1. Symposium on
Microstructure, Texture and
Formability in Metal Alloys
16th & 17th September 2015
The University of Pretoria and the University of Cape Town present

2. 1
Microstructure, Texture and Formability of Metal Alloys
NOTES
Contents
GENERAL INFORMATION....................................................................... 3
INTRODUCTION TO THE SYMPOSIUM:.....................................................3
CONFERENCE ORGANISERS ......................................................................4
USEFUL INFORMATION & CODE OF CONDUCT ........................................5
CONFERENCE NAME TAG .....................................................................5
ORAL SESSIONS.....................................................................................5
INTERNET AND WI-FI DURING THE SYMPOSIUM .................................5
DRESS....................................................................................................5
DISCUSSION TIME AT END OF EACH SESSION ......................................5
USEFUL NUMBERS ................................................................................5
BIO’S OF KEYNOTE SPEAKERS:............................................................... 6
DR OLAF ENGLER ......................................................................................6
DR STEFAN ZAEFFERER .............................................................................7
PROF HAMISH FRASER..............................................................................8
SYMPOSIUM VENUE ............................................................................. 9
SYMPOSIUM DINNER.......................................................................... 10
PROGRAMME..................................................................................... 11
ABSTRACTS......................................................................................... 13
BIMODAL MCROSTRUCTURE DEVELOPPEMENT IN TWO-PHASE Ti ALLOY
DURING ROLLING SIMULATION..............................................................14
CHARACTERISING THE MICROSTRUCTURE AND TEXTURE EVOLUTION
DURING STECKEL MILL ROLLING SIMULATION OF AISI 430 FSS.............16
DENSIFICATION MECHANISMS AND MICROSTRUCTURAL EVOLUTION
DURING SPARK PLASMA SINTERING OF Ti6Al4V ALLOY.........................18
THE EFFECT OF INTERMEDIATE ANNEALING ROLLING AND DIRECT
ROLLING ON THE MICROTEXTURE OF AISI 436 FERRITIC STAINLESS STEEL
................................................................................................................20
THE EFFECT OF CAST STRUCTURES ON THE EVOLUTION OF TEXTURE
DURING PROCESSING OF AISI 433 FERRITIC STAINLESS STEEL...............22

3. 2
Microstructure, Texture and Formability of Metal Alloys
NOTES
THE INFLUENCE OF HOMOGENISATION CONDITIONS ON GALLING
RESISTANCE OF ALUMINIUM ALLOY AA3104.........................................23
INFLUENCE OF MICROSTRUCTURAL FEATURES ON THE TENSILE
BEHAVIOR OF TI-6AL-4V ALLOY..............................................................25
INTRODUCTION TO EBSD AND DATA PROCESSING FOR
MICROSTRUCTURAL AND TEXTURE ANALYSIS .......................................27
INVESTIGATION OF MICROSTRUCTURE AND PHASE FRACTION CHANGES
IN 2304 DUPLEX STAINLESS STEEL (DSS); INFLUENCE OF HOT WORKING
CONSTANTS ............................................................................................29
MICROSTRUCTURAL ANALYSIS OF A 12% Cr STEEL AT SUBGRAIN LEVEL
USING EBSD ............................................................................................31
MIPAR™: 2D AND 3D MICROSTRUCTURAL CHARACTERIZATION
SOFTWARE DESIGNED FOR MATERIALS SCIENTISTS, BY MATERIALS
SCIENTISTS..............................................................................................33
TEXTURE CAPABILITIES AT NECSA’S DIFFRACTION FACILITIES ...............35
NOTES ................................................................................................ 37
NOTES ................................................................................................ 38

4. 3
Microstructure, Texture and Formability of Metal Alloys
NOTES
GENERAL INFORMATION
INTRODUCTION TO THE SYMPOSIUM:
The Centre for Materials Engineering at the University of Cape Town and
the Department of Materials Science and Metallurgical Engineering at the
University of Pretoria are jointly hosting this symposium on Microstructure,
texture and formability of metal alloys.
The purpose of this event is to improve the level of metals characterisation
expertise in South Africa with particular emphasis on aspects of
crystallographic texture and its impact on metals processing and
formability.
The theme of the symposium will be led by three international experts, who
will each focus on a particular metal alloy group. The metal groups will
include aluminium alloys, steels and stainless steels, and titanium alloys.
The programme will then include the presentation of pertinent research
activities from our local universities and research institutions. The research
activities will be further contextualized by contributions from our local
metals producing industries.
It is the intention that the symposium will be instructional and principally
aimed at advancing the knowledge of our local postgraduate students and
researchers. In this regard, we have planned that the experts will provide
comprehensive overviews of well-established characterisation techniques,
as well as the presentation of novel, state of the art analytical tools for
crystallographic texture measurement. Moreover, they will focus attention
on the critical microstructure evolution that occurs during metal processing
and metal forming as it relates to the development of texture and the
impact that texture has on controlling mechanical properties and
formability. The student presentation sessions will close with a 5 minute
discussion time, where the audience, led by the session chair and the
keynote speakers, will offer constructive advice and guidance to help the
students improve and lead their work in the direction of a publication in a
high level journal.

5. 4
Microstructure, Texture and Formability of Metal Alloys
NOTES
CONFERENCE ORGANISERS
Sarah George
University of Cape Town
Charles Siyasiya
University of Pretoria
Robert Knutsen
University of Cape Town
Waldo Stumpf
University of Pretoria

6. 5
Microstructure, Texture and Formability of Metal Alloys
NOTES
USEFUL INFORMATION & CODE OF CONDUCT
CONFERENCE NAME TAG
You will receive an official conference name tag when you register. Please
wear this name tag at all times. This will be used for entry into lunch, as
well as entry onto the transport to the conference dinner.
ORAL SESSIONS
Presentations are limited to 20 minutes. All presentations must be pre-
loaded onto the supplied PC in the venue before the start of the session.
INTERNET AND WI-FI DURING THE SYMPOSIUM
The conference venue has wi-fi. Please enquire at registration for the
details.
DRESS
The delegates are expected to respect a smart casual business attire dress
code for the symposium.
DISCUSSION TIME AT END OF EACH SESSION
There will be 5-10 minutes allocated time at the end of each session for a
discussion on the talks that were presented in that session. If you have any
advice for the presenters regarding their content, or any further questions,
please use this time to address the presenters.
USEFUL NUMBERS
Conference organiser Sarah George (UCT) 084 2833296
Conference organiser Charles Siyasiya(UP) 082 3617764
Rondebosch police station 021 685 7345

7. 6
Microstructure, Texture and Formability of Metal Alloys
NOTES
BIO’S OF KEYNOTE SPEAKERS:
DR OLAF ENGLER
The keynote presentation contribution of
Dr Engler will focus on microstructure and
texture development during aluminium
processing, which will be underpinned by
participation from our local Hulamin Rolled
Products industry.
Dr Olaf Engler is a well-established and
internationally acclaimed physical
metallurgist, who has worked extensively in
the field of aluminium alloys and stainless steels. He was instrumental in
the early day in contributing towards the development of the electron
backscattered diffraction techniques and he spent several years working on
aspects of texture characterisation at the Los Alamos National Laboratory
in the USA. He has worked extensively in industry, particularly within the
production and development of aluminium alloys and he is presently the
senior scientist within the research and development division of Hydro
Aluminium Rolled Products in Germany. He has published over 100
reviewed papers, contributed substantially to international conferences
and contributed to the publishing of 13 books, book chapters and
monographs.

