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The E ff ect of Phase Morphology and
Volume Fraction of Retained Austenite on
the Formability of Transformation
Induced Plasticity Steels
by
Jeff rey Benjamin Rutter Lawrence
A thesis submitted to theDepartment of Mechanical and Materials Engineering
in conformity with the requirements for
the degree of Masters of Applied Science
Queens University
Kingston, Ontario, Canada
January 2010
Copyright Jeff rey Benjamin Rutter Lawrence, 2010
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Abstract
Transformation induced plasticity (TRIP) steels are a class of steels with exceptiona
formability properties, due mainly to the presence of meta-stable retained austenite
which transforms to martensite under loading, locally hardening the steel. The vol
ume fraction and mechanical stability of the retained austenite play an important role
in producing the high formabilities of TRIP steels. In this thesis, two separate mor
phologies of retained austenite, equiaxed versus lamellar, have been produced throug
thermo-mechanical processing of a single common TRIP steel chemistry. The she
formability characteristics of these two microstructures were examined, with varyin
volume fractions of retained austenite, through uniaxial tensile and in-plane plane
strain (IPPS) testing.
It was found that higher levels of retained austenite produced better formabilityproperties for both microstructures and strain paths. In uniaxial tension it was seen
that the the lamellar microstructure attained higher strains at maximum load, and
exhibited more sustained instantaneous n values than the equiaxed structure, despite
having a lower volume fraction of retained austenite.
IPPS testing was performed using an optical measurement of local strain and a
comparative forming limit based on diff
erences in strain rate between a developingneck and the surrounding material. It was found that the lamellar microstructure
performed better than the equiaxed microstructure for this strain path, achieving
higher strains before reaching the comparative forming limit.
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For Rachel, without whose love and support this thesis would not have been possible.
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Acknowledgements
I would like to acknowledge:
My supervisors Dr. Keith Pilkey and Dr. Doug Boyd, for their great help, sup
port, expertise, and friendship which was invaluable in all aspects of completin
this thesis.
Dr. Thomas Krause of RMC for the very generous use of equipment as well a
his time and knowledge of magnetic testing methods, along with Brian Judd
for his technical help and support of the same.
Jasmine Chiang whose exemplary work as a summer student greatly helped in
the completion of this thesis.
Dr. Alison Mark whose Ph.D work this thesis extends, for her availability in
answering questions about all things TRIP.
The Mechanical and Materials Engineering Department technicians at Queens
uUiversity, namely Charlie Cooney, Chris Gabryel, Joyce Cooley and the staff
of the Mechanical Engineering machine shop for their technical skill and patien
explanations.
CANMET for manufacturing and supplying the steel used in this thesis.
The nancial support of Queens University, the STELCO graduate fellowshi
and AUTO21 Canadian automotive research group.
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Contents
Contents i
List of Figures v
List of Tables ix
List of Nomenclature xi
Chapter 1: Introduction 1
1.1 Motivation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1
1.2 Research Objective . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2
Chapter 2: Literature Review 4
2.1 TRIP Steel Microstructure . . . . . . . . . . . . . . . . . . . . . . . . 4
2.1.1 Chemistry . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5
2.1.2 Processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7
2.1.3 The Bainite Transformation . . . . . . . . . . . . . . . . . . . 10
2.1.4 Retained Austenite Stability . . . . . . . . . . . . . . . . . . . 17
2.2 TRIP Steel Properties . . . . . . . . . . . . . . . . . . . . . . . . . . 222.2.1 Tensile Properties . . . . . . . . . . . . . . . . . . . . . . . . . 22
2.2.2 Hardening Rate . . . . . . . . . . . . . . . . . . . . . . . . . . 23
2.2.3 Sheet Formability Testing . . . . . . . . . . . . . . . . . . . . 24
2.2.4 IPPS Testing . . . . . . . . . . . . . . . . . . . . . . . . . . . 27
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Contents v
2.2.5 TRIP Steels and Forming Path . . . . . . . . . . . . . . . . . 28
2.3 TRIP Steel Characterization . . . . . . . . . . . . . . . . . . . . . . . 30
2.3.1 X-Ray Diff raction . . . . . . . . . . . . . . . . . . . . . . . . . 30
2.3.2 Magnetic Saturation . . . . . . . . . . . . . . . . . . . . . . . 322.4 Literature Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . 34
Chapter 3: Experimental Methods 36
3.1 Heat Treatments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
3.1.1 The As-Received Condition . . . . . . . . . . . . . . . . . . . 36
3.1.2 TRIP Steel Heat Treatments . . . . . . . . . . . . . . . . . . . 37
3.2 Magnetic Saturation Measurements . . . . . . . . . . . . . . . . . . . 433.2.1 Magnetic Theory . . . . . . . . . . . . . . . . . . . . . . . . . 43
3.2.2 Magnetic Sample Preparation . . . . . . . . . . . . . . . . . . 43
3.2.3 Magnetic Testing Conditions . . . . . . . . . . . . . . . . . . . 44
3.3 X-Ray Diff raction Measurements . . . . . . . . . . . . . . . . . . . . 52
3.3.1 X-Ray Diff raction Theory . . . . . . . . . . . . . . . . . . . . 52
3.3.2 X-Ray Diff raction Sample Preparation . . . . . . . . . . . . . 56
3.4 Optical Microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56
3.4.1 Optical Sample Preparation . . . . . . . . . . . . . . . . . . . 57
3.5 Uniaxial Tensile Testing . . . . . . . . . . . . . . . . . . . . . . . . . 58
3.5.1 Sample Preparation and Dimensions . . . . . . . . . . . . . . 58
3.5.2 Testing Equipment and Conditions . . . . . . . . . . . . . . . 59
3.5.3 Data Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . 60
3.6 In-Plane Plane-Strain Testing . . . . . . . . . . . . . . . . . . . . . . 60
3.6.1 Sample Preparation and Geometry . . . . . . . . . . . . . . . 63
3.6.2 Testing Equipment and Conditions . . . . . . . . . . . . . . . 64
3.6.3 Data Analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . 65
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Contents vi
Chapter 4: Results 67
4.1 Characterization and Validation of the Magnetic Determination of Re-
tained Austenite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67
4.1.1 Vibrating Sample Magnetometer Setup and Results . . . . . . 674.1.2 X-Ray Diff raction Results . . . . . . . . . . . . . . . . . . . . 67
4.1.3 Optical Analysis . . . . . . . . . . . . . . . . . . . . . . . . . 69
4.2 Heat Treatments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70
4.2.1 Equiaxed and Lamellar Microstructures . . . . . . . . . . . . . 70
