatures, another way to
tailor electrical conductivity of alumina – CNT nanocomposites is to use longer
processing durations. In figure 7.5, alumina grains were slightly coarsened using
longer SPS processing time at 1400 oC. The grain growth in alumina – 5 wt% is not
significant (figure 5.11). In this context, this route can be adopted to slightly increase
the electrical conductivity of these nanocomposites. For example, increasing
processing time by 13 minutes resulted in a 12% increase in the electrical
conductivity. However, using higher sintering temperatures (figure 7.4) proved to be
145
more effective in increasing the electrical conductivity as compared to using longer
dwell times (figure 7.5).
Figure 7.5. Electrical conductivities of alumina – 5 wt% nanocomposites as the
function of SPS processing durations. CNTs were supplied by Nanocyl, Belgium.
7.3.3. Electrical conductivity as a function of temperature
Figure 7.6 shows the electrical conductivity of alumina – CNTs nanocomposites as a
function of temperature. It should be noted that CNTs are the only conductor in the
nanocomposite. MWNTs show metallic behaviour [158]. This is evident from figure
7.6, where the electrical conductivity is decreasing with the rise in temperature. At a
particular temperature, CNTs started oxidizing and a sharp decrease in the
nanocomposite’s electrical conductivity was observed. For instance, alumina – 5 wt%
CNT nanocomposite (SPSed at 1400 oC/ 100 MPa/ 3 minutes) showed an abrupt
146
decrease in electrical conductivity at 449 oC, where as alumina – 5 wt% CNT
nanocomposite (SPSed at 1400 oC/ 100 MPa/ 20 minutes) showed such behaviour at
496 oC. This is due to higher oxidation resistance in large grained nanocomposites. In
large grained nanocomposites, there is a small volume of grain boundary and it is
difficult for more oxygen to diffuse in and react with the CNTs. In this way, the
electrical conductivity of alumina – CNT nanocomposites can be improved.
Alumina grain size
o
Electrical Conductivity (S/m)
496 C
701 nm
653 nm
364 nm
o
471 C
o
449 C
Temperature (oC)
Figure 7.6. Electrical conductivities of alumina – 5 wt% nanocomposites as the
function of temperature. The grain sizes were coarsened by using longer processing
durations. Heating rate: 2 oC/ minute.
7.4. Conclusions
The electrical conductivity of alumina – CNT nanocomposites is four times higher as
compared to alumina – carbon black nanocomposites for the same mass content of
147
carbon due to the fibrous nature and high aspect ratio of CNTs. The conductive
network structure allowed the percolation of CNTs at low volume fractions [96] and
thus increased the electrical conductivity compared to alumina – carbon black
nanocomposites. The electrical conductivity of alumina – CNT nanocomposite
increased with increasing grain size due to the concentration of CNTs in fewer
conductive paths. Because of the electronic properties of CNTs, the electrical
conductivity of alumina – CNT nanocomposites varied with temperature for different
grain sizes.
148
Chapter 8.
Oxidative Stability of Ceramic – CNT
Nanocomposites
8.1. Introduction
Ceramics are used in high temperature applications [279]. In order to get most out of
CNTs in ceramic – CNT nanocomposites, it is necessary to retain their structure and
properties, and avoid degradation. The potentially improved properties of
nanocomposites can then be explored. However, because of their low oxidation
temperature resistance (~500
o
C) of carbon [354], using ceramic – CNT
nanocomposites at high temperatures is an obstacle for their commercial success.
Most of previous reports about oxidation resistance of CNTs are based on polymer
based CNT composites [355-357]. For improving electronic and electrical properties,
ceramic layers were coated on CNTs [234, 358-361]. However, no comments were
presented on the stability against oxidation of the coated CNTs. Wang et al. [362]
coated 10 nm of silicon layer and reported an improvement of 105 oC in oxidation
resistance. Li et al. [363] reported that the increase in Ni catalyst during CVD growth
improves the oxidation resistance of multiwall CNTs. The catalyst composition had a
significant effect on the CNT structure and stability, and is important for large-scale
CNT synthesis [363].
To date, there are no reports on the effect of the ceramic matrices and ceramic
coatings on the oxidation resistance of CNTs. Coating CNTs by atomic layer
149
Chapter 8. Oxidative Stability of Ceramic – CNT Nanocomposites
deposition (ALD) provides shielding that improves oxidative stability of the
encapsulated CNTs. This chapter reports the high temperature protection effect of
alumina coating on CNTs. It also reports the effect of grain boundary area on the
oxidative stability of CNT in alumina – CNT nanocomposites.
MWNTs supplied by Nanocyl, Belgium and NanoDynamics, USA were used in this
chapter. The preparation method is described in chapter 4 and 5. All nanocomposites
were fully dense (rel. theoretical density: ~100, table 8.1). The oxidation resistance of
CNTs in different systems was characterised by Thermo Gravimetric Analysis (TGA)
as described in section 3.3.7.
Table 8.1. Alumina – 5 wt% CNT nanocomposites SPSed for this chapter.
Alumina
coating
Weight
%
CNT supplier
SPS conditions
Rel. theoretical
density (%)
No
5
Nanocyl, Belgium
~100
No
5
Nanocyl, Belgium
~100
8.3.1. Oxidation of CNTs
Figure 8.1 shows the oxidation behaviour of CNTs obtained from different suppliers.
477 oC is the onset of oxidation (intersection point of first two slopes in figure 8.1) for
CNTs provided by NanoDynamics, USA and 507 oC is the onset of oxidation for
CNTs provided by the other source. CNTs provided by NanoDynamics, USA have
larger average diameter (~20 nm) as compared to the CNTs provided Nanocyl,
150
Belgium (~12 nm). CNTs provided by NanoDynamics, USA showed slightly less
oxidation resistance as compared to CNTs provided by Nanocyl, Belgium. The
oxidative stability of CNTs is influenced by defects [363] and nanotube diameter
[363, 364]. Oxygen molecules react easily with larger surface areas, resulting in
decreased oxidative stability of CNTs (supplied by NanoDynamics, USA). The
oxidation of CNTs is not rapid and acute like combustion, which is also evident in
other report [365]. Because the kinetic energy of oxygen varies with temperature,
there is not a critical temperature when the oxidation of CNTs starts [365] as shown in
figure 8.1. During the initial stage of TGA, all samples showed a slight mass loss due
to the presence of amorphous carbon, as reported in the literature [363, 366]. In the
second stage of TGA, the curve slope remains almost the same in the definite
temperature range for both types of CNTs. In the third stage of TGA, there was no
weight gain observed during thermal treatment, since no oxidation of the impurities
occurred. The weight loss for both types of CNTs was not 100 % due to the presence
of impurities (section 5.3.5).
o
477 C
507 oC
Figure 8.1. TGA of raw CNTs from different suppliers. Heating rate: 5 oC/ minute.
151
8.3.2. Oxidation of alumina coated CNTs
To improve the oxidation resistance of CNTs, the CNTs were coated with alumina
[250]. The oxidative stability of CNTs was distinctly improved due to the protective
alumina coating (figure 8.2).
o
553 C
o
477 C
530
o
C
Figure 8.2. TGA of uncoated CNTs and alumina coated CNTs. CNTs were supplied
by NanoDynamics, USA. Heating rate: 5 oC/ minute.
The onset oxidisation temperature for sample coated for 50 ALD cycles is now as
high as 553 oC in air atmosphere, which is 76 oC higher than that of uncoated CNTs.
