1.2714 - Ruhr-Universität Bochum

Transcription

Interface Characterization, Mechanical Properties and Chemical
Interdiffusion Behavior of Hot Direct Extruded Tool Steel Powder
Coatings on Low Alloyed Steel Substrates
Dissertation
zur
Erlangung des Grades
Doktor-Ingenieur
der
Fakultät für Maschinenbau
der Ruhr-Universität Bochum
von
Pedro Augusto de Souza e Silva
aus Belo Horizonte (MG), Brasilien
Bochum, Nov. 2008
Interdiffusion Behavior of Hot Direct Extruded Tool Steel Powder
Coatings on Low Alloyed Steel Substrates
Dissertation
zur
Erlangung des Grades
Doktor-Ingenieur
der
Fakultät für Maschinenbau
der Ruhr-Universität Bochum
von
aus Belo Horizonte (MG), Brasilien
Bochum, Nov. 2008
Dissertation eingereicht am: 25.11.2008...............
Tag der mündlichen Prüfung: 03.02.2009................
(erst bei Druck der Pflichtexemplare)
Erster Referent: Prof. Anke R. Pyzalla..........................
Zweiter Referent: Prof. W. Theisen.........................
Interdiffusion Behaviour of Hot Direct Extrusion of Tool Steel
Powder Coatings on Low Alloyed Steel Substrates
Abstract
In this work, low alloyed steel bars were co-extruded with pre-sintered tool steel powders with
or without the addition of fused tungsten carbides (W2C/WC) as hard particles. During the hot
extrusion process of these massive and powdery materials, an extrudate is formed consisting of a
completely densified wear resistant coating layer and a bulk steel bar as the tough substrate core.
The microstructures at the interfaces between the steel substrate cores and the wear resistant
coatings were characterized by optical and scanning electron microscopy (OM and SEM) in
combination with electron backscatter diffraction (EBSD), energy dispersive X-ray analysis (EDX) and
electron-probe microanalysis (EPMA). Hardness maps and profiles, as well as tensile tests of
miniaturized samples were performed to obtain mechanical properties. Concentration profiles were
calculated using the software DICTRA showing a good agreement with the experimental findings. In
the materials combination K+1.2714 - where K is a gas-atomized cold work tool steel powder
X220CrVMo13-4 (2.39 C, 12.56 Cr, 1.10 Mo, 3.69 V, 0.37 Mn, 0.55 Si) and 1.2714 is a nickel alloyed
hot work tool steel 55NiCrMoV7 (0.56 C, 1.15 Cr, 0.46 Mo, 0.08 V, 0.75 Mn, 0.29 Si, 0.11 Cu, 1.74 Ni)
- carbon diffuses into the substrate material against the concentration gradient due to a higher activity,
leading to an increase of carbide volume fractions close to the interface. The mechanical tests show a
brittle fracture region with high hardness localized about 50µm away from the interface in the coating
material.
The investigation of the microstructure at the interface between coating and substrate of four
hot extruded rods with different coatings revealed the influence of hard particle (HP) addition on the
formation of M7C3 and MC carbides in the coating. It could be verified that the interface region is free
of retained austenite which was expected to be present locally due to an enrichment of carbon at the
interface between substrate and coating.
A combination of experimental measurements (EPMA) and diffusion calculations (DICTRATM)
was carried out in order to investigate the effect of hard particle addition and its dissolution, as well as
the formation of M6C carbides, on the properties of two different PM tool steel coatings hot extruded
with a 1.2714 steel bar. A carburization effect resulting from the W2C hard particles is responsible for
an increase of the 1.2344 steel matrix hardness. The 1.2344 steel is a gas-atomized hot work steel
powder X40CrMoV5-1 (0.40 C, 5.04 Cr, 1.34 Mo, 0.97 V, 0.30 Mn, 0.19 Si, 0.10 Ni). The mechanical
properties of the interface region between coating matrix and substrate are influenced by chemical
interdiffusion of carbon and other alloying elements occurring during hot extrusion.
1I
Acknowledgements
The biggest challenge of my Life started in the summer of 2003, in Berlin, during my IAESTE
internship. The first person I´d like to thank is Dr. Alice Bastos, the brazilian who introduced me to Dr.
Haroldo Pinto and started everything. Haroldo is from the same city as me, Belo Horizonte (MG) in
Brazil. Fortunately, the same soccer team: Clube Atlético Mineiro, the famous GALO. He was moving
from Berlin to Vienna and asked if I knew someone interested in a PhD at the TU Wien, Austria. I
answered “me”, asking when I’d start. I thank you forever for this question. Another important person
in Berlin was Dimas Souza, a brazilian living in the same house as me. He became a loyal and trustful
friend. Thanks, man! Back to Brazil, I had to finish my engineering course and pray for the good news
coming from Vienna. Seven months later I was there starting a PhD in Materials Science without any
knowledge or even had read a paper before. Prof. Pyzalla would be my supervisor and Mr. Pinto my
guide, in the scientific and real life. We survived… both.
Mr. Rodrigo Coelho (Mancebo) and Daniel Silveira (KXA) were and still are my brothers here,
there and everywhere. The group was small, but started to grow. Karolina, Heinz Kaminski, Leozinho,
Mr. Reza, Leo Agudo, Augusta, Moscicki, Fede Sket, Claudia, Murilo and Carla. Most of them are
here now, finishing their dissertations. The argentinians Warcho, Cecília, Guille and hombre Gato, and
the crazy Cynthia were also essential during extra-scientific activities. I thank you all for the moments
we lived in Vienna, Düsseldorf and around Europe. I’ll never forget you guys.
Vienna changed my life completely. I’ve arrived single, without experience in materials science
and to speak German was a dream. I’ve left with a bit of knowledge and married with the Love of my
Life, Juliana Lachini. I believe that Vienna’s role was only this, to meet her and start a new life. Juju,
without you, your help, impulse and our mutual admiration I could never finish this task. I love you,
Amore! But thanks to the South American Mafia, speak German is a dream… still.
In Vienna I’ve started working with TiAl alloys, but in Düsseldorf at the Max-Planck-Institut für
Eisenforschung, the MMC processed by hot extrusion became part of my life. Mr. Pinto said goodbye,
professionally speaking, and Dr. Sebastian Weber said hallo! New city, institute, project, and group
leader. Dr. Weber already got his place in Heaven after being so kind, polite, smart, patient and loyal
during these “hot extrusion, MMC, DICTRA” days. I’ll be forever grateful. Fr. Adrian, our lovely
secretary, appeared in my life. Thank you for everything. And the group continued to grow. The
“República Sulamericana” increased with new (old) faces like Pedro Brito, an old buddy from PUC-MG
and now a great friend. Later on, appeared José Garcia, Maria Maccio, Laís Mujica, David Rojas,
Orlando Prat and Mauro Martin, most of them became my collegues at the Ruhr-Universität Bochum.
The shy-in-the-beginning and now tricky girl Maitena came from France. Dr- Alex Kostka from Poland
is another person who I must say “thank you very much” for all efforts and patience during the EBSD
experiments and also for the fruitful discussions about Formula 1 and Moto GP. To the colleagues
Frank, Jürgen, Hauke, Preilowski, Kryz and Prof. Borbely, and to the IAESTE students Raphael,
Andrea, Lucas and Daniel “Zero Dois”, many thanks! People from other groups like Srdjan, Clara, Luiz
Eleno and Pati Llorente were more than important during this period. Luiz became a brother, forever!
Furthermore, I’d like to gratefully acknowledge the financial support of the DGM within the
project “Strangpressen von Pulverkapseln mit Hartphasen/Metallmatrix-Verbunden auf Fe-Basis”
(DFG-Projekt-Nr.: TH531/3-1&-2 and RE688/58-18.2). For the execution of all the extrusion trials at
the Extrusion Research and Development Center of the Technical University of Berlin I thank Dr. Ing.
K.B. Müller and Prof. Reimers. In addition I’d like to thank Mr. Bialkowski and Monika Nellessen from
the MPIE for the great care and support during sample preparation, Dr. M. Palm and Mrs. I. Wossack
for performing the EPMA measurements and Mr. O. Prat for the fruitful discussions and support during
DICTRATM calculations. Prof. Inden is also gratefully acknowledged for his invaluable help during
diffusion calculations and his kindness. Johnny, the man from Ghana who cleaned your offices almost
everyday and learnt a few words in Portuguese was also an important person for me during these
researching days. To Andreas, who fixed my bike many times, the “Italian” guy from Sri-Lanka, Katja,
Tao and Prof. Fromeyer for his example I also say “thank you”. The people from Bochum were also
important, especially Markus Karlsohn and Arne Röttger for the cooperation. Prof. Theisen is also
acknowledged for his help and cooperation during the project.
Finally, to my lovely wife Juliana, father Guido, mother Jane for the patience and love, sister
Helena + Leocádio, my brothers Lucas, Marcos and Daniel with their respective partners Paulinha and
Candice, my aunt Vânia, cousins Analaurinha + Miminho, Licao, Jota + Samhila = Larota, Juliana’s
parents Jonas and Marisa Meu Amor and my brother and sisters-in-law Leopoldo, Déborah, Maraia +
Henrique and the lovely Matheus, and the essential women Silva, Rosemary and Fatima, thank you all
for your Faith, Patience, Support and Love during these years, months and days of my existence.
Prof. Anke Pyzalla, without your unconditional support, loyalty and especially patience, this
PhD would ever be possible. Thanks a lot!!!
II2
Table of Contents
___________________________________________________________________
Abstract
…………………………………………………………………………… I
Acknowledgements……………………………………………………………………II
Table of Contents …………………………………………………………………… III
List of Abbreviations
…………………………………………………………… VI
Chapter 1: Introduction …………………………………………………………… 1
1.1 Scientific background
1.1.1 Wear Resistance and Casting Restrictions
1.1.2 Hot Isostatic Pressing (HIP)
1.1.3 Hot Direct Extrusion
1.2 Processes
1.3 Aims
1.4 Outline
Chapter 2: Experimental Details………..……………………………....……..… 9
2.1 Materials selection
2.1.1 Coating Matrices
2.1.2 Steel Substrate Bars
2.2 Heat treatments
2.3 Characterization techniques
2.3.1 Microstructure
2.3.1.1 Optical Microscopy (OM)
2.3.1.2 Scanning Electron Microscopy (SEM)
2.3.1.3 Electron Probe Micro Analysis (EPMA)
2.3.1.4 Electron Backscattered Diffraction (EBSD)
2.3.2 Mechanical properties
2.3.2.1 Tensile Tests
2.3.2.2 Hardness Tests
2.3.3 Diffusion calculations (DICTRATM / ThermoCalcTM)
2.3.3.1 ThermoCalcTM
2.3.3.2 DICTRATM
III3
2.4 Proposed literature
Chapter 3: Interface Characterization and Mechanical Properties of the Cold
Work Steel Coating (K) Co-Extruded on a 1.2714 Steel Substrate
………………………………………………………..…………………. 22
3.1 Introduction
3.2 Materials Processing and Experiments
3.2.1 Metal Matrix Substrate and Coating
3.2.2 Hot Extrusion Process
3.2.3 Metallography and Microscopy
3.2.4 Heat Treatment
3.2.5 Electron Probe Micro Analysis (EPMA)
3.2.6 Diffusion Calculations with DICTRATM
3.2.7 Mechanical Properties
3.3 Results and Discussion
3.3.2 Measured Element Distributions
3.3.3 Calculated Element Profiles
3.3.4.2 Micro Hardness
Chapter 4: Correlation between Interface Microstructures and Mechanical
Properties of Co-Extruded Layered Structures…………………. 39
4.1 Introduction
Chapter 5:
Microstructure Characterization by EBSD/EDX Focusing on the
Influence of Hard Particle Addition and the Formation of Retained
Austenite …..………………………………………………………….. 46
5.1 Introduction
5.2.1 Metal Matrix Substrate and Powder Steel Coatings
5.2.2 Hard Phases and Sample Designation
4
IV
5.3.2 Retained Austenite (RA or γ-Fe) and Vanadium Carbides (VC)
Chapter 6: Influence of Hard Particle Addition and Chemical Interdiffusion
Investigated by Diffusion Calculations on the Mechanical
Properties .………………………………………………………..…… 58
6.1 Introduction
6.2.1 Materials Processing
6.2.3 Hardness Measurements
6.3.1 Influence of FTC on Hardness
6.3.2 Dissolution of W2C in the 1.2344 Coating and Formation of M6C
6.3.3 Influence of W2C Hard Particle on 1.2344 Coating Matrix
6.3.4 Interactions of the Coatings WW1 and KW1 with the 1.2714 Steel
Substrate
Chapter 7: Conclusions ………………………………………………………….. 72
7.1 General Remarks
7.2 Specific Remarks
7.3 Outlook and Future Works
7.4 Suggestions for Industrial Applications
Chapter 8: References……………………………………………………………… 79
Curriculum Vitae …………………………………………………………………… 81
V5
List of abbreviations
___________________________________________________________________
EDM
Electro Discharge Machining
OM
Optical Microscopy
SEM
Scanning Electron Microscopy
SE
Secondary Electron
BSE
Backscattered Electron
EDX
Energy Dispersive X-ray Spectroscopy
EPMA
Electron Probe Microanalysis
WDS
Wave Length Dispersive Spectometer
EBSD
Electron Backscatter Diffraction
OIM
Orientation Image Microscopy
EX
as EXtruded
HT
Heat Treated
QT
Quenched and Tempered
DICTRATM
DIffusion
Controlled
TRAnsformation
software
(ThermoCalc
AB,
Stockholm, Sweden)
TC
ThermoCalcTM software (ThermoCalc AB, Stockholm, Sweden)
HIP
Hot Isostatic Pressing
MMC
Metal Matrix Composite
PM
Powder Metallurgy
FTC
Fused Tungsten Carbide (W2C/WC)
HP
Hard Particles
RA
Retained Austenite
Rp0,2
Yield strength
Wt. %
Weight percent
VI
6
Introduction
1
___________________________________________________________________
1.1
Scientific background
1.1.1 Wear Resistance and Casting Restrictions
The necessity of high abrasion resistant materials in applications, for instance
in the mining and cement industry, led to the development of metal matrix
composites (MMC) produced by powder metallurgy (PM) to overcome the restrictions
of casting. For many purposes the wear resistant material is not necessary or even
useful for the whole tool, but only the near-surface region. Such a layered structure
necessitates the cladding of the wear resistant material onto a dissimilar substrate.
Metal matrix composites (MMC) produced by powder metallurgy and based on heat
treatable steel matrices with or without dispersed hard particles exhibit a higher
resistance in certain wear applications compared to conventional materials [1], e.g.
white cast iron. In cast materials the evolution of the microstructure is mainly
dominated by the alloy system and the formation of non stoichiometric mixed phases.
1.1.2 HIP
In recent years, several concepts have been developed, one of them being
cladding by hot isostatic pressing (HIP) [2], for producing thick wear resistant
coatings on tough substrate materials. To obtain PM-MMC, a steel metal powder
serving as matrix material is mixed with hard phase particles (HP) and consolidated
within a gas-tight capsule in a HIP furnace. So far, hot isostatic pressing (HIP) has
been the usual method for producing low alloyed steel rods clad with thick wear
resistant layers of MMCs.
1
For this purpose, typically pre-alloyed tool steel powders showing a small
sintering activity are used as matrix materials exhibiting powder grain sizes ranging
from 40µm to 150µm. A powder mixture is filled into a gas-tight capsule and
consolidated by HIP. This way of consolidation leads to several limitations of the
process: the welded capsule has to be completely gas-tight to ensure densification of
the material. Furthermore, the maximum size of a component is limited by the size of
available HIP furnaces.
A high resistance against abrasive wear has been achieved by the
development of metal matrix composites with coarse carbide hard phases [3]. These
hard phase reinforced steel composites so far could only be clad onto a steel
substrate by hot isostatic pressing [2]. Compared to hot isostatic pressing hot
extrusion of the mixture of steel powder and hard phase powder onto a steel
substrate appears beneficial with respect to production costs, product size, and
versatility. This production process will be described in details in section 1.2.1 of this
work.
1.1.3 Hot direct extrusion
Hot direct extrusion has been recently introduced as a novel process for the
production of low alloy steel rods clad with MMCs as high wear resistant coatings [4].