4.2.2 Investigation of the Bainite Hold . . . . . . . . . . . . . . . . 72
4.2.3 Equiaxed Tensile Samples . . . . . . . . . . . . . . . . . . . . 75
4.2.4 Equiaxed IPPS . . . . . . . . . . . . . . . . . . . . . . . . . . 77
4.2.5 Lamellar Tensile . . . . . . . . . . . . . . . . . . . . . . . . . 77
4.2.6 Lamellar IPPS . . . . . . . . . . . . . . . . . . . . . . . . . . 82
4.3 Tensile Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83
4.3.1 Equiaxed Stress Strain Analysis . . . . . . . . . . . . . . . . . 83
4.3.2 Lamellar Stress Strain Analysis . . . . . . . . . . . . . . . . . 87
4.3.3 Instantaneous n Values . . . . . . . . . . . . . . . . . . . . . . 904.4 In-Plane Plane-Strain Testing . . . . . . . . . . . . . . . . . . . . . . 93
4.4.1 Strain in Cells . . . . . . . . . . . . . . . . . . . . . . . . . . . 94
4.4.2 Comparative Forming Limits . . . . . . . . . . . . . . . . . . . 95
Chapter 5: Discussion 99
5.1 Characterization of TRIP steels . . . . . . . . . . . . . . . . . . . . . 99
5.1.1 Use of Magnetic Saturation for the Measurement of Volume
Percent of Retained Austenite . . . . . . . . . . . . . . . . . . 99
5.1.2 X-Ray Diff raction Measurements of Volume Percent of Retained
Austenite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101
5.1.3 Microscopy and the Determination of TRIP Steel Microstructures103
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Contents vii
5.2 Microstructural Evolution . . . . . . . . . . . . . . . . . . . . . . . . 104
5.2.1 Initial Microstructures . . . . . . . . . . . . . . . . . . . . . . 104
5.2.2 The Inter-critical Hold . . . . . . . . . . . . . . . . . . . . . . 105
5.2.3 The Bainite Hold Curve . . . . . . . . . . . . . . . . . . . . . 1065.2.4 Temperature and Heating Rate Dependance of Microstructure 108
5.3 Role of Retained Austenite in Formability . . . . . . . . . . . . . . . 109
5.3.1 Retained Austenite in Uniaxial Tension . . . . . . . . . . . . . 109
5.3.2 Eff ect of Retained Austenite on Material Hardening . . . . . . 109
5.3.3 Retained Austenite in IPPS . . . . . . . . . . . . . . . . . . . 110
5.4 Aff ect of Phase Morphology on Performance of TRIP Steel . . . . . . 111
5.4.1 Diff erences in Morphology . . . . . . . . . . . . . . . . . . . . 111
5.4.2 Diff erences in Mechanical Properties . . . . . . . . . . . . . . 112
5.4.3 Diff erences in Hardening Behavior . . . . . . . . . . . . . . . . 113
5.4.4 Diff erences in IPPS . . . . . . . . . . . . . . . . . . . . . . . . 113
Chapter 6: Conclusions and Recommendations 115
6.1 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 115
6.1.1 Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . 115
6.1.2 Mechanical Properties . . . . . . . . . . . . . . . . . . . . . . 116
6.1.3 Characterization . . . . . . . . . . . . . . . . . . . . . . . . . 117
6.2 Recommendations . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117
6.2.1 Microstructure . . . . . . . . . . . . . . . . . . . . . . . . . . 118
6.2.2 Mechanical Properties . . . . . . . . . . . . . . . . . . . . . . 118
6.2.3 Characterization . . . . . . . . . . . . . . . . . . . . . . . . . 118
References 120
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List of Figures
1.1 Ranges of UTS and total elongation for high-strength steel (HSS) and
advanced high-strength steels (AHSS), showing the benets of TRIP
steels over HSLA and DP steels . . . . . . . . . . . . . . . . . . . . . 3
2.1 The aff ect of alloying elements on the thermal processing of TRIP steels 6
2.2 Schematic representation of the thermo-mechanical treatments applied
to TRIP steels, from hot- or cold-rolled material to the two-stage heat-
treatment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8
2.3 Typical SEM micrograph of equiaxed TRIP structure, phase A is
austenite, B is bainite and F is inter-critical ferrite . . . . . . . . . . . 9
2.4 Sketch of the evolution of the martensite-based TRIP structure froma) martensite laths to b) partially recovered laths + carbon moving
to boundaries, c) recovered ferrite laths + austenite and d) recovered
ferrite + austenite with bainite . . . . . . . . . . . . . . . . . . . . . 11
2.5 SEM of martensite-based TRIP steel, light phase is austenite/martensite,
dark phase is ferrite/bainite . . . . . . . . . . . . . . . . . . . . . . . 12
2.6 The growth of a bainite sheaf from primary nucleation at a grain
boundary (top) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
2.7 Representative T 0 and Ae3 lines for a given steel . . . . . . . . . . . . 16
2.8 Theoretical relation between the grain size and thermal stability of
retained austenite grains . . . . . . . . . . . . . . . . . . . . . . . . . 19
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List of Figures ix
2.9 Volume fraction of retained austenite measured during tensile testing
for four austenite morphologies . . . . . . . . . . . . . . . . . . . . . 21
2.10 Representative stress-strain curves of DP, TRIP and HSLA steels . . 23
2.11 Comparison of instantaneous n values calculated for DP, TRIP andHSLA steels . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25
2.12 FLCs of HSLA, mild and DP steels, showing strain paths from uniaxial
to equi-biaxial tension . . . . . . . . . . . . . . . . . . . . . . . . . . 26
2.13 Volume fraction of retained austenite with increasing major strain for
TRIP steel tested in several strain paths . . . . . . . . . . . . . . . . 29
2.14 Magnetic hysteresis curve for an austenitic stainless steel after under-
going 55% reduction in thickness to produce 75 vol.% martensite . . . 33
3.1 Schematic graph of the two-stage heat treatment used to produce the
equiaxed microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . 39
3.2 Coupon dimensions for tensile sample heat treatments showing position
of tensile bar and magnetic testing discs, all dimensions in mm. . . . 40
3.3 Coupon dimensions for IPPS sample heat treatments showing position
of IPPS sample and magnetic testing discs, all dimensions in mm. . . 41
3.4 Schematic graph of the three-stage heat treatment used to produce the
lamellar microstructure. . . . . . . . . . . . . . . . . . . . . . . . . . 42
3.5 VSM showing the diff erent parts of the machine, the sample is vibrated
vertically through the magnetic eld. . . . . . . . . . . . . . . . . . . 45
3.6 Magnetic hysteresis loops for pure ferrite (calculated) and sample TE100-
1 with 13.1 vol.% retained austenite. . . . . . . . . . . . . . . . . . . 473.7 Second derivative of the magnetic hysteresis loop of the equilibrium-
cooled sample showing peak values at the inection point of the rst
derivative. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49
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List of Figures x
3.8 Curve-tting results for sample TE100-1 showing the negative ap-
proach to saturation. . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.9 Full XRD prole for sample AH60 with 6.4 vol.% R.A. . . . . . . . . 53
3.10 Full XRD prole for sample AM200 with 11.2 vol.% R.A. . . . . . . . 543.11 Deconvolution of the double copper ferrite (200) peak for sample AL200.