8. 7
Microstructure, Texture and Formability of Metal Alloys
NOTES
DR STEFAN ZAEFFERER
The keynote presentation contribution of Dr
Stefan Zaefferer will focus on microstructure and
texture development during steel and stainless
steel processing, which will be underpinned by
participation from our local Columbus Stainless
Steel and Arcelor Mittal industries.
Dr Stefan Zaefferer studied physical metallurgy
and metal physics at the Technical University of
Clausthal in Germany, and spent time at
Université Paris-Sud, Orsay in Paris and Kyoto University in Japan, where he
focussed mainly on the deformation and recrystallization mechanisms of
metals with fcc crystal structures using TEM. Since 2000 he is head of the
research group Microscopy and Diffraction at the Max-Planck-Institut für
Eisenforschung GmbH (MPIE). In this time he has mainly dealt with the
development of electron diffraction techniques in SEM and TEM and
applied these methods to the investigation of various metallic and
intermetallic materials. Currently, Dr Zaefferer tries to focus his research
activities on projects where environmental protection and sustainability
play an important role, e.g. materials for thermal solar power plants,
photovoltaic materials, and light, abundant and recyclable structural
materials.
In 2009 he passed his habilitation at the RWTH Aachen and teaches there,
together with colleagues from the RWTH and the MPIE, a self-developed
master course “Microstructures, Microscopy and Modelling“. Dr Zaefferer
is also a guest professor at various international universities (currently
Vienna, Vancouver and Melbourne). Dr Zaefferer has published extensively
and his work can be accessed from his website:
http://www.mpie.de/2973988/employee_page?c=3079071&employee_id
=40057

9. 8
Microstructure, Texture and Formability of Metal Alloys
NOTES
PROF HAMISH FRASER
The keynote presentation contribution of
Professor Hamish Fraser will focus on
microstructure and texture development
during titanium alloy processing, which will be
underpinned by participation from
representative from the Titanium Centre of
Competence (a current DST initiative).
Professor Hamish Fraser is a well-established
and internationally acclaimed physical
metallurgist, who has worked extensively in the field of titanium alloys. He
presently serves as Director of the Centre for the Accelerated Maturation
of Materials (CAMM) at OSU. He has been a member of the National
Materials Advisory Board and the US Air Force Scientific Advisory Board. He
has consulted for a number of national laboratories and several industrial
companies. He is a Fellow of TMS, ASM, IOM3
(UK), and MSA. He has
published over 380 papers in scholarly journals, and given over 280 invited
presentations. He has graduated 48 doctoral students and 36 students
graduating with the degree of M.S. His work is based on research involving
the development of advanced methods of materials characterization
(involving high resolution and analytical electron microscopy), materials
processing, and microstructure/property relationships. Dr. Fraser has an
active research program in the development of new and improved
materials, including: advanced materials characterization, direct 3-D
microstructural representation, modelling microstructure/properties in
light alloys, with an emphasis at present on titanium alloys and more
recently high entropy alloys.

10. 9
Microstructure, Texture and Formability of Metal Alloys
NOTES
SYMPOSIUM VENUE
BELMONT SQUARE CONFERENCE CENTRE
Belmont Road, Rondebosch
Tel: 021 685 2625 Fax: 086 523 8262
Email: info@belmontsquare.co.za
Website: www.belmontsquare.co.za
GPS 33°57'34''S 18°'28'25''E

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Microstructure, Texture and Formability of Metal Alloys
NOTES
SYMPOSIUM DINNER
The symposium dinner will take place at The Harbour House, in Kalk Bay.
Kalk Bay Harbour, Kalk Bay
021 788 4134/5
kalkbay@harbourhouse.co.za
There will be a bus to transport you to and from the venue. The bus will
leave directly after the close of day 1 from the Belmont Conference Centre,
and will drop you off at the same place at about 22h30.
Please use a taxi or Uber to get home from the venue, as it is not safe to
walk too far late at night.
There will be a free 30 minutes once we get to Kalk Bay for you to walk along
the harbour wall and admire the view before we must be seated for dinner.
Please note: Beer, cider and softdrinks can be ordered at the restaurant,
and wine will be provided on the tables. Any other drinks will be at your
own expense.

12. 11
Microstructure, Texture and Formability of Metal Alloys
NOTES
PROGRAMME
WEDNESDAY 16 SEPTEMBER 2015
TIME
8.00-8.30 REGISTRATION (FOYER)
8.30-8.45 WELCOME ADDRESS: PROF ROBERT KNUTSEN (HEAD OF MECHANICAL ENGINEERING,
UNIVERSITY OF CAPE TOWN)
8.45-9.45 KEYNOTE SPEAKER: PROF STEFAN ZAEFFERER
Characterisation of metal microstructures by orientation measurements; from grain boundary
engineering to bulk plasticity
SESSION 1 – CHAIRED BY CHARLES SIYASIYA
PRESENTER TITLE
9.50-10.10
MAUMELA
The effect of cast structure on texture, ridging and formability in AISI 433 ferritic
stainless steel
10.10-10.40
MASINDI
Microstructure evolution of AISI 430 ferritic stainless steel during steckel mill hot
rolling
10.40-11.00 DICKS Introduction to EBSD and data processing for microstructural and texture
analysis – Part 1
11.00-11.10 SESSION DISCUSSION
11:10-11:30 - TEA
SESSION 2 – CHAIRED BY WALDO STUMPF
PRESENTER TITLE
11.30-11.50 BILL Investigation of microstructure and phase fraction changes in 2304 duplex
stainless steel (DSS); Influence of hot working constants
11.50-12.10 DEYZEL Microstructural analysis of a 12% Cr steel at subgrain level using EBSD
12.10-12.30
MATTHEWS
Columbus stainless steel – An overview for stainless steel research in South
Africa
12.30-12.40 SESSION DISCUSSION
12:40-13:40 - LUNCH
SESSION 3 – CHAIRED BY SARAH GEORGE
PRESENTER TITLE
13.40-14.40 KEYNOTE SPEAKER: PROF HAMISH FRASER
The role of texture in influencing the mechanical properties of titanium alloys (including
representation and interpretation of non-cubic HCP type texture components in titanium alloys).
14.40-15.00 KAFER The influence of temperature and strain rate on the texture development of
commercial
Ti-6Al-4V alloy
15.00-15.20 SOSA MIPAR™: 2D and 3D microstructural characterization software designed for
materials scientists, by materials scientists
15.20-15.30 SESSION DISCUSSION
15.30-15.50 – TEA
15:50-16:00 CLOSE OF DAY 1: INSTRUCTIONS FOR DINNER

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Microstructure, Texture and Formability of Metal Alloys
NOTES
PROGRAMME
THURSDAY 17 SEPTEMBER 2015
TIME
SESSION 1 – CHAIRED BY ROBERT KNUTSEN
PRESENTER TITLE
9.00-9.20
RASIAWAN
Investigation of the fine grain heat affected zone of repair welded creep aged
X20 CrMoV12-1
9.20-9.40 UANANISA The effect of intermediate annealing rolling and direct rolling on the micro-
texture of an AISI 436 ferritic stainless steel
9.40-9.50 SESSION DISCUSSION
9.50-10.50 KEYNOTE SPEAKER: DR OLAF ENGLER
Controlling texture during commercial sheet production of aluminium alloys and the importance
of textures in influencing formability.
10:50-11:10 - TEA
SESSION 2 – CHAIRED BY CHARLES SIYASIYA
PRESENTER TITLE
11.10-11.30
MAGIDI
The influence of homogenisation conditions on galling resistance of aluminium
alloy AA3104
11.30-11.50 MASETE Influence of microstructural features on the tensile behavior of TI-6AL-4V alloy
11.50-12.10 VILANE Title
12.10-12.30 SHIRRAN Hulamin Rolled Product – An overview for aluminium research in South Africa
12.30-12.40 SESSION DISCUSSION
12:40-13:40 - LUNCH
SESSION 3 – CHAIRED BY SARAH GEORGE
PRESENTER TITLE
13:40-14:00 VENTER Texture capabilities at NECSA’s diffraction facilities
14.00-14.20 CAIN Title
14.20-14.40 ADEGBENJO Densification mechanisms and microstructural evolution during spark plasma
sintering of
Ti-6Al-4V alloy
14.40-15.00
MUTOMBO
Bimodal microstructure development in two-phase Ti alloy during rolling
simulation
15.00-15.10 SESSION DISCUSSION
15.10-15.30 - TEA
15.30-15.45 CONFERENCE CLOSING: PROF ROBERT KNUTSEN (HEAD OF MECHANICAL ENGINEERING,
UNIVERSITY OF CAPE TOWN)

14. 13
Microstructure, Texture and Formability of Metal Alloys
Symposium on
Microstructure, Texture and Formability
of Metal Alloys
2015
ABSTRACTS
FOR ORAL PRESENTATIONS