In contrast to HIP the hot extrusion process allows the cost efficient production of
long products with full density of the coating and comparable properties and is thus
considered as a possible alternative for producing semi-finished parts with a large
aspect ratio. Furthermore it provides a full densification of the material by the
hydrostatic pressure in the die of the extrusion press. It was recently found out that
hot direct extrusion is a feasible and cost efficient process for the production of these
PM-MMCs with tool steel matrices.
Therefore a novel manufacturing route via direct hot extrusion of bulk steel
bars and pre-sintered tool steel powders, partly mixed with hard particles, was
developed [4]. The different steps of this process are comparable to those of HIP
processing. The powder or powder mixture is filled into a capsule of large wall
thickness, which is evacuated, sealed, pre-heated in a furnace for several hours and
subsequently pressed. During pre-heating sintering of the powder particles takes
place, influencing the deformation behavior of the capsule in the extrusion press [5].
2
The basic concepts of this process and resulting materials are described in detail
here [6].
. The extrusion trials were performed at the Extrusion Research &
Development Center of TU Berlin and cylindrical rods consisting of claddings of
either steel MMCs with hard phases or tool steel on lower alloyed steel substrates
were successfully produced [3].
Further investigations considered direct hot extrusion as an alternative to HIP
cladding for producing thick wear resistant coatings on low alloyed substrates.
Therefore, a massive steel bar was inserted into the capsule and the retained cavity
filled with powder. The co-extrusion of these massive and powdery materials leads to
a complete densification and the formation of a tough substrate coated with a thick
wear resistant layer. This production process will be described in details in section
1.2.2 of this work.
1.2
Processes
HIP is used to eliminate porosity from cast or sintered components and
consolidate encapsulated powders to provide fully dense materials with excellent
properties. Searching for ways of improving the mechanical properties of a material,
especially in critical, highly stressed applications and abrasive environments, the use
of HIP as a method for producing components from different powdery materials has
become well established. Dissimilar materials can be clad together to produce
unique, cost effective components.
In order to reduce the porosity of metals, improve the mechanical properties
and workability, a container with a powder mixture of, e.g. tool steel powder, is
subjected to high temperatures after vacuum is used to remove air and moisture from
the powder. After sealing the container, an inert gas is applied in high pressure to the
material from all directions in such a way that no chemical reactions occur. This
results in the removal of internal voids and creates a strong mechanical bonding
throughout a homogeneous material, uniform grain size and full density.
During the HIP process, internal voids are eliminated, clean and solid bonds
are created and a fine, uniform microstructure is produced. Welding or casting
3
materials do not possess these characteristics. Another advantage is the
improvement of fatigue strength due to the nearly complete elimination of internal
voids and microporosity through a combination of plastic deformation, creep and
diffusion bonding.
Initial advantages are the reduction of micro-shrinkage and the consolidation
of powder metals, ceramic composites and metal cladding. HIP is also used as part
of a sintering process (powder metallurgy) and for manufacturing metal matrix
composites (MMC).
The HIP process has the ability to create near-net shaped parts that require
little machining using 80-90% of the purchased material. This clear advantage in
comparison with conventional manufacturing methods reduces costs and machining
time significantly. Almost all shapes and sizes can be produced by HIP, including
cylindrical billets and solid shapes with complex external geometry and shape. This
manufacturing
process
enables
the
production
of
materials
from
metallic
compositions that are difficult or even impossible to forge or cast.
The HIP process is not only used for densifying castings, but in other areas
like powder metal consolidation and diffusion bonding of dissimilar materials. Besides
porosity elimination from welding, casting and sintered materials, three other areas
can be identified for the application of HIP: claddings, production of near-net shaped
parts and consolidation of powder metals (PM).
A common application of HIP in the production of abrasion and wear resistant
materials takes place by cladding as a selective bonding of hardfacing materials into
various substrate surfaces. The basic idea is to coat a less expensive material with a
thin or a thick layer of powdery metal, depending on the application, creating a
protection on its wear surface. The reduction of costs results from applying expensive
wear resistant materials only where they are necessary. An increase in the wear
resistant properties also occurs without wasting valuable resources. Another aspect
of cladding is that incompatible materials such as metals, intermetallics and ceramic
powders can be bonded together.
A wide range of different applications, including chemical processing,
petroleum, medical and automotive industries, makes use of HIP products. A higher
freedom of design when compared with forging and casting satisfies product and
project engineers. By eliminating shrinkage, porosity defects and reaching
4
mechanical property requirements, HIP was developed and now provides the
possibility to produce fine grained materials combined with the desired high density.
Nowadays, the HIP process is used also in the aerospace industry in the
production of rocket engines, satellites and aerospace airframe castings.
As with any technology, awareness by industry is the key to growth. Costs
also pay a key role in the development and popularization of any production process.
As an established technique to clad parts with thick abrasion resistant layers, the HIP
process unfortunately had already reached limitations on the size of the equipments.
Besides, this process is cost intensive and the necessity for near-net shaped parts
avoided the acceptance of this method in a wider range of applications.
Knowledge of several extrusion techniques, such as direct, indirect, cold, hot
and hydrostatic allowed the development of a novel and alternative method to clad
rods with coatings based on a metal matrix composite (MMC) by hot direct extrusion.
Basically, this method forces a metal or an encapsulated metal powder to flow
from a container through a cavity determined by a die, and a mandrel fixed to the
press ram. The main advantage of a powder extrusion is to achieve a desired shape
which cannot be easily obtained by conventional methods.
Many variations of extrusion methods are encountered. In our case,
considering the movement of the extrusion with respect to the ram, the die remains
stationary and the ram moves towards it. In this configuration, the extrusion process
is called direct. The press is positioned horizontally and a hydraulic drive performs
the extrusion using hydrostatic pressure enabling the consolidation of the material to
full density.
The extrusion of metal powders possesses several characteristics that make
this method a powerful manufacturing technique:
•
the possibility to create shapes and/or forms from materials that are difficult or
impossible to obtain by casting or forging;
•
the capacity to produce wrought structures without sintering or other thermal
treatments;
•
mechanical properties are improved due to minimisation of segregation and
microstructural refinement obtained from powder processing;
5
•
the dissipation of one material in another because of the extrusion of powder
mixtures;
•
smaller extrusion pressures, wider temperature and ram speed ranges for
powder extrusion in comparison of extruded cast billets.
Other important advantages of the hot extrusion process are the possibility to
manufacture complex cross-sections and to process brittle materials, due to the fact
that only compressive and shear stresses are acting in the materials during
processing.
Thus, hot direct extrusion appears as a promising and cost-efficient
manufacturing route allowing the production of Fe-base metal matrix composites
(MMC) and an alternative to HIP, specially suited for the production of long semifinished products. The need of full density and the low sintering activity does not
allow a sintering route for these materials. Capsules similar to the ones used in HIP,
filled with steel metal powder or a powder mixture with hard particles, are hot direct
extruded after being closed, evacuated, sealed and pre-heated for several hours in a
furnace to processing temperature. During the pre-heating state, sintering of the steel
powder occurs influencing the deformation behaviour of the capsule. Due to the hot
extrusion, the powder mixture is consolidated and bonded to the substrate.
The costs for machinery and its maintenance are the main disadvantages of
this production method. The process starts to be economically feasible when
producing between several kilos to many tons, depending on the extruded material.
1.3
Aims
Within this work, different combinations of substrates and coating materials
were investigated. The main focus of this work is the characterization of the coatingsubstrate interface region formed by different configurations of PM steel coatings and
steel substrates and the applied heat treatment.
A systematic characterization of the microstructures formed in the interface
region as well as the formation of these microstructures and their effect on the
mechanical properties was performed.
The diffusion mechanisms between the different coatings and substrates were
analyzed and correlated with mechanical properties such as hardness and yield
6
strength. The influence of alloying elements, formation of carbides, and dissolution of
hard particles in the coating matrix and the presence of retained austenite were also
investigated and elucidated.
Furthermore, the influence of heat treatments and the hard particle additions in
the coating matrices were studied in detail and related with the diffusion processes in
the interface region between coating and substrate, especially carbon diffusion.
1.4
Outline
This dissertation is structured in seven chapters, starting with this
Introduction, in the following sequence:
In Chapter 2 a detailed description of the experiments is presented, including
materials selection, preparation for the extrusion trials, sample preparation and a
brief citation about the equipments and instruments used for the characterization of
the specimens.
This work presents different characterization methods and investigates
different materials combinations using several approaches. To facilitate reading, a
brief introduction followed by experimental details, results and discussions
corresponding to each set of analyzed specimens and investigations are organized
separately in the Chapters 3, 4, 5 and 6. These chapters correspond to original
publications already submitted or published in scientific journals during the course of
this work.
In Chapter 3 the simplest materials combination (without hard particles in the
coating) is characterized, mechanical properties are presented and measured
element distributions are compared to calculated concentration profiles. Carbon plays
a key role for the diffusion processes influencing the mechanical properties locally.
Chapter 4 presents the comparison between three materials combinations
including a second coating steel powder and the addition of hard particles coextruded with the same steel substrate of 1.2714. Characterization techniques using
EBSD and differences in carbon activity in the coatings and substrates are correlated
with mechanical properties revealing a carburization/decarburization effect and its
impact on the mechanical properties.
7
Chapter 5 depicts the work using EBSD carried out in the investigation of four
different configurations of hot extruded bars, adding now a second steel substrate.
The influence of hard particle addition on the formation of M7C3 and MC carbides in
the coating matrix started to be understood and the presence of retained austenite at
the interface region is investigated.
In Chapter 6 a specific effect is deeply analyzed: the hard particle addition
and, especially, its dissolution on the properties of two different PM tool steel
coatings hot extruded with a 1.2714 steel bar. The effect of carburization from the
W2C hard particles, formation of M6C carbides and how the chemical interdiffusion of
carbon and other alloying elements during heat treatment influence the mechanical
properties are revealed.
Chapter 7 shows general and specific concluding remarks. An outlook for
future works and industrial applications is given for the investigated specimens
produced by the novel manufacturing process.
Chapter 8 contains the literature used as references during the execution of
this dissertation.
8
Experimental Details
2.1
Materials selection
The main conditions to be fulfilled by the materials selected to produce a
successful hot extruded bar are:
•
a high fracture toughness substrate core and;
•
a wear resistant coating layer.
2.1.1 Coating matrices
The two selected tool steel powder coatings are the 1.2380 and the 1.2344
tool steels. The gas-atomized steel powder of 1.2380 (X220CrVMo13-4) is a
ledeburitic cold work steel and was selected as the metal coating considering the
wear resistance of the matrix material. In the quenched and tempered conditions its
microstructure is formed by tempered martensite as well as chromium-rich M7C3 and
vanadium-rich MC carbides increasing its hardness. This steel is commonly used for
die cutting tools.
The gas-atomized hot work tool steel powder 1.2344 (X40CrMoV5-1) exhibits
a high wear and thermal shock resistance in a temperature range of 400-700°C, as
well as a high level of toughness and ductility. Depending on the heat treatment
applied, it can reach a typical hardness of 50-56 HRC containing virtually no carbides
in its martensitic microstructure. This steel can be used as a standard material for hot
forming and extrusion tools, forging dies, pressure casting tools, hot shear knives
and also as tools for the plastic industry.
9
2.1.2 Steel substrate bars
As a substrate steel bar, the hot work steel 1.2714 (55NiCrMoV7) and the nonalloy structural steel S355 were chosen. The nickel alloyed hot work tool steel 1.2714
was selected because of its good hardenability, tempering resistance and
dimensional stability as well as very good strength and excellent toughness. The
main applications are for dies, tools for rod and tube extrusion, forming dies and
plastic moulds. The steel S355 was chosen due to its low cost, good weldability, cold
formability and a high fatigue limit.
2.2
Heat treatments
The major intention of heat treating the extruded bars is to obtain a sufficient
hardness and toughness of the coating material. However, using extended heat
treatment times it was possible to change the chemical gradient at the interface
region in both directions, coating and substrate, as well as the microstructure directly
at the interface region.
As an example, depending on the carbon activity, diffusion from the substrate
to the coating may occur. After quenching, this was supposed to increase the amount
of retained austenite at the interface, especially on the coating side, influencing the
mechanical properties locally.
2.3
Characterization techniques
All specimens analyzed in this work were cut parallel to the extrusion direction
by electro discharge machining (EDM) in order to reduce the influence of cutting on
the microstructure and to keep the interface region parallel to the extrusion flow.
There is a wide range of aspects in the analysis of the microstructure,
including size, shape, and orientation of grains, chemical composition and the
correlations of these characteristics with the physical properties like yield strength,
hardness, ductility, fracture toughness and chemical properties like the diffusion
behaviour.
Regarding the microstructural investigations conducted in this work, the
following micro-analytical tools were used: optical microscopy (OM), scanning
10
electron microscopy (SEM), electron probe micro analysis (EPMA) and electron
backscattered diffraction (EBSD).
In order to highlight how these tools were applied during this work, examples
of some microstructural investigations are showed. In section 2.4, more details of the
equipment and the physical fundaments behind it are presented as “proposed
literature”.
2.3.1.1 Optical Microscopy (OM)
The first analyzes of the specimens were conducted by OM aiming to compare
the different microstructures formed and to measure the width of the interface region
between coating and substrate. Polished and eventually etched surfaces were used
to assess the basic characteristics of the microstructures showing grain and phase
boundaries, size of α-martensitic laths as well as pores and/or cracks, especially in
the interface region. This task was conducted using a Leica DM4000 optical
microscope.
Sample preparation
A successful investigation in the materials science field is directly related to a
smart and useful specimen preparation. The etching methods and the correct
chemical agent and time also play a key role to obtain satisfactory results in an OM
analysis.
This
work
was
conducted
applying
standard
metallographic
sample
preparation for all specimens:
-
grinding manually with water from 54µm to 15µm in special discs designed for
hard metals. This step makes the surface as flat as possible and removes the
first scratches;
-
polishing in rotating machines with diamond suspension fluids from 3µm to
0.25µm;
-
short final polishing with a finely napped disc using ethanol for cleaning
purposes;
-
cleaning specimens with ethanol in an ultrasonic bath between each polishing
step;
Nevertheless, each materials combination investigated presented its particular
problems of specimen preparation. The coating materials were always harder than
11
the substrates and the interface region is a mix between them. When one side was
perfectly polished, the other still needed more polishing time. To find a balance
between them was the most difficult task during sample preparation.
All specimens were etched with Nital 3% in order to reveal grain and phase
boundaries and the α-martensitic laths of the substrate steels.
Examples of the investigated microstructures
OM is a fast and easy technique and was applied in this work to characterize
different materials combinations and to obtain important information from the
investigated microstructures.
In the interface region between coating and substrate, the initial analysis
focused on the width of the different interfaces formed and on the formation of pores
and cracks. Figure 1 shows the limits of the interface region formed between the coextruded WW1 steel coating and the S355 structural steel. The analysis of not
expected α-martensitic laths close to the interface, the orientation of W2C/WC
particles and chromium- and vanadium-rich carbides (Fig. 1d) is an example of
features observed using OM.
a)
S355
WW1
b)
S355
WW1
WC/W2C
WC/W2C
12
c)
S355
WW1
d)
S355
WW1
WC/W2C
WC/W2C
Figure 1: Examples of the use of OM in this work: a) an overview of the microstructure formed in the
materials combination KW1+S355, b) and c) analysis of the interface region width and pore formation,
d) interface region morphology and a W2C/WC hard particle close to the interface region, on the
coating side (right hand site).
In addition, OM is suitable to correlate the microstructural morphology with
local mechanical properties. In Figure 2 the correlation between hardness
indentations and the microstructure is depicted revealing a higher hardness in the
FTC particle core followed by the η-carbide diffusion seam and the coating matrix.
a)
b)
WC/W2C
WC/W2C
10 µm
10 µm
Figure 2: Hardness indentations around the W2C/WC particle revealing a higher hardness in the FTC
particle core followed by the η-carbide diffusion seam and the coating matrix. The WW1+1.2714
specimens were a) hardened, and b) hardened and tempered.