Once the oxidation started, the degradation was also reduced to 0.41%/oC, compared
to 0.92 %/oC for uncoated CNTs. The degradation process was delayed because it
became more difficult for oxygen molecule to approach CNTs after coating. By
analyzing the third stage (> 700 oC, figure 8.2), it is possible to quantify the mass
content of the alumina on the coated CNTs (figure 8.3a). The coated CNTs with 25
152
ALD cycles had a residual mass of 56.4 % of CNTs of the original mass and the 50
ALD cycles had a 44.9 % residual mass. A thicker alumina coating could further
inhibit the oxidisation of CNTs but it may decrease the mechanical properties of
CNTs by making the coated CNTs brittle. After cooling down the TGA furnace, white
coloured alumina nanotubes (figure 8.3b) were left in the platinum pan, which were
previously surrounding CNTs. This could be a route to mass-produce alumina
nanotubes (figure 8.4).
(a)
(b)
Figure 8.3. Platinum pan used for TGA: a) alumina coated (50 ALD cycles) CNTs
before oxidation; and b) alumina nanotubes left after oxidation of coated CNTs. CNTs
were supplied by NanoDynamics, USA.
153
(a)
(b)
Figure 8.4. Alumina nanotube left after the oxidation of coated CNTs. CNTs were
supplied by NanoDynamics, USA: a) at lower magnification; and b) at higher
magnification.
8.3.3. Oxidation of alumina – CNT nanocomposites
Dense 5 wt% CNT (uncoated) dispersed alumina nanocomposites with different grain
sizes were fabricated using SPS. In figure 8.5, the onset oxidisation temperature for
sample sintered at 1800 oC is 588 oC in air, which is 81 oC higher than that of raw
CNTs. Once the oxidation started, the degradation was 0.026%/oC, which is 97% less
than that of raw CNTs. Oxidative reactivity in these nanocomposites is influenced by
the grain boundary area. The sample sintered at 1200 oC and 1800 oC have grain sizes
of 0.15 μm and 0.66 μm (figure 5.10b), respectively. The larger grain size material
therefore had a smaller total area of grain boundaries, making the ingression of
oxygen slower and increasing the oxidation resistance. Thus the oxidative stability of
alumina – CNT nanocomposites can be tailored by changing the grain boundary area.
154
o
588 C
o
547 oC
507 C
Figure 8.5. TGA of raw CNTs and SPSed alumina – 5 wt% uncoated CNT
nanocomposites. CNTs were supplied by Nanocyl, Belgium.
Heating rate: 5 oC/ minute.
8.4. Conclusions
It is necessary to preserve the chicken wire hexagonal distribution of CNTs in
ceramics for high-temperature applications. It was observed that TGA is a good tool
to evaluate the mass content of CNTs in the coated CNT and SPSed nanocomposites.
Coating CNTs by ALD provides shielding that improves oxidative stability of the
encapsulated CNTs. A thicker alumina coating could further inhibit the oxidisation of
CNTs and enhance the thermal stability of CNTs. Coarser grained materials have
higher oxidation resistance of CNT due to the presence of fewer grain boundaries.
The oxidative stability of alumina – CNT nanocomposites can be tailored by changing
the grain boundary area. SPS is a good processing method to modify the grain size
(and grain boundary area) of alumina – CNT nanocomposites.
155
Chapter 8. Oxidative stability of Ceramic – CNT nanocomposites
Chapter 9.
Mechanical Properties of Ceramic – CNT
Nanocomposites
9.1. Introduction
Ceramics are inherently brittle. Zhan et al. [138, 140] SPSed alumina – SWNT
nanocomposites and reported a threefold improvement in the fracture toughness
(section 2.3.1). Since then, a number of research groups have investigated various
routes to fabricate strong and tough ceramic – CNT nanocomposites [110-112, 134,
137, 141-145, 153, 161-163, 168, 192, 332]. A detailed literature review is presented
in chapter 2. However, to date, the role of CNTs on the mechanical properties of
ceramic nanocomposites remains inconclusive. This is because of the usage of
different types and compositions of CNTs, inhomogeneous distribution of CNTs and
differences in experimental conditions. In many reports [106, 109, 125, 131, 134,
332], the presence of good quality CNTs was not confirmed after conventional
processing at high temperatures. This is also one of the reasons for variation in the
mechanical properties of CNT reinforced ceramic nanocomposites. The structural
stability of CNTs after sintering was confirmed in chapter 6. In this chapter, alumina,
alumina – carbon black and alumina – CNTs nanocomposites were prepared using
SPS. Hardness, indentation fracture toughness, flexural modulus, flexural strength and
thermal shock resistance (qualitative assessment) were analysed. Good interfacial
bonding is required to achieve load transfer across the CNT-matrix interface, a
condition necessary for
improving the mechanical properties of ceramic
nanocomposites [176]. The toughening enhancement produced by ceramic coated
156
Chapter 9. Mechanical Properties of Ceramic – CNT Nanocomposites
CNTs was much better than that resulted from application of either CNTs or
nanometre ceramic powders individually [187]. To improve interfacial bonding,
alumina coated CNTs were used. The mechanical properties of nanocomposite
reinforced with alumina coated CNTs are compared with nanocomposite reinforced
with CNTs. The outcomes of these comparisons are discussed in relation with their
microstructures.
MWNTs (supplied by NanoDynamics, USA and Nanocyl, Belgium) and CB (supplied
by Cabot, USA) were used to prepare alumina nanocomposites. All coated CNTs
were coated for 50 cycles of ALD, which resulted in ~5-10 nm coating of alumina
(figure 3.1e). All materials were prepared as explained in chapter 4 and 5. Alumina
and alumina – carbon nanocomposites were sintered at 1800 oC/ 100 MPa/ 3 minutes.
To exclude the effect of porosity on the mechanical properties, all specimens were
fully densified (rel. theoretical density: ~100%) using SPS. The sintering conditions
used for all samples were the same for the sake of comparison. The high residual
stresses at the surface of alumina grains may cause the nanotube/ matrix interface
debonding [196]. To avoid the influence of residual stresses on the mechanical
properties, a relatively slow cooling rate (~50 oC/ minute) was adopted.
Vickers hardness and indentation fracture toughness (IFT) were measured as
described in section 3.3.5 using 2.5 kg load. For surface roughness analysis, all
specimens were ground using SiC paper and diamond polished down to 1 μm. The
data presented for Vickers hardness and IFT is the average values obtained from at
least ten indentations on the same specimen.
157
9.2.1. Flexural bending
Rectangular bars of dimensions ~ 2 x 1.5 x 25 mm were cut using a high speed
diamond cutter. The flexural modulus was measured by non-destructive 3-point
bending method using Dynamic Mechanical Analyzer (DMA Q800, TA Instruments).
A small load (up to 18 N) was applied and deflection of the beam was measured. The
flexural modulus was calculated by measuring the slope of the load-deflection curve.
The technique is quite accurate and reliable. Gou [367] reported similar values of
flexural modulus for different ceramic materials using an Impulse Excitation
Technique (IET) and flexural bending method.
Flexural strength was evaluated by 3-point bending in accordance with ASTM
C1161-02c [368]. For achieving good results, a special jig was used (figure 9.1). The
bending tests were performed at room temperature on an Instron 6025 using a load
cell of 1 kN. A cross-head speed of 0.2 mm/ minute was used for all specimens.
Sample dimensions were ~ 2 x 1.5 x 25 mm. All edges and corners on these samples
were chamfered on 4000 grit size SiC grinding paper in order to avoid stress
concentration points (figure 9.2). Configuration A [368] was used for all specimens.
The flexural strength of the specimens was determined from the failure load and the
geometry of the test piece using equation 9.1.
S=
3PL
2bd 2
Equation 9.1
Where, S = flexural strength, P = maximum loading level, L = 20 mm, outer
(support) span, b = specimen width and d = specimen thickness. The data presented
for fracture strength and flexural modulus is the average values obtained from at least
ten test specimens of the same composition. Fractographic analysis was performed on
selected fractured samples after flexural testing to identify the cause of failure. The
origin of the fracture was identified by carefully monitoring the surfaces under an
optical microscope and a field emission SEM.