15. 14
Microstructure, Texture and Formability of Metal Alloys
NOTES
BIMODAL MCROSTRUCTURE DEVELOPPEMENT IN TWO-PHASE Ti ALLOY
DURING ROLLING SIMULATION
K. Mutombo1,2
, C. Siyasiya2
, W.E. Stumpf2
1
MSM/ Light Metals, Council for Scientific and Industrial Research (CSIR), Pretoria,
South Africa,
2
Department of Materials Science & Metallurgical Engineering, University of
Pretoria, South Africa
The two-phase (α+β) Ti alloys allow a
mixture of both α and β-phase at room
temperature. They contain alloying
elements such as Al, Ga, Ge, Sn, and Zr
as α-stabilizers and V, Mo, Nb, Ta, Cr,
Mn, Fe, Cu, Ni as β stabilizers. The
Ti6Al4V alloy is the most common two-
phase α/β Ti alloy, accounting for more
than 50% of the total market. A two–
phase Ti alloy can exhibit a variety of
microstructures, such as α/β colony
lamellar structure developed during
slow cooling from high-temperature
single-phase β field. This microstructure
leads to moderate strength and good
fatigue crack growth resistance but low
ductility. In contrast, a microstructure
comprising globular α-phase in a matrix
of β or transformed β possesses a better
balance of strength and ductility and it is
thus more desirable for forming process
such as forging and extrusion because of
its superplasticity property. In this
regard, a two-phase Ti alloy has
attracted many researchers to
investigate its complexity in terms of
microstructural evolution, phase
transition and morphology, and
deformation behavior during hot
compression and rolling1,2,3
. Therefore,
the effect of temperature, strain rate,
strain and rolling passes on
development of α-globular grains in a
transformed β matrix needs to be
further investigated.
The two phase Ti alloy, specifically
Ti6Al4V alloy grade 5, in the annealed
condition, was supplied by BAOTI
Titanium industry as a billet of 65 mm
diameter and 2540mm length. The
rolling simulation was performed in the
Gleeble(TM)
1500D thermo-mechanical
simulator. Cylindrical Ti6Al4V samples
were wire cut and machined to 10mm
diameter and 15mm height. Hot rolling
simulation was carried out on these
cylindrical samples in the temperature
range of 750-900o
C, at a strain rate of
0.001-5s-1
and to a true strain ranging
from 0.3 to 1.4. Tungsten carbide (WC)
anvils with a silica barrier around them
were used to minimize the temperature
gradient along the length of the sample.
The metallographic observation was
performed on the inverted LeicaTM light
optical microscope equipped with an
automatic image analyzer Image-
ProPLUS™ 5.1 and IMAGEJ™, and a
Scanning Electron Microscope (SEM)
equipped with Energy Dispersive X-ray
spectroscopy (EDX). The grain size, α/β
colony size, α lath thickness, and the
volume fraction of α and β-phase was
determined. The spheroidization of α/β

16. 15
Microstructure, Texture and Formability of Metal Alloys
NOTES
lamellae was then quantified. Ti6Al4V
samples, in deformed conditions were
cut to about 3mm thick and 15mm
diameter to isolate the area of interest
during ion milling. The ZeissTM Focused
Ion Beam/Scanning Electron Microscopy
(FIBSEM) was used for cutting and
thinning of Transmission Electron
microscopy (TEM) specimens at a
specific location. The TEM micro-
specimen was transferred to a copper
grid for further TEM analysis. The HR-
TEM JEOLTM JEM-2100 operating at
200kV with LAB6 filament and Tecnai
TM G2 F20 X-Twin MAT with field
emission gun was used to identify the
phases and investigate the
microstructural change during the
rolling simulation process. X-Ray
Diffraction (XRD) analysis was
performed on deformed samples to
investigate the microstructural change,
the presence and phase transition
during rolling.
The resulting bimodal microstructure, as
shown in Fig.1 is characterized by α-
globules in the matrix of transformed β.
A large amount of deformation
happened in the β or (α+β) grains, where
the deformation initiated at the α/β
phase boundaries, as revealed in Fig.2.
References
1. Zherebtsov, S., et al. (2011).Acta
Mater. 59, 4138–50.
2. Collings, E.W. (1984) The Physical
Metallurgy of Titanium Alloys. USA:
American Society for Metals..
3. Semiatin, S.L. (2005) ASM Handbook,
14A , Metalworking: Bulk Forming. ASM
Int. 552–613; 651–99.
Figure 1. Spheroidized α-phase and
transformed β matrix developed
during the rolling of Ti6Al4V alloy.
Figure 2. Rolled Ti6Al4V,
containing α, β and α'-phase, and
deformed areas.
Corresponding author: kmutombo@csir.co.za

17. 16
Microstructure, Texture and Formability of Metal Alloys
NOTES
CHARACTERISING THE MICROSTRUCTURE AND TEXTURE EVOLUTION
DURING STECKEL MILL ROLLING SIMULATION OF AISI 430 FSS
R.R. Masindi, S.L. George and R.D. Knutsen
Centre for Materials Engineering, Dept. Mechanical Engineering, University of
Cape Town, South Africa
Ferritic stainless steels (FSS) can exhibit
ridging parallel to the rolling direction
when subjected to tensile forces or deep
drawing. Ridging is caused by
inhomogeneous plastic flow, as a result
of the spatial distribution of texture
components. Although AISI 430 is a
ferritic stainless steel, it is in the dual
ferrite-austenite phase during
conventional Steckel mill hot rolling.
Consequently, the microstructural
evolution is influenced by both phase
balance and imposed thermo-
mechanical processing conditions. The
Steckel mill rolling process comprises an
odd number of roll passes, usually five,
seven or nine, at a temperature within
the dual ferrite-austenite phase region
throughout the rolling process. The
strain rates and interpass times increase
with each pass. The final texture
distribution and microstructure
evolution results from the processing
parameters and microstructural
development that occurs through
deformation, as well as both dynamic
and static restoration mechanisms. The
dominant mechanism for
microstructural evolution in ferritic
stainless steels during this type of
thermo-mechanical processing is
dynamic recovery1
.
The objective of this study is to simulate
the first three passes of the Steckel mill
hot rolling processing in order to
characterise the microstructure
evolution and texture development of
AISI 430 FSS. Steckel mill hot rolling
parameters are simulated using plane
strain compression testing in the
Gleeble 3800 thermo-mechanical
facility. The samples were quenched to
room temperature after each
deformation pass. The effect of the
accumulated deformation on the
microstructural evolution after each roll
pass is investigated using both light
microscopy, Scanning Electron
Microscopy (SEM) and Electron
Backscatter Diffraction (EBSD).
Figure 1: EBSD Euler maps of three
consecutive Steckel mill passes.
The microstructures and flow stress
curves reveal that dynamic recovery is
the dominant restoration mechanism in
AISI 430 stainless steel. Fig. 1 shows the
microstructure after three consecutive
simulated Steckel mill passes. The
microstructure consists of bands of

18. 17
Microstructure, Texture and Formability of Metal Alloys
NOTES
ferrite and martensite, the latter
representing the prior austenite regions.
The microstructure analysis after each
pass showed that the ferrite subgrain
size is generally much smaller in regions
immediately adjacent to the prior
austenite phase as a result of strain
partitioning to the ferrite by the harder
austenite phase at rolling temperatures.
Therefore, the deformation conditions
and volume fraction of austenite during
hot rolling influence the subgrain
structure and strored energy
accumulation in the ferrite. The degree
of stored energy may influence
recrystallization in the ferrite. Fig. 2
shows the texture development in the
ferrite and prior austenite regions after
each successive pass. The ODFs in Fig. 2
show that the ferrite grains consisted of
α and γ fiber textures. The clustering in
the ODFs show that, as the amount of
strain increases (i.e. increase in pass
number), the {001}<110> and
{111}<112> texture components
become more prominent, although
there is a much greater strengthening of
the deformation {001}<110>
component. The prior austenite regions
also show similar texture components to
the ferrite grains, but in this case the γ
fiber texture {111}<112> becomes more
prominent with increasing strain.
The simulation of the first three passes
of the Steckel mill roll schedule has
demonstrated the polarization in
texture development between the
primary ferrite and prior austenite
regions in the dual phase microstructure
evident in AISI430 FSS. The
microstructure evolution within the
ferrite grains is attributed to recovery as
the dominant mechanism which is
evident in the strong development and
retention of the {001}<110> fiber
texture.
References
1. Hinton, J.S. and Beynon, J.H.,
(2007) ISIJ, 47(10), 1465-1474.
2. Jonas, J.J., (2009), Springer
Corresponding author: msnrab001@myuct.ac.za
Figure 2: ODFs at ϕ2 = 45° showing texture development during the simulated
Steckel mill hot rolling process in the ferrite and prior austenite regions.