2.3.1.2 Scanning Electron Microscopy (SEM)
SEM analysis combines high magnification, larger focus depth, greater
resolution and easiness of sample observation making this instrument one of the
most used in materials research today. Analysis of the microstructural morphology,
failure mode, and chemical composition by EDX could be performed using the SEM.
13
The study of fracture and failure were carried out on rough surfaces from tensile test
specimens right after the performed test. Moreover, hardness tested samples were
also investigated with SEM in order to correlate local mechanical properties with the
microstructure.
These tasks were conducted using a Jeol JSM-6500F field emission
microscope and a Jeol JSM-6490 tungsten filament instrument, both equipped with
the EDAX-TSL EBSD software and a Zeiss Neon 40 field emission microscope
equipped with the Hikari EDAX-TSL EBSD software.
Sample preparation
The specimens were prepared using the same receipt already described for
OM investigations. Rough surfaces designed for fractography analysis did not require
any special preparation.
The power of this tool in microstructure analysis can disclose details not
clearly revealed with OM. The formation and chemical composition of carbides,
presence of pores and the interface width could be better seen using this technique
combined with EDX.
As
an
example
of
applying
the SEM in the microstructural
characterization, Figure 3 shows the shape, size and orientation of the chromium-rich
M7C3 carbides and vanadium-rich MC carbides close to the interface region and to
the FTC hard particles.
a)
b)
WC/W2C
1.2714
KW1
1.2714
KW1
14
c)
d)
WC/W2C
M6C
WC/W2C
1.2714
KW1
1.2714
KW1
Figure 3: Examples of the use of SEM in this work: a) an overview of the microstructure formed in the
materials combination KW1+1.2714, b) reveals an interface region width of ~15µm, c) and d)
formation of Cr7C3 and VC carbides in the coating steel KW1 (right hand side).
Another example using SEM/EDX showed in Figure 4 reveals the chemical
composition of the M7C3 type chromium- and MC type vanadium-rich carbides
formed in the 1.2380 (K) coating steel microstructure.
b)
a)
11µm
µm
1 µm
c)
V
1 µm
Cr
d)
Fe
1 µm
Figure 4: EDX map showing the chemical composition of the carbides formed in the materials
combination KW1+1.2714: a) SEM micrograph, b) chromium map, c) vanadium map, d) iron map.
15
Moreover, SEM proved to be an interesting research tool applied for the
analysis of failure mechanisms. A brittle failure mechanism by cleavage fracture
mode is depicted in a fractography (Fig. 5) in the coating side of the materials
combination K+1.2714.
Figure 5: Fractography showing a brittle failure mechanism by cleavage fracture mode after tensile
test in the materials combination K+1.2714.
2.3.1.3 Electron Probe Micro Analysis (EPMA)
An Electron Probe Microanalyser (EPMA) is an instrument to determine and
analyze the chemical composition of materials in a non-destructive way. The surface
of a sample is scanned with an electron beam and the signals coming from it are
collected, working similar to a SEM. The beam makes the elements from the sample
surface to emit X-rays with a characteristic energy being detected wavelengthdispersive by the system.
The chemical composition is determined by comparing the intensities of
characteristic X-rays from the sample material with intensities from standard
16
compositions. The determination of any variation in chemical composition in a
material can be easily performed.
Sample preparation
OM investigations, but with a final step polishing manually with diamond suspension
spray down to 0.25µm.
2.3.1.4 Electron Backscattered Diffraction (EBSD)
EBSD is a microstructural-crystallographic technique which, combined with
SEM, allows the investigation of the crystallographic orientation of a microstructure
by the analyses of the diffraction of backscattered electrons. This tool can be used to
index and identify crystal systems, crystal orientation, phase identification, grain
boundary and morphology studies.
A polished and flat sample is inserted into the SEM chamber and tilted to 70°
towards the camera. The stationary electron beam hits the sample surface, the
information of the crystal structure being analyzed is detected on a fluorescent
screen as a diffraction pattern (Kikuchi pattern), satisfying Bragg conditions. The
diffraction patterns are indexed according to the Miller indices and used to identify
phases, to reveal crystal orientations and grain boundaries, and, when combined with
EDX, to measure chemical compositions.
The result of a scanned area of interest is the formation of orientation maps
and a qualitative and quantitative representation of the microstructure processed by
the OIM software. Maps of the microstructure, charts and plots can be produced
showing easily the grain morphology, orientation, phase- and grain-boundaries.
Sample preparation
OM investigations. However, for an EBSD analysis the specimen’s surface must be
as flat as possible avoiding any roughness and micro-scratches. In order to reach
this condition, that task was conducted applying the same standard metallographic
sample preparation described for OM investigations adding important steps
according to the following receipt:
17
-
grinding manually with water from 54µm to 15µm in special discs designed for
hard metals. This step makes the surface as flat as possible and removes the
first scratches;
-
polishing in rotating machines using the Struers disc Dur with high pressure (5
bar) using a 3µm diamond suspension fluid and blue lubricant during 10-15
min. The machine should rotate in the opposite direction of the disc;
-
polishing manually with high pressure using a 1µm diamond suspension fluid
and blue lubricant during 1-2 min in a special disc;
-
polishing manually with high pressure using a 0.25µm diamond suspension
spray and pink lubricant during 30”-1’ in a special disc;
-
final polishing (30”–1’) with the finely napped disc Chem from Struers using
SiO2 suspension, soap and few drops of water;
-
final polishing (30”) with the finely napped disc Chem from Struers using
ethanol for cleaning purposes and to remove SiO2 small particles from the
coating surface;
-
cleaning specimens with ethanol in an ultrasonic bath between each polishing
step;
The main advantage of an investigation using EBSD is the wide range of
information obtained from a scanned area, which is not readily acquired by using
conventional methods. The changes in the crystallographic orientation created by the
hot extrusion in the KW1 coating matrix and substrate steel are shown in Figure 6 as
an example of the application of this technique.
1.2714
KW1
Figure 6: EBSD scan of the materials combination KW1+1.2714 showing the crystallographic
orientation of the grains. The interface region is in the centre of the micrograph.
18
The reliability of the interface region depends on the inter-atomic bonding
developed during the pre-heating, the hot extrusion process, heat treatment and aircooling. The quality of the interface region was determined by mechanical tests, such
as tensile tests and hardness measurements.
2.3.2.1 Tensile tests
The weakest link within the chain “coating layer-substrate-coating layer” could
be determined. The specimens were clamped for mounting in the testing facility. The
miniaturized tensile test specimens (Fig. 7) were extracted by spark erosion and
grinded on the top and bottom side removing oxides and smoothing the surface. The
tests were performed according to the DIN 10002-2 standard.
a)
b)
interface
substrate
coating
Figure 7: a) Macroscopic view of the cross section showing the external capsule, coating, substrate
(~8mm thickness) and the interface region, and b) sketch of the miniaturized specimens taken from
the cross section of an extruded bar used for the tensile tests.
2.3.2.2 Hardness tests
The hardness tests were carried out on polished specimens and the
indentations were done perpendicular to the extrusion direction according to the DIN
50359-1 (1997) Universal hardness standard and ASTM E384-99 Vickers microhardness standard. Line profiles and 2D-maps were measured.
19
2.3.3 Diffusion calculations (DICTRATM / ThermoCalcTM)
To support the experimental results of the heat treatments, diffusion
calculations using the software DICTRATM were performed focusing on the
interdiffusion of alloying elements at the interface region during processing and heat
treatment. Detailed information of the latest versions and about the software
company can be found at http://www.thermocalc.se.
2.3.3.1 ThermoCalcTM
The Thermo-CalcTM software is widely spread around the world and has
numerous users, probably being the most frequently used thermodynamic simulation
software worldwide. The software not only performs standard equilibrium calculations
and calculation of thermodynamic quantities based on thermodynamic databases,
but is also equipped with some unique features in special modules for special types
of calculations for the advanced user.
This work was performed with the TCC version R from June 20th 2007 using
the TCFE4 database and, to obtain mobility data, the MOB2 Thermo-CalcTM
database [10] was used.
2.3.3.2 DICTRATM
DICTRATM is the pioneering software for accurate simulations of diffusion in
multi-component alloy systems. DICTRATM is coupled with Thermo-CalcTM for
necessary thermodynamic calculations and has been applied to numerous problems
of practical and scientific interest.
This work was performed with the version 2.4 from June 20th 2007 using the
same thermodynamic and mobility databases used in Thermo-CalcTM.
2.4
Proposed literature
Materials Information and Heat Treatment
•
Berns, H. and Theisen, W. – Ferrous Materials – Steel and Cast Iron, 2008.
•
Jones, R.M. - Mechanics of Composite Materials, 1999.
•
ASM Handbook, Vol. 4 – Heat Treating, ASM International, 1991.
•
Heat Treating Processes and Related Technology, ASM International, 1995.
•
Guidelines for Heat Treating of Steel, ASM International, 1995.
20
Characterization techniques
•
Brandon, D., Kaplan, W.D. – Microstructural Characterization of Materials, Second Ed. –
John Wiley & Sons Ltd., 2008.
•
ASM Handbook, Vol. 1 – Properties and Selection Irons Steels and High Performance
Alloys, ASM International, 1990.
•
ASM Handbook, Vol. 8 – Mechanical Testing and Evaluation, ASM International, 2000.
•
ASM Handbook, Vol. 9 – Metallography and Microstructures, ASM International, 1985.
•
ASM Handbook, Vol. 10 – Materials Characterization, ASM International, 1986.
•
ASM Handbook, Vol. 11 – Failure Analysis and Preventions, ASM International, 2002.
•
ASM Handbook, Vol. 12 – Fractography, ASM International, 1987.
21
Interface Characterization and Mechanical
Properties of the Cold Work Steel Coating (K)
Co-Extruded on a 1.2714 Steel Substrate [19]
3.1
Introduction
Materials with high resistance against abrasive wear are of interest for many
tool applications e.g. in mining industry. A special issue is the cladding of these
materials to low alloyed substrates for new protection purposes. A novel
manufacturing route via hot direct extrusion of bulk steel bars and pre-sintered tool
steel powders was applied. In this manner, wear resistant claddings of PM tool steels
on steel substrates were obtained. A further development of the hot extrusion of
abrasion resistant coatings was achieved by inserting a massive steel bar into the
capsule and filling the retained cavity with powder. The co-extrusion of this massive
and powdery materials leads to a complete densification and the formation of a tough
substrate coated with a thick wear resistant layer. Depending on the materials
combination and the heat treatment, differences in the formation of the interface
substrate-coating can be determined.
The characterization of the interface region, taking processing parameters into
account, is the focus of this chapter. The microstructures at the interface between the
steel substrate cores and the wear resistant coating were characterized by means of
22
optical and scanning electron microscopy (OM and SEM/EDX) and electron-probe
microanalysis (EPMA). Hardness maps and profiles, as well as tensile tests on
miniaturized samples were performed to obtain mechanical properties. Concentration
profiles were calculated using the software DICTRATM showing a good agreement
with the experiments. Carbon diffuses against the concentration gradient due to a
higher activity into the substrate material leading to an increase of carbide volume
fractions close to the interface region. The mechanical tests show a brittle fracture
region with high hardness localized about 50µm away from the interface region in the
coating material.
3.2
Materials Processing and Experiments
3.2.1 Metal Matrix Substrate and Coating
A gas-atomized cold work tool steel powder X220CrVMo13-4 (1.2380) was
selected as the metal coating considering the wear resistance of the matrix material.
A hot work steel bar made of 55NiCrMoV7 (1.2714) with a diameter of 30mm was
chosen as the substrate core for the clad rods. The chemical compositions of the
steel metal matrix powders and the substrate core are shown in Table 1. Information
about the physical properties of the steels is given in Table 2. The steel matrix
powder grain sizes are below 200µm with most of the powder grains exhibiting sizes
between 40µm and 80µm, which is typical for gas atomized powders.
The steel 1.2380 is a ledeburitic cold work steel with a high wear resistance
and is often used for die cutting tools. Its microstructure in the quenched and
tempered conditions consists of tempered martensite as well as chromium-rich M7C3
and vanadium-rich MC carbides increasing its hardness. The nickel alloyed hot work
tool steel 1.2714 was chosen as the substrate core due to its good hardenability,
tempering resistance and dimensional stability as well as very good strength and
excellent toughness. It is also used for dies, tools for rod and tube extrusion, forming
dies and plastic moulds.
Table 1: Chemical composition of the substrate core (1.2714) and the coating steel powder (1.2380)
Material
Chemical composition [Wt.-%]
C
Cr
Mo
V
Mn
Si
Cu
Ni
Fe
1.2380
2.39
12.56
1.10
3.69
0.37
0.55
-
-
bal.
1.2714
0.56
1.15
0.46
0.08
0.75
0.29
0.11
1.74
bal.
23
Table 2: Physical properties of the substrate core (1.2714) and the coating steel powder (1.2380)
Material
α [10-6 K-1]
Hardness
Density
TA
[HRC]
[g/cm3]
[°C]
100°C
600°C
X220CrVMo13-4 (1.2380)
54 – 63
7,60
1050 – 1150
12,2
13,9
55NiCrMoV7 (1.2714)
52 – 58
7,84
850 – 1000
12,5
14,3
The clad rods were produced putting the hot work steel bar as the substrate
material into large capsules (Ø = 78mm, l = 200mm) made of a commercial austenitic
stainless steel (X5CrNi18-10, 1.4301) (Fig. 8). The surrounding space was filled with
the steel powder (1.2380) being pre-compressed to tap density by vibration. The
capsules were evacuated, sealed by TIG welding, and subsequently pre-heated at
1150°C for two hours. To reduce the friction between the rod and the die, the hot
capsules were rolled in glass powder, which acts as a lubricant that solidifies during
cooling down to room temperature and adheres as a solid layer on the rods. Finally,
the capsules were put into the preheated extrusion container (480°C) and extruded
with a ram speed of 36mm/s and a pressing ratio of 5.2:1 into rods with a diameter of
35mm (Fig. 9). Due to the hot extrusion the steel powder is consolidated and bonded
to the massive substrate material while the substrate itself is also deformed during
the process. In the end, an extruded bar with approximately Ø 35mm is formed
consisting of a tough core and a wear resistant layer of several millimeters in
thickness (Fig. 10).
Figure 8: Stainless steel capsules (Ø = 78 mm, l = 200 mm) with evacuation sockets filled with hot
work steel bars and 1.2380 tool steel powder for hot extrusion.
24
Figure 9: Sketch of the direct extrusion process.
b)
a)
interface
substrate
coating
Figure 10: a) extruded bars (~Ø = 35 mm) after hot extrusion and b) macro view of a cross section
showing, substrate, coating (~8mm thickness) and capsule.
The extrusion parameters of the specimens under investigation are
comparable to those of a previous work [4]. All extrusion trials were performed on the
8MN horizontal extrusion press at the Extrusion Research and Development Center
of the Technical University, Berlin. The extrusion press is equipped with load cells to
record the total force (Ft), die force (Fd) and friction force (Ff) during extrusion (Ft = Fd
+ Ff). The microstructure of the product can be influenced by variation of the
extrusion parameters temperature, extrusion speed and product shape, as well as by
the choice of lubricant and the extrusion method itself (direct, indirect, with mandrel
or porthole die) [7].
Microstructural examination was carried out by optical microscopy (OM) and
scanning electron microscope (SEM). For OM and SEM samples were cut parallel to
25
extrusion direction by electro discharge machining (EDM) to minimize the influence of
cutting on the microstructure. All specimens were ground and polished with great
care using diamond paste down to 1µm grade in order to avoid causing particle
damage in this stage. For OM and SEM the specimens were etched, when
necessary, with Nital 3%.
The production of high wear-resistant materials requires knowledge of the
hardening and tempering behavior to reach full secondary hardness. The extruded
bars were austenitized at 1070°C for thirty minutes, quenched in air to room
temperature and afterwards tempered at 520°C two times for two hours being cooled
in air between each step. This condition is called QT.
In order to compare the heat treatment effect on the diffusion mechanisms a
second procedure was carried out. An as-extruded sample was put in a furnace with
argon atmosphere for eight hours at 1150°C and cooled in air to room temperature,
totalizing 10 hours of heat treatment. After that, the hardening and tempering steps
were done in the same manner as described above for the specimens pre-heated for
two hours at 1150°C. This condition is denominated HT and the as-extruded state,
EX.