The original data is shown with the two Pearson VII curves, back-
ground radiation level, the total curve t and the residuals (dotted) . 55
3.12 Sub-sized tensile sample geometry, sheet thickness of 1 mm, all dimen-
sions in mm. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59
3.13 Position of the grid dots on the surface of the IPPS samples, dots were
centered in the major and minor directions . . . . . . . . . . . . . . . 61
3.14 Experimental setup for IPPS testing, showing digital camera, Instron
with wide grips and digital clock . . . . . . . . . . . . . . . . . . . . . 62
3.15 Post-processed image of deformed grids showing necked and un-necked
rows . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63
3.16 Geometry of IPPS samples, all dimensions in mm, sheet thickness is
1 mm . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64
4.1 a) Longitudinal and b) transverse section of sample AM200. Etched
with picric/sodium metabisulte, martensite/austenite is light, fer-
rite/bainite dark. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71
4.2 Micrographs of sample TL300-2 in a) longitudinal and b) transverse di-
rections. Etched with picric/sodium metabisulte, martensite/austenite
is light, ferrite/bainite is dark. . . . . . . . . . . . . . . . . . . . . . . 724.3 The volume percent of RA produced at bainite hold temperatures of
350, 400 and 450 C, as a function of bainite hold time . . . . . . . . 74
4.4 Volume percent of RA as a function of bainite hold time for TE series
heat treatments, employing a bainite hold temperature of 450 C . . . 76
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List of Figures xi
4.5 Volume percent of RA as a function of bainite hold time for IE series
heat treatments, employing a bainite hold temperature of 450 C . . . 79
4.6 Volume percent of RA as a factor of bainite hold time for TL series
heat treatments, employing a bainite hold temperature of 450
C . . . 814.7 Typical engineering stress-strain curves for TE series samples . . . . . 84
4.8 Strain at UTS and volume percent RA as a function of bainite hold
time for TE series heat treatments, employing a bainite hold temp of
450 C . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 86
4.9 Typical engineering stress-strain curves for TL series heat treatments 88
4.10 Strain at UTS and volume percent RA for TL series heat treatments 89
4.11 Instantaneous n as it changes with strain for three TE series heat
treatments. TE30-1 - Green, TE100-3 - Blue, TE1800-1 - Red . . . . 91
4.12 Instantaneous n as it changes with strain for three TL series heat treat-
ments. TL60-1 - Green, TL100-2 - Blue, TL1800-1 - Red . . . . . . . 92
4.13 Strain proles for the a) beginning, b) middle and c) end of the IPPS
testing of sample IE100-3 . . . . . . . . . . . . . . . . . . . . . . . . . 94
4.14 Curves of the diff erence in strain rate for sample IE100-3 betweennecked and un-necked rows at ve locations across the sample width . 96
4.15 Strain at relative necking limit for IE and IL series samples plotted
against the volume percent of RA in each sample . . . . . . . . . . . 98
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List of Tables
3.1 Chemistry of the TRIP steel, as received in weight percent, balance is
iron . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37
3.2 The measured ferrite and austenite peaks with their corresponding
expected diff raction angles . . . . . . . . . . . . . . . . . . . . . . . . 53
4.1 Volume percent retained austenite as measured by VSM and XRD . . 68
4.2 A-series heat treatment conditions and volume percent retained austen-
ite, as measured by magnetic saturation . . . . . . . . . . . . . . . . 74
4.3 Bainite hold times and measured volume percent retained austenite for
equiaxed tensile samples. . . . . . . . . . . . . . . . . . . . . . . . . . 75
4.4 Bainite hold times and average volume percent of retained austenitemeasured magnetically for equiaxed tensile samples. The A and B
samples were taken from either side of the IPPS sample as shown in
Figure 3.3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78
4.5 Bainite hold times and volume percent of retained austenite measured
magnetically for lamellar tensile samples. . . . . . . . . . . . . . . . . 80
4.6 Lamellar IPPS samples with volume percent of retained austenite mea-
sured magnetically on either side of the IPPS sample as shown in Fig-
ure 3.3, bainite hold time is 100 seconds. . . . . . . . . . . . . . . . . 83
4.7 Tensile properties of TE samples . . . . . . . . . . . . . . . . . . . . 85
4.8 Tensile properties of lamellar tensile samples . . . . . . . . . . . . . . 88
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List of Tables xiii
4.9 Retained austenite and strain at the local necking condition for equiaxed
and lamellar IPPS samples . . . . . . . . . . . . . . . . . . . . . . . . 97
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List of Nomenclature
Ae3 Equilibrium carbon concentration in austenite
Ao Original cross-sectional area
d Atomic spacing
e Engineering strain
F app Applied force
H Applied magnetic eld
I hklx Integrated intensity of the hkl plane of phase x
k Hardening constant
l Gauge lengthlo Original gauge length
l Change in gauge length
M Magnetic moment
M s Magnetic saturation
n Strain hardening exponent
n i Instantaneous n
Rhklx Theoretical intensity ratio for the hkl plane of phase x
S Engineering stress
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List of Nomenclature xv
T 0 Thermodynamic maximum carbon concentration in austenite for the
bainite transformation
True stress
True strain
Diff raction angle
Wavelength
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Chapter 1
Introduction
1.1 Motivation
Increasing oil prices and environmental consciousness in consumers, along with n
environmental legislation have made fuel effi
ciency a top priority of the automotiveindustry. A straightforward means to increasing a vehicles fuel efficiency is to reduce
vehicle weight, which often means using new materials to produce lighter vehic
parts without sacricing strength or performance. A reduction of 100 kg in th
average weight of a gasoline passenger car has been found to lead to a decrease in f
consumption of 0.35 liters per 100 km or 700 liters over the vehicles life cycle [1]. The
automotive industry has been looking towards plastics, composites and light-weigh
alloys to replace traditional materials, while at the same time using small amounts o
advanced steels to replace larger pieces made of traditional steels. Automotive bod
parts are commonly stamped out of sheet steels, forming complex structural member
from thin sheet. If the strength of the steel were increased, thinner sheet could b
used to provide the same strength at a lower weight. It has been estimated that by
replacing conventional steel with advanced high-strength steels the overall weight o
a standard family car could be reduced by 118 kg [2]. This could result in even largertotal weight savings in the vehicle as the size of many components is dependent o
the vehicle weight, so that reducing the overall weight of the vehicle allows addition
weight savings to be made elsewhere. Assuming one liter of gasoline contains 2
kg of CO2 [3], the savings in emissions over the life cycle of a car could be as hig
1
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Introduction 2
as 1916 kg of CO2 . Conventional high-strength steels do not possess the formability
required to create many of these complex sheet parts using a reduced sheet thicknes
Therefore, the challenge is to design a steel which exhibits both high strength fo
weight reduction and good formability for forming complex structural componentsMulti-phase TRansformation Induced Plasticity aided steels, or TRIP steels,
are a class of advanced high-strength steels that potentially have the high strength
and good formability needed for signicant weight reductions in automotive applic
tions. These steels possess a multi-phase microstructure which includes a meta-stab
retained austenite phase which is capable of producing the TRIP eff ect, along with
inter-critical ferrite and bainitic ferrite. When the steel is deformed the austenite
transforms into martensite, which helps to absorb energy and locally harden the
material. This hardening eff ect, combined with the volume increase of the marten-
site transformation, acts to resist necking in the material and postpone failure in
sheet forming operations. TRIP steels also incorporate the advantages of the com
posite structure of a dual-phase steel, containing strong austenite/martensite phases
amongst the weaker ferrite and bainite.
The macroscopic benets of the TRIP eff ect can be demonstrated by comparingdiff erent steels in terms of their ultimate tensile strength and total elongation. A
representative plot is shown in Figure 1.1 where the TRIP steels exhibit a much
higher total elongation (which is loosly related to formability) than comparable high
strength low-alloy (HSLA) or dual-phase (DP) steels.
1.2 Research Objective
The microstructure of a TRIP steel can greatly alter its mechanical properties. The
most important aspects of the TRIP steel microstructure are the volume percent, size
and morphology of the retained austenite phase, as these properties directly aff ect
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Introduction 3
Figure 1.1Ranges of UTS and total elongation for high-strength steel (HSS) andadvanced high-strength steels (AHSS), showing the benets of TRIPsteels over HSLA and DP steels [4]
the austenite-to-martensite transformation when the steel is deformed [5, 6]. Previous
research has investigated a variety of diff erent phase morphologies and the resulting
mechanical properties in TRIP steels. In this thesis, the eff ect of austenite phase
morphology on TRIP steel formability is examined by producing two distinct microstructures of retained austenite with a common steel chemistry. To accomplish thi
objective two groups of steel samples were heat-treated to obtain similar surroundin
matrices of ferrite and bainite, but diff ering morphologies of retained austenite. It
was found that the two steel microstructures had very diff erent properties, both under
uniaxial tensile and in-plane plane-strain testing, and that the mechanical properties
of both steels relied heavily on the volume fraction and morphology of the retaineaustenite phase.
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Chapter 2
Literature Review
TRIP steels present a complex class of materials, where the mechanical propertie
are aff ected by an evolving, multi-phase microstructure. In order to understand the
microstructure and its evolution through processing and service, a range of character
ization techniques are needed to analyze these steels. By considering the contributio
of each microstructural component to the overall mechanical behavior of the materia
it is possible to design TRIP steels with a microstructure optimized for the mechan
ical properties needed to solve real-world problems. In this study, the evolution an
morphology of two separate TRIP microstructures are analyzed, and their mechanica
properties are evaluated for their use in sheet forming operations. Hence, there ar
three main areas of interest: the microstructure, the mechanical properties and the
characterization of the TRIP steels.
2.1 TRIP Steel Microstructure
The microstructure of TRIP steels revolves around the presence of meta-stable re
tained austenite which produces the TRIP eff ect. There are many compositions and
microstructures which allow the austenite to be stabilized at room temperature. In
this study, low-alloy TRIP steels with austenite stabilized by inter-critical and bainite
holds are considered.