158
Figure 9.1. Special jig for improved 3 –
Figure 9.2. Chamfered edges to avoid
point flexural testing.
stress concentration points on corners.
Sample thickness: 1.5 mm.
Sample: Fractured alumina surface.
9.2.2. Thermal shock resistance during SPS (observation)
A SPS furnace was used to compare the thermal shock resistance of alumina and
alumina – carbon nanocomposites. This analysis is possible by rapidly cooling a thin
sample after sintering. Dried alumina powder (0.07 cm3) was poured into a carbon die
and cold pressed at 0.62 MPa for 5 s before sintering. Nanocomposite discs (thickness
0.2 mm and diameter 20 mm) were prepared by SPS. A pressure of 100 MPa was
applied concurrently with the heating (rate 300 oC/ minute). The sample was sintered
at 1800 oC. After sintering, temperature was dropped to 450 oC using a cooling rate of
100 – 300 oC/ minute. A constant load of 100 MPa was used while cooling.
Application of pressure while cooling is necessary in order to remove heat from the
system. The carbon die was given sufficient time for cooling. The sintered material
was then carefully taken out from the die and analysed for cracks.
The same
procedure was followed for alumina – carbon nanocomposites.
159
9.3.1. Surface finish
Figure 9.3 shows the surface finish of different compositions of alumina – CNT
nanocomposites prepared using CNTs supplied by Nanocyl, Belgium. All samples
were polished in the same manner.
(a)
(b)
(c)
(d)
Figure 9.3. Polished surfaces of alumina – CNT (Nanocyl, Belgium) nanocomposites
containing: (a) no CNTs, thermally etched at 1500 oC for 10 minutes; (b) 2 wt% (~4.7
vol%) CNTs; (c) 5 wt% (~11.2 vol%) CNTs; and (d) 10 wt% (~21 vol%) CNTs.
It is difficult to achieve good surface finish for alumina – CNT nanocomposites.
Because of the strong covalent bonding in alumina and due to lubricating nature of
carbon, there was poor interfacial adhesion at the ceramic/ CNT interface. On
160
polishing, alumina grains were selectively plucked out (figure 9.3b, 9.3c and 9.3d)
because of poor interfacial adhesion. This is very obvious in nanocomposite
containing 10 wt% CNT (figure 9.4). The arrows show the location of alumina grain
before plucking (figure 9.4). Such surface will result in poor mechanical properties
because of significant number of surface flaws.
Figure 9.4. Polished surface of alumina
Figure 9.5. Alumina coated CNT (50
– 10 wt% CNT nanocomposite. Arrows
ALD cycles) encapsulated in an alumina
showing alumina grain ex-sites. CNTs
grain of nanocomposite. CNTs were
were supplied by Nanocyl, Belgium.
supplied by NanoDynamics, USA.
Apart from surface flaws, high weight content of CNTs (10 wt%) yields an
inhomogeneous dispersion of CNTs and a porous microstructure. Trapped gases in
agglomerates reduce the mechanical and electrical properties of the nanocomposite
[112, 177]. In this context, the maximum amount of CNT in alumina – CNT
nanocomposites should be restricted to 5 wt%. However, when alumina coated CNTs
were used, CNTs were located inside the grains (figure 9.5) and much better surface
finishes could be achieved.
161
9.3.2. Vickers hardness
Table 9.1 shows the grain size and Vickers hardness of alumina and alumina – carbon
nanocomposites. Alumina is one of the hardest structural ceramics. The Vickers
hardness of the nanocomposites decreased monotonously with increasing carbon
content. It is quite well known that the hardness of alumina increases with decreasing
grain size [369]. However, this is not true for alumina – carbon nanocomposites.
Table 9.1. Vickers hardness of alumina and alumina – carbon nanocomposites.
§
Material
Alumina coating
Grain size (µm)
Vickers hardness (GPa)
Alumina
-
43.20
18.9 ± 0.3
Alumina – 2 wt% CNT§
-
1.81
10.1 ± 0.3
Alumina – 2 wt% CNT#
-
1.79
10.2 ± 0.3
~5-10 nm
3.23
14.9 ± 0.2
-
0.66
12.0 ± 0.3
~5-10 nm
2.34
18.6 ± 0.2
Alumina – 5 wt% CB
-
4.67
11.0 ± 0.3
CNTs were supplied by Nanocyl, Belgium
#
CNTs were supplied by NanoDynamics, USA
The presence of a soft phase (hardness of MWNT in radial axis ~6 – 10 GPa [370])
eases the penetration of diamond indenter during testing. Such effect can be estimated
by hardness rule-of-mixtures for composites [371] as indicated in equation 9.2:
H vc = H vm (1 - v f )+ H vf v f
Equation 9.2
Where, H vc = hardness of composite, H vm = intrinsic hardness of the matrix phase,
v f = volume fraction of the filler phase and H vf = intrinsic hardness of the filler
phase. It must be noted here, that equation 9.2 is a useful tool for approximation and
162
does not consider the dispersion and quality of CNTs. The reduced hardness values of
the nanocomposites are due to the presence of very soft phase at the grain boundaries,
which is nullifying the effect of fine grains. However, there may be some contribution
coming from grain size when hardness values of alumina – 5 wt% CNT and alumina –
5 wt% CB nanocomposites are compared.
The presence of alumina coating increased the hardness of alumina – CNT
nanocomposites. The alumina coated CNTs are located inside the grains (figure 9.5
and figure 9.6a) unlike uncoated CNTs, which form a network of entangled CNTs on
the grain boundaries (figure 9.6b). This gives the flexibility of making CNT based
intragranular nanocomposites (figure 9.6a) and intergranular nanocomposite (figure
9.6b) or both. In alumina coated CNT nanocomposites, no soft lubricating phase is
present at the grain boundaries. Higher frictional forces between alumina grains are
resisting the penetration of indenter and grain sliding that is improving the hardness of
the nanocomposite.
(a)
(b)
Figure 9.6. Different types of fracture modes in alumina – 2 wt% CNT
nanocomposites; (a) intragranular fracture mode in coated CNT nanocomposite (50
ALD cycles, NanoDynamics, USA); and (b) intergranular fracture mode in CNT
nanocomposite (Nanocyl, Belgium).
163
9.3.3. Indentation Fracture Toughness (IFT)
Table 9.2 shows the indentation fracture toughness (IFT) for different materials.
Mechanical properties are influenced by the type of carbon nanotubes used in ceramic
nanocomposites [107]. Zhan et al. [140] reported a value of 9.7 MPa√m (measured by
IFT method) for alumina – 10 vol% SWNT. In the current study MWNTs were used.
The internal shells of MWNT are unable to bond to the alumina matrix and therefore
tensile loads are carried entirely by the external shell [140]. An increase in IFT could
be obtained even by improving the quantity of carbon nanotubes. A 78%
improvement in the IFT was found for alumina – 5 wt% CNT nanocomposite as
compared to alumina. Figure 9.7a shows alumina having longer and wider radial
cracks as compared to alumina – 5 wt% CNT nanocomposite (figure 9.7b). This is
due to the presence of strong entangled network of CNTs at the grain boundaries
(figure 9.6b). Such a phenomenon was not observed in alumina – CB nanocomposite
(figure 5.9b).
Table 9.2. Indentation fracture toughness of different materials.
§
Material
Alumina coating
Indentation Fracture Toughness (MPa√m)
Alumina
-
3.3 ± 0.2
Alumina – 2 wt% CNT§
-
5.7 ± 0.1
-
5.6 ± 0.3
~5-10 nm
5.5 ± 0.5
-
5.9 ± 0.3
~5-10 nm
5.8 ± 0.2
-
3.4 ± 0.3
#
164
(a)
(b)
Figure 9.7. Vickers indent after applying 2.5 kg load in: (a) alumina; and (b) alumina
– 5 wt% CNT nanocomposite. CNTs were supplied by NanoDynamics, USA.