19. 18
Microstructure, Texture and Formability of Metal Alloys
NOTES
DENSIFICATION MECHANISMS AND MICROSTRUCTURAL EVOLUTION
DURING SPARK PLASMA SINTERING OF Ti6Al4V ALLOY
A.O. Adegbenjo, E. Nsiah-Baafi, S. Diouf, M.O. Durowoju, M.B. Shongwe, M.
Ramakokvhu and P.A. Olubambi
Institute for NanoEngineering Research, Tshwane University of Technology, PMB
X680, Pretoria, South Africa
Ti6Al4V is the most widely used α+β
titanium alloy due to its good
comprehensive mechanical
performance1
. Among the production
processes commonly associated with
this alloy, powder metallurgy has been
the most recent focus, owing to the
ability to produce near net shaped parts
with complex geometry, with the added
advantage of reduced production costs.
Notwithstanding the known benefits of
this type of technique, there is limited
literature available relating to its
densification and microstructural
evolution during sintering. Spark Plasma
Sintering (SPS), though a relatively new
technique has attracted much interest in
research activities over conventional
sintering methods due to the high
heating rates, short sintering cycles and
low sintering temperatures utilized
during the process. This present study
investigates the densification
mechanisms and microstructural
evolution during the spark plasma
sintering of pre-alloyed Ti6Al4V (100
mesh) powders.
Sintering experiments were performed
in the 650 °C –850 °C temperature
range, at constant heating rate of 100
°C/min, an applied pressure of 50 MPa
and a holding time of 5 minutes. The
densification mechanism was
investigated using the sintering profile
obtained from the SPS system. The
experimental density of the sintered
samples was measured using the
Archimedes’ principle. Microstructural
characterization of the as-polished
sintered samples was done using
Scanning Electron Microscopy (SEM)
and Energy Dispersive x-ray
Spectroscopy (EDS). Polished samples
were etched with Kroll’s reagent (6 ml
HF, 12 ml HNO3 in 150 ml H2O).
The sintering profile seen in Figure 1,
shows near full densification and
maximum shrinkage was achieved at
850 °C. The measured density increased
from 85.33 % at 650 °C to 99.55 % at 850
°C, representing a 14.22 % enhancement
in densification with the increased
sintering temperature. At 650 °C (Fig. 2),
a porous structure with the pores (dark
spots indicated by arrow) situated at the
grain boundaries was observed. The
basket weave (α+β acicular structure)
characteristic of Ti6Al4V alloys were
observed at both 650 and 750 °C, and
can be identified in the images in Figure
2. The proportion of each phase was
dependent on the sintering
temperature. The α phase was more
dominant at lower temperature and

20. 19
Microstructure, Texture and Formability of Metal Alloys
NOTES
owing to the fact that α is less dense
than the β phase, coupled with the high
level of porosity evident in the
microstructure after sintering at that
temperature, explains the lower density
of the sample at 650 °C. A reduction in
pore size and pore density was observed
with increased temperature. The
samples sintered at 850 °C, show very
few pores leading to an improvement in
density. The β phase was more
prominent at high temperature with
plates of the α phase lamella situated
within the β matrix. Diffusion bonding
at particle boundaries in addition to
other diffusivity paths was enhanced at
high temperatures resulting in a re-
orientation of the Ti6Al4V particles
thereby closing up pores due to
boundary sliding and grain rotation, thus
leading to an enhanced material
density2, 3
.
References
1. Long, Y. et al. (2013) Mater Sci Eng A,
585, 408-414.
2. Ergǘl, E. et al. (2009) Powder Metall,
52(1), 65-71.
3. Teber, A. et al. (2012) Int. J Refract
Met Hard, 30, 64-70
Figure 1. Densification profile of sintered samples at different temperatures
Figure 2. SEM micrographs of sintered samples at different temperatures
Corresponding author: waleeleect@gmail.com

21. 20
Microstructure, Texture and Formability of Metal Alloys
NOTES
THE EFFECT OF INTERMEDIATE ANNEALING ROLLING AND DIRECT
ROLLING ON THE MICROTEXTURE OF AISI 436 FERRITIC STAINLESS STEEL
H.J. Uananisa1
, C.W. Siyasiya1
, W.E. Stumpf1
and J. Papo2
1
Dept. Material Science & Metallurgical Engineering University of Pretoria, South
Africa
2
Advanced Materials Division, Mintek, Randburg 2125, South Africa
The need to reduce weight while
maintaining good mechanical properties
in materials used in the automotive
industry has resulted in an increased
interest in steel alloys. The Ferritic
Stainless Steel (FSS) family has been
widely used in this industry, with the AISI
436 type increasingly being used for
automotive trims and mufflers for
exhaust systems. However, there have
been reports of some poor surface
roughening of this material during deep
drawing, with tearing and/or cracking
occuring. This has been attributed to
some differences in the localized
mechanical properties as a result of
grain clusters in the rolled and annealed
material. The main objective of this
study is to optimize the process
parameters within the cold rolling
process, thereby eliminating the
intermediate annealing stage, and
consequently reducing processing costs
while achieving improved formability
and surface roughness properties in the
AISI 436 FSS.
AISI 436 FSS samples were used for a
production trial operation, in which two
processing routes, namely Direct Rolling
(DR) and Intermediate Rolling (IR) were
followed. This study focuses on cold-
rolled and annealed samples, with
thicknesses ranging from 0.46mm to
0.50mm. Each sample was firstly
analyzed under an optical microscope to
determine the grain morphology and
grain-size measurements using ‘Image J’
line-intercept method. The samples
were then prepared for Electron Back
Scatter Diffraction (EBSD) analysis using
a Jeol JSM-IT300LV Scanning Electron
Microscope (SEM). Orientation
Distribution Functions (ODFs) were
developed using the series expansion
method (lmax=22). The mean R-values
(Rm) and planar anisotropy values (Δr)
for the various samples were measured
after 10% strain along the longitudinal
(0o
), transverse (90o
), and diagonal (45o
)
directions. The Rm and Δr values were
then calculated using some standard
equations.
Planar anisotropy (Δr) gives an
indication of the amount of necking or
earing that will occur on the edges of the
deep-drawn items1
. Moreover, separate
samples of the steel, which were then
deep-drawn into lug-nut covers through
a multi-stage forming process, were also
analyzed for surface roughness. Sample
‘D1’ was intermediate rolled (IR) prior to
the deep-drawing process, while ‘D2’
was directly rolled (DR). This analysis
was used to investigate the effect of the
production processing route on the
surface roughness of the steel, and the

22. 21
Microstructure, Texture and Formability of Metal Alloys
NOTES
differences between the two processing
routes are clearly visible in Fig.1. The
results obtained showed an even
distribution of the γ-fiber texture,
orientation of grains in the {111} plane,
for both processing routes, i.e. texture
banding was not observed in the two
routes. Moreover, the process route did
not seem to have a significant influence
on the average grain-size, although an
effect on grain-size distribution was
observed, with some grain-size
clustering clearly visible in some
instances (see Figure 2). This grain
clustering phenomenon was seen by
Knutsen and Wittridge2
where the
influence of grain-size banding is
identified and associated with
differential yielding under tension,
hence distortion of the surface. The
Lankford values (Rm) were calculated for
both D1 and D2 and gave similar values
for the two routes, recording average
values of about 1.9 and 1.6 for DR and IR
respectively. However, the IR route
showed a significantly lower planar
anisotropy, as evidenced by an average
∆r value of 0.5 compared to an average
∆r value of -1.4 for the DR route, hence
suggesting superior formability and
drawability of the IR route over the DR
route.
References
1. Maruma, M.G., Siyasiya, C.W. and
Stumpf, W.E. (2003), SAIMM, 113,115-
120.
2. Knutsen, R.D. and Wittridge, N.J.
(2002), Materials Science and
Technology 18, 1279-1285.
Figure 1. SEM images of D1 and D2 after the stage 2 of forming process.
Figure 2. Inverse pole figure maps //ND for the stage 1 of the deep-drawing.
Corresponding author: h.uananisa@gmail.com
001
111
101