4
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2.1.1 Chemistry
There are several ways in which to chemically stabilize austenite at room temperature
The simplest approach is through the addition of nickel as in austenitic stainless steel
The mechanical advantages of the TRIP eff ect were in fact rst observed in nickel-
bearing meta-stable austenitic steels in 1967 by Merz et al. [7]. Transformation due
to plasticity in austenitic steels is still a major area of interest [8], but the high nickel
contents of these steels make them very expensive and so the stabilization of austenit
using low levels of alloying elements was developed out of research involving sim
dual-phase steels. In low-alloy TRIP steels the austenite is stabilized through it
carbon content. The carbon content necessary to stabilize austenite is approximately1 wt.% [9]. However, carbon content this high prevents the steel from being easily
welded, so the overall carbon content of TRIP steels is limited to 0.2 wt.% [10] so that
it can be widely used in industry. A complex heat-treatment is therefore employe
to concentrate the carbon within the austenite.
A common TRIP steel chemistry also contains small additions of other element
to both help in stabilizing the austenite and to aid in the creation of microstructures
which partition carbon into the austenite. The most common additions in TRIP steels
are 1.5 wt.% of both silicon and manganese. The manganese directly stabilizes th
austenite by lowering the martensite start temperature; it will also lower the start-
temperature for the formation of cementite, preventing pearlite from forming durin
cooling [10, 11]. The silicon content of the steel is also important as silicon is highly
insoluble in cementite. The silicon will therefore act to greatly delay the formation
carbides, especially during the bainite transformation, as time must be given for thsilicon to diff use away from the bainite grain boundaries before cementite can form
[10, 12]. The eff ects of each of these elements, as they apply to TRIP steel processing
are shown in Figure 2.1.
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Figure 2.1The aff ect of alloying elements on the thermal processing of TRIPsteels [10]
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The use of silicon in TRIP steels does have one major drawback. It cause
serious problems when applying a zinc coating to the surface, as is the case for
standard galvanneal process used in the production of automotive sheet steel [13].
This problem has lead researchers to experiment with other alloying elements in placof the silicon, such as aluminum and phosphorus [14]. Additional alloying elements,
such as vanadium, may also be included to precipitation strengthen the ferrite matrix
[15].
2.1.2 Processing
Specic processing routes are required to concentrate carbon into the austenite inorder to stabilize it at room temperature. The most common method for accom
plishing this task in a TRIP steel is to start with a hot-rolled steel, followed by col
rolling in order to deform the microstructure and impart the potential energy needed
for efficient recrystallization. The steel is then re-heated to the inter-critical region
where the steel recrystallizes, growing small grains of austenite and ferrite. In th
inter-critical region the austenite can be enriched in carbon up to the eutectoid chem
istry (approximately 0.8 wt.% carbon) by holding it at the bottom of the temperature
range. Depending on the amount of prior deformation and the holding temperature
this recrystallization to ferrite and austenite can be very fast. Long hold times are
avoided in this step to prevent detrimental grain growth. A schematic of the common
two-stage TRIP heat treatment is shown in Figure 2.2.
After inter-critical annealing the steel is then cooled to the bainite region wher
it is held for a short length of time. This bainite hold is necessary to further enricthe carbon content of the austenite. The steel is then quenched to room tempera-
ture to halt the bainite reaction. The bainite grows into the inter-critical austenite,
partitioning carbon to the austenite according to the transformation described in Sec
tion 2.1.3. Usually the bainite reaction is accompanied by the formation of carbides
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around the bainitic ferrite, but the addition of silicon prevents this occurrence, and
so the carbon is partitioned directly into the austenite [16]. The length of the bainite
hold time must be carefully chosen as holds which are too short will not allow enou
bainite to grow and the austenite will not be stabilized by enough carbon; hence, thaustenite will transform to martensite upon quenching to room temperature. If the
bainite hold time is too long the high-carbon austenite will decompose into carbide
reducing the volume percent of austenite retained at room temperature [16].
Figure 2.2Schematic representation of the thermo-mechanical treatments applied to
TRIP steels, from hot- or cold-rolled material to the two-stageheat-treatment [12]
The heat treatment described above will produce a structure consisting of ferrite
bainite and austenite, with some martensite or cementite depending on the bainite
hold conditions. The ferrite will appear as equiaxed grains that were formed durin
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the inter-critical hold. The inter-critical austenite grains will be partly consumed by
bainitic ferrite which has grown into them during the bainite hold. The austenite
which is retained to room temperature appears in small equiaxed or blocky grain
at the boundaries of the ferrite or bainite grains [17]. This structure is shown inFigure 2.3.
Figure 2.3Typical SEM micrograph of equiaxed TRIP structure, phase A isaustenite, B is bainite and F is inter-critical ferrite [17]
It is possible to produce diff erent microstructures from the one depicted in Fig-
ure 2.3 by using the same two-stage heat-treatment but starting with a diff erent
microstructure. For example, Sugimoto et al. started the heat treatment with a mi-
crostructure of ne martensite [6]. In this work, the TRIP steel was rst heated up to
the fully austenitic region, and held at this temperature for sufficient time to produce
a structure of only austenite. The steel was quenched directly to room temperature
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in order to produce ne martensite, and then re-heated up to the inter-critical tem-
perature so that the two-stage TRIP heat treatment could be applied. The highly
disordered martensitic structure produces a similar driving force for recrystallizatio
as the plastic deformation from cold rolling in the usual TRIP process. When thmartensite is heated to the inter-critical region it transforms into a mixture of ferrite
and austenite. Moreover, the high-carbon martensite partitions over short ranges,
and so the ferrite and austenite form alternating plates from the original repeating
lath structure of the martensite [6]. When the steel is subsequently cooled to the
bainite region, the bainite grows into the austenite, stabilizing it with carbon. The
nal microstructure consists of a matrix of ferrite grains or sub-grains arranged i
thick plates. The retained austenite exists in thin plates between these larger areas
of ferrite and bainite [6]. This microstructure was also investigated in the Ph.D.
work of Mark [5]. Its evolution is sketched in Figure 2.4, and a representative SEM
micrograph is shown in Figure 2.5.
The two TRIP steel processing routes described above were examined in thi
thesis. Each is referred to by the resulting morphology of the retained austenite, suc
that the common, two-stage heat-treatment TRIP steel microstructure is referred toas equiaxed and the more complex martensite-based microstructure is referred t
as lamellar for the thin, plate-like structure of the retained austenite.
2.1.3 The Bainite Transformation
Of the microstructural evolutions in TRIP steel processing, the bainite hold is the
most complex and also the most important in the stabilization of the austenite at roomtemperature. The transformation of austenite into bainite is a complex phenomenon
which is a matter of much research and some controversy. In theory, the reactio
is a non-equilibrium process that produces many ne lenticular sub-units of super
saturated ferrite which arrange themselves into needle-like structures referred to a
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Figure 2.4Sketch of the evolution of the martensite-based TRIP structure from a)martensite laths to b) partially recovered laths + carbon moving toboundaries, c) recovered ferrite laths + austenite and d) recovered ferrite+ austenite with bainite [5]
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Figure 2.5SEM of martensite-based TRIP steel, light phase is austenite/martensite,dark phase is ferrite/bainite [5]
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sheaves of bainite. Bainite will nucleate at an austenite grain boundary and grow
into the grain. The carbon content of the ferrite is higher than equilibrium, while th
remaining carbon is either partitioned to the surrounding austenite or forms carbides
which surround the bainite sheaves. The process by which the bainite nucleates angrows into this structure has been a matter of considerable debate. The crux of this
debate has been whether bainite grows athermally or through a diff usion process.
The original theory of diff usional bainite growth was developed by Heheman et al.
in 1972 [18] and there continues to be kinetic models of bainite growth developed
based on diff usion processes [19]. These models, however, do not account for the
observed kinetics of the bainite transformation which occurs at speeds which ar
orders of magnitude greater than what has been predicted from diff usion models
[20]. In the diff usionless process proposed by Bhadeshia [21] the bainite nucleates in
equilibrium, grows diff usionlessly and the excess carbon is subsequently partitioned
to the surrounding matrix. This model, based on the thermodynamics of the bainite
reaction, accounts for the observed chemistry and shape of the bainite grains and
forms the basis of the model adopted in this thesis.