3.2.5 Electron Probe Micro Analysis (EPMA)
As the as-extruded rods consist of two different tool steels, diffusion at the
interface region driven by chemical composition and resulting activity gradients of the
alloying elements can be expected. The change in the concentration gradients for
each element was investigated by several line profiles and element mappings
performed on a JEOL model JXA-8100 instrument using a wave-length-dispersive
spectrometer for electron-probe micro analysis (WDS-EPMA) operated at an
acceleration voltage of 15kV and a probe current of 20nA. The electron beam was
set to perform line-scans of 200µm length being symmetric with respect to the
interface region, starting on the coating side and going towards the substrate
material, perpendicular to the extrusion direction. All specimens for EPMA were
mechanically ground and polished using diamond paste till 0.1µm.
26
For calculating diffusion profiles between the cold work tool steel clad to the
hot work tool steel substrate, the software package DICTRATM [8] was used.
DICTRATM stands for DIffusion Controlled TRAnsformation and it is based on a
numerical
solution
of
the
multi-component
diffusion
equations
and
local
thermodynamic equilibrium at the phase interfaces. The program is suitable for
treating e.g. moving boundary problems as well as growth, dissolution and
coarsening of particles in a matrix phase.
DICTRATM considers a system as divided into regions and/or cells. In this
study, the calculations were carried out isothermally at 1150°C using only one region
with a size of 16 mm, according to the macroscopic dimensions of the extruded bars
(Fig. 10). This region was symmetrically divided into two parts, coating and substrate,
by defining concentration profiles for each element using the heavy-side step function
hs(x). For setting appropriate starting conditions, the equilibrium state was calculated
for both steels at T=1150°C and p=101325 Pa with Thermo-CalcTM using the TCFE4
database and considering all alloying elements given in Table 1. The hot work tool
steel is fully austenitic at this condition, while the ledeburitic cold work steel 1.2380
exhibits an austenitic matrix in equilibrium with M7C3- and MC-carbides (Table 3).
Both types of carbides were included in the DICTRATM simulation using the included
model for dispersed phases and setting their volume fractions again using the
function hs(x). To account for the influence of the dispersed carbides on diffusion, a
labyrinth factor of f2, with f being the volume fraction of the matrix was introduced [9].
A grid consisting of 150 points and a higher point density towards the interface region
was defined while the simulation time was set to 36.000s (10h). For obtaining
mobility data the Thermo-CalcTM MOB2 database [10] was used.
To evaluate the bond strength between the substrate and the wear resistant
coating, tensile tests with miniaturized specimens (Fig. 11) were performed at room
temperature using a Zwick/Roell Z100 testing machine and a cross-head speed of
0.5mm/min. Tensile specimens with a length of 30mm, a gauge length of 16 mm and
a cross section of 1.5x2mm were machined by EDM perpendicular to the extrusion
direction. In order to remove the EDM surface layer and to reduce the surface
roughness, all specimens were polished with 6µm diamond paste.
27
To characterize the mechanical properties at the interface region, micro
hardness measurements were performed using the Fischerscope H100 equipment
and a load of 0.1N. An area of 1000x100µm in size, symmetric with respect to the
interface region, was defined and measured with a point distance of 10µm. The
Universal hardness values were calculated from the force-indentation curves
according to DIN 50359-1 and plotted two-dimensionally. Additionally, Vickers
hardness profiles with a load of 0,3kg (ASTM E 384-99) were also measured.
Figure 11: Sketch and dimensions of the miniaturized specimens used on the tensile tests.
3.3.
Results and Discussion
The macroscopic view of the as extruded bar (Fig. 10b) reveals a defect free
coating of about 8mm in thickness. This result is in agreement with further
investigations on the extrusion of wear resistant metal matrix composites [4, 5]. It
could be shown that a full densification of tool steel powders is possible by hot
extrusion. In the current investigation, the inserted substrate material does not
derogate the densification behaviour of the powdery layer. Besides, the diameter of
the massive hot work tool steel, serving as substrate material, is reduced from 30mm
to 16±0.3mm [11].
An overview of the microstructure at the interface region in the quenched and
tempered (QT) condition is depicted in Fig. 12a-12b. The substrate is fully martensitic
28
while the wear resistant layer of the steel 1.2380 is made up of a tempered
martensitic matrix with embedded globular iron-chromium- (M7C3 or M23C6) and
vanadium-carbides (MC). A characterisation of the phases present in these materials
was performed in earlier works using synchrotron radiation [5]. While phase- and
grain-boundaries can be clearly identified in both materials (Fig. 12c, 12d), the
interface region appears as a bright un-etched part of the microstructure with an
apparent width of 10-20µm. Even at high magnifications no pores could be identified
at the interface region, thus, high bond strength could be anticipated.
coating
substrate
a)
coating
substrate
b)
coating
substrate
c)
coating
substrate
d)
Figure 12: a, b) OM images of the interface region between 1.2714 and 1.2380 (K), sample quenched
and tempered. c, d) SEM images of the interface region between 1.2714 and 1.2380 (K) in the asextruded state.
3.3.2 Measured element distributions
The distributions of the most important alloying elements at the interface
region between the steel substrate 1.2714 and the wear resistant layer of the coating
1.2380 are depicted in Fig. 13a in the as-extruded state (EX), after 2h pre-heating at
29
1150°C. Strong signals from chromium and vanadium can be attributed to the
corresponding carbides. However, chromium is also dissolved in the vanadium-rich
MC-carbides and as well vanadium is dissolved in the iron-chromium carbides
confirming the results from the Thermo-CalcTM calculations (Table 3). The interface
region towards the cold work tool steel is determined by the presence of the
aforementioned carbides. Towards the substrate a layer enriched in silicon can be
determined while nickel diffuses from the substrate into the coating.
Table 3: Activities, equilibrium phases and corresponding compositions of 1.2380 and 1.2714 at
T=1150°C and p=101325 Pa calculated with Thermo-Calc using the TCFE4 database
Material
Chemical composition [Wt.-%]
C
Cr
Mo
V
Mn
Si
Cu
Ni
Fe
0.9
7.63
0.77
0.67
0.39
0.65
-
-
bal.
8.74
46.44
1.54
7.75
0.32
-
-
-
15.82
14.43
7.13
60.09
0.02
-
-
-
bal.
0.56
1.15
0.46
0.08
0.75
0.29
0.11
1.74
bal.
1.2380
FCC_A1#1
(Austenite)
M7C3
FCC_A1#2
(MC)
1.2714
FCC_A1#1
(Austenite)
Material
Activities [Dimensionless]
1.2380
1,86e-02
4,83e-04
5,94e-05
3,82e-06
2,82e-06
3,60e-08
-
-
1,62-03
1.2714
1,97e-02
7,81e-05
3,34e-05
4,80e-07
5,12e-06
1,67e-08
2,75e-05
2,14e-05
1,73e-03
The interface region average widths measured by EPMA line profiles (Fig. 14)
are listed on Table 4. In the element maps (Fig. 13) it is not obvious, but on the
EPMA line profiles (Fig. 14) the element distribution reveals a wider interface region
on the HT sample when compared with QT due to the higher diffusion activity as a
consequence of a longer heat treatment. It is worth to mention that the nickel
concentration increased up to approximately 0,5% within the coating, close to the
interface region.
30
a) As extruded (EX)
extrusion
direction
10µm
SEM
Cr
V
Si
Ni
b) Heat-treated for 2h at 1150°C (QT)
extrusion
direction
10µm
SEM
Cr
V
Si
Ni
c) Heat-treated for 8h at 1150°C (HT)
extrusion
direction
10µm
SEM
Cr
V
Si
Ni
Figure 13: Element distribution maps in the coating (left hand side), interface region (middle) and
substrate (right hand side) – a) EX, b) QT and c) HT.
In the quenched and tempered sample (Fig. 13b) a band like structure of
carbides is formed. A band of iron-chromium carbides is formed at the interface
region, separated by a band of vanadium-rich carbides. At this position, the volume
fraction of vanadium-rich carbides is increased while any iron-chromium carbides can
be found.
After an additional tempering of 8h at 1150°C followed by a QT heat treatment,
the band like structure of carbides cannot be found anymore (Fig. 13c). However, a
zone enriched in vanadium-rich carbides being free from iron-chromium carbides can
be detected at the interface region. Comparing the EPMA results for the chromium
and vanadium signals, a shift of the iron-chromium carbides towards the coating and
31
an enrichment of vanadium-rich carbides at the interface region can be detected. The
silicon signal in the analysed region shows an even distribution regarding the matrix
concentration.
a) 10
Cr QT
Cr HT
b)
Wt.%
Wt.%
V QT
V HT
2,5
8
6
3,0
2,0
1,5
4
1,0
2
0,5
0,0
0
80
100
120
140
160
80
180
100
c)
120
140
160
180
Distance [µm]
Distance [µm]
0,8
Si QT
Si HT
0,7
d) 2,0
Ni QT
Ni HT
Wt.%
Wt.%
1,5
0,6
0,5
0,4
1,0
0,5
0,3
0,2
80
0,0
100
120
140
160
Distance [µm]
180
60
80
100
120
140
160
180
Distance [µm]
Figure 14: EPMA line-scans showing a) chromium, b) vanadium, c) silicon and d) nickel profiles
between the coating (left hand side), interface region (middle) and substrate (right hand side), samples
quenched and tempered (QT) and heat treated (HT) at 1150°C. Between the dotted lines is the wider
interface region of sample HT.
3.3.3 Calculated element profiles
In Fig. 15 the calculated diffusion profiles for chromium, vanadium, silicon and
nickel are depicted. The diffusion range for each element was determined by defining
a deviation of 5% from the initial value. Comparing the values with those measured
by EPMA a good agreement can be noticed, except for silicon and molybdenum
(Table 4). Diffusion in the coating material takes place slower than in the substrate
due to the effect of the dispersed carbides. While chromium (Fig. 15a) and nickel
(Fig. 15d) exhibit a regular interdiffusion profile, vanadium and, in particular, silicon
are enriched at the interface region between the steels (Fig. 15b, 15c).
32
Table 4: Interface region average width per element according to EPMA and DICTRATM
As extruded (EX)
QT [7.200s]
8h at 1150°C + QT (HT) [36.000s]
EPMA
DICTRATM
~ 39 µm
~ 55 µm
~ 85 µm
~ 28 µm
~ 31 µm
~ 60 µm
~ 65 µm
~ 19 µm
~ 25 µm
~ 76 µm
~ 56 µm
~ 168 µm
Ni
~ 18 µm
~ 23 µm
~ 26 µm
~ 44 µm
~ 57 µm
Mo
~ 12 µm
~ 22 µm
~ 53 µm
~ 54 µm
~ 107 µm
EPMA
EPMA
Cr
~ 21 µm
~ 24 µm
V
~ 22 µm
Si
TM
DICTRA
Diffusion profiles for the carbon distribution in the FCC matrix were calculated
because we reckon that, due to the large difference in the carbon content between
the two steels, the EPMA measurements of the carbon content were not precise
enough to be shown. In Fig. 15e the carbon distribution over the full range of 16mm
is depicted. Due to the difference in carbon activity (Table 3), the substrate material
is decarburized over a range of about 5mm while the coating is enriched in carbon. A
steep increase in carbon content to a value of 1.3 wt.% was calculated for the
interface region. At a smaller length scale of 500µm (Fig. 15f), the profile in the
carbon content of the FCC matrix can be seen more clearly. Carbon is significantly
enriched at the interface region in a range of 10µm after 7.200s and 25µm after
36.000s. This leads to the assumption of a high content of retained austenite in this
region even after double tempering. Investigations with electron back scatter
diffraction (EBSD) are under examination to investigate the crystal structure of the
matrix at the interface region.
The DICTRATM model for diffusion in dispersed systems considers only
diffusion in the matrix, but not in the dispersed phases [13, 14]. However, the
dispersed phases, in this case the carbides in the coating, are taken into account for
the equilibrium calculations. Therefore, the evolution of phase fractions can be
analysed with respect to the heat treatment time. In Fig. 16 the profiles of the carbide
volume fractions over the full range of 16mm are depicted. The results exhibit an
increase in the volume fraction of the M7C3 and MC carbides at the interface region.
At a smaller length scale, the profiles can be seen more clearly (Fig. 16c). The
overlapping of both profiles shows that the maximum in M7C3 content is shifted to the
left, towards the coating, while the MC content increases close to the interface
region. These results are in agreement with the findings from the EPMA
measurement for the specimen heat treated additionally for 8h (Fig. 13c).
33
a
b
c
d
e
f
Figure 15: Concentration profiles after 2h (7.200s) and 10h (36.000s) calculated with DICTRATM:
a) chromium, b) vanadium, c) silicon, d) nickel and e, f) FCC matrix. The interface region is at the
distance of 8mm.
34
a
b
c
d
Figure 16: Calculated profiles of a) M7C3, and b) MC. Volume in the initial state and c) after 2h
(7.200s) and d) 10h (36.000s) at 1150°C.
The results of the tensile tests (Fig. 17a) disclose continuous yielding and
average yield strength of 800±110MPa, being in agreement with the yield strength of
the substrate material after the aforementioned heat treatment. The average ultimate
tensile strength (UTS) is 1300±150MPa. As expected, this materials combination
shows brittle fracture with a plastic strain of about 2,6%. Fracture occurs in the
coating material close to the interface region, as can be seen in OM pictures taken
from the fractured specimens (Fig. 17b, 17c) and on the SEM image of the fractured
surface (Fig. 17d). Rests of the coating can be detected on the substrate, indicating a
35
fracture in the coating and not along the interface region. This result argues for the
good bonding between the co-extruded materials. Assuming a high amount of
retained austenite at the interface region and taking the shift of the M7C3 carbides
into account, a fracture occurring in the coating is self-evident.
Elastic deformation occurs in both materials. As soon as the yield strength of
the substrate material is reached, plastic deformation takes place localized in the
coating 1.2380. Due to work hardening, the local yield strength of the coating
material close to the interface region is reached. Plastic deformation of this material
at room temperature is not possible due to the high volume fraction of carbides acting
as defects for crack initiation. Thus, fracture occurs in the region with the highest
volume fraction of carbides and higher hardness, located about 50µm away from the
initial interface region.
a)
b)
1400
Stress [MPa]
1200
1000
800
600
400
c)
200
0
0
1
2
3
4
5
6
7
8
9
Strain [%]
c)
d)
coating
substrate
Figure 17: a) Tensile test curves, b, c) LOM pictures showing where the fracture occurs and d) SEM
image of the fractured surface on the coating (d). The tests were performed on the sample QT.
36
3.3.4.2
Micro hardness
The 2D-microhardness map is shown on Fig.18. The effect of the heat
treatment is clearly observed on the higher hardness of the specimen QT due to less
decarburization, especially on the interface region. Coloured in yellow and orange we
can observe, respectively, the chromium and vanadium-rich carbides dispersed on
the coating and a high hardness concentration region shifted to the left, right beside
the interface region. In accordance with the results from the tensile tests, the fracture
area has approximately 80µm in width and is localized about 50µm away from the
interface region. In order to compare, Fig. 19 depicts a hardness profile of the
materials combination on the as-extruded state (EX) and the specimen quenched
and tempered (QT). A decrease of hardness near the interface region on the
substrate and the increase on the coating is also observed. These results are also in
good agreement with the investigations using DICTRATM and EPMA.
HUplast 0.1/20 [N/mm2]
Figure 18: Micro hardness map showing the coating (left hand side), interface region (middle, between
dashed lines) and substrate (right hand side), sample quenched and tempered (QT).
37
substrate
coating
Figure 19: Micro hardness profile showing the decrease of hardness near the interface region on the
substrate and the increase on the coating side. The interface is in the middle at value zero.
38
Correlation between Interface Microstructures
and Mechanical Properties of Co-Extruded
Layered Structures [20]
4.1 Introduction
The aim of this chapter is a characterization of the microstructures at the
interface between the steel substrate cores and the wear resistant coatings using
optical and scanning electron microscopy with energy dispersive X-ray analyses and
electron backscatter diffraction tools. The results of the investigations reveal that
carbon diffusion against the concentration gradient influences the microstructure and
mechanical properties, such as hardness and fracture toughness, at the interface
region of three different combinations of wear resistant coatings co-extruded with the
same steel substrate.