23. 22
Microstructure, Texture and Formability of Metal Alloys
NOTES
THE EFFECT OF CAST STRUCTURES ON THE EVOLUTION OF TEXTURE
DURING PROCESSING OF AISI 433 FERRITIC STAINLESS STEEL
M. Maumela1,2
, C.W. Siyasiya1
and W.E. Stumpf1
1
Dept. Materials Science and Metallurgical Engineering, Faculty Engineering, Built
Environment & Information Technology, University of Pretoria
2
Mintek, Randburg, South Africa
433 ferritic stainless steel (FSS) alloy is a
modification of the 430 alloy through
the addition of Al. The 433 alloy is
generally susceptible to ridging during
deep drawing operations. Ridging is a
surface defect characterized by the
formation of narrow ridges parallel to
the rolling direction of the sheet which
impair the aesthetics of the product.
Since the ridging defect is cause by
banding of micro-texture, it can be
minimized or eliminated by optimizing
the homogeneously distributed γ-fibre
texture, which is referred to all
orientations of {111}//ND [2].
Therefore, this work explores the role of
the cast structure (i.e. columnar versus
equiaxed grain structure) on the
evolution of crystallographic texture.
433 material, with different starting
grain structures, namely columnar and
equiaxed, were lab rolled in order to
simulate the last two rough and last
three Steckel hot rolling passes. This was
followed by 62% cold rolling and
annealing at 900o
C for 180 seconds.
Texture was measured using Scanning
Electron Microscopy Electron
Backscattered Diffraction (SEM-EBSD),
ridging properties were assessed using
Taylor Hobson profilometer the
Ultimate Tensile Strength (UTS) and
elongation were measured using a
tensile load applied along the RD with a
crosshead speed of 2 mm/minute in an
Instron Universal Tensile testing
machine of 10 ton capacity according to
the ASTM A 370 standard. Figure 1
shows EBSD Orientation Distribution
Function (ODF) diagrams for the cold
rolled and annealed materials of an (a)
equiaxed and (b) columnar cast
structure. The starting equiaxed grain
structure developed the homogenous γ-
fibre texture, {111}//ND, while the
columnar grain structure developed
heterogeneous γ-fibre, {001} < 110 > ,
and Goss {110}[001] texture
components. The desired homogeneous
γ-fibre texture is owed to the as-cast
random texture of the equiaxed grains.
The dominant Cube texture in the
columnar materials is known to be
responsible for the subsequent poor
texture after hot rolling, cold work (CW)
and annealing processes. Raabe et al.
found that heterogeneous texture
distribution impairs the forming
properties, while homogeneous γ-
texture improves them1
.
Figure 2 shows the inverse pole figure
(IPF) and the corresponding EBSD map.
The samples with the starting structure
of columnar grains resulted in, not only
poor texture, but clustering of grains in
the form of bands, which is associated

24. 23
Microstructure, Texture and Formability of Metal Alloys
NOTES
with ridging. Contrary to this, the same
band-like clusters were not observed in
the sample with the equiaxed starting
grain structure. However, there was no
significant difference in the average
grain size for the two materials, i.e.
average grain size of 33 μm and 38 μm,
for the equiaxed and columnar starting
structures respectively. The UTS for both
materials was found to be
approximately 590 MPa and the
elongation was found to be 25% and
28% for columnar and equiaxed starting
structures respectively. The ridging
resistance was found to be higher for
the equiaxed sample than the columnar
samples, which was attributed to the
respective recrystallization textures.
References
1. Raabe, D., et al. (2005). Steel
Research Int. 76 no.11.
2. Engler, O, and Huh, M.Y. (2001)..
Material Science and Engineering A
3008, pp 74-87.
Figure 1. EBSD textures ODF section 𝜑2 = 45° after 62% CW and annealing at 900 o
C
for 180 s, (a) equiaxed and (b) columnar.
Figure 2. IPF after 62% CW and annealing at 900 o
C for 180 s, (a) equiaxed and (b)
columnar.
Corresponding author: maumela.britz@gmail.com

25. 23
Microstructure, Texture and Formability of Metal Alloys
NOTES
THE INFLUENCE OF HOMOGENISATION CONDITIONS ON GALLING
RESISTANCE OF ALUMINIUM ALLOY AA3104
L.T. Magidi, S.L. George, and R.D. Knutsen
Centre for Materials Engineering, Dept. Mechanical Engineering, University of
Cape Town, South Africa
Aluminium alloy AA3104 is commonly
used for the manufacturing of beverage
can bodies. This alloy has excellent
formability and strength properties. The
evolution of the AA3104’s
microstructure and intermetallic
particles during thermo-mechanical
processing has a direct impact on quality
parameters. These parameters
influence the formability of the material
during beverage can deep drawing and
wall ironing. The parameters are earing,
tear-off and galling resistance1,2
. During
homogenisation of the cast ingot, an
intermetallic particle phase
transformation from the orthorhombic
β-Al6(Fe,Mn) phase to the harder cubic
α-Al15(Fe,Mn)3Si2 phase occurs2
. The
presence of the α-phase intermetallic
particles is crucial for galling resistance,
thus improving the formability of the
material. For ideal galling resistance the
material requires 1-3% total Volume
Fraction (VF) of intermetallic particles,
50% of which should be the harder α-
phase.
The homogenisation practice variables,
such as temperature and time, as well as
the effect of the intermetallic particle
VFs with the correct β to α ratio are
investigated. Phase transformation and
VF analyses were performed on samples
extracted from the edge and center of
the AA3104’s direct-chill cast ingot. The
samples were homogenised using a two-
step practice, where the primary
temperature step was varied between
560°C and 580°C, and the secondary
step was performed at 520°C.Two
dimensional intermetallic particle
analysis can be misleading, therefore
three dimensional analysis is necessary.
A particle extraction process was utilised
to characterize the three-dimensional
particle morphology and the relative VF
using x-ray diffraction (XRD). The
process was designed by the SINTEF
Group of Norway and has been used by
a few researchers to extract
intermetallic particles from aluminium
alloys3,4
. This process dissolves the Al-
matrix using dry butanol at 135°C, while
leaving the intermetallic particles
unaffected. The phases are
distinguished based on morphological
identification using Scanning Electron
Microscopy (SEM) and compositional
identification using Energy Dispersive
Spectroscopy (EDS). The Si content is
used to differentiate between the two
phases. Figure 1 shows that the darker
β-Al6(Fe,Mn) phase is more geometric in
shape, while the lighter α-
Al15(Fe,Mn)3Si2 phase comprises isolated
areas that contain the Al-matrix within
the particle centers Chinese-script like).
XRD patterns also show the presence of

26. 24
Microstructure, Texture and Formability of Metal Alloys
NOTES
β and α-phase particles within the
homogenised samples at the edge and
center. Image analysis and particle
extraction results indicate that
homogenisation at 560°C yields a higher
intermetallic particle VF near the edge
and a lower VF at the center.
Homogenisation at the higher
temperature (580°C) however, yields an
even distribution of intermetallic
particle VF at the edge and center.
Results show that homogenisation at
560°C yields a more ideal phase balance
between the β and α-phases, in so much
that there is approximately 50% of the
harder α-phase. The homogenisation
treatment at 560°C would therefore be
an appropriate temperature to improve
galling resistance., although the varied
particle distribution may affect other
parameters, such as particle stimulated
nucleation during subsequent hot and
cold rolling.
References
1. Marshall, G.J. (1996) Materials
Science Forum. 217-222, 19-30.
2. Kamat, R.G. (1996) JOM. 48(6), 34-38.
3. Simensen, J.C., Fartum, P. and
Anderson A. (1984) Fresenius Z anal
Chem. 319, 286-292.
4. Kumar, S. et al. (2014) Metallurgical
and Materials Transactions A. 45A 2842-
2854.
Figure 1. Secondary SEM image showing the difference in β-Al6(Fe,Mn) and α-
Al15(Fe,Mn)3Si2 particles at the ingot center, homogenised at 560/520°C.
Corresponding author: mgdliv002@myuct.ac.za