Figure 9.8a shows the sub-surface damage caused by 2.5 kg loading on alumina and
alumina – 5 wt% uncoated CNT. It was not possible to observe the same sub-surface
optical effect for alumina – CNT nanocomposite because of the opaque nature of
sample (figure 9.8b).
(a)
(b)
Figure 9.8. Sub-surface damage analysis after Vickers indentation. Optical
micrograph (dark field image) showing sub-surface cracking in: (a) alumina; and (b)
alumina – 5 wt% uncoated CNT (Nanocyl, Belgium) nanocomposite.
165
Sub-surface damage was analyzed by observing the cross-section of indented
specimen under SEM (figure 9.9). The sample was carefully fractured to observe the
cross-section. As compare to alumina (figure 9.9a), no major damage was observed
for alumina – 5 wt% CNT nanocomposite (figure 9.9b). Significant redistribution of
stresses by the strong entangled network of CNTs was the main reason for improved
sub-surface damage resistance [127]. It prevented the formation of cracks that were
observed for brittle alumina. These properties are important for ceramics used in
contact-mechanical applications such as bearings, valves, nozzles, seals, wear parts,
armour and prostheses [127].
(a)
(b)
Figure 9.9. Sub-surface damage analysis after Vickers indentation. Cross-section of:
(a) alumina, showing major crack; and (b) alumina – 5 wt% uncoated CNT (Nanocyl,
Belgium) nanocomposite, showing no major sub-surface damage.
In uncoated CNT nanocomposites, fracture occurred along the grain boundaries
(intergranular) and not within the alumina grains (transgranular). In coated CNT
nanocomposite (figure 9.6a), the fracture path was intergranular and transgranular.
Increasing the quantity of CNTs in alumina – coated CNT nanocomposite decreased
the IFT further. Here transgranular fracture indicates weak grains and this is due to
poor contact between coated CNTs and alumina grains. The objective of using
alumina coated CNTs in alumina for improved IFT was not successful.
166
9.3.4. Flexural modulus
Table 9.3 shows the flexural modulus for different materials. Previous publications
[107, 116, 133, 200] reported a decrease in flexural modulus of ceramic – CNT
nanocomposites due to poor density and damage caused to CNTs while processing.
Chapter 5 confirms the improved density of raw CNT reinforced nanocomposite and
chapter 6 confirms the structural stability of CNTs after SPSing. The decrease in
flexural modulus is due to grain size refinement, lubricating nature of CNTs and poor
adhesion between CNTs and alumina grains. The argument is also valid for alumina –
CB nanocomposite. These properties may be useful in applications where mechanical
flexibility and good electrical conductivity is desired. A detailed discussion on the
applications of ceramic based CNT nanocomposites is presented in appendix B. In
coated CNT nanocomposites, CNTs were individually encapsulated in coarse alumina
grains (figure 9.5). Unlike uncoated CNT nanocomposite, no soft lubricating phase
was present between alumina grains that resulted in slightly higher flexural modulus
of coated CNT nanocomposites.
Table 9.3. Flexural modulus of different materials.
#
Material
Alumina coating
Grain size (µm)
Flexural modulus (GPa)
Alumina
-
43.20
308 ± 2.9
-
1.81
234 ± 1.9
~5-10 nm
3.23
263 ± 3.9
-
0.66
165 ± 2.1
~5-10 nm
2.34
185 ± 3.1
-
4.67
201 ± 1.3
167
9.3.5. Flexural strength
The variability of the flexural strength of a ceramic can be analysed in a number of
ways. The most commonly used is Weibull statistical analysis (appendix C). Table 9.4
shows the fracture strength for different materials.
Table 9.4. Flexural strength of different materials.
#
Material
Alumina coating
Grain size (µm)
Flexural strength (MPa)
Alumina
-
43.20
241 ± 8.9
-
1.81
253 ± 9.3
~5-10 nm
3.23
278 ± 6.9
-
0.66
261 ± 11.2
~5-10 nm
2.34
289 ± 7.3
-
4.67
233 ± 12.3
According to the Griffith criterion, the fracture stress is expected to decrease with
increasing grain size [279]. The inherent flaw size is increases with the increase in
grain size. Minor improvement in flexural strength was observed for alumina –
uncoated CNT nanocomposites as compared to alumina. The grain size of alumina – 5
wt% CNT is >43 times smaller than alumina, but only 15% improvement in flexural
strength was observed (table 9.4). There is no major effect of the fine grains on the
168
fracture strength of alumina – uncoated CNT nanocomposites. The effect of fine grain
size is very significant in monolithic alumina. A 21% improvement in fracture
strength was reported in monolithic alumina when grain size was reduced from 8 to 2
microns [372].
Fracture strength depends on the flaw geometry and dimensions (equation 9.3), i.e.
larger the flaw size, lower the fracture strength [373].
K =  c
Equation 9.3
Where, K = stress-intensity factor,  = fracture strength,  = dimensionless
geometric constant and c = flaw (including surface flaws) size. There was no
significant reinforcing effect of uncoated CNTs on the fracture strength of alumina.
Poor grain boundary adhesion and surface flaws (figure 9.3b and 9.3c) are responsible
for poor fracture strength of alumina – uncoated carbon nanocomposites.
Using alumina coated CNTs as reinforcement for alumina proved slightly effective
(table 9.4). As explained in section 9.3.2 and 9.3.3, in alumina – coated CNT (figure
9.6a), fracture occurred in an intergranular and transgranular fashion. A mixed inter/
transgranular fracture mode is a representation of improved fracture strength [147].
All samples failed forming a compression curl on the loading surface, which is a
representation of low energy failure (figure 9.10). It was observed that all samples
failed by a surface flaw located on the tensile surface. Figure 9.11 shows alumina –
CNT nanocomposites mounted in a way that the two surfaces are in direct contact. A
typical half-penny shaped region at the centre can be observed in figure 9.11a. Figure
9.11b shows the failure originated from sub-surface agglomerate/ porosity. It is
difficult to identify the fracture path because of the higher sub-surface damage
resistance of alumina – CNT nanocomposites (figure 9.9b).
169
(a)
(b)
Figure 9.10. Low energy failure in 3 – point flexural testing: (a) Optical micrograph
(side view) of alumina – 5 wt% uncoated CNT nanocomposites; and (b) schematics as
per ASTM C1161-02c [368].
(a)
(b)
Figure 9.11. Fractured surfaces of alumina – uncoated CNT nanocomposites. Two
tensile surfaces are mounted together. Fracture origin can be identified. CNT
concentration: (a) 2 wt%; and (b) 5 wt%.
170
9.3.6. Thermal shock resistance (qualitative assessment)
Table 9.5 and figure 9.12 show different materials after rapid cooling from 1800 oC.
Thermal shock resistance of ceramics improve with decreasing grain size and
porosity. All materials were fully densified before rapid cooling. Apart from alumina
– CNT nanocomposites, alumina and alumina – CB nanocomposites were found
fractured after rapid thermal shocking (table 9.5).
Table 9.5. Comparison of thermal shock resistance for different materials.
§
Alumina – 2
Alumina – 2
wt% CB
wt% CNT§
43.20
5.48
1.81
Fracture status (Cooling rate: 100 oC/ minute)
NC
NC
NC
PC
NC
NC
CR
CR
NC
Materials/ Properties
Alumina
Grain size (μm)
NC: Not cracked, PC: Partially cracked and CR: Cracked into two or more fragments.
(a)
(b)
(c)
Figure 9.12. Effect of rapid cooling (300 oC/ minute) from 1800 oC after sintering.