A gas-atomized cold work tool steel powder X220CrVMo13-4 (1.2380) and a
gas-atomized hot work steel powder X40CrMoV5-1 (1.2344) were selected as metal
matrices for the coatings because of their good hardenability and high wear
resistance. A hot work steel bar 55NiCrMoV7 (1.2714) with a diameter of 30mm was
39
chosen as the substrate core also due to the good hardenability, strength and
toughness. The chemical compositions of the metal matrix of the coatings and of the
substrate are shown in Table 5.
Table 5: Chemical composition of the coating steel powders and substrate core
Designation
1.2380 (K)
1.2344 (W)
1.2714
Coating
Substrate
Composition [wt.-%]
C
Cr
Mo
V
Mn
Si
Cu
Ni
2,39
12,56
1,10
3,69
0,37
0,55
-
-
0,40
5,04
1,34
0,97
0,30
0,19
-
-
0,56
1,15
0,46
0,08
0,75
0,29
0,11
1,74
Fe
Bal.
The three materials combinations investigated are:
-
K+1.2714 where the cold work tool steel powder 1.2380, here denominated K
(kalt = cold), is hot extruded on the steel 1.2714 as substrate.
-
WW1+1.2714 stands for the hot work steel 1.2344, W as hot (warm), as
coating on the same substrate. W1 means the addition of 10 vol.% of WC/W2C
(fused tungsten carbides, FTC) particles into the coating.
-
KW1+1.2714 extruded with 10 vol. % of WC/W2C (FTC) is also analysed.
The extrusion process and the process parameters of the investigated
specimens were described previously [19]. Microstructure examination was carried
out using optical microscopy (OM), scanning electron microscopy (SEM) and electron
backscatter diffraction (EBSD). The samples were cut parallel to extrusion direction
by electro discharge machining (EDM) in order to minimize the influence of cutting on
the microstructure. All specimens were ground and polished down to 1µm grade. For
OM and SEM the specimens were etched with Nital 4%. For EBSD a final polishing
step using colloidal silicon oxide was necessary.
The extruded bar WW1+1.2714 was austenitized at 1050°C for thirty minutes,
quenched in oil to room temperature and afterwards tempered at 570°C two times for
two hours being cooled in air between each step. Specimens K+1.2714 and
KW1+1.2714 were austenitized at 1070°C for thirty minutes, quenched in air to room
temperature and afterwards tempered at 520°C three times for two hours being
cooled in air between each step. In order to investigate the effect of the heat
treatment (HT) on the diffusion mechanisms, a second austenitizing treatment of
eight hours at 1150°C was carried out on the K+1.2714 specimen in the as-extruded
state. Considering the first two hours of pre-heating before the extrusion process and
40
not considering the annealing time, the heat treatment in this case has 10 hours in
total [19]. After the prolonged high temperature heat treatment of K+1.2714, it was
hardened and tempered in the same way as described above for the same
specimen.
In order to evaluate the bond strength between the substrate and the wear
resistant coating, tensile tests with miniaturized specimens were performed at room
temperature using a Zwick/Roell Z100 testing machine and a cross-head speed of
0.5mm/min. For characterizing the mechanical properties at the interface, micro
hardness measurements were performed using the Fischerscope H100 equipment
with a load of 0.1N for 20s per point. An area of 260x130µm in size, symmetric with
respect to the interface region between substrate and coating of the extrudates, was
defined and micro hardness was measured with a point distance of 10µm. Due to the
small load used in the measurement, any influence of work hardening was detected
with respect to the small point distance used for the indentations. Additionally,
Vickers hardness profiles with a load of 0,3kg (ASTM E 384-99) were determined.
4.3
An overview of an extruded bar is depicted in Fig.20a. The coating thickness is
approximately 8mm for all extrudates. Full densification of the steel powders of the
coating both with and without hard particles occurred [3, 4]. Microscopy showed that
the FTC hard particles do not seem to harm the densification and the bonding
between 1.2380 coating steel and the 1.2714 substrate, even if a WC/W2C particle is
located exactly at the interface region. Further, the interface region between coating
and substrate is free of defects, which is a necessity for good bonding between the
MMC cladding and the steel substrate.
The steel matrix of the coating WW1 is almost carbide free (Fig. 20b). Both the
matrices of the cold work steel K (Fig. 20c) and KW1 with FTC (Fig. 20d) consist of
tempered martensite with embedded globular chromium carbides (M7C3) and
vanadium carbides (MC). The interface region appears as a band between coating
and substrate with an average width of 15-20µm. Phase boundaries can be
recognized and, even at higher magnifications, a defect free interface region could be
found.
41
a)
b)
interface
WW1
1.2714
substrate
coating
20 µm
c)
K
1.2714
d)
KW1
1.2714
20 µm
20 µm
Figure 20: a) Macro view of the cross section showing substrate, coating (~8mm thickness) and
external capsule. SEM images of the interface region between the substrate steel 1.2714 on the left
hand side and the coating steel powders on the right hand side: b) WW1, c) K and d) KW1. The
extrusion direction is parallel to the interface region.
In the EBSD image of the combination KW1+1.2714 (Fig. 21) Cr7C3 carbides
with an average size of 2-2.5µm can be detected in the metal matrix of the coating.
The shape of these Cr7C3 carbides varies between ellipsoid and globular. Ellipsoidal
Cr7C3 appear aligned with respect to the extrusion direction. Beside the Cr7C3
carbides a small amount of VC with an average size of 0.8-1.2µm is dispersed within
the martensitic matrix of the coating.
42
KW1
1.2714
5µm
Figure 21: EBSD scan of KW1+1.2714 showing the α-martensite substrate matrix (yellow) and the
KW1 coating matrix. The coating contains Cr7C3 carbide particles (green) and vanadium carbides as
the small and dispersed particles colored in red. The extrusion direction is parallel to the interface
region.
Yield strength and ultimate tensile strength (UTS) values (Table 6) for the
three extrudates in the hardened and tempered condition (hot extruded after 2h at
1150°C + annealing afterwards) were determined as average values of at least five
specimens each. A typical stress versus strain curve of each substrate/coating
combination and a sketch of the micro tensile test specimen are shown in Figure 22.
Analyses of the tensile curves (Fig. 22a) reveal that plastic deformation occurs
in the substrates of the specimens K+1.2714 and KW1+1.2714. In case of the hot
work tool steel coating on WW1+1.2714 the specimens do not show plastic
deformation.
Table 6: Results of the tensile tests (values are an average of five micro tensile tests)
Ultimate Tensile Strength (UTS)
Combination
Yield Strength [MPa]
WW1+1.2714
---
715 ± 100
---
K+1.2714
810 ± 140
1300 ± 150
3.5
KW1+1.2714
245 ± 5
415 ± 30
~8
[MPa]
Plastic Strain [%]
43
Stress [MPa]
a)
1400
b)
CW1+1.2714
C+1.2714
HW+1.2714
1200
1000
800
600
400
200
0
0
1
2
3
4
5
6
7
8
9
10
Strain [%]
Figure 22: a) Micro tensile test curves and, b) sketch of the miniaturized specimens used for the
tensile tests.
Fracture occurs on the coating side of the specimens close to the interface
region in all specimens (Fig. 23). Parts of the coating material remain on the
substrate indicating that fracture occurs within the coating in a distance of about
50µm and not right at the interface. This confirms the good bonding between coating
and substrate that was indicated by microscopy images of the interface region. The
low yield strength and UTS values for KW1+1.2714 cannot be explained yet.
K
a)
b)
KW1
Figure 23: SEM image of a fractured surface: a) K+1.2714 and b) KW1+1.2714. The coating is located
on the right hand side.
Micro hardness profiles (Fig. 24a) and micro hardness maps (Fig. 24b, 24c)
show a decrease of the substrate hardness in the region near the interface. In
contrast to the hardness decrease of the substrate a pronounced increase in
hardness appears close to the interface region in the coating in all three different coextruded materials combinations. This peak in hardness of the coatings is located
44
between 25µm and 50µm from the interface region. This distance to the interface
region corresponds to the area where fracture occurred in the tensile tests.
The micro hardness profile and the micro hardness maps further reveal a
softening of the substrate in case of K+1.2714 and WW1+1.2714, but not for
KW1+1.2714. A carburization of the coating and a decarburization of the substrate
are expected due to a difference in the carbon activity for the combination K+1.2714
[19].
950
Hardness [HV0,05]
a) 900
850
WW1
b)
CW1+1.2714
C+1.2714
HW1+1.2714
c)
K
d)
KW1
800
750
700
650
coating
600
550
substrate
2440
3014
500
-500 -400 -300 -200 -100
0
100 200 300 400 500
Distance from interface [µm]
3588
4162
HUplast 0.1/20
[N/mm2]
4736
5310
5884
6458
7032
7606
8180
Figure 24: a) Micro hardness profile of the three combinations. Micro hardness maps: b)
WW1+1.2714, c) K+1.2714 and d) KW1+1.2714. The coating is located on the right hand side.
45
Microstructure Characterization by EBSD/EDX
Focusing on the Influence of Hard Particles
Addition and the Formation of Retained Austenite
5.1
Introduction
Hot direct extrusion has been established as a successful method to clad low
alloy steel substrates with tool steels or metal matrix composites (MMC). Here low
alloyed steel bars were co-extruded with pre-sintered tool steel powders with or
without the addition of tungsten carbides (W2C/WC) as hard particles. During the hot
extrusion process, an extrudate is formed consisting of a wear resistant coating layer
and a bulk steel bar as the substrate core. The microstructure at the interface region
between coating and substrate of four hot extruded rods with different coatings was
characterized using optical and scanning electron microscopy (OM and SEM) in
combination with electron backscatter diffraction (EBSD) and energy dispersive X-ray
analysis (EDX).
Electron backscatter diffraction (EBSD) in the SEM here appears as
particularly well suited method to study the influence of hard particle addition on the
formation and distribution of chromium- and vanadium-rich carbides. Calculations
performed using the program DICTRA [8] revealed a strong carbon enrichment of the
46
cladding at the interface to the low alloyed steel substrate. Thus, EBSD here was
also employed to check for the presence of retained austenite.
The investigations revealed the influence of hard particle (HP) addition on the
formation of M7C3 and MC carbides in the coating. They further showed that the
interface region is free of retained austenite which was expected to be present due to
an enrichment of carbon at the interface between substrate and coating.
5.2.
5.2.1 Metal Matrix Substrates and Powder Steel Coatings
A gas-atomized cold work tool steel powder X220CrVMo13-4 (1.2380) and the
gas-atomized hot work tool steel powder X40CrMoV5-1 (1.2344) were used as
coating materials for the PM-MMCs. A hot work steel bar made of 55NiCrMoV7
(1.2714) and a non-alloy structural steel S355 with a diameter of 30mm were chosen
as the substrate materials for the clad rods. The chemical composition of the tool
steel powders and the substrates is given in Table 7.
An earlier work [19] gives more details about the steel powder 1.2380 as well
as the substrate steel 1.2714. The steel powder X40CrMoV5-1 is based on the hot
work tool steel 1.2344 exhibiting a high wear and thermal shock resistance in a
temperature range of 400-700° C, as well as a high level of toughness and ductility. It
can be heat treated to a typical hardness of 50-56 HRC and contains virtually no
carbides in its martensitic microstructure, depending on the heat treatment applied.
The substrate made of the non-alloy structural steel S355 was selected due to its low
cost, good weldability and cold formability as well as a high fatigue limit.
Table 7: Chemical composition of the coating steel powders and substrate cores
Designation
1.2380 (K)
1.2344 (W)
S355
1.2714
coating
substrate
Concentration (wt.-%)
C
Cr
Mo
V
Mn
Si
Cu
Ni
2,39
12,56
1,10
3,69
0,37
0,55
-
-
0,40
5,04
1,34
0,97
0,30
0,19
-
-
0,20
0,15
0,03
-
1,38
0,22
0,30
0,12
0,56
1,15
0,46
0,08
0,75
0,29
0,11
1,74
Fe
Bal.
47
5.2.2 Hard Phases and Sample Designation
Coarse hard particles (HP) of monolithic WC/W2C (fused tungsten carbide,
FTC) with a density of 16.5 g/cm3, a micro hardness of 2600 HV0.05 and size of 100150µm were added to the matrix powders to improve the wear resistance in abrasive
environments. FTC offers a good ratio of hardness to fracture toughness [7].
Powders containing 10 vol.% of hard particles (HP) were blended in a shaker-mixer
for 1h. The steel 1.2380 with 10 vol.% of WC/W2C (FTC) has a theoretic density of
8,49 g/cm3 and the hot work steel 1.2344 with 10 vol.% of FTC one of 8,65 g/cm3.
The materials combination of the hot work steel powder 1.2344 as a coating
and the steel bar made of 1.2714 as a substrate is designated as WW1+1.2714. “W”
stands for hot work steel and “W1” is the reference for 10 vol.% of FTC. The cold
work tool steel powder 1.2380 coated on the 1.2714 steel is called KW1+1.2714 in
this publication. Letter K is the cold work steel. Finally, the powder 1.2380 as a
coating in combination with the structural steel S355 results in KW1+S355. The
configuration of each specimen is given in Table 8.
Table 8: Designation of each MMC specimen configuration
Coating
Substrate
W (1.2344) + 10% W2C/WC = WW1
K (1.2380)
=K
WW1 + 1.2714
1.2714
K (1.2380) + 10% W2C/WC = KW1
K (1.2380) + 10% W2C/WC = KW1
Materials Combination
K + 1.2714
KW1 + 1.2714
S355
KW1 + S355
The rods cladded with MMCs were produced putting a steel bar as the
substrate material (1.2714 or S355) into capsules (Ø = 78mm, l = 200mm) made of a
commercial austenitic stainless steel (X5CrNi18-10, 1.4301). The surrounding space
was filled with a mixture of steel powder (1.2380 or 1.2344) and 10% of FTC particles
being pre-compressed to tap density by vibration. The capsules were evacuated,
sealed by TIG welding, and subsequently preheated at 1150°C for 2h. To reduce the
friction between the rod and the die, the hot capsules were rolled in glass powder,
which acts as a lubricant that solidifies during cooling down to room temperature and
adheres as a solid layer on the rods. Finally, the capsules were put into the
preheated extrusion container (480°C) and extruded with a ram speed of 36mm/s
and a pressing ratio of 5.2:1 into rods with a diameter of 35mm. These parameters
48
were based from sintering and dilatometric deformation tests as well as earlier
investigations of Al alloys with high Si content [21]. Due to the hot extrusion, the steel
powder is consolidated and bonded to the massive substrate material while the
substrate itself is also deformed during the process. Finally, an extruded bar with
approximately Ø 35mm is formed consisting of a tough core and a wear resistant
layer of several millimeters in thickness.
The extruded bar WW1+1.2714 was austenitized at 1050°C for thirty minutes,
quenched in oil to room temperature and afterwards tempered at 570°C two times for
two hours being cooled in air between after each step. Specimens K+1.2714,
KW1+1.2714 and KW1+S355 were austenitized at 1070°C for thirty minutes,
quenched in air to room temperature and tempered at 520°C three times for two
hours also being cooled in air after each step. These heat treatments were performed
in order to reach full secondary hardness aiming at the production of high wearresistant materials.
The microstructural examination was carried out by scanning electron
microscope (SEM). SEM samples were cut parallel to extrusion direction by electro
discharge machining (EDM) to minimize the influence of cutting on the
microstructure. A macroscopic view of the cross section is depicted in Figure 25a. All
specimens were ground and polished using diamond paste down to 1µm grade in
order to avoid causing particle damage in this stage. The SEM specimens were
etched, when necessary, with Nital 3%.
EBSD specimens were also ground and polished using diamond paste down
to 0.25µm in a rotating polishing machine using high pressure. This step was used in
order to avoid reliefs in the interface region caused by a different behavior of coating
and substrate material during polishing. Several different procedures were tried till
the best receipt was encountered. A final polishing step using colloidal silicon oxide
(SiO2) was necessary aiming to obtain a surface as flat as possible and high quality
diffraction patterns.