27. 25
Microstructure, Texture and Formability of Metal Alloys
NOTES
INFLUENCE OF MICROSTRUCTURAL FEATURES ON THE TENSILE
BEHAVIOR OF TI-6AL-4V ALLOY
S. Masete1,2,
, K. Mutombo1,2
, C. Siyasiya2
and W.E. Stumpf 2
1
MSM /Light Metals, Council for Scientific and Industrial Research, (CSIR),
Pretoria, South Africa
2
Department of Materials Science and Metallurgical Engineering, University of
Pretoria, South Africa
Aerospace industries use materials with
a combination of high strength to weight
ratio such as titanium alloys. The most
common alloy is Ti-6Al-4V which has
exceptional mechanical properties, low
density and good corrosion resistance.
This alloy is used for turbine engine parts
and structural components in aircraft.
Heat treatment is critical as it affects the
microstructure and mechanical
properties. Tensile property data of Ti-
6Al-4V Solution Treated (ST) above the
-transus and Water Quenched (WQ),
Air Cooled (AC) or Furnace Cooled (FC)
has been reported1,2,3
. The ST/WQ Ti-
6Al-4V has a higher Yield Strength (YS)
and Ultimate Tensile Strength (UTS)
than ST/AC or ST/FC Ti-6Al-4V. The
higher YS and UTS are due to the
presence of fine  (martensite) needles
in ST/WQ while ST/FC has coarser /
laths2
. Tensile elongation was lowest for
ST/WQ (high volume fraction of )
samples while highest after ST/FC (high
volume fraction of )2,3
. Ductility
improved with increasing  content as
slip occurs easily than with , hence
ST/FC gives relatively high tensile
elongation1
. The strain hardening
exponent and strain hardening rate was
reported for Ti-6Al-4V that was ST below
the -transus temperature (955C)5
. The
strain hardening rate was higher for the
ST/WQ sample than for ST/AC due to the
higher dislocation density5
. The strain
hardening exponent for ST/WQ sample
could not be determined due to a very
short plastic deformation regime5
.
Morita et al.6
reported data on Ti-6Al-4V
alloy ST below the -transus
temperature (930C), WQ and aged at
480C-680C. The strain hardening of
ST/WQ Ti-6Al-4V was high while
increasing ageing temperature reduced
strain hardening6
. Solution treating Ti-
6Al-4V below the -transus temperature
has been shown to produce lower
strength materials than ST above the -
transus temperature1
. With limited data
on the influence of microstructure on
the strain hardening exponent and
strain hardening rate for ST above the -
transus temperature, this work seeks to
investigate the optimization of the
microstructure and the corresponding
tensile behavior of the Ti-6Al-4V alloy.
A fully martensitic morphology resulted
from ST(1050C, 30mins)/WQ and
consisted of  and -phase on Grain
Boundaries (GB). A colony structure
produced by a ST(1050C, 30mins)/FC
consisted of colonies of / laths with
-phase on GB. The partial 

28. 26
Microstructure, Texture and Formability of Metal Alloys
NOTES
morphology, generated by ST/WQ
followed by ageing (900C, 30mins/FC),
consisted of a mixture of  and fine /
laths with GB partially transformed from
 to -phase. Tensile properties were
determined using Instron 1342/H1314
universal tester at 0.5mm/min
according to ASTM E8. The 0.2% YS,
UTS, percentage elongation (% EL) and
percentage reduction in area (% RA)
were determined as well as the strain
hardening exponent and strain
hardening rate. Samples were cut,
ground, polished and etched using
Kroll’s reagent. Microstructural analysis
was performed using a Leica DMI5000M
inverted Optical Microscope (OM) with
image analysis software (Image Pro Plus)
and a Jeol JSM-6510 Scanning Electron
Microscopy (SEM) equipped with X-ray
electron diffraction spectrometry (EDX).
The fully  samples had high YS and UTS
(Table. 1) due to fine  needles and the
lower ductility (% EL,% RA). The lamellar
samples contained coarser / laths and
had a lower YS and UTS but increased
ductility in comparison. The YS and UTS
of the partial  sample lies between the
fully  and lamellar sample. The strain
hardening exponent and strain
hardening rate of the fully 
microstructure were high. The ratio of
YS/UTS was used as an indirect measure
of strain hardening, where high ratio
indicates high strain hardening6
. The
current study shows ST/WQ having high
YS/UTS ratio while ST/FC had the lowest
ratio. The interactions of dislocations
with each other and with the fine 
needles would be higher in a ST/WQ
leading to dislocation pile up and
entanglement hence increasing strength
by strain hardening mechanism4
. Thus
showing that there is a strain hardening
increase with increasing volume fraction
of fine  in the microstructure.
References
1. H. Reham Reda, Adel Nofal, (2013) J.
Metall. Eng., vol. 2, no. 1, pp. 48–54
2. M. T. Jovanović, S. Tadić, S. Zec, Z.
Mišković, and I. Bobić, (2006) Mater.
Des., vol. 27, pp. 192–199.
3. P. Guo, Y. Zhao, W. Zeng, and Q.
Hong, (2013) Mater. Sci. Eng. A, vol.
563, pp. 106–111.
4. R. K. Gupta, C. Mathew, and P.
Ramkumar, (2015) Front. Aerosp.
Eng., vol. 4, no. 1, pp. 1–13.
5. B. D. Venkatesh, D. L. Chen, and S.
D. Bhole, (2009) Mater. Sci. Eng. A,
vol. 506, no. 1–2, pp. 117–124.
6. T. Morita, K. Hatsuoka, T. Iizuka, and
K. Kawasaki, (2005 Mater. Trans.,
vol. 46, no. 7, pp. 1681–1686.
Table 1. Average YS and UTS values at various heat treatments.
ST 1050°C, 30min,
WQ
ST 1050°C, 30min,
WQ, 900°C, 30min, FC
ST 1050°C, 30min, FC
YS (MPa) 905 820 810
UTS (MPa) 1040 925 875
Corresponding author: SMasete@csir.co.za

29. 27
Microstructure, Texture and Formability of Metal Alloys
NOTES
INTRODUCTION TO EBSD AND DATA PROCESSING FOR
MICROSTRUCTURAL AND TEXTURE ANALYSIS
K.G. Dicks
Oxford Instruments NanoAnalysis, Halifax Road, High Wycombe, Bucks HP12 3SE
Electron BackScattered Diffraction
(EBSD) is a highly established technique,
frequently complimenting Energy
Dispersive Spectroscopy (EDS). EDS
provides elemental information;
quantitative analyses and maps of
element distributions.
Whilst element associations can be
shown, the technique cannot
unequivocally identify compounds and
thus the ability to discriminate certain
phases may be compromised. Also the
interaction volume from which x-rays
are obtained is relatively large.
Therefore the spatial resolution may be
limited unless low accelerating voltages
are employed in the SEM.
Conversely, EBSD can identify
compounds or ‘phases’; i.e. regions that
may be either chemically or
crystallographically distinct, or both. It
does so by using electron diffraction to
collect and identify diffraction patterns,
rather than just elemental information
carried in the x-ray yield.
Electron diffraction patterns carry a
wealth of information, which can be
used to create a very diverse range of
plots, maps and output. These can be
compared with and compliment other
techniques such as X-Ray Diffraction
(XRD). EBSD is particularly convenient
and accessible, being performed on the
familiar platform of an SEM, which also
delivers high quality images for
reference with the material’s
microstructure. This is a very distinct
advantage over XRD for crystallographic
texture analysis as XRD has no reference
image, and no way of relating the
texture information with
microstructural features. Thus EBSD is a
highly visual as well as quantitative
technique, as a reference image of the
microstructure is provided.
EBSD also has the advantage of having
far better spatial resolution than EDS.
This is because the escape depth of
diffracted electrons is on a nanometre
scale, whereas the interaction depth of
EDS may be on a micron scale. EBSD can
also be performed on TEM samples
which further improves the spatial
resolution by ~10X.
These factors make EBSD a very
attractive proposition in its own right
and particularly in association with EDS.
EBSD is being increasingly adopted to
compliment EDS, XRD and TEM studies,
and is the ‘go to’ technique for
quantitative microstructure analysis.
EBSD is a standardised method for grain
size determination.
EBSD facilitates measures usually
acquired by x-ray diffraction and TEM
studies, including: Phase identification,
Phase or compound mapping,

30. 28
Microstructure, Texture and Formability of Metal Alloys
NOTES
crystallographic texture evaluation and
quantification, grain size and
morphology, grain boundary analysis,
strain and phase transformations. EDS
and EBSD information can be acquired
simultaneously and the resulting dataset
‘mined’ in post processing to reveal
measures of interest.
The presentation will give an overview
of the technique in combination with
EDS and then go on to show aspects of
post processing EBSD data, including
texture analysis.
Corresponding author: Keith.Dicks@oxinst.com