Bainite initially nucleates at the boundaries of prior austenite grains [21]. Thenuclei will form on pre-existing aws at the grain boundaries such as dislocation ne
works. These areas are highly disordered, and this high free energy allows the nucle
to form with almost any carbon concentration, allowing larger changes in free energ
at nucleation. From a given nucleus the bainite grows quickly and diff usionlessly in
a martensite-like transformation to form a small sub-unit. This sub-unit is made up
of super-saturated ferrite, and its growth must be accommodated plastically by thesurrounding austenite, given that the specic volume of ferrite (BCC) is higher tha
that of austenite (FCC). Plastic accommodation of the bainite sub-unit limits its
growth as the austenite strain-hardens around it. The geometric mismatch between
the parent austenite and the growing bainite results in growth which is more easil
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accommodated in certain crystallographic directions than others. As with marten-
site, this anisotropic strain limits the growth of the sub-units perpendicular to the
habit plane, which in turn results in their lenticular shape, the aspect ratio of which
depends on the formation temperature as well as the carbon and manganese contentof the steel [21].
The newly formed bainite sub-unit is supersaturated in carbon, having the same
carbon content as the parent austenite due to its diff usionless growth. The tem-
perature at which bainite forms is sufficient for this carbon to partition out of the
supersaturated ferrite into the surrounding austenite, but is not sufficient to allow
diff usion of substitutional alloying elements. The carbon, therefore, partitions out o
the sub-units into the surrounding austenite shortly after formation. Being a diff u-
sion process, the concentration and distance of carbon partitioning in the austenite i
temperature dependent. The end result is austenite enriched in carbon, and bainite
with a chemical composition closer to that of equilibrium ferrite.
Once a bainite sub-unit has formed, the total surface area of the austenite grain
boundary has been eff ectively increased by the presence of the boundary between the
newly formed sub-unit and the parent austenite. New bainite sub-units can thereforenot only form on the existing austenite grain boundaries but also at the tip of the
newly-formed sub-units, as shown in Figure 2.6. This additional nucleation site gives
rise to the sheaf structure of the bainite sub-units and eff ectively increases the number
of possible nucleation sites for sub-units as the reaction progresses.
The bainite transformation will end well before the austenite completely trans-
forms as a result of two factors: the thermodynamic driving force for nucleatiodecreases as the carbon content of the remaining austenite increases, and the bai
nite must be plastically accommodated by the austenite which is increasingly strain
hardened by the transformation. For reasonably large grains of austenite the ther-
modynamic limit is more important [22]. The thermodynamic limitations on bainite
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Figure 2.6The growth of a bainite sheaf from primary nucleation at a grainboundary (top) [21]
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growth place a maximum on the possible carbon content of austenite. This maximum
is temperature dependent; that is, decreasing temperatures increase the driving force
of the bainite reaction, and will complete with smaller volume fractions of austen
ite, but the austenite will have a higher carbon content. The line which denes thmaximum carbon content for the austenite at any given temperature is known as
the T 0 line. As can be seen in Figure 2.7, this line lies well below the equilibrium
concentration of the austenite dened by the Ae3 line.
Figure 2.7Representative T 0 and Ae3 lines for a given steel [21]
The transformation from austenite into bainite is a very complex process, gov
erned by non-equilibrium thermodynamics, having multiple transformation steps and
having several ways in which the prior formation of bainite aff ects further transforma-
tion. The microstructures which result from this process in TRIP steels are equally
complex, having several phases of diff erent properties. The structure is made even
more complex by the eff ect of the dynamic martensitic transformation of the retained
austenite which is based on many factors including grain size, stress and strain state
and the composition of the surrounding matrix.
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2.1.4 Retained Austenite Stability
The stability of retained austenite refers to its resistance to transformation with stress,
strain and temperature, so that for a given temperature more stable retained austenite
will transform to martensite at higher stresses and strains. In this work the stability
of austenite will refer to its resistance to transformation with the global strain applied
to the steel, as opposed to the micro-mechanical state of the retained austenite grain
themselves. There are several factors which aff ect the stability of retained austenite.
The most important of these is the concentration of stabilizing alloying elements in th
austentite, usually carbon, silicon and manganese [23]. Of these, carbon is the most
important element aff ecting the stability of austenite in low-alloy TRIP steels [12].Other important factors which control the stability of the retained austenite are the
size of the austenite grains, the nature of the surrounding phases and the morpholog
of the retained austenite. In a study by Reisner et al. [24] it was found that austenite
grains with very low levels of carbon (
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stable under deformation, with grains larger than approximately 1m transforming
too readily to help improve ductility. The relationship between size and stability o
retained austenite grains has been investigated theoretically by Wang and Van Der
Zwaag [23]. As the steel is deformed, dislocations build up in the austenite formingpotential nucleation sites for the martensite transformation. For a given potential
nucleation density small grains of austenite have fewer sites for the nucleation o
martensite, and so are less likely to transform than larger grains with more nucleation
sites available [23]. This relationship is shown below in Figure 2.8 which shows the
theoretical aff ect of grain size on the martensite start temperature (Ms ), a measure
of the thermal stability of the austenite grains. It has also been suggested that the
increase in stability seen in small grains of austenite is partly due to the restriction i
martensite variants that can be selected for transformation in small grains of austenite
[17]. Only certain martensite variants will be favoured by the stress-strain state of the
grain and some of these variants will be geometrically impossible to accommodate
One of the more complex contributions to austenite stability comes from the
interaction of the austenite with the surrounding matrix. The size, shape and com-
position of the phases surrounding the retained austenite grain play a major role inthe global level of strain required for transformation. Thinking of the multiphas
structure of the TRIP steel as a composite material, it is clear that the mechanical
properties of the surrounding matrix will aff ect the stress and strain carried by the
austenite at a given strain level and therefore its transformation behaviour [26]. The
strain and stress state of inidividual TRIP phases has been studied by Jacques et al.
[17] using in-situ neutron diff
raction and in-situ SEM tensile tests. It was found thateach phase carries a diff erent level of stress and strain under loading, with martensite
being the strongest followed by bainite and austenite, while ferrite was the weakes
If the austenite phase is surrounded by soft ferrite it experiences more of the globa
stress and strain than if it were surrounded by bainite. This was found by Jacques e
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Figure 2.8Theoretical relation between the grain size and thermal stability of retained austenite grains [23]
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al. [16] where TRIP steels with retained austenite which had a lower level of carbo
exhibited improved apparent stability due to the stronger surrounding matrix.
The nal source of martensite stabilization is the eff ect of retained austenite phase
morphology, which is the topic of this thesis. Retained austenite has been produceand observed to exist in several diff erent morphologies within TRIP steels. The
simplest of these are the small islands present in a steel with a classic two-stage he
treatment as described in Section 2.1.2. In addition to this morphology researchers
have observed retained austenite as thin layers between bainite sheaves [23] or as a
thicker lamellar shape as produced by Sugimoto et al. [6]. It has been found that
diff erent morphologies of retained austenite have diff erent mechanical stabilities, as
found by Timokhina et al. [26]. It was found that austenite existing as layers between
bainite sheaves was much more stable than the retained austenite existing between
bainite grains. A direct measurement of the transformation behaviour of diff erent
austenite morphologies was presented by Mark [5] using in-situ neutron diff raction
experiments. These experiments provided a direct measure of the volume fraction o
retained austenite which existed in the samples at several strain levels during tensil
loading. Mark investigated four separate austenite morphologies, two fully bainitiTRIP steels (coarse and ne), one traditional equiaxed structure and one lamellar
structure base on the work of Sugimoto et al. [6]. It was found that the ne and
coarse bainitic TRIP steels had the most stable morphology of retained austenite
showing almost no transformation through the testing. The least stable morphology
was found to be the equiaxed samples which transformed quickly at low strains. Th
lamellar structure showed intermediate stability and steadily transformed throughoutthe entire test. These results are summarized in Figure 2.9 which shows the measured
volume fraction of retained austenite with increasing strain for all four austenit
morphologies.