Microstructure characterisations were carried out by electron back scattering
diffraction (EBSD) using a Zeiss Neon 40 field emission scanning electron
49
microscope equipped with the Hikari EDAX/TSL EBSD system and a energy
dispersive X-ray analysis (EDX).
5.3
All interface region between coatings and substrates are free of macroscopic
defects. Figure 25 shows the interface between the coating of 1.2380 (approximately
8mm in thickness) and the 1.2714 substrate bar (Ø = 16mm). A full densification of
the powder mixtures both without and with 10 vol.% FTC after hot extrusion is
observed for all specimens (Fig.26a-d). This reveals that the insertion of a massive
substrate bar does not deteriorate the densification of the PM-coatings.
The long axes of the FTC hard particles present in the coating of three of the
specimens (Fig. 26b-d) are oriented in extrusion direction. Around each of the FTC
particles the formation of a diffusion seam of η-carbides (M6C), which has already
been identified in HIPped microstructures [22], is clearly recognizable. FTC particles
are even found directly at the interface to the substrate. These particles bond to the
low alloy steel substrate without any pore formation.
The etched microstructures show slight differences between the two
substrates, 1.2714 and S355, close to the interface. The main difference between the
substrate steel S355 and 1.2714 close to the interface is the size of the α-martensite
laths (Fig. 26d). Typically, the microstructure of S355 is ferritic/perlitic due to the low
amount of alloying elements. Here, however, martensite appears in the substrate at
the interface to the coating. Generally, the formation of martensite in S355 is possible
only, if a very fast cooling is applied after the austenitization. The presence of αmartensite laths close to the interface in this material is likely to be the result of
alloying elements diffusing from the coating KW1 to the S355 substrate.
Corresponding diffusion calculations with DICTRATM [8] result in an element level
within the substrate of 0.39 wt% of C, 0.20 wt% of Cr, 0.10 wt% of V and 1,37 wt% of
Mn at a distance of 20µm from the interface. In particular, the increased carbon level
is responsible for the formation of martensite.
50
b)
a)
8mm
16mm
EBSD
scans
interface
substrate
substrate
coating
coating
capsule
Figure 25: a) Macroscopic view of the cross section showing the external capsule, substrate, coating
(~8mm thickness) and the interface region, b) sketch indicating the location of the performed EBSD
scans. The interface region is parallel to the extrusion direction.
a)
1.2714
K
KW1
b)
1.2714
WC/W2C
51
WW1
c)
1.2714
d)
KW1
S355
WC/W2C
WC/W2C
Figure 26: Optical micrographs of the interface region between coating (right hand side) and substrate
(left hand side): a) K + 1.2714, b) KW1 + 1.2714, c) WW1 + 1.2714, and d) KW1 + S355. The interface
region is parallel to the extrusion direction.
Figure 27a-d shows the microstructures and the interface regions (in-between
the white lines) of the four hot extruded clad steel bars. These four materials were
selected to account for an addition of hard particles (KW1+1.2714 and K+1.2714), a
change of the tool steel powder of the coating (KW1+1.2714 and WW1+1.2714) and
a change of the substrate material (KW1+1.2714 and KW1+S355).
In specimen K+1.2714 (Fig. 27a), containing the same substrate but with a
different coating tool steel, which was already described in detail in an earlier work
[19], the interface width is 10-15µm. The microstructure of the substrate material is
martensitic while in the coating three different phases are present: tempered
martensite, MC and M7C3 carbides.
In Fig. 27b the wear resistant coating of the steel 1.2380 consists of a
tempered martensitic matrix with embedded iron-chromium- (M7C3 or M23C6) and
vanadium-rich carbides. Phase-boundaries can be recognized and no pores are
observed anticipating high bond strength. The interface region is the unetched band
between coating and substrate with an average width of 11-15µm. A hard particle of
FTC is situated directly at the interface between substrate and coating. It exhibits a
diffusion seam of M6C as well as a full bonding to the substrate material. The final
polishing preparation step using SiO2 had left small white particles incrusted in the
coating microstructure. After different polishing methods, these features were not
visible anymore.
In the specimen WW1 + 1.2714 (Fig. 27c), the width of the interface region is
10-15µm. The α-martensitic matrix in the 1.2714 substrate can be clearly recognized.
52
Around the FTC particle with a core consisting of W2C/WC, a diffusion seam of M6C
(η-carbide) is visible.
In Fig. 27d the interface microstructure of KW1 cladded on S355 is depicted.
An average interface width of approximately 19-22µm was determined. This
materials combination exhibits the largest interface region compared to the others.
K
a)
1.2714
b)
KW1
M 6C
1.2714
10µm
c)
1.2714
M 6C
WW1 d)
S355
10µm
KW1
M 6C
WC/W2C
10µm
10µm
Figure 27: SEM micrographs in a higher magnification showing the interface region between coating
(right hand side) and substrate (left hand side): a) K + 1.2714, b) KW1 + 1.2714, c) WW1 + 1.2714
and d) KW1 + S355. The interface region is parallel to the extrusion direction.
53
a)
1.2714
α-Fe
VC
Cr7C3
b)
1.2714
K
5µm
KW1
5µm
5µm
c)
1.2714
WW1
5µm
d)
S355
KW1
5µm
Figure 28: EBSD phase maps: a) K+1.2714, b) KW1+1.2714, c) WW1+1.2714 and d) K+S355. The
substrate is in the left hand side and the coating in the right hand side. The extrusion direction is
parallel to the interface region.
The position of the selected areas from which the EBSD scans were taken is
depicted Figure 25b. The respective EBSD phase maps corresponding to the
interface regions of each analysed sample are presented in Figure 28. In Figure 28a,
the EBSD phase map of the materials combination K+1.2714 is depicted. The cold
work steel coating forms chromium- and vanadium-rich carbides with an average size
of 0.3–2µm and 1–3.5µm with a volume fraction of 2.6% and 26.9%, respectively.
The major axis of larger chromium-rich M7C3 carbides, exhibiting an elongated and
54
ellipsoid shape, is aligned with respect to the extrusion direction. This tendency to
align in extrusion direction is also observed for the particles of FTC [5]. The extrusion
direction is the one with minimum resistance to material flow. Due to a high
deformation degree associated with a radial gradient in the extrusion velocity, the
hard particles in a larger extent, and the chromium-rich carbides in a smaller one, are
forced to orient in the extrusion direction during the hot extrusion process.
A higher concentration of vanadium-rich MC carbides is located close to the
interface region. This finding was reported earlier [19] after an EPMA analysis of the
same specimen and is related with interdiffusion processes between substrate and
coating during the processing at high temperatures. In the steel substrate 1.2714, the
α-martensitic laths are perpendicularly oriented with respect to the extrusion
direction. The α-martensite laths become smaller, randomly oriented and deformed
close to the interface region. This behaviour is observed for all substrates of the four
different materials combinations.
The main difference between specimen K+1.2714 and KW1+1.2714 (Fig. 28b)
is the addition of 10% of FTC into the cold work tool steel coating. Comparing both,
the presence of W2C/WC particles is responsible for an increase in the M7C3 volume
fraction (~31.3%) and a decrease in the content of MC particles (~ 1.5%). The size
range of these carbides is reduced to 2.0–2.5µm for the chromium-rich and 0.8 –
1.2µm for the vanadium-rich carbides.
The higher concentration of VC close to the interface region observed in the
coating K (Fig. 28a) does not occur in the same coating with hard particles addition,
KW1 (Fig 28b). The presence of FTC particles influences the MC distribution and the
amount of these carbides in the α-martensitic matrix of the coating steel 1.2380.
What is similar in the microstructure of these two cold work tool steel coatings is the
shape and orientation of the chromium-rich carbides – globular and ellipsoid, aligned
with respect to the extrusion direction.
The steel matrix of the coating WW1 (Fig. 28c) contains some vanadium-rich
carbides with FCC structure due to the heat treatment and the reaction of the matrix
steel 1.2344 with the particles of FTC. The MC carbides exhibit a globular shape and
are randomly dispersed at a small volume fraction of 0.7% and an average size of
0.5 – 1.0µm. The substrate steel 1.2714 in this materials combination shows the
same behaviour when compared to KW1+1.2714, but contains smaller α-martensite
laths.
55
The same coating, KW1, was hot extruded with a different substrate, the nonalloy structural steel S355 (Fig. 28d). In this combination, the globular shaped M7C3
carbides have a size of 0.8–2.5µm and the smallest volume fraction (~ 17.8%)
compared to the other coatings with KW1. The chromium-rich carbides are not
aligned in the extrusion direction and the ellipsoidal shape rarely occurs. A higher
concentration of MC close to the interface occurs in a similar way as observed in the
materials combination K+1.2714. The average size range of these carbides is the
same as for combination KW1+1.2714, but the volume fraction is higher (~ 2.4%). A
few microns away from the interface region, the steel substrate S355 possesses
much bigger laths of α-martensitic in comparison to the substrate 1.2714.
5.3.2 Retained austenite (RA or γ-Fe) and Vanadium carbides (VC)
Earlier investigations [19], supported by diffusion calculations, lead us to the
assumption of a high content of retained austenite (RA) in the interface region for the
materials combination K+1.2714 due to and enrichment of carbon at the interface. As
γ-Fe and MC carbides have the same FCC crystal structure and space group
(Fm3m), an EBSD measurement combined with an EDX analysis was carried out for
the specimen KW1+1.2714 (Fig. 29) in order to differentiate these phases and verify
the hypothesis. The quality of this EBSD measurement is not as high as the already
presented due to a larger step size, but has enough information to allow a precise
identification of either γ-Fe (RA) or vanadium-rich FCC carbides.
The phase map (Fig. 29a) shows particles with FCC structure in red and the
chromium-rich carbides in green with a particle distribution and size already
described. The circles are indicating the position of several FCC particles. The same
positions are then indicated in Fig. 29b for the vanadium map, in Fig. 29c for the
distribution of chromium and finally, in Fig. 29d, for the distribution of iron. Strong
vanadium signals are measured for all places indicated as FCC particles in Fig. 29a.
The chromium signals correspond to the Cr7C3 carbides indicated in the same figure.
A depletion of the iron signal corresponding to the positions of the FCC particles is
shown in Fig. 29d. Another important issue is that vanadium is dissolved in the ironchromium carbides as well as chromium is dissolved in the vanadium-rich MC
carbides. This finding was reported in a previous contribution [19], referring to the
materials combination K+1.2714.
56
Thus, the amount of retained austenite in this materials combination can be
considered negligible according to the EBSD analysis.
a)
1.2714
α-Fe
VC
Cr7C3
KW1
3µm
b)
V
3µm
c)
Cr
3µm
d)
Fe
3µm
Figure 29: EBSD phase and EDX maps of specimen KW1+1.2714: a) phase map, b) vanadium map,
c) chromium map and d) iron map. Circles are indicating the position of several FCC particles. The
substrate is in the left hand side and the coating in the right hand side, with the extrusion direction
being parallel to the interface region.
57
Influence of Hard Particles Addition and
Chemical Interdiffusion Investigated by Diffusion
Calculations on the Mechanical Properties
6.1
Introduction
It was recently found out that hot direct extrusion is a feasible and cost
efficient process for the production of PM-MMCs with tool steel matrix. In the next
step of our investigations, low alloyed steel bars were co-extruded with pre-sintered
tool steel powders with the addition of tungsten carbides (W2C/WC) as hard particles.
During the hot extrusion process of these massive and powdery materials, an
extrudate is formed consisting of a completely densified wear resistant coating layer
and a bulk steel bar as the tough substrate core.
This work combines experimental measurements (EPMA) and diffusion
calculations (DICTRATM) to investigate the effect of hard particle addition and its
dissolution, as well as the formation of M6C carbides on the microstructure and
properties of two different PM tool steel coatings hot extruded with a 1.2714 steel
bar. Current investigations focus on the influence of fused tungsten carbide (FTC)
particles added to the coatings for increasing the abrasive wear resistance as well as
a modification of the matrix material for the coating.
58
As shown in [23, 24], diffusion of alloying elements during the processing and
heat treatment has an essential influence on the development of the interface and
the failure mode occurring in mechanical tests. The particles of FTC are suggested to
affect particularly the carbon concentration by reacting with the tool steel matrix of
the MMC coating. The investigations and results presented here are thus focusing on
the influence of the FTC addition on the microstructure.
6.2
6.2.1 Materials Processing
A gas-atomized cold work steel powder X220CrVMo13-4 (1.2380) and a gasatomized hot work steel powder X40CrMoV5-1 (1.2344) were used as coating
materials for the PM-MMCs. Bars with a diameter of 30mm made of the hot work
steel 55NiCrMoV7 (1.2714) were chosen as substrates for the clad rods. The
chemical compositions of the steels are given in Table 9. An earlier work gives more
details about the coating based on 1.2380 as well as the substrate steel 1.2714 [19].
The second coating investigated here is based on the hot work steel 1.2344.
Coarse particles (100-150µm) of monolithic WC/W2C (fused tungsten carbide,
FTC) were added to the steel powders to improve the wear resistance of the resulting
PM-MMCs. FTC was chosen due to its good ratio of hardness to fracture toughness
[7]. Powder mixtures containing 10 vol.% of hard particles (HP) were blended in a
shaker-mixer for 1h.
The materials combination of the hot work steel powder 1.2344, mixed with
particles of FTC, with the steel bar 1.2714 as substrate is designated “WW1+1.2714”.
“W” states for the hot work steel and “W1” is a reference for 10 vol.% WC/W2C
(FTC). The mixture of the cold work steel 1.2380 with FTC coated on the substrate of
1.2714 is called “KW1+1.2714” with the letter “K” referring to the cold work steel.
The rods cladded with MMCs were produced by putting the substrate material
into capsules (Ø = 78mm, l = 200mm) made of a commercial austenitic stainless
steel (X5CrNi18-10, 1.4301). The surrounding space was filled with the powder
mixture (10% of FTC with 1.2380 or 1.2344) and pre-compressed to tap density by
vibration. The capsules were evacuated, sealed by TIG welding, preheated at
1150°C for 2h and extruded with a ram speed of 36mm/s and a pressing ratio of
59
5.2:1. Due to the hot extrusion, the steel powder is consolidated and bonded to the
massive substrate material while the substrate itself is also deformed during the
process.
The hot extrusion was followed by a heat treatment to assure a high wear
resistance and sufficient toughness of the compound. The extruded bar
WW1+1.2714 was austenitized at 1050°C for thirty minutes, quenched in oil to room
temperature and tempered at 570°C two times for two hours being cooled in air
between each step. The material KW1+1.2714 was austenitized at 1070°C for thirty
minutes, quenched in air to room temperature and tempered at 520°C three times for
two hours also being cooled in air between each step. Finally, an extruded bar with a
diameter of approximately 35mm is formed consisting of a tough core and a wear
resistant layer of several millimeters in thickness [19].
Table 9: Chemical composition of the steels used for coatings and substrates
ASTM designation
1.2380 PM
1.2344 PM
1.2714
coating
substrate
Typical composition (wt.-%)
C
Cr
Mo
V
Mn
Si
Cu
Ni
2,39
12,56
1,10
3,69
0,37
0,55
-
0,30
0,40
5,04
1,34
0,97
0,30
0,19
-
0,10
0,56
1,15
0,46
0,08
0,75
0,29
0,11
1,74
Fe
Bal.
The microstructural examination was carried out by optical microscopy (OM)
and scanning electron microscope (SEM). Samples were cut parallel to the extrusion
direction by electro discharge machining (EDM) to minimize the influence of cutting
on the microstructure. All specimens were ground and polished using diamond paste
down to 1µm grade in order to avoid particle damage. For OM and SEM the
specimens were etched with Nital 3%.
As the co-extruded rods consist of two different tool steel powders in the
coating and another steel in the substrate, diffusion at the interfaces driven by
differences in chemical composition and resulting activity gradients can be expected.
Additionally, the particles of FTC added to the coatings are supposed to influence the
element levels of the MMC matrix. The changes in concentration across the interface
were investigated for each element by several line profiles measured with a wavelength-dispersive spectrometer for electron-probe micro analysis (WDS-EPMA, JEOL
60
JXA-8100). The instrument was operated at an acceleration voltage of 15kV and a
probe current of 20nA. The electron beam was set to perform line scans of 200µm
length being symmetric with respect to the interface region, starting on the coating
side and going towards the substrate material, perpendicular to the extrusion
direction.