Disk diameter: 20 mm; (a) alumina; (b) alumina – 2 wt% CB nanocomposite; and (c)
alumina – 2 wt% uncoated CNTs (Nanocyl, Belgium) nanocomposite.
171
The results can be explained by Hasselman’s theory on the thermal shock resistance
of brittle ceramics [374]. According to the theory, the residual strength of material
after water quenching is an important index of the thermal shock resistance. The
thermal stress fracture resistance parameter R was used to estimate the thermal stress
fracture resistance of materials [374]:
Tmax =
 f (1- )
E
R  K IC 2
Equation 9.4
Equation 9.5
Where, Tmax = maximum quenching temperature before fracture,  f = flexural
strength,  = Poisson ratio, E = elastic modulus,  = coefficient of thermal
expansion and K IC = fracture toughness. The coefficient of thermal expansion of
CNT nanocomposites is quite high when compared to their matrices [196, 375]. It is
clear in table 9.6 that alumina – CNT nanocomposites have higher indentation
fracture toughness, higher flexural strength, finer grains, lower elastic modulus and
higher coefficient of thermal expansion. As per equation 9.4 and 9.5, these
mechanical properties are dominating in increasing the thermal shock resistance of
alumina – CNT nanocomposites. This is a comparative assessment that infers the
alumina – CNT nanocomposites have better thermal shock resistance as compared to
alumina and alumina – CB nanocomposites.
172
#
-
-
-
-
18.9
3.3
308
241
Cracked
Vickers hardness (GPa)
Indentation fracture toughness (MPa√m)
Flexural modulus (GPa)
Fracture strength (MPa)
Thermal shock resistance analysis†
§
5.48
43.20
Grain size (μm)
-
-
Alumina coating (nm)
-
233
201
3.4
11.0
4.67
-
Alumina –
5 wt% CB
Cracked
Alumina –
2 wt% CB
Alumina
Materials/ Properties
-
261
165
5.9
12.0
0.66
-
Alumina –
5 wt% CNT
Cooling rate: 300 oC/ minute
-
Not
cracked§
†
278
263
5.5
14.9
3.23
~5-10
Alumina –
2 wt% CNT
253
234
5.6
10.2
1.81
-
Alumina –
2 wt% CNT
Table 9.6. Summary of mechanical properties of alumina and alumina – carbon nanocomposites#.
-
289
185
5.8
18.6
2.34
~5-10
Alumina –
5 wt% CNT
9.4. Summary of mechanical properties
173
9.5. Conclusions
Raw CNTs form strong entangled networks around alumina grains that is useful in
making flexible (low flexural modulus) and tough ceramic nanocomposites. CNTs
reduced sub-surface damage in alumina – CNT nanocomposites as compared to
alumina. Poor hardness was observed for alumina – uncoated CNTs nanocomposites,
because of the soft lubricating nature of CNTs. A marginal improvement in flexural
strength was observed for alumina – uncoated CNT nanocomposites as compared to
alumina. Uncoated CNTs reduced the grain size of alumina nanocomposite (as
compared to alumina) which proved beneficial for fracture strength and thermal
shock. However, for flexural strength, refining grain size for monolithic alumina is
more effective [372] as compared to carbon addition. Wide scattering in fracture
strength data (low Weibull moduli, appendix C) was observed for alumina – uncoated
CNTs nanocomposites due to poor surface finish as opposed to alumina – coated CNT
nanocomposites. Owing to the poor surface finish, it was difficult to analyse the
intrinsic contribution of uncoated CNTs on the fracture strength of alumina – CNT
nanocomposites.
Effective bonding of CNTs with the matrix plays a vital role in the mechanical
properties of CNT reinforced ceramics. The effectiveness of using alumina coated
CNTs for improved toughness of the nanocomposite was not very significant. Poor
grain boundary adhesion reduced the strength of nanocomposite’s grains that resulted
failure in transgranular mode. Alumina coated CNTs may be useful for applications
where improved IFT as well as good hardness is desired for ceramic nanocomposite.
A good interfacial adhesion is required to increase the stress transfer ability in
alumina – CNT nanocomposites. This will be the subject of future research. Alumina
coated and uncoated CNTs showed higher mechanical reinforcement effectiveness in
alumina nanocomposite as compared to carbon black. This is because of their
geometry and outstanding mechanical properties (table 1.1).
174
Future Work
Conclusions
1.
In the literature, the role of CNTs on the properties of ceramic matrices remains
inconclusive. This is because of the usage of different types and compositions of
CNTs, inhomogeneous distribution of CNTs and difference in experimental
conditions.
2.
The selection of the ultrasonication medium for homogenous dispersion of CNTs
is very important for the final properties of the nanocomposite. From dispersion
stability observations and agglomerate size measurements, it is clear that DMF
produces fine and stable CNT and alumina dispersions. Faster re-agglomeration
was observed in ethanol as compared to DMF for pristine and coated CNTs.
Coating CNTs with alumina reduced the re-agglomeration rate. No evidence of
agglomeration and a good distribution of the CNTs was observed in FE-SEM
micrographs of the SPSed samples when they were mixed with alumina in DMF.
3.
Nanocomposites prepared using DMF dispersions showed better dispersions,
better mechanical properties and higher electrical conductivity as compared to
those prepared using ethanol dispersions. Therefore, it is concluded that DMF is
an ideal dispersant for making homogeneous and agglomerate-free slurries by any
type of colloidal processing.
4.
The addition of CNTs or carbon black to alumina significantly increases its
sintering rate. The sintering temperature required to achieve full densification of
alumina – CNT nanocomposites was reduced by 500 oC as compared to alumina.
5.
CNTs form entangled networks at the grain boundaries, which produce
significant grain growth retardation. Using this effect, an alumina /
nanocomposite laminate structure with a grain size difference of about ten times
was successfully co-sintered.
175
Conclusion and Future Work
6.
CNTs were found to be well preserved in alumina after being SPSed up to
1900 oC/ 100 MPa/ 3 minutes. CNT may not be suitable for ceramics that require
high processing temperatures and longer processing durations. Structural
preservation of CNTs in ceramic nanocomposites depends on the nature of
ceramic, SPS temperature and dwell times.
7.
The electrical conductivity of alumina – CNT nanocomposites is four times
higher as compared to alumina – carbon black nanocomposites for the same mass
content of carbon due to the fibrous nature and high aspect ratio of CNTs. The
electrical conductivity of alumina – CNT nanocomposite increased with
increasing grain size due to the concentration of CNTs in fewer conductive paths.
8.
Fabricating ceramic – CNT nanocomposite by ALD coating of the CNTs and SPS
densification provide shielding that improves the oxidative stability of the
encapsulated CNTs. In ceramic – CNT nanocomposites, coarser grained materials
have higher oxidative resistance as compared to the finer ones, due to the
presence of fewer grain boundaries. The oxidative stability of alumina – CNT
nanocomposites can be tailored by changing the grain boundary area.
9.
CNTs in alumina reduce indentation sub-surface damage as compared to alumina.
The effectiveness of using alumina coated CNTs for improved toughness of the
nanocomposite was not very significant. However, from a mechanical properties
point of view, alumina coated CNT nanocomposite is a good compromise
between hard (alumina) and tough (uncoated CNT nanocomposite) materials.
10. Owing to the poor surface finish, it was difficult to analyse the intrinsic
contribution of uncoated CNTs on the fracture strength of alumina – CNT
nanocomposites. This was not true for coated CNT nanocomposites where
marginal and significant improvements in fracture strength and Weibull modulus
were observed respectively.
176
11. Alumina coated and uncoated CNTs showed higher mechanical reinforcement
effectiveness in alumina nanocomposite as compared to carbon black. This is
because of the geometry and outstanding mechanical properties of CNTs.
12. A qualitative assessment indicated that alumina – CNT nanocomposite has better
thermal shock resistance as compared to alumina and alumina – CB
nanocomposites.