31. 29
Microstructure, Texture and Formability of Metal Alloys
NOTES
INVESTIGATION OF MICROSTRUCTURE AND PHASE FRACTION CHANGES IN
2304 DUPLEX STAINLESS STEEL (DSS); INFLUENCE OF HOT WORKING
CONSTANTS
O.M. Bill, C.W. Siyasiya and W.E. Stumpf
Department of Materials Science and Metallurgical Engineering, University of
Pretoria, 002 Pretoria, South Africa
Duplex steels are increasingly becoming
more attractive in many industrial
applications more than their single phase
steel counterparts; austenitic and ferritic
steels. Duplex steels constitute both the
austenitic and ferritic phases. However,
the presence of these two phases with
different deformation behaviours
complicates the processing of these steels,
often resulting in edge cracking and poor
surface finish 1,2
. Thus, there is a need to
understand the hot working behaviour of
duplex steels in order to address the
problems discussed above. Research has
also indicated that the ductility of duplex
steels is dependent on the deformation
conditions, behaviour of constituent
phases and volume fraction of austenite in
ferrite3,4,5
. However, past research has
neglected the fact that volume fraction of
austenite does not remain constant during
hot working; neither does the distribution
nor the orientation.
The current research aims at investigating
the evolution of microstructure and phase
fraction due to hot deformation conditions
in a 2304 duplex stainless steel through
isothermal hot compression tests.
The isothermal hot compression
deformation conditions were varied from
800 o
C - 1050 o
C with true strain levels of
0.2 - 0.6 and strain rates of 0.1 s-1
, 1 s-1
, 5 s-
1
and 10 s-1
. The microstructural evolution
was studied using optical microscopy and
Electron Backscatter Diffraction (EBSD).
The microstructural changes were rotation
and alignment, fragmentation and the
change in size as well as morphology and
distribution of the austenite phase as the
deformation conditions were varied within
the temperature and strain ranges above.
The volume fraction of austenite was
observed to increase from 30% to 49% at
850 o
C and 28% to 39% at 1050 o
C, as the
strain rate increased from 0.1 s-1
to 10 s-1
,
see Figure 1. The volume fraction of
austenite was also observed to increase
with decrease in temperature. No increase
in the volume fraction of austenite was
observed during the interpass times. EBSD
results also confirmed that
recrystallization was strongly inhibited
during deformation and prevailed during
interpass times. This can be a confirmation
that an increase in volume fraction of
austenite and recrystallization are two
energy competing processes.
These findings will be taken into account in
the development of a model for the
prediction of the microstructural evolution
of lean duplex stainless steels in the
finishing mill.

32. 30
Microstructure, Texture and Formability of Metal Alloys
NOTES
References
1. Iza-Mendia, A. et al. (1998) Metall Mater
Transactions A, 2975-2985.
2. Dehghan-Manshadi, A., Barnett, M.
and Hodgson, P. (2007) Mater Sci Tech,
23, 1478-1484.
3. Hoffmann, W.A.M. and Balancin, O.
(1998) Informacion Tecnologica, 9, 11-16.
4. Reis, G., Jorge, J.A. and Balancin, O.
(2000) Materials Research, 3, 31-35.
5. Momeni, A. and Dehghani, K. (2011)
Mater Sci Eng A, 528, 1448-1454.
Figure 1. Microstructure changes of 2304 DSS deformed at 850 o
C: (a) 0.1 s-1
(b) 10 s-1
Corresponding author: moshebill80@gmail.com

33. 31
Microstructure, Texture and Formability of Metal Alloys
NOTES
MICROSTRUCTURAL ANALYSIS OF A 12% Cr STEEL AT SUBGRAIN LEVEL
USING EBSD
G. Deyzel1,2
and J.E. Westraadt1,2
1
Physics Department, 2
Centre for HRTEM, Nelson Mandela Metropolitan
University, Port Elizabeth
Creep-strength-enhanced ferritic (CSEF)
steels are widely used in fossil fuel
plants1
. They have complex
microstructures and contain different
obstacles to dislocation motion. These
obstacles included the grain boundaries
(GB) and subgrain boundaries (SGB) that
are impenetrable by mobile
dislocations. The general microstructure
consists of prior-austenite grains that
are divided into packets and further into
martensite block and lath SGB, which
are elongated in shape. In addition, a
dislocation substructure of small-angle
GB is introduced during heat treatment
as a result of the rearrangement of
dislocations during annealing2
. The short
width of the elongated SGB plays an
important role in the creep
strengthening of the material, which is
quantified by a sub-boundary hardening
stress value2
. The SGB coarsen during
long-term creep3
resulting in a decrease
in the creep resistance. By obtaining an
Electron Back Scattered Diffraction
(EBSD) orientation map, one is able to
quantitatively investigate the different
SGB in an area.
The aim of this study was to determine
the SGB width using EBSD for a CSEF
steel in the as-received and creep-aged
conditions.
The material used for this study was an
X20CrMoV11-1 (12% Cr) stainless steel.
A weldment of new X20 steel welded
onto a creep damaged X20 steel was
analysed. The samples were polished
down to 0.25µm and etched with
alkaline colloidal silica solution.
Orientation maps were acquired using a
Nordlys HKL system fitted to a JEOL JSM
7001F SEM. A 15kV accelerating
voltage, 4nA probe current and
specimen tilt of 70° from the horizon,
was used for EBSD analysis. EBSD maps
were obtained with magnifications of
x250 and x500 in combination with a
0.2µm step size. For the subgrain
analysis, both the Grain Reconstruction
(GR) and Linear Intercept (LI) methods
from the HKL EBSD software was used. A
critical misorientation angle of 5º with
boundary completion down to 0° was
used to reconstruct the grains.
Figure 1 provides the Inverse Pole Figure
(IPF) map from the bulk region of the
new X20 steel weldment. An
enlargement of the rectangular section
highlighted in Figure 1 is given in Figure
2. The prior austenite GB and elongated
martensite block and lath subgrains are
clearly visible. Table 1 shows the
measured elongated SGB short widths
(λSGB) for the bulk region of the new and
damaged X20 steel weldments. It can be

34. 32
Microstructure, Texture and Formability of Metal Alloys
NOTES
concluded that the short width values
obtained from the LI method are closer
in accordance with previous results3
than that obtained from the GR method.
The results from the GR method are
dependent on the critical misorientation
angle used to reconstruct the grains.
Subgrains with a critical misorientation
less than 5º are grouped into one grain,
resulting in an over estimation of the
SGB size. Furthermore, assumptions
made regarding the aspect ratio and 2D
projection of the 3D subgrains will
influence the results. Therefore the LI
method seems to be more effective in
measuring the elongated SGB short
width. Further work will include the use
of Transmission Electron Microscopy
(TEM), Transmission Kikuchi Diffraction
(TKD) and High-Resolution (HR) EBSD to
provide the necessary spatial and
angular resolution needed to resolve the
dislocation substructure.
References
1. Bhadeshia, H.K.D.H. (2001) ISIJ Int. 41,
626-640.
2. Holzer, I. (2010) PhD. Dissertation,
Graz University of Technology, Austria.
3. Aghajani Bazazi, A. (2009). Dr.-Ing.
Dissertation, University of Bochum,
Germany
Table 1. Measured elongated SGB short widths (λSGB) for the new and creep-
damaged X20 steel.
Corresponding author: s210059257@nmmu.ac.za
Figure 1. EBSD IPF_Z0
orientation map for
bulk region of new X20
steel. The solid line
indicates a prior
austenite GB.
Figure 2. Enlarged rectangular
section extracted from Figure 1.
The solid line ellipse encloses a
single block and the dashed line
ellipse encloses a single
martensite lath within the block