6.2.3 Hardness Measurements
To determine the influence of the FTC addition on the matrices of the MMCs,
hardness measurements were performed using a Vickers indenter and a load of
0,5kg. The PM steels 1.2344 and 1.2380 with and without the addition of FTC
particles were measured choosing 10 separate points randomly distributed. In the
coatings containing FTC, the points were carefully chosen between the hard particles
not taking them into account for the hardness values. A region in the middle of the
cross-section of the rod was selected to avoid possible diffusion influences in
hardness from the capsule material and from the substrate.
Due to the large size of the FTC particles and their comparatively low volume
fraction of 10 vol. %, a large mean free path between them results. In combination
with the small load of the hardness measurement, an influence of hard particles
located under the surface is unlikely to occur. This assumption is partly reflected in
the low values of the standard deviation of the hardness values.
For calculating diffusion profiles between the different coatings and the
substrates, the software package DICTRATM [8] was used. With this software
diffusion-controlled transformations are treated on the basis of the following
fundamental concepts [25]:
•
The movement of a phase interface is controlled by the mass balance
obtained from the fluxes of the elements diffusing across the interface region,
•
Diffusion is considered in terms of mobility and true thermodynamic forces,
i.e., gradients in chemical potential. The thermodynamic functions are
calculated with Thermo-CalcTM which runs as a subroutine to DICTRATM;
• A local equilibrium is assumed to exist between phase interfaces.
For mobility data the database MOB2 [10] was used covering a vast number of
elements and phases [26-28]. Additionally, the database TCFE4 was used in
61
Thermo-CalcTM and DICTRATM providing the thermodynamic data. The calculations
were carried out isothermally at 1150°C, considering a one dimensional setup of
16mm in size according to the macroscopic dimensions of the extruded bars (Fig.
30). This region was symmetrically divided into two parts, coating and substrate, by
defining concentration profiles using the heavy-side step function hs(x) (Fig. 31).
Equilibrium states for the steels at T=1150°C and p=101325 Pa calculated with
Thermo-CalcTM [29] software using the TCFE4 database are given in Table 10 and
were used as the starting values for the DICTRATM calculations.
The hot work steels 1.2714 and 1.2344 are fully austenitic at this condition,
while the ledeburitic cold work steel 1.2380 exhibits an austenitic matrix in equilibrium
with M7C3- and MC-carbides (Table 10). Both types of carbides were included in the
DICTRATM simulation using the model for dispersed phases and setting their volume
fractions again using the function hs(x). A grid consisting of 150 points and a higher
point density towards the interface was defined while the simulation time was set to
7.200s (2h).
Table 10: Equilibrium phases and corresponding compositions of 1.2380, 1.2714 and 1.2344 at
T=1150°C and p=101325 Pa calculated with Thermo-Calc using the TCFE4 database
Material
Chemical composition [wt.-%]
C
Cr
Mo
V
Mn
Si
Cu
Ni
Fe
FCC_A1#1 (Austenite)
0.9
7.63
0.77
0.67
0.39
0.65
-
-
bal.
M7C3
8.74
46.44
1.54
7.75
0.32
-
-
-
FCC_A1#2 (MC carbide)
15.82
14.43
7.13
60.09
0.02
-
-
-
bal.
0,40
5,04
1,34
0,97
0,30
0,19
-
0,10
bal.
0,56
1,15
0,46
0,08
0,75
0,29
0,11
1,74
bal.
1.2380
1.2344
1.2714
62
Figure 30: Macroscopic view of the cross section showing the external capsule, substrate, coating
(~8mm thickness) and the interface region.
coating
substrate
interface
16 mm
Figure 31: One dimensional setup of the DICTRATM calculation with a cell size of 16mm according to
the macroscopic dimensions of the extruded bars. Calculations were carried out isothermally at
1150°C, applying a combined model for moving boundaries and dispersed systems.
As the particles of FTC (WC/W2C) cannot be directly considered
simultaneously within one region in the DICTRATM calculations, their dissolution was
calculated separately for the 1.2344 matrix focusing on the analysis of carbon and
tungsten diffusion into the austenitic matrix.
According to the sketch presented in Figure 32, a spherical particle of W2C
with a radius of r=75µm and HCP type in a fully austenitic matrix at 1150°C was
considered as a basic setup for the dissolution calculations. The element levels of the
FCC matrix were set according to 1.2344 (Table 9) while for the initial composition of
W2C the stoichiometric values for W and C were used.
63
The formation of M6C within the diffusion zone between the W2C particle and
the FCC matrix cannot be treated as a layer around the W2C particle as suggested in
Fig. 32a, since this situation would require diffusion through M6C. However, so far
there are no diffusion data available within M6C. Consequently, M6C is treated as a
so-called “diffusion-none” phase. A way of circumventing this problem is offered by
introducing M6C as a so-called spheroidal phase. In this case M6C is introduced as
spherical particles around the W2C carbide. This allows diffusion to be treated within
the region between the particles. The composition of the spheroidal η-carbide phase
was given as the equilibrium composition according to Thermo-CalcTM calculations
and an initial mole fraction of zero for the whole FCC region. During the diffusion
calculations, the mole fraction of M6C is allowed to change according to the
equilibrium calculated locally at each grid point. This allows for the possibility of the
formation of a discontinuous seam of M6C particles as schematically shown in Fig.
32b.
a)
“diffusion-none“ phase
b)
η-carbide (M6C)
η-carbide (M6C)
W2C
hard particle
W2C
hard particle
FCC
spheroidal phase
FCC
Figure 32: Basic setup for the dissolution calculation of a spherical particle of W2C with a radius of
r = 75µm and HCP type in an austenitic matrix of 1.2344: a) M6C (η-carbides) diffusion seam defined
as a closed “diffusion-none” phase around the W2C hard particles, b) defining M6C as a spheroidal
phase not considering diffusion through the η-carbide, but only its thermodynamic stability locally at
the grid points.
6.3
A microscopic view of the interfaces of the two different compounds is
depicted in Figure 33. On the left hand side, the substrate material of 1.2714 with a
64
fully martensitic microstructure can be seen. On the right hand side, the coating of
either 1.2344 or 1.2380 as matrix material with embedded particles of FTC is
presented. The matrix of KW1 additionally contains a high volume fraction of
carbides dispersed in a martensitic microstructure, while the one of WW1 is fully
martensitic. In Figure 33 the seams of M6C formed around each particle of FTC can
be clearly recognized. These seams assist the bonding of the hard particles in the
coating matrix but, due to their brittleness and comparatively low hardness, are an
undesirable microstructural constituent in wear resistant MMC. A diffusion controlled
interaction of the hard particles with the coating matrix can be anticipated. Thus, the
influence of the addition of hard particles on the properties of the coating matrices
was analyzed by measuring the Vickers hardness in the hardened as well as in the
hardened and tempered condition.
WW1 b)
a)
1.2714
M 6C
KW1
1.2714
M 6C
WC/W2C
WC/W2C
Figure 33: Optical micrographs of the interface region between coating (right hand side) and substrate
(left hand side): a) WW1 + 1.2714; b) KW1 + 1.2714.
6.3.1 Influence of FTC on Hardness
The results of the hardness measurements are 580±14 HV5 for the PM steel
1.2344 and 850±25 HV5 for the matrix of the WW1 coating after hardening from
1050°C in oil. The presence of the FTC particles is responsible for the significatively
higher hardness of the coating matrix in the as-quenched state.
After an additional tempering (2 x 2h at 570°C, air), the hardness of the PM
steel 1.2344 without particles of FTC is 525±8 HV5 and 600±15 HV5 for the matrix of
the WW1 coating. Also in the hardened and tempered condition, the interaction of
FTC particles with the surrounding steel matrix leads to an increase of the hardness.
Carbon diffusion from the W2C/WC particles into the coating matrix during heat
65
treatment is likely to be the main source of this increase in hardness by changing the
carbon concentration locally. Thus, a DICTRATM calculation was performed with the
ambition to simulate the dissolution of FTC followed by the carburization of the steel
matrix.
6.3.2 Dissolution of W2C in 1.2344 and Formation of M6C
The dissolution of a W2C hard particle in a matrix of 1.2344, calculated with
DICTRATM, is depicted in Figure 34. The initial tungsten concentration profile and the
one after 2h (7.200s) is presented in Figure 34a. In the first seconds of the
calculation, the FTC particle starts to grow and then a shrinkage process is observed.
After 2h (7200s) a shrinkage of the hard particle by ~0,4µm took place, leading to the
dissolution of tungsten to the steel matrix of 1.2344. At the initial interface between
W2C/WC and 1.2344 coating matrix (x=75µm), a peak in the tungsten concentration
is present, being related to the formation of the M6C phase. Furthermore, a diffusion
profile of tungsten towards the steel matrix of 1.2344 is generated exhibiting a width
of about 15µm. Calculations for longer times resulted in a continuous dissolution
process of the FTC particle and tungsten diffusion to the coating.
In Figure 34b the most important result of the calculation is presented, the
formation of a seam of the M6C carbide. As already pointed out, the model for
dispersed systems used for this calculation allows the formation of a discontinuous
seam of M6C. In other words, the mole fraction around the FTC particle increases by
the formation of M6C particulates and not by the formation of a closed layer as shown
in Figure 32a. At a position where the mole fraction of M6C almost reaches the value
1, a closed layer is formed. This effect can be seen in Figure 34b at the global
coordinate x=75µm. In the surrounding of this closed layer, a lower mole fraction of
M6C is present. The extension of the range incorporating M6C is about 1,5µm after a
calculation at 1150°C for 5min (600s), 5,5µm after 2h (7.200s) and 7µm after 10h
(36.000s).
The order of magnitude but not the absolute values are in agreement with
measurements performed with a conventionally hot isostatic pressed (HIP) material
of FTC particles in a matrix of 1.2380 (Figure 5c). This HIP material is an MMC
successfully used in industrial wear resistant applications and thus well suited for
comparison with the hot extruded materials investigated here. Particularly for
extended tempering times, the widths of the diffusion seams found experimentally
66
are larger than the calculated ones. This deviation could result from the different steel
matrix materials in the measurement and the simulation, but will also be an effect of
the simplified one-dimensional approach used in DICTRATM, considering W2C
instead of FTC and not considering diffusion along grain boundaries.
A solution reaction occurs at the interface between the FTC particle and the
1.2344 matrix. The interdiffusion of elements between matrix and FTC particles,
including its grain boundaries, is maintained by the concentration gradients [30]. This
reaction results in the formation of a diffusion seam of M6C and the partial
transformation of the FTC particle from WC to W2C, explained by the equation WC Ù
W2C + C [31]. This could be a continuous process which causes instability of the WC
particles and dissolution of carbon into new carbides or in the matrix. This continuous
source of carbon is able to carburize the matrix increasing its hardness. However, a
comparison of carbon contents was made in three different matrix alloys [32]
indicating that the carbon content does not play an important role in the formation of
M6C and W2C carbides and in the dissolution of WC.
b)
a)
W2C particle
M6C
W2C
1.2344 matrix
(coating)
1.2344 matrix
(coating)
M6C
67
Figure 34: Concentration profile of tungsten and content of M6C at the interface between W2C and
1.2344 after 2h (7.200s) at 1150°C calculated with DICTRATM: a) Tungsten concentration between the
W2C spherical particle and the matrix of 1.2344 (the step profile represents time = 0s); b) Mole fraction
of M6C showing the formation of this carbide in a range of about 1,5µm after a calculation for 5min
(600s), 5,5µm for 2h (7.200s) and ~7µm for (36.000s), as well as the formation of a closed layer for
the global coordinate x=75µm at the value 1.0 of the M6C mole fraction. c) Mean widths of M6C
diffusion seams measured by image analysis around particles of fused tungsten carbide in 1.2380
processed by hot isostatic pressing at 1150°C for 4h and subsequent tempering at the same
temperature for 2h, 4h and 8h.
The M6C carbide has been previously identified in the Co-W-C system [31,
33], with “M” consisting of the elements W, Cr and Mo. To compare the results of the
DICTRATM calculations with experimental ones, the element levels after 5min, 2h and
10h at 1150°C within the closed seam of M6C at the global coordinate x=75µm were
taken from DICTRATM and compared to EDX measurements. For this purpose,
several EDX measurements were performed in a WW1 coating heat treated during
2h to obtain mean values of the concentrations of Fe, Mo, Cr, W and V in the M6C
diffusion seam (Table 11). The calculated values after 2h differ significantly from the
measured ones. Only iron has calculated values higher then the EDX measurements.
All the others showed smaller volume fractions calculated with DICTRATM.
Comparing the FTC particles added in two different coating steels, WW1 and KW1,
the M6C carbide has roughly the same composition. The W2C/WC particles differ on
the Fe and Mo content, but showed similar values for Cr, W and V.
68
Table 11: EDX measurement values of the FTC particles added in two different coating steel matrices,
in vol. fraction %:
WW1
EDX (measured)
W2C/WC
DICTRA
M6C
2h (7.200s)
KW1
TM
(calculated)
M6C
EDX (measured)
W2C/WC
5min (300s)
2h (7.200s)
10h (36000s)
M6C
2h (7.200s)
Fe
0,63
21,71 ± 0,75
36,05 ± 3,32
30,60 ± 4,53
25,75 ± 0,78
1,29 ± 0,9
21,83 ± 0,59
Mo
0,25
0,49 ± 0,24
0,16 ± 0,10
0,25 ± 0,06
0,36 ± 0,02
0,08 ± 0,1
0,45 ± 0,21
Cr
0,20
2,82 ± 0,09
0,70 ± 0,18
0,68 ± 0,18
0,72 ± 0,11
0,29 ± 0,2
3,74 ± 0,06
W
98,80
73,88 ± 0,88
58,30 ± 6,2
65,10 ± 7,50
70,25 ± 1,5
98,14 ± 1,1
72,62 ± 0,6
V
0,12
1,11 ± 0,12
0,05 ± 0,04
0,07 ± 0,02
0,15 ± 0,06
0,19 ± 0,13
1,36 ± 0,18
6.3.3 Influence of W2C Hard Particle on 1.2344 coating matrix
In Figure 35a the evolution of the integrated mass fractions of carbon in the
1.2344 steel coating reacting with the FTC particle is depicted overlapped with the
mole fraction of M6C after a 2h (7.200s) calculation. At the global coordinate x=75µm
a carburization of the coating matrix is clearly observed in a range of about 85µm.
Calculations for 5min (600s) and 10h (36.000s) revealed a carburization range of
55µm and 60µm, respectively. Tungsten diffuses into the coating matrix, but in a
lower extent compared to carbon (Fig. 35b). The integral content of tungsten in the
FCC matrix is constantly increasing showing that W diffuses easily through the M6C
diffusion seam into the coating matrix [30].
a)
b)
W2C
1.2344 matrix
(coating)
Figure 35: a) Mass fraction of carbon (left hand side y axis) and M6C mole fraction (right hand side, y
axis) showing the carburization effect of the coating matrix starting at the global coordinate x = 75µm,
and b) tungsten mass fractions in grams in the coating matrix as a function of time after 2h (7.200s)
calculated with DICTRATM.
69
The effect of tungsten and particularly carbon diffusion towards the coating
matrix of 1.2344 is an increase in hardness after austenitizing and quenching the
material due to carburization, as already mentioned in section 6.3.1. The results of
the calculation are, thus, in agreement with the experimental findings.
6.3.4 Interactions of the coatings WW1 and KW1 with the steel substrate of 1.2714
The schematic drawing presented in Figure 31 shows the calculation setup
based on the real dimensions of the extruded bars. In order to illustrate the real
conditions, WC/W2C particles are shown in the scheme, but not considered in the
calculations as explained in section 6.2.4. Thus, the diffusion calculations are only
considering the diffusion between the substrate materials 1.2714 and the two
different steels used for the coatings, 1.2344 and 1.2380 not taking into account the
carburization effect of the particles of FTC. This effect was analysed separately and
already explained for the 1.2344 coating.