Future Work
1.
CNTs form a strong entangled network around alumina grains. Using longer
CNTs will result in forming much stronger network that will be beneficial for
grain growth retardation and creep resistance at high temperatures. The influence
of quality of CNTs (particularly aspect ratio) on the sintering, grain growth and
properties should be studied.
2.
Breakdown of the conductive networks through deformation or micro-cracking
leads to a large change in the electrical conductivity of alumina – CNT
nanocomposites. The electrical properties could be used to monitor the damage in
sintered nanocomposites. The electrical conductivity could be analysed during
mechanical testing to monitor damage evolution.
3.
The increased electrical conductivity (section 7.3.1) is attributed to the presence
of undamaged CNTs, due to the utilisation of SPS technique that allowed lower
sintering temperatures and shorter sintering times. Other conventional techniques
(hot-pressing and pressureless sintering) should be compared with SPS to gain
deeper understanding.
177
4.
Alumina – CNT nanocomposites were prepared using 2 – 10 wt% (~ 4.7 – 21
vol%) of CNTs. As compared to alumina and coated CNT nanocomposites, a
poor surface finish was observed for uncoated CNT nanocomposites. Better
properties are likely to be achieved at lower volume contents of about 0.5 – 1.5
wt%. The influence of using lower concentrations of CNTs should be studied.
5.
The potential of using different types of CNTs (i.e. DWNT and SWNT) for
improved mechanical properties of ceramic nanocomposites should be
investigated.
6.
The
improvement
in mechanical properties
of alumina coated CNT
nanocomposites was not significant. A better coating like monazite [376] should
be used in future studies. Monazite is already well used to improve the interfacial
properties of ceramic – ceramic fibre nanocomposites [376]. Coated CNTs are
much easier to disperse than the uncoated CNTs and the coating may
significantly improve the strength and toughness of the nanocomposite.
7.
CNTs have outstanding thermal properties. Thermal conductivity and thermal
shock resistance of ceramic – CNT nanocomposites should be investigated.
8.
The potential of various prospective applications (appendix B) will be analysed in
detail, particularly heating element (section B.2.4) and percolating nano-porosity
(section B.2.5).
178
Appendix A. Properties of DMF
Appendix A.
Properties of DMF
A.1. Introduction
N,N-Dimethylformamide (DMF) is a colourless organic liquid, miscible with water
and other organic solvents. It is uniquely versatile and powerful solvent that has a
wide liquid range, good chemical and thermal stability, a high polarity, and a wide
solubility range for both organic and inorganic compounds [377]. The principal
stabilising mechanism for DMF is electrostatic, i.e. the overlap of similarly charged
electric double layers [378]. DMF increases the stability of CNTs by dispersal of
charge [378, 379].
Martelli [380] reported that DMF affects living organs (kidneys and livers) and
digestive tract of pigs when given orally or subcutaneously. Liver disturbances,
stomach complaints, headache, loss of appetite and nausea were found in workers
subjected to less than 20 ppm vapours [377].
A.2. Chemistry of DMF
DMF (figure A) is a derivative of formamide, the amide of formic acid. It is a polar
aprotic solvent, which has dipole moments that help to solvate cations by electron
donation from an oxygen atom [379]. The dipole moments of DMF are several times
as large as that of water. DMF lacks the ability to form hydrogen bonds because all
hydrogen atoms are bonded with carbon [379]. Some of the properties of DMF are
given in table A.
179
Represents resonance of
double bond
Figure A. Chemical formula of DMF, C3H7ON.
Table A. Properties of DMF [377, 381, 382].
#
Properties
Values/ description
Molecular weight
73.1
Boiling point
153 oC
Melting point
-60.4 oC
Flash point
67 oC
Viscosity (at 20 oC)
0.92 cP
Viscosity (at 40 oC)
0.74 cP
Surface tension (at 20 oC)
0.0368 N/m
Surface tension (at 40 oC)
0.0344 N/m
Solubility in water
Infinite
Critical volume
0.25 litres/mol
Dipole moment
3.86 D
Dielectric constant (at 25 oC)
36.71
Electrical conductivity (at 25 oC)
6 x 10-6 S/m
Toxicity threshold limit#
10 ppm or 30 mg/m3
In vapour form
180
Appendix B.
Future of Ceramic – CNT Nanocomposites
B.1. Introduction
The possibilities for the application of CNTs are intriguing and challenging. CNT
based nanocomposites may significantly increase in the near future [383]. Apart from
making nanocomposites, the main avenues of potential applications of CNTs are:
ultimate reinforcement fibres, conducting nanowires, field emitters, nano-tools,
energy storage and energy conversion devices, sensors, drug delivery, medical
diagnostics and cancer therapy [303, 384]. At the moment, commercial usage of
CNTs is quite limited i.e. AFM tips [385] and specialized sports equipments (figure
B.1), like bicycle frames [386], baseball bats [387], tennis rackets [388], ice hockey
sticks [389].
Figure B.1. Different sports goods based on non-ceramic – CNT nanocomposites
[386-389].
181
Appendix B. Future of Ceramic – CNT Nanocomposites
For the development of commercial products, research into ceramic – CNT
nanocomposites is at a very preliminary stage. A commercial product based on
ceramic – CNT nanocomposite is not available yet in the market. This chapter
discusses the future prospects of ceramic – CNT nanocomposites and the issues
relevant to their potential commercial application.
B.2. Applications of ceramic – CNT nanocomposite
CNTs can be added in ceramics for the following reasons:
1. As a sintering aid
2. To control microstructure
3. To improve electrical and thermal conductivity
4. To improve mechanical properties
The first three reasons are the most promising ones. In particular, CNTs are one of the
potential reinforcements for ceramic matrices in microelectronics devices, microwave
devices, consumer products, medical devices, batteries, solid oxide fuel cells,
chemical sensors, gas turbine engines, high-temperature reactors and structural
components that are exposed to high-temperature and aggressive environments [390].
Some of the potential applications of ceramic – CNT nanocomposites are discussed
below:
B.2.1. Conductive ceramic – CNT nanocomposites
Ceramic composites with tailorable electrical conductivity have many industrial
applications. Dense electrically conductive ceramics are used for static charge
dissipation, lightning protection, ceramic heaters, electric discharge machining
182
(EDM), electromagnetic interference shielding in electronic, mechanical, structural,
chemical, and vacuum applications and anti-static floor tiles [324]. In particular,
ceramic alumina with added electrically conductive fillers has been used to fabricate
substrates for handling semiconductor wafers that require static protection.
In contrast, porous electrically conductive ceramics have applications in highperformance radiative heaters, filters for the aeration of liquids, ceramic foam heaters,
and exhaust traps for automotive applications, as well as for the combustion of diesel
soot and the non-catalytic oxidation of noxious gases [324].
B.2.2. Electric Discharge Machining (EDM)
Ceramic materials with precise and complicated shapes can be manufactured by
electric discharge machining. EDM requires materials with a low electrical resistivity,
below ~1 Ωm [26]. Hence insulating ceramics, like alumina cannot be electric
discharge machined. Adding CNTs to alumina significantly reduces the electrical
resistivity (section 7.3.1) and this enhances the electric discharge machinability of
alumina. This can be further employed to other insulating matrices, e.g., boron nitride,
zirconia, silicon nitride. CNTs have a low oxidation temperature (~ 477 oC, section
8.3.1). If CNTs are not required in the microstructure, ceramic – CNT nanocomposite
can then be very easily converted into monolithic ceramic completely by a heat
treatment (up to 700oC, section 8.3.3).
B.2.3. Special purpose electrodes
CNTs can replace graphite in Carbon Ceramic Electrodes (CCEs) for molecule
sensing [391] and water treatment in aggressive environments [392]. This is due to
their large surface area and superior electrochemical properties [393]. Other
183
applications of ceramic – CNT nanocomposites include ceramic microelectrodes for
biomedical applications [394], where electrical conductivity is the main concern.