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Figure 2.9Volume fraction of retained austenite measured during tensile testing forfour austenite morphologies [5]
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While the work of Mark and others has shown that there is clearly a relationship
between the morphology and austenite stability, the reasons for this improved stability
are complex and not well dened. Large changes in the processing conditions a
needed to produce these diff erent austenite morphologies and so the chemistry, sizeand surrounding matrix of the retained austenite is necessarily also changed. A
shown above, any of these factors can greatly inuence the stability of the austenit
and so it is not clear what direct aff ect a simple change in morphology would have
on the stability of the retained austenite. Timokhina et al. [26] suggest that the
diff erences in stability seen with diff erent morphologies in their study is due mainly
to the location of the austenite grains in the matrix and their carbon content. In
Marks work [5] the high stability of the bainitic TRIP structures is thought to arise
from the presence of the strong bainite surrounding the austenite grains and also
from the small size of the austenite grains. The equiaxed structure is not constraine
in this way, and transforms readily, while the lamellar structure transformation is
slightly retarded by the surrounding ferrite phase producing behaviour intermediate
between the equiaxed and bainitic structures. A full discussion of the mechanisms
martensite nucleation and all of the factors aff ecting the TRIP transformation can befound in Marks thesis [5].
2.2 TRIP Steel Properties
2.2.1 Tensile Properties
TRIP steels show major improvements over dual phase and other high-strength steelin their behavior under uniaxial tension [27]. The dynamic transformation of TRIP
steels, along with their multi-phase structures, allow them to attain very high uniform
elongations, over 20%, and also very high ultimate tensile strengths (UTS) up to
1000MPa [28]. These properties make TRIP steels very attractive for high-strength
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forming operations. A comparison of stress-strain curves for TRIP versus dual phas
(DP) steels is plotted in Figure 2.10.
Figure 2.10Representative stress-strain curves of DP, TRIP and HSLA steels [4]
TRIP steels exhibit both an increased elongation and a good UTS, due in part to
the sustained hardening of the material that helps suppress the onset of necking [12].
There are several aspects of a stress-strain curve which reveal important informatio
about a given materials formability. The strain at maximum load of a steel present
the onset of plastic instability or diff use necking and is a weak predictor of the forma-
bility of the material, but provides a simple measure which is easily determined fro
the stress-strain data. Better indicators of formability, such as the instantaneous
hardening exponent, can also be calculated from the stress-strain data.
2.2.2 Hardening Rate
The rate at which a material hardens, and how this rate changes with strain, is very
important with respect to its forming behavior. The hardening of a material is often
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modeled as a simple power-law relationship, such that:
= k n (2.1)
where k is the hardening constant, and n is the strain-hardening exponent of the
material. Equation 2.1 is a simplication of the hardening behavior of a material,
but it applies quite well to many metal alloys. A more informative measure of
materials hardening behavior is its instantaneous hardening rate, a value that can
be calculated by taking natural logarithms of Equation 2.1 to isolate n. If the stress-
strain data is taken as a series of discrete data points, and the hardening constant k is
assumed to be constant from point to point, then the instantaneous strain-hardening
exponent (n i ) of the material at a given stress ( m ) and strain (m ) is dened by:
n i =ln m
( m 1)
ln m( m 1)
(2.2)
A material that exhibits a sustained, high value of ni is expected to be resistant to
plastic instability and the onset of necking, producing greater uniform elongations ana more formable material. This is one area in which the martensitic transformation
of TRIP steels greatly aids in their mechanical properties. TRIP steels exhibit more
sustained levels of hardening than DP and other types of steels, as is evident from
the n i values plotted in Figure 2.11.
2.2.3 Sheet Formability Testing
In sheet forming operations the most important quality of a metal is how far it can
be stretched along any given strain path before it fails. Failure in sheet forming i
dened by the onset of local necking. Local necking appears in a sheet material as
line across the surface of the sheet which has experienced higher strains in the plan
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Figure 2.11Comparison of instantaneous n values calculated for DP, TRIP andHSLA steels [4]
of the sheet than the surrounding material and has thinned through the thickness of
the sheet. Considering uniaxial tension, local necking is a condition which follows tonset of diff use necking or plastic instability, ie. after the material is strained past the
point of maximum applied load. If the material is strained beyond this point there wi
eventually develop an orientation in the sheet which undergoes smaller and smalle
increments of strain, approaching zero. This direction of zero strain will dene th
local neck in the material. It is possible to calculate the strain at which local neckin
starts by assuming stress states and hardening behaviours in the material, but in
practice the onset of local necking is usually determined experimentally by deformin
a sheet in a specic strain path and then dening a criterion (usually visual) fo
the onset of local necking. Once this criterion is reached the local major and min
strains within the necked (failed) and un-necked (safe) regions are measured. Whe
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this test is repeated for many diff erent strain paths and the results are plotted in
terms of major and minor strain, the highest safe and lowest failed strains will den
an upper boundary of safe forming. This method of measuring sheet formability w
introduced by Keeler [29] and Goodwin [30] in the 1960s and is widely used andunderstood in industry. This diagram is known as a forming limit diagram (FLD
and the upper boundary is known as the forming limit curve (FLC). An example FLD
with FLCs of HSLA, mild and DP steels is shown in Figure 2.12.
Figure 2.12FLCs of HSLA, mild and DP steels, showing strain paths from uniaxialto equi-biaxial tension [4]
The FLD in Figure 2.12 shows the forming limits for strain paths from negative
minor strains on the left (approaching uniaxial tension) through in-plane-plane strain
at the vertical axis to equi-biaxial tension on the right. The forming limits of a meta
generally follow the shape shown in Figure 2.12, highest in uniaxial tension, lowest
in in-plane plane-strain and increasing again in equi-biaxial tension.
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2.2.4 IPPS Testing
The in-plane plane-strain (IPPS) forming path is the strain path of greatest interest
for materials in sheet forming applications because it typically produces the lowe
level of formability. IPPS represents a forming path where strain is applied to th
major direction of the material and strain in the orthogonal minor direction is con
strained to zero, usually due to the geometry and gripping of the sample. The sampl
is therefore free to deform in the major direction and through the thickness, denin
a plane of plane-strain. Given that the failure of a sheet material is dened by th
visible necking or thinning of the material, and this strain path poses the lowest re
sistance to necking, IPPS testing measures the lower bound of sheet formability foa given material. This low resistance to necking is due to the fact that one of th
criteria for the formation of a neck is the development of a zero-strain direction in th
material along which a neck can form. In other strain paths this zero-strain directio
takes some time to develop, however in IPPS this criterion is automatically satise
by the strain state of the material (zero strain in the minor direction), and so material
thinning and necking occur at lower strains as can be seen from Figure 2.12.
The development of a local neck can either be dened by the thinning of the m
terial (such as the visible formation of a neck), or through the major strain behaviou
of the material. As the material thins locally to form a neck, the material inside the
necked region will have an increasing strain rate compared with the material outsid
the neck which will have a strain rate which decreases towards zero. If a criterion f
local necking is established as a limit on the diff erence between these two strain rates
then the initiation of a local neck can be determined directly from the local majostrain rates of the material.
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2.2.5 TRIP Steels and Forming Path
As discussed above, the lowest limit on TRIP steel forming is along the in-plan
plane-strain path. This reduction in forming limit has been observed to be more
severe in TRIP steels than other grades. The forming limits of several high-strengt
steels were investigated by Bleck et al. [31] who found that TRIP steels showed
the lowest forming limit of all the steels tested in IPPS, but had more comparabl
levels of forming in stretching (tension-tension) and uniaxial tension. It is suggeste
with support from Mamalis et al. [32], that the TRIP steels performed particularly
poorly in IPPS because the full benet of the dilatational martensitic transformation
is dependent on the hydrostatic stress component of the loading.Not only does the strain path aff ect the usefulness of the retained austenite
transformation, the strain path will also aff ect how quickly the retained austenite
transforms during loading. In order to build a multi-scale mechanical model of TRI
steel behaviour Jacques et al. [17] measured the change in volume fraction of retained
austenite as the steel was tested on several strain paths. It was found that the
retained austenite transformed at larger strains in uniaxial and equi-biaxial tension,
but transformed much more quickly for samples nearer to the plane strain condition
This result is shown for one TRIP steel sample below in Figure 2.13.