35. 33
Microstructure, Texture and Formability of Metal Alloys
NOTES
MIPAR™: 2D AND 3D MICROSTRUCTURAL CHARACTERIZATION
SOFTWARE DESIGNED FOR MATERIALS SCIENTISTS, BY MATERIALS
SCIENTISTS
J.M. Sosa1
, D.E. Huber1
, B.A. Welk1
, H.L. Fraser1
1
Center for the Accelerated Maturation of Materials, Department of Materials
Science and Engineering, The Ohio State University, Columbus
Stereology, the science of estimating
three-dimensional quantities from two-
dimensionally acquired measurements,
has historically been the sole technique
for microstructural quantification1
. Over
the last decade and a half, 3D
characterization has begun to replace
stereology with direct-3D
quantification. As data acquisition
techniques continue to advance, the
need for more materials science-
orientated analytical 2D and 3D
software has become evident.
This led to the development of a
comprehensive software suite known as
MIPAR™ (Materials Image Processing
and Automated Reconstruction)2
.
MIPAR was written and developed
within MATLAB™, but is deployable as a
standalone cross-platform application.
MATLAB’s powerful 2D and 3D
processing libraries have greatly
contributed to and accelerated MIPAR’s
development. MIPAR is more an
application environment than a single
program. With a total of five
applications, it was designed to handle
all post-acquisition stages of 3D
characterization: alignment, pre-
processing, segmentation, visualization,
and quantification, as well as provide a
powerful platform for materials science-
oriented 2D image analysis. Images of
MIPAR’s three most commonly used
applications are shown in Fig. 1.
While direct-3D quantification offers
several advantages over stereology such
as the absence of sectioning variation
and superior quantification of complex
shapes, it is not without limitations.
With good reason, the representative
volume or area element (RVE or RAE)
has been an increasingly popular topic of
study and discussion3
. However, a
definition of the target quantification
precision is often left out of RVE/RAE-
related discussions. Therefore, this
paper will present the use of MIPAR,
together with statistical tools such as
random sampling and bootstrapping, to
establish quantitative relationships
between sampled volume/area size and
measurement precision for a variety of
microstructural metrics. In addition to
exploring the influence of sectioning
variation on various stereological
metrics, direct-3D quantification can
either validate or invalidate
stereological assumptions. A common
stereological metric is the mean linear
intercept. Measured from a series of
random lines placed within a segmented
microstructure, the mean linear
intercept has been employed to
estimate three-dimensional quantities
such as the mean diameter of spheroidal

36. 34
Microstructure, Texture and Formability of Metal Alloys
NOTES
precipitates and mean width of plate-
like features4
. In both cases, the
constitutive equations rely of several
assumptions regarding the shape and
size distribution of the intercepted
features. For features in α+β titanium
microstructures, MIPAR has been used
to explore the validity of these
assumptions and determine the
sensitivity of stereological quantification
to deviations from such assumptions.
The efficacy of these characterization
efforts was critically dependent on
segmentation. Therefore, a strong focus
has been placed on developing a
method of objectively quantifying
segmentation quality. This method,
reliant on the similarity metric of mutual
information, has been integrated into
MIPAR’s Image Processor and examples
of its application will be presented.
References
1. Russ, J.C. and DeHoff, R.T. (2000)
Practical Stereology. Denmark, Kluwer
Academic Pub.
2. Sosa, et al.(2014) Integr. Mater.
Manuf. Innov. 3(10), 18.
3. Shan, Z. and Gokhale, A.M. (2002),
Computational Materials Science. 24(3),
361–379.
4. Collins, et al. (2009) Mater. Sci. Eng. A.
508(1-2), 174–182.
Corresponding author: sosa.12@buckeyemail.osu.edu
Figure 1. Images of three applications used for 3D characterization in MIPAR™ where
(a) reveals the Image Processor, (b) the Batch Processor, and (c) the 3D Toolbox

37. 35
Microstructure, Texture and Formability of Metal Alloys
NOTES
TEXTURE CAPABILITIES AT NECSA’S DIFFRACTION FACILITIES
A.M. Venter1
, D. Marais1
, T.P. Ntsoane1
1
Research and Development Division, Necsa Limited SOC, Pretoria, South Africa
Radiation diffracted from
crystallographic lattice planes gives
direct information of the crystallite
orientations (texture)1
. Of these, X-ray
and neutron diffraction are well-
established methods enabling
determination of the macrotexture, as
opposed to individual orientations
measured with electron back-scatter
diffraction techniques to give the
microtexture.
The materials science diffractometer
MPISI at the SAFARI-1 research reactor
of Necsa has been designed to meet the
demands of non-destructive residual
stress and texture determination by
neutron diffraction methods. This
instrument is complemented with an X-
ray diffraction equivalent instrument,
Bruker D8 Discover. Both instruments
are equipped with the required sample
manipulation stages, such as Eulerian
cradles, area detectors and advanced
data processing software such as
MulTex2
and popLA3
to enable multi-
reflection pole figure compilation and
further processing towards the
determination of orientation
distribution functions. Data are
outputted in a three column format, i.e.
I, φ, χ.
In principle, pole figure measurement
methodology from neutron and X-ray
scattering data are very similar. X-ray
photons scatter through an
electromagnetic interaction with the
electron charge cloud of the material,
whilst neutrons are scattered by a
nuclear interaction with the nuclei. In
general, the interaction strength of X-
rays with matter is directly related to the
atomic number of the materials,
whereas neutron scattering lengths are
approximately equal in magnitude for
most atoms. Apart from the interaction
strength differences, penetration
depths are also atomic species
dependent, as summarised in Table 1.
Table 1: Neutron and X-ray scattering.
Half attenuation lengths correspond to
the thickness of material that reduces
the intensity by 50%.
Element
Neutron
scattering
length(bc)
[fm=10-15m]
X-ray
scattering
length,
[fm=10-15m]
Half attenuation
lengths
1.8Å
neutron
s
[mm]
1.5Å
X-rays
[μmm]
Mg 5.375 0.246 43 100
Al 3.449 0.258 66 52
Ti -3.37 0.451 12 11
Fe 9.54 0.565 6 15
Co 2.49 0.594 2 14
Ni 10.3 0.624 3 16
Zr 7.16 0.884 24 8
Hf 7.77 1.715 1 3
W 4.86 1.762 5 2
U 8.417 2.172 9 1
X-rays effectively give the near surface
texture (a few tens of micrometres),
whereas thermal neutrons can
penetrate at least 1000 times deeper to
give the global texture over the full

38. 36
Microstructure, Texture and Formability of Metal Alloys
NOTES
sample volume illuminated. Specifically,
neutron techniques provide information
on the bulk texture in the volume of
material illuminated. Complementary
information is thus possible. Depth-
resolved texture is achieved by
extracting representative samples from
thick-walled components such as plates
and pipes by slitting at each depth. In
addition, due to their phase sensitivity,
texture of constituent phases in multi-
phase materials can be measured. This
makes the techniques applicable to
fields of study that include metals
forming and processing, geological and
cultural heritage applications. Fig. 1
shows bulk pole figures measured with
neutron diffraction of ferritic stainless
steel specimens subjected to different
heat-treatment parameters after rolling.
Significant texture development is
evident after different heat treatment
conditions, i.e Sample 1 vs Sample 1 pole
figures. This variation in texture could
lead to anisotropic properties and
physical phenomena such as earing,
warping and general dimensional
changes and distortions.
This work addresses the
complementarity of neutrons and X-
rays. The diffraction techniques could
add significant value to issues of
microstructure characterisation and
texture measurement.
References
1. Olaf Engler and Valerie Randle (2010)
Taylor and Francis Group
2.https://www.bruker.com/products/x-
ray-diffraction-and-elemental-
analysis/x-ray-diffraction/xrd-
software/overview/multex.html
3. http://www.ccp14.ac.uk/ccp/web-
mirrors/popla/orgs/mst/cms/poplalapp
.html
Figure 1: Pole figure representations measured with neutron diffraction of selected
reflections from two ferritic stainless steel samples subjected to different heat
treatments
Corresponding author: Andrew.venter@necsa.co.za
Sample 1 Sample 2
(110) (200) (211) (110) (200) (211)
Min = 0.04
Max = 3.18
Norm = 1065
Min = 0.02
Max = 2.76
Norm = 641
Min = 0.42
Max = 2.01
Norm = 870
Min = 0.26
Max = 2.07
Norm = 518
Min = 0.24
Max = 2.12
Norm = 300
Min = 0.68
Max = 1.47
Norm = 389
Min: 0.04
Max: 3.18
0.43
0.82
1.22
1.61
2.00
2.39
2.79
Min: 0.02
Max: 2.76
0.37
0.71
1.05
1.39
1.74
2.08
2.42
Min: 0.42
Max: 2.01
0.62
0.81
1.01
1.21
1.41
1.61
1.81

39. 37
Microstructure, Texture and Formability of Metal Alloys
NOTES
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Microstructure, Texture and Formability of Metal Alloys
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