B.2.4. Heating elements
As a result of outstanding electrical and thermal properties of CNTs, ceramic – CNT
nanocomposites can be used as heating elements (figure B.2). To date, Zhan et al.
[97] reported the highest electrical conductivity for ceramic – CNT nanocomposite,
(3345 S/m), which is quite high as compared to some of the commercially available
SiC heating elements [395]. To prevent oxidation of CNTs, these heating elements
would need to be covered with a protective ceramic layer. Other areas of interest that
might require high electrical and thermal conductivity of ceramic – CNT
nanocomposites include electrical contacts, electrical switches, electromagnetic
interference shielding, electronic devices, thermal plates and electrical/ thermal
pastes.
(a)
(b)
(c)
Figure B.2. Heating element based on alumina – CNT nanocomposite. Sample
diameter is 20 mm: (a) alumina – 5 wt% CNT (Nanocyl, Belgium); (b) demonstration;
and (c) after oxidation, white area shows oxidation of CNTs.
184
B.2.5. Porous structures
CNTs exist at the grain boundaries (section 5.3.3). Oxidising CNTs leaves a
nanoscale, porous network in the ceramic matrix (figure B.3). Because of the low
percolation threshold for CNTs, a nanoporous material could be produced with a
small content (<1 vol%) of porosity. These materials could be used as nano-filters,
bio-scaffolds, thermal insulations and water purification membranes [392].
(a)
(b)
Figure B.3. CNT network on the grain boundaries of alumina grains. (a) Before
oxidation; and (b) After oxidation of CNTs, percolating porosity.
B.2.6. Ceramic armour
Adding CNTs to ceramic armour (e.g., silicon carbide, alumina, boron carbide,
titanium diboride etc.) may improve their ballistic properties. One of the key
properties required for improved ballistic resistance of ceramic armour is fracture
toughness [396]. Ceramic – CNT nanocomposites have demonstrated an increased
sub-surface damage resistance and fracture toughness (section 9.3.3). Apart from
ductility, adding CNTs to ceramics can also improve their sinterability and refine their
microstructure (section 5.3.1 and 5.3.3), which is beneficial for strength of ceramic
nanocomposites as observed in many reports [26, 102, 110, 111, 118-120, 128].
185
B.2.7. Functional Gradient Materials (FGMs)
Functional Gradient Materials (FGMs) are required for the application demanding
different and unique set of properties in different areas of a designed component.
Grain-size FGMs are very advantageous for bio-medical applications [328]. An
alumina / nanocomposite laminate structure was fabricated in this work (section 5.3.4)
to demonstrate the grain refinement effect of the CNTs and their ability to produce
grain size laminated materials.
B.2.8. Nanoceramics
By reducing the grain size to nanoscale, the nanocrystalline monoliths exhibit
strength, hardness and wear resistance, in contrast to those achieved by conventional
composite formation [272]. In section 5.3.4 (figure 5.12-5.14), sub-micron alumina
after oxidizing CNTs from the alumina – CNT nanocomposite was prepared. It is a
lengthy process to make nanoceramics. However, further optimisation of the process
may lead to an advance, commercial way of preparing monolithic nanoceramics using
CNTs.
B.3. Challenges in the development of ceramic – CNT nanocomposite
Section B.2 describes some of the applications that may emerge for ceramic – CNT
nanocomposites. However, for the commercial success of these applications, there are
some immediate obstacles that have to be overcome.
186
B.3.1. CNT related issues
The commercial success of any technology relies on the availability of materials at a
reasonable price. Very little SWNT material is available and most investigators have
been limited to working with milligram quantities. This may be problematic in the
development of ceramic – SWNT nanocomposites, since bench top processing may
not always be representative of larger scale processing [66]. However, the situation
for MWNTs is different as industrial quality MWNTs is available in bulk. The 2006
global production capacity of MWNTs was more than 300 tons/year with the potential
to grow at a significant rate [92]. International manufacturer Bayer Material Science
AG has an annual capacity of 60 tons, and is planning to scale up to 3000 tons by
2010 [61]. Availability in bulk and low cost are the prime factors for the commercial
use of MWNTs rather than SWNTs. But unbundled SWNTs have several advantages
over MWNTs, i.e. enabling lower percolation threshold for electrical conductivity,
reducing the required loading levels in nanocomposites [303]. SWNTs may be more
important is microelectronics, where the quantities is much less of an issue.
Within the time frame of years, prices have been predicted to decrease between 10 to
100 times, depending on the type of nanotubes [22]. Already, Bayer Material Science
AG has started providing industrial grade MWNTs as cheap as £ 0.05/ gram [397].
Future market developments of ceramic – CNT nanocomposites will be fuelled by the
declining prices of high quality CNTs.
Standards for terminology and testing are required to improve understanding and
confidence. They are very important for new materials and designs, as they aid in
determining the reliability and effectiveness of the new products. A nanotechnology
standards debate has been going since 2004. Some organisations, like ASTM, ANSI,
IEEE, BSI and ISO have set some terminologies and standards for nanotechnology
[398]. However, there are no standardized protocols for the evaluation of the quality
of CNTs and their nanocomposites, which is one of the pre-requisites for the
commercial success of ceramic – CNT nanocomposites.
187
Other related issues are the de-bundling of SWNTs, toxicology and the health hazards
associated with CNT exposure.
B.3.2. Understanding of nanocomposites
CNTs exist in different forms, such as CNF, SWNT, DWNT and MWNT. A
comparison needs to be done to see the effectiveness of the different types of CNTs
for ceramic – CNT nanocomposite. A deeper understanding is required of the
fundamental mechanisms associated with the properties of ceramic – CNT
nanocomposites. In this work, the interfacial bonding was improved by using alumina
coated CNTs. A constitutive modelling approach has to be followed to study the
interface, load transfer mechanisms and fracture mechanics of CNT reinforced
ceramics. The interfacial adhesion to the matrix [399], dispersion and stress transfer
[31] of polymer – CNT nanocomposites can be enhanced by chemical
functionalisation of CNTs. However, the role of functionalised CNTs on the
properties and sintering of ceramic – CNTs nanocomposite has not yet explored.
B.4. Conclusions
The future for ceramic – CNT nanocomposites is very bright, especially for
applications concerned with the electrical and thermal properties. However, research
into ceramic – CNT nanocomposites is at very early stage. CNT reinforced ceramic
nanocomposites are not available as a product in the market. Much of the global
research interest is in the development for polymer – based nanocomposites, which
has resulted in some limited commercial success of polymer – CNT nanocomposites.
Apart from a deeper understanding of CNT nanocomposites, CNT related issues must
be resolved for the substitution of CNT based technology in the real world
188
Appendix C.
Weibull Statistical Analysis
Weibull analysis is the empirical statistical distribution function describes the scatter
in strength values of ceramic materials. It is used to assign mechanical properties to
brittle materials in probabilistic terms, and to define design requirements in terms of
strength and reliability [400]. To increase the confidence level, the data sample has to
be sufficiently large (generally ≥10), as a small uncertainty in Weibull modulus can
result in large uncertainties in the survival probabilities. Metals typically have values
of m ~100, while traditional ceramics have values ~5 and engineering ceramics are in
the range 10-25 [400]. The survival probability, i.e., the fraction of samples that
would survive a given stress level can be given by equation C1 [279].
   m 
S p = exp     
   o  
Equation C1
Where S p = survival probability,  = maximum design stress,  o = normalising
parameter (when lnln(1/S) = 0) and m = Weibull modulus.