As discussed in Section 2.1.4, the nature of the TRIP transformation means that
there is an optimal level of austenite stability which allows TRIP steels to resis
failure and prolong ductility. As shown here, the transformation rate of the austenite
is strongly dependant on the strain path. There is therefore no one optimal level o
austenite for a TRIP steel, but rather the optimal austenite stability will depend onthe strain path applied.
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Literature Review 29
Figure 2.13Volume fraction of retained austenite with increasing major strain forTRIP steel tested in several strain paths [17]
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Literature Review 30
2.3 TRIP Steel Characterization
One of the most important aspects of a TRIP steel microstructure is the volume
fraction of retained austenite. While there are several characterization techniques
available to perform this analysis, such as metallography, neutron diff raction and
dilatometry. The two characterization methods employed in this study were X-ray
diff raction and magnetic saturation analysis.
2.3.1 X-Ray Di ff raction
X-ray diff raction is currently the most commonly used method for determining the
volume percent of retained austenite in TRIP steels, as the equipment is widely
available [33]. X-ray diff raction can distinguish the phases in a material by their
variation in atomic spacing. When a beam of X-rays interacts with the surface laye
of a material it will be diff racted at a specic angle depending on the wave-length of
the X-rays ( )and the atomic spacing of the material (d) as dened by Braggs Law:
= 2d
sin (
). (2.3)
For a given wavelength of X-rays, the angle at which a diff raction peak is detected
is related to the atomic spacing of the corresponding phase by Equation 2.5, and the
intensity of the peak is related to the volume fraction of that phase. The volum
fraction of retained austenite in a TRIP steel can therefore be determined from the
intensity of the austenite and ferrite peaks in the material as described in Jatczak [34].
Bainite and martensite have very similar atomic spacings, such that their diff ractionpeaks lie on top of those for ferrite; hence a comparison between the two sets
peaks provides an accurate volume percent of the austenite versus all of the othe
phases. In rolled sheet material there is always the chance that the material has
developed a preferred crystallographic texture, which means it will not diff ract from
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Literature Review 31
all orientations equally. All available orientation peaks of each phase must therefor
be averaged in order to account for any texture within the sample. The intensity o
each peak is divided by its theoretical intensity and averaged with all of the othe
peak intensities. The volume fraction of a specic phase, namely austenite in thcase, can be found using [34]:
V a =1
n An0 (I hA kl/R hA kl))
1
n An0 (I hA kl/R hA kl) +
1
n F n0 (I hF kl/R hF kl)
(2.4)
where nA and nF are the numbers of austenite and ferrite diff raction peaks, I A and
I F are the integrated peak intensities, and RA and RF are the theoretical intensities
of these peaks, which are also listed by Jatczak [34]. Equation 2.4 assumes that the
carbide volume fraction is zero and only the ferrite/bainite/martensite and austenite
peaks are measured.
If the X-ray source is not entirely monochromatic, multiple peaks will appear fo
each lattice spacing. X-rays created from copper radiation will present two separat
peaks, Cu-K 1 and Cu-K 2 . These peaks will lie within a few degrees of each other,
separated by = 2(
)tan (2.5)
where = 0.00382 and = 1.54178 . Therefore, the peaks must be decon-
voluted into two separate peaks in order to properly curve-t them and determine
the integrated intensity. A method for this deconvolution is outlined by Gupta [35]
who employed a Pearson VII t and the assumption that the Cu-K 2 peak has a
maximum intensity of half of the Cu-K 1 peak and a peak shift of , but is otherwiseidentical. Following this approach if the Cu-K 1 curve can be represented by the
equation f (x), then the entire intensity prole is t to the equation:
y(x) = f (x) + 0 .5f (x ) + g(x). (2.6)
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Literature Review 32
The Pearson VII t denes the intensity prole of a peak as:
f (x) = a 1 + (x d)2
b2
m
. (2.7)
where (d, a) are the co-ordinates of the peak maximum, exponent m is the shape
factor and b is proportional to the full-width at half-maximum. When m = the
curve is Gaussian, when m = 1 the shape is Cauchy. The full equation of the curve,
assuming a linear change in background radiation, is:
y(x) = a 1 + (x d)2
b2
m
+ a2
1 + (x d )2
b2
m
+ ( px + q ). (2.8)
The intensity y(x) can be t to the experimental data using the variables a, b, m, d, p
and q [35]. In this way the separate peaks of the copper radiation can be separated
and curve t, eliminating noise and background levels of radiation in the experiment
data and allowing for easier integration of the individual peaks.
2.3.2 Magnetic Saturation
The second technique used to characterize the volume percent of retained austenite i
TRIP steels was based on a measurement of magnetic saturation. The magnetic sat
uration is the maximum magnetic response of a material under an applied magneti
eld. With an increasing magnetic eld the degree to which the material is magne
tized increases, at rst linearly, and then eventually approaching a maximum. If the
eld is then decreased back to zero and reversed, the sample will again pass throug
a linear region before approaching the negative magnetic saturation, as seen in Fig
ure 2.14 which illustrates the hysteresis curves produced in using magnetic saturation
for measuring martensitic transformations in austenitic steels. This reverse respons
is off set from the positive direction to create a hysteresis loop for the material. The
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Literature Review 33
positive off set of the magnetization at zero applied eld is the property which allows
the creation of permanent magnets.
Figure 2.14Magnetic hysteresis curve for an austenitic stainless steel afterundergoing 55% reduction in thickness to produce 75 vol.% martensite [8]
The curve in Figure 2.14 can be modeled as two distinct parts, the linear low-
eld response, and the approach to saturation [33]. According to Berkowitz [36], the
approach to saturation will follow the equation:
M = M s 1 aH bH 2
cH 3 . . . . (2.9)
where M is the magnetic response, M s is the magnetic saturation, H is the applied
eld and a, b and c are material constants.
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Literature Review 34
The magnetic saturation is a material property that varies with temperature,
crystal structure and material chemistry [37], and as such, can be used to diff er-
entiate the volume fraction of crystal structures within a material. This is a wel
established technique and is commonly applied to materials which contain severaphases of known magnetic saturations by stepping the sample through temperature
ranges and measuring the total magnetic saturation of that sample. Since each crys
tal structure has a magnetic saturation that varies with temperature, the magnetic
contribution of each phase, and therefore its volume percent, can be calculated [37].
In the case of TRIP steels this process can be simplied by assuming that the
material only consists of two phases, ferrite and austenite. Given that the austen-
ite is non-magnetizing (para-magnetic) the magnetic response is entirely due to th
response of the ferrite. The presence of martensite and cementite in the structure
does create a source of uncertainty; however, the volume fraction of these phases
low, and their magnetic saturation is close to that of ferrite, so their aff ect on the
retained austenite measurements should be minimal [33]. If the magnetic saturation
of pure ferrite is known for the TRIP steel chemistry, then the volume percent re
tained austenite can be directly determined from the ratio of the measured magneticsaturation to that of pure ferrite. The magnetic saturation of a given sample can be
calculated by tting Equation 2.9 to the measured hysteresis curve and then solving
for M s [33].
2.4 Literature Summary
TRIP steels have great potential for use in high-strength forming operations. Their
formation and the basis of their excellent mechanical properties are, however, ver
complex. To this point low alloy TRIP steels have been formed and tested wit
a variety of chemistries and microstructures. Alloying elements will greatly aff ect
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Literature Review 35
the formation and retention of austenite in the steels and the resulting mechanical
properties [10, 14], especially in the case of carbon, silicon and manganese. Through
the use of the bainite transformation these steels can be produced with a carbon
enriched and meta-stabilized austenite phase [12]. The volume fraction of this meta-stable phase is one of the most important factors aff ecting the formability of the
steel [10] and can be characterized through optical, X-ray and magnetic methods
[33, 34, 38].
The formability of TRIP steels is aff ected by the transformation characteristics
of the retained austenite phase, which is in turn aff ected by the austenite chemistry,
steel morphology and the stress/strain state applied to the material [26]. It has been
found that TRIP steels with lamellar morphologies have austenite with a greater
resistance to transformation. This extra resistance could be most benecial to the
steel in strain paths, such as plane-strain, which show a transformation of retained
austenite at lower strain level [17].