The line profiles of the major alloying elements were measured by EPMA and
compared with the results of the diffusion calculations. The extension of the interface
region was determined for each element separately from the calculated and
measured profiles. For this purpose, a deviation of 5% from the initial value was
defined in order to determine the diffusion range for each element. The results for the
interface widths are given in Table 12. Comparing the values with those measured by
EPMA a good agreement can be noticed, except for chromium and vanadium in the
materials combinations WW1+1.2714 (Fe) and KW1+1.2714 (Ni).
Table 12: Average interface width per element according to EPMA and DICTRATM, in µm
WW1+1.2714
EPMA
KW1+1.2714
TM
DICTRA
EPMA
K+1.2714 [19]
TM
DICTRA
EPMA
DICTRATM
Fe
86
63
57
44
---
---
Cr
86
58
75
48
24
39
Ni
42
39
53
35
23
26
V
104
61
50
29
28
31
For all alloying elements, the interface region in WW1+1.2714 is wider than in
KW1+1.2714. The reason for this is the larger diffusivity of elements in the compound
with the hot work steel coating as diffusion takes place in a one-phase austenitic
microstructure and is not hindered by the presence of dispersed carbides.
70
In Figure 36, the measured and calculated element profiles for the compound
consisting of a coating of 1.2344 with a steel substrate of 1.2714 are depicted. The
calculated element profiles are in good agreement with the micro-probe analyses.
Vanadium and chromium show depletion in element content from the coating towards
the substrate, corresponding to the concentration gradients. The opposite occurs for
nickel and iron. The measured profile of vanadium shows a peak concentration of
this element of about 1.0 wt% on the coating side, which is not reproduced by the
calculation.
a)
W DICTRA
W EPMA
b)
W DICTRA
W EPMA
1,0
0,8
4
3
V [wt.%]
Cr [wt.%]
5
substrate
coating
0,6
0,4
substrate
coating
2
0,2
1
-100
-50
0
50
100
0,0
-100
-50
Distance [µm]
c)
1,6
d) 95
Fe [wt.%]
Ni [wt.%]
1,4
1,2
1,0
substrate
coating
0,8
50
100
Distance [µm]
W DICTRA
W EPMA
1,8
0
0,6
W DICTRA
W EPMA
94
93
substrate
coating
92
0,4
0,2
91
0,0
-100
-50
0
50
Distance [µm]
TM
Figure 36: Results of DICTRA
100
-100
-50
0
50
100
Distance [µm]
calculations compared with EPMA line-scans. The concentration
profiles are between the coating (left hand side), interface (middle at value zero) and substrate (right
hand side) for the coating of 1.2344 (W) and the substrate of 1.2714. The profiles are from a)
chromium, b) vanadium, c) nickel, and d) iron.
The results for the compound consisting of a coating of 1.2380 with a steel
substrate of 1.2714 are presented in Figure 37. The agreement of calculation and
measurement is worse compared to the aforementioned compound, especially in the
interface region and for the iron profile in the coating side. This might be related to
71
the more complex coating, consisting of a matrix phase and two different kinds of
carbides. In dispersed systems, the DICTRATM model takes into account only
diffusion in the matrix, not considering the dispersed phases. The chromium- and
vanadium-rich carbides in the coating of 1.2380 are taken into account only for the
equilibrium calculations, instead. An enrichment of the chromium level close to the
interface towards the coating is observed in Figure 37a. This enrichment is related to
a higher concentration of M7C3 carbides. In Figure 37b an enrichment of the
vanadium content is also observed corresponding to the VC concentration close to
the interface. The same is found experimentally for vanadium and the vanadium-rich
MC carbide [19] and was investigated using EBSD [34].
a)
14
K DICTRA
K EPMA
12
b)
5
K DICTRA
K EPMA
4
8
V [wt.%]
Cr [wt.%]
10
substrate
coating
6
3
substrate
coating
2
4
1
2
0
-100
0
-50
0
50
100
-100
-50
1,8
50
100
Distance [µm]
Distance [µm]
c)
0
K DICTRA
K EPMA
d)95
K DICTRA
K EPMA
Fe [wt.%]
Ni [wt.%]
1,5
1,2
substrate
coating
0,9
90
substrate
coating
85
0,6
80
0,3
-100
-50
0
50
Distance [µm]
100
75
-100
-50
0
50
100
Distance [µm]
Figure 37: Results of DICTRATM calculations compared with EPMA line-scans. The concentration
profiles are between the coating (left hand side), interface (middle at value zero) and substrate (right
hand side) for the coating of 1.2380 (K) and the substrate of 1.2714. The profiles are from a)
chromium, b) vanadium, c) nickel, and d) iron.
72
Conclusions
This chapter gives an overview of the scientific investigations carried out in
this work. Specific remarks related to the Chapters 3 to 6 are presented separately.
Suggestions for future works and developments are also addressed in the wide field
of abrasion/wear resistant materials and metal matrix composites (MMC) produced
by hot direct extrusion.
The focus of this work was the characterization of the interface microstructure,
its correlation with mechanical properties and the chemical interdiffusion behaviour of
different materials combinations produced by hot direct extrusion.
7.1
General remarks
The aim of this work is attested by the successful hot direct extrusion of two
different PM tool steel powder coatings and a massive tool steel bar producing thick
wear resistant coating layers with a high fracture toughness core.
Several characterization methods were used focusing on the interface region
between coating and substrate giving special attention to:
•
changes in the microstructure and the influence on mechanical properties of
different materials configurations;
•
influence of alloying elements;
•
heat treatment parameters;
73
•
addition of hard particles to the coating matrix;
•
presence of retained austenite;
•
diffusion and dissolution processes.
The main challenges of the work presented here were the diffusion
calculations and sample preparation, especially for the EBSD measurements.
Differences, e.g. in hardness, between two dissimilar materials, coatings and
substrates, made this task very difficult and time consuming.
The DICTRATM calculations proved to be a powerful tool combined with EPMA
measurements facilitating and improving the characterization works, the entire
comprehension of the interface region and a complete assessment of this area giving
quantitative and qualitative reliable data.
7.2
Specific remarks
Chapter 3
Interface Characterization and Mechanical Properties of the Cold
Work Steel Coating (K) Co-Extruded on a 1.2714 Steel Substrate
[19]
The microstructure and mechanical properties of the materials combination
K+1.2714 were investigated using electron microscopy, EPMA (electron probe microanalysis), hardness and tensile tests. Diffusion calculations were performed using the
software DICTRATM.
The obtained results revealed that by co-extrusion of preheated capsules filled
with tool steel powder and a massive tool steel bar, thick wear resistant coating could
be successfully produced showing a pore-free and complete densification of the
microstructure. The element profiles calculated with DICTRATM and measured using
EPMA are in good agreement with experimental results, except for silicon and
molybdenum. An enrichment of the carbon matrix content was calculated for the
coating side of the interface region, as well as a carbon uphill diffusion against the
concentration gradient revealed by DICTRATM calculations due to a higher carbon
activity in the substrate material. EPMA measurements showed an increase in the
carbide volume fractions of Cr7C3 and VC close to the interface with the highest
74
concentration of VC carbides at the interface and of Cr7C3 shifted towards the
coating.
Analysing the mechanical properties in the tensile tests, continuous yielding
and poor plastic elongation occurs while the measured yield strength is in agreement
with the values expected for the substrate material. Hardness profiles correlated with
the fractography indicate the brittle fracture region shifted approximately 50µm from
the interface within the coating side with highest volume fraction of carbides and
increased hardness.
The increase of carbon content in the coating side is certainly one of the
factors responsible for the good bonding between coating and substrate in this
materials combination. One consequence of the interdiffusion of alloying elements is
the concentration of VC and Cr7C3 carbides, both measured and calculated, near the
interface region. On the other hand, after analysing the mechanical properties, the
coating matrix has a high hardness and yield point, but a desired degree of ductility is
not achieved due to the high volume fraction of carbides working as defects for
cracking initiation. These factors fulfil the conditions for a fracture to occur. The
reason for this could be an austenitizing temperature a little higher than enough for
the K+1.2714 materials combination, resulting in loss of ductility and toughness.
Chapter 4
Correlation between Interface Microstructures and Mechanical
Properties of Co-Extruded Layered Structures [20]
The microstructures and mechanical properties of the material combinations
K+1.2714, KW1+1.2714 and WW1+1.2714 were investigated using electron
microscopy, hardness and tensile tests. Diffusion calculations were performed using
the software DICTRATM
The mechanical properties of the interface region were tested using micro
tensile tests and fracture of the specimens occurred within the hardened region of the
coatings. While the difference in carbon activity results in softening of the substrate
close to the interface, the hardness of the coating increases near the interface due to
carburization.
Carbon diffusion and carbon activity from coating to substrate is the most
important feature influencing hardness and ductility in this comparison between three
different materials combinations. The hardening of the coatings and the softening of
75
the substrates close to the interface region are directly related with carburization and
decarburization, both governed by carbon interdiffusion.
Chapter 5
Microstructure Characterization by EBSD/EDX Focusing on the
Influence of Hard Particles Addition and the Formation of Retained
Austenite
Four different materials configurations, K+1.2714, KW1+1.2714, WW1+1.2714
and KW1+S355, of hot extruded rods using two different tool steel powders, with or
without hard particle (HP) addition, clad to two different steel substrate materials
were characterized by EBSD and EDX with special focus on the specimen
KW1+1.2714.
The influence of hard particle addition on the formation of chromium- and
vanadium-rich carbides and the presence of retained austenite (RA) at the interface
region were investigated by a combined EBSD/EDX measurement and elucidated.
The structural steel S355 extruded with the KW1 coating showed the widest interface
region after hot extrusion and heat treatment. The reason for this could be a higher
gradient in activity of the alloying elements between coating and substrate. An
accumulation of vanadium-rich MC carbides close to the interface was found for the
materials K+1.2714 and KW1+S355 but not for KW1+1.2714. Retained austenite
could not be found in the KW1+1.2714 materials combination even though a
pronounced enrichment of carbon at the interface was anticipated [19].
Chapter 6
Influence of Hard Particles Addition and Chemical Interdiffusion
Investigated by Diffusion Calculations on the Mechanical
Properties
The influence of an addition of particles of fused tungsten carbide on the
microstructure of two different tool steels was investigated. Diffusion processes
occurring during processing of particle reinforced MMC by hot extrusion were taken
into account performing DICTRATM calculations, compared to experimental results of
hardness and EPMA measurements. The dissolution of a spherical particle of W2C in
a matrix of the steel 1.2344 was calculated and its influence on the element levels in
76
the steel analyzed. Two effects found experimentally, the shrinkage of the FTC hard
particle and the diffusion controlled formation of a seam of M6C, could be simulated.
The latter were compared to measured widths of diffusion seams around FTC
particles in 1.2380, showing the same order of magnitude, but larger widths in reality
for extended tempering times. Furthermore, an increase in hardness is observed in
the matrix of 1.2344 due to carburization as a result of carbon diffusion from the W2C
hard particles. This result is in agreement with the calculated one. Diffusion profiles
across the interface of substrates and coatings were measured by EMPA and are in
good agreement with the calculated results, particularly for the materials combination
WW1+1.2714. During processing and subsequent heat treatments, chemical
interdiffusion of alloying elements, predominantly of carbon, takes place, influencing
the mechanical properties at the interface of substrate and coating locally.
7.3
Outlook and future works
The investigations and results presented in this work improve the
understanding of diffusion processes in multi-component Fe-base alloy systems and
its influence on mechanical properties.
Questions regarding residual stresses and TEM (transmission electron
microscopy) remain unanswered and would help for a more detailed comprehension
of the chemical and mechanical behaviours in the interface region.
In Chapter 3, further investigations and/or DICTRATM calculations could be
carried out in order to understand why a higher carbon activity occurs in the
substrate, where carbon concentration is far smaller than in the coating. More
mechanical tests might be important for a better understanding of the carbide
influence, e.g. on the mechanical bonding.
In Chapter 4, the influence of other alloying elements could be investigated in
detail as well as mechanical tests in specimens after different heat treatment times. A
ratio between the width of the interface and the location of the fracture region could
be estimated and a comparison between the different heat treatment times would be
done.
In Chapter 5, a complete comparison between the other specimens using
EBSD might be done in order to improve the understanding of the influence of hard
77
particles addition on the formation of chromium- and vanadium-rich carbides and the
presence of retained austenite (RA). A higher activity gradient was identified in the
structural steel S355 used as a substrate, but the reason for that is not fully
understood.
In Chapter 6, the same calculations for the dissolution of a W2C spherical
particle could be performed in a matrix of the steel 1.2380 (K) and the differences
compared with the steel 1.2344 (W). The influence of chemical interdiffusion of
carbon and other alloying elements on the mechanical properties of the 1.2380
matrix could also be analyzed and compared with the 1.2344 matrix.
Furthermore, carbon activity plays an important role influencing the
mechanical properties of coatings and substrates in all the investigated materials
combinations. A measurement or a simulation of carbon activity and the influence of
each alloying element at given temperatures and/or pressures would be a very
important value to help understanding its effects.
Following this path, another important question is how the diffusion of alloying
elements influences the consolidation and bonding of the powdery coating to the
substrate steel bar during the hot direct extrusion process.
The material flow during hot extrusion is also very important for process
optimizations and industrial applications. Studies on distribution of flow were
performed in the early 1930’s and end of 1960’s [36 - 38] and reassessed in the
1990’s [6], but in small scale experiments with cylindrical tin samples. Modelling
efforts in material flow may become an important feature to predict the microstructure
formation and the influence of hard particle addition on the quality of the hot extruded
rods.
7.4
Suggestions for industrial applications
A laboratory sample of a typical industrial application is depicted in Fig. 38: a
cylindrical roller using the KW1+1.2714 materials combination to be used in the
mining and cement industry in areas where abrasion and wear resistance are
necessary [39]. This roller has a wear resistant coating co-extruded with a high
fracture toughness core made of a cheap steel bar. Pin-on-disc wear tests in the
KW1 coating matrix resulted in an increase in the wear resistance in comparison with
78
traditional cast iron. The direct consequence in an industrial plant is an increase in
the equipment’s availability and in the production as well as a reduction of preventive
and corrective maintenance costs.
10mm
10mm
Figure 38: Typical industrial application of a hot direct extruded bar using the KW1+1.2714 materials
combination [39] developed to withstand wear and abrasion: a cylindrical roller to be used in the
mining and cement industry.
Normally, hot extrusion can be used to produce complex shapes even using
metals which are difficult to form. In addition, small lot sizes can be produced
economically. Hot extruded profiles offer the benefit of different material thicknesses
within one profile cross-section, the possibility to use them in highly sensitive areas,
where the special profiles must withstand specific demands of temperature, pressure,
aggressive media or hygienic requirements. Typical parts produced by extrusions are
trim parts used in automotive and construction applications, window frame members,
railings and aircraft structural parts.
Nowadays, rigid–viscoplastic finite element programs have been developed as
simulation tools for analyzing the extrusion of steel profiles [39]. For large industrial
production of extruded rods with wear resistant coating, the effects of various design
parameters such as flow guide and the die type could be simulated aiming for an
optimized production process.
79
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81
Curriculum Vitae
PERSONAL INFORMATION
Last Name
DE SOUZA E SILVA
First Name
PEDRO AUGUSTO
Nationality
Brazilian
Date and place of birth
January, 29th 1974, Belo Horizonte (MG)
EDUCATION AND
PROFESSIONAL EXPERIENCE
Since March 2006
PhD student at the Max-Planck Institute for Iron
research, Material Diagnostics and Steel
Technology Department, Düsseldorf – Germany
June 2004 to Feb. 2006
PhD student at the Institute of Materials Science
and
Technology,
Vienna
University
of
Technology, Vienna – Austria
Aug. 2000 to May 2004
Assistant Mechanical Engineer / Technical
Assistant at Minerações Brasileiras Reunidas
S.A. – MBR (mining company), Nova Lima (MG)
– Brazil
Feb. 1997 to April 2000
Technical-Commercial Budget Assistant at
Milplan Engenharia e Montagens (erection
works), Belo Horizonte (MG), Brazil
Feb. 1997 to Dec. 2003
Mechanical Engineering course (night shift) at
the Pontifícia Universidade Católica – PUC-MG,
Belo Horizonte (MG), Brazil
82

1.2714 - Ruhr-Universität Bochum

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