The Weibull modulus is relatively low for alumina, alumina – uncoated CNT and
alumina – CB nanocomposites (table C). For alumina – carbon nanocomposites, this
is attributed to the poor surface finish as reported in section 9.3.1. Higher Weibull
189
Appendix C. Weibull Statistical Analysis
moduli were observed for coated CNT nanocomposites due to good surface finish
(comparatively) that resulted in less variation in the fracture strength data (table C).
Table C. Weibull modulus (fracture strength) of different materials.
Material
Alumina coating
Weibull modulus
Alumina
-
9.05
-
9.26
~5-10 nm
13.77
-
7.38
~5-10 nm
13.35
-
6.52
CNTs and CB were supplied by NanoDynamics, USA and Cabot, USA respectively.
Tables C1 – C6 and figures C1 – C6 show Weibull statistical analysis for different
materials used in this study. The reason that -lnln(1/S) is plotted rather than lnln(1/S)
is aesthetic, such that the high survival probabilities appear on the upper left-hand
sides of the plots [279]. Note for Weibull modulus analysis, S j was calculated using
equation C2.
 j - 0.3 
S j = 1- 
 N +0.4 
Equation C2
Where S j = survival probability of jth sample, j = sample rank and N = total no. of
samples tested, which was 10 for this study.
190
Table C1. Weibull statistical analysis (fracture strength) for alumina.
Flexural Strength, 
(MPa)
ln (  )
Rank, j
S
-lnln (1/S)
204.97
5.32
1
0.93
2.66
205.37
5.32
2
0.84
1.72
227.80
5.43
3
0.74
1.20
228.76
5.43
4
0.64
0.82
229.59
5.44
5
0.55
0.51
242.87
5.49
6
0.45
0.23
253.66
5.54
7
0.36
-0.03
254.91
5.54
8
0.26
-0.30
259.50
5.56
9
0.16
-0.59
300.95
5.71
10
0.07
-0.99
-lnln(1/S)
m = 9.05
ln (σ)
Figure C1. Weibull plot for data shown in table C1. Weibull modulus: 9.05.
191
Table C2. Weibull statistical analysis (fracture strength) for alumina – 5 wt% CB
nanocomposite. CNTs were supplied by NanoDynamics, USA.
(MPa)
ln (  )
Rank
S
-lnln (1/S)
178.30
5.18
1
0.93
2.66
195.24
5.27
2
0.84
1.72
201.08
5.30
3
0.74
1.20
212.29
5.36
4
0.64
0.82
217.99
5.38
5
0.55
0.51
247.98
5.51
6
0.45
0.23
250.89
5.53
7
0.36
-0.03
254.82
5.54
8
0.26
-0.30
264.32
5.58
9
0.16
-0.59
307.43
5.73
10
0.07
-0.99
-lnln(1/S)
m = 6.52
ln (σ)
192
Table C3. Weibull statistical analysis (fracture strength) for alumina – 2 wt%
uncoated CNT nanocomposite. CNTs were supplied by NanoDynamics, USA.
(MPa)
ln (  )
Rank
S
-lnln (1/S)
199.68
5.30
1
0.93
2.66
229.12
5.43
2
0.84
1.72
230.64
5.44
3
0.74
1.20
244.74
5.50
4
0.64
0.82
247.80
5.51
5
0.55
0.51
255.90
5.54
6
0.45
0.23
257.63
5.55
7
0.36
-0.03
282.87
5.64
8
0.26
-0.30
287.27
5.66
9
0.16
-0.59
293.05
5.68
10
0.07
-0.99
-lnln(1/S)
m = 9.26
ln (σ)
193
Table C4. Weibull statistical analysis (fracture strength) for alumina – 2 wt% coated
CNT nanocomposite. CNTs were supplied by NanoDynamics, USA.
(MPa)
ln (  )
Rank
S
-lnln (1/S)
244.47
5.50
1
0.93
2.66
257.36
5.55
2
0.84
1.72
262.30
5.57
3
0.74
1.20
269.21
5.60
4
0.64
0.82
270.61
5.60
5
0.55
0.51
273.56
5.61
6
0.45
0.23
284.13
5.65
7
0.36
-0.03
295.84
5.69
8
0.26
-0.30
303.32
5.71
9
0.16
-0.59
315.21
5.75
10
0.07
-0.99
-lnln(1/S)
m = 13.77
ln (σ)
194
Table C5. Weibull statistical analysis (fracture strength) for alumina – 5 wt%
uncoated CNT nanocomposite. CNTs were supplied by NanoDynamics, USA.
(MPa)
ln (  )
Rank
S
-lnln (1/S)
184.20
5.22
1
0.93
2.66
239.14
5.48
2
0.84
1.72
242.34
5.49
3
0.74
1.20
249.53
5.52
4
0.64
0.82
257.28
5.55
5
0.55
0.51
268.82
5.59
6
0.45
0.23
277.22
5.62
7
0.36
-0.03
281.40
5.64
8
0.26
-0.30
297.28
5.69
9
0.16
-0.59
310.85
5.74
10
0.07
-0.99
-lnln(1/S)
m = 7.38
ln (σ)
195
Table C6. Weibull statistical analysis (fracture strength) for alumina – 5 wt% coated
CNT nanocomposite. CNTs were supplied by NanoDynamics, USA.
(MPa)
ln (  )
Rank
S
-lnln (1/S)
248.95
5.52
1
0.93
2.66
261.19
5.57
2
0.84
1.72
267.76
5.59
3
0.74
1.20
290.76
5.67
4
0.64
0.82
291.28
5.67
5
0.55
0.51
292.07
5.68
6
0.45
0.23
303.40
5.72
7
0.36
-0.03
304.21
5.72
8
0.26
-0.30
308.41
5.73
9
0.16
-0.59
323.59
5.78
10
0.07
-0.99
-lnln(1/S)
m = 13.35
ln (σ)
196
References
Appendix D.
Recent Publications
(on CNT nanocomposites)
1.
F. Inam, H.X. Yan, T. Peijs, M.J. Reece, "The sintering and grain growth behaviour of
ceramic – carbon nanotube nanocomposites", Journal of the American Ceramic Society,
Ready for submission.
2.
F. Inam, H.X. Yan, M.J. Reece, T. Peijs, “Stability of multiwall carbon nanotubes in sintered
ceramic nanocomposite”, Advances in Applied Ceramics, Accepted (2009).
3.
J. Dusza, G. Blugan, J. Morgiel, J. Kuebler, F. Inam et al., “Hot pressed and Spark Plasma
Sintered zirconia / carbon nanofibre composites”, Journal of the European Ceramic Society,
Article in Press (2009). DOI: 10.1016/j.jeurceramsoc.2009.05.030
4.
F. Inam, H.X. Yan, T. Peijs, M.J. Reece, "Electrically conductive alumina – carbon
nanocomposites prepared by Spark Plasma Sintering", Journal of the European Ceramic
Society, Article in Press (2009). DOI: 10.1016/j.jeurceramsoc.2009.05.045
5.
F. Inam, H.X. Yan, M.J. Reece, T. Peijs, "Dimethylformamide: an effective dispersant for
making ceramic – carbon nanotube composites", Nanotechnology, Vol. 19, No. 19 (2008)
195710 (5 pages)
6.
F. Inam, H.X. Yan, R. Zhang, D. Hua, M. Reece, T. Peijs, "Firing up on all cylinders: Carbon
nanotube based nanocomposites", Materials World, Vol. 15, No. 10 (2007) pp. 24-25
7.
F. Inam, T. Peijs, "Re-aggregation of Carbon Nanotubes in two-component epoxy system",
Journal of Nanostructured Polymers and Nanocomposites, Vol. 2, No. 3 (2006) pp. 87-95
8.
F. Inam, T. Peijs, "Transmission light microscopy of Carbon Nanotubes – epoxy
nanocomposites involving different dispersion methods", Advanced Composite Letters, Vol.
15, No. 1 (2006) pp. 7-15
197
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