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Chapter 3
Experimental Methods
3.1 Heat Treatments
3.1.1 The As-Received Condition
The material for this study was produced at CANMET-MTL. As received, the stee
contains roughly 0.17 wt.% carbon, 1.5 wt.% manganese and 1.5 wt.% silicon. T
full chemistry of the steel is shown in Table 3.1. The steel was received in a hot
rolled condition. It was rolled in two stages, from 127 mm to 50.8 mm, and then fro
50.8 mm to 5 mm with a reheat in between. The steel was rolled rst through eigh
passes from 127 mm to 50.8 mm and from 1220 C to 920 C. The steel was then
reheated to 1200
C and rolled through seven passes to a thickness of 5 mm and anal temperature of 825 C. The nal sheet, excluding the rounded tails, measured
aproximately 600 mm by 180 mm. Sections of approximately 10 cm by 9 cm were th
cut from the uniformly rolled area of the sheet. The mill scale of these smaller piec
was removed from the surface of the material using a rotary grinder. The sample
were then cold-rolled in the same direction as the hot-rolling, to a nominal thickne
of 1 mm, by approximately twenty passes through the experimental rolling mill a
Queens University. The resulting sheet was used to make heat treatment coupons
tensile samples, and in-plane plane-strain samples.
36
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Experimental Methods 37
Table 3.1Chemistry of the TRIP steel, as received in weight percent, balance isiron
C Si Mn Al Ti P S O N
0.17 1.53 1.50 0.03 0.021 0.007 0.005 0.002 0.0048
3.1.2 TRIP Steel Heat Treatments
In this study, two heat treatment schedules were used to make two distinct TRIP
steel morphologies: one with equiaxed grains of retained austenite and one wit
lamellar grains of retained austenite. These two microstructures shall be referred toas equiaxed and lamellar, respectively. In both of these treatments the steel i
heated to the inter-critical region to form a mixture of austenite and inter-critical
ferrite, and is then cooled to the bainite temperature range where it is held, allowin
some of the austenite to transform into bainite. This transformation enriches the
remaining austenite with carbon, stabilizing it to room temperature. The steel is
then quenched in water. The resulting microstructure consists of inter-critical ferrite
bainitic ferrite and retained austenite. If the austenite is not sufficiently stabilized
some high-carbon martensite may also be present in the microstructure. Cementite
will also form at the edges of the bainite sheaves if the bainite hold time is sufficient
to allow for the diff usion of silicon through the carbon-rich austenite away from the
bainite interfaces.
The diff erence between the equiaxed and the lamellar microstructures stems from
the starting condition of the material. To form the equiaxed microstructure, the steelundergoes the above treatment from the as-cold-rolled condition, while for th
lamellar microstructure the cold-rolled steel is rst heated to the fully-austenitic rang
and then quenched in water to form martensite before undergoing the intercritical and
bainite holds.
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Experimental Methods 38
For all heat treatments the steel was heated to the inter-critical or fully-austenitic
range using an electrical resistance air furnace. The bainite holds were performe
in a 101.6 mm salt pot containing molten NaCl. All quenching was performed b
immersing the sample in still, room-temperature water.Initial heat treatments were performed to investigate the optimal conditions for
the bainite hold for the equiaxed microstructure. Heat-treatment coupons, measuring
25.4 by 50.8 mm were cut out of the cold-rolled sheet using a hydraulic shear. The
coupons were heated to an inter-critical temperature of 750 C and held for 5 minutes.
The coupons were then held in the bainite region at either 350, 400 or 450 C, for
times varying from 30 to 1800 seconds and then quenched to room temperature.
schematic graph of this process is shown in Figure 3.1. The heat-treated coupons
were analyzed optically, magnetically and using X-ray diff raction to determine the
volume fraction of retained austenite in the steel, and the morphology of the variou
phases.
From these heat treatments it was found that holding the steel at 450 C pro-
duced the largest volume percent (14.6 vol.%) of retained austenite in the equiaxe
microstructure. This temperature was then adopted for the bainite holds in all furtherheat treatments.
In order to create tensile samples with an equiaxed structure and varying levels
of retained austenite, heat-treatment blanks measuring 100 by 16 mm were cut from
cold-rolled sheet. These blanks were then held in the inter-critical region at 750 C
for ve minutes. This was followed by a bainite hold at 450 C for times of 30, 60,
100, 300 and 1800 seconds and a nal quench in water. The blanks were then cut insub-sized tensile specimens and magnetic testing specimens using a water-jet cutte
as shown in Figure 3.2. The magnetic test revealed that the peak level of retained
austenite, 14.1 vol.%, was created through a bainite hold of 100 seconds. Blank
measuring roughly 50 by 90 mm were cut from the cold-rolled sheet and used
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Experimental Methods 39
Figure 3.1Schematic graph of the two-stage heat treatment used to produce theequiaxed microstructure.
create in-plane plane-strain samples. These samples were heated to the inter-critica
range as above, and then held at 450 C for 30, 60 and 100 seconds. The blanks werethen cut into in-plane plane-strain samples, and magnetic testing samples as shown
in Figure 3.3. The level of retained austenite in these samples increased to a peak of
11.4 vol.% at the 100 second hold.
To create samples with lamellar microstructures the steel was rst cold-rolled
to a thickness of one millimeter. The cold-rolled sheet was then cut into tensile an
in-plane plane-strain blanks measuring 100 by 16 mm and 50 by 90 mm, respectivelEach blank was placed in a circulating air furnace at 950 C for 1000 seconds, and then
quenched in water. Following this treatment, the tensile blanks were heated to 750 C
and held in the inter-critical region for ve minutes. The samples were subsequent
moved to the salt pot for bainite holds at 450 C lasting 60, 100, 300 and 1800 seconds
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Experimental Methods 40
Figure 3.2Coupon dimensions for tensile sample heat treatments showing positionof tensile bar and magnetic testing discs, all dimensions in mm.
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Experimental Methods 41
Figure 3.3Coupon dimensions for IPPS sample heat treatments showing position of IPPS sample and magnetic testing discs, all dimensions in mm.
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Experimental Methods 42
and then quenched in water. A schematic graph of this process is shown in Figure 3.4.
Sub-sized tensile samples and magnetic testing samples were cut out of the blank
using a water-jet cutter as shown in Figure 3.2. The samples were characterized
magnetically and optically, and mechanically tested under uniaxial tension. Fromcharacterization of the magnetic samples it was found that the maximum volume
fraction of retained austenite of 9.53 vol.% was produced with a bainite hold of 10
seconds.
Figure 3.4Schematic graph of the three-stage heat treatment used to produce thelamellar microstructure.
The in-plane plane-strain blanks were taken from the as-quenched condition and
placed in a circulating air furnace at 750
C for ve minutes, and then moved to asalt pot for a bainite hold of 100 seconds at 450 C followed by a quench in water.
In-plane plane-strain and magnetic testing samples were cut out of the blanks using
a water-jet cutter as shown in Figure 3.3.
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Experimental Methods 43
3.2 Magnetic Saturation Measurements
3.2.1 Magnetic Theory
Measuring the magnetic properties of a material provides a wealth of informatioabout its structure, chemistry and composition. If the phases present in a material
have diff erent magnetic properties it is possible to calculate the volume fraction o
each phase present by measuring the magnetic response. In steels, the ferrite crysta
structure is very strongly ferro-magnetic (an increasing eld will produce an increa
ing magnetic response) whereas austenite is paramagnetic and does not respond t
magnetic elds. For a ferromagnetic material, such as ferrite, when the applied mag
netic eld becomes large enough to change all of the magnetic orientation of th
domains, the material will approach a magnetic saturation point, where an increase
in the applied eld does not increase the magnetic response. This saturation poin
is reliant only on the temperature, crystal structure, chemistry and volume of the
material. Assuming that the sample contains only ferrite and austenite, the volume
fraction of ferrite in a sample can be directly related to the magnetic saturation that
the steel achieves as a fraction of the magnetic saturation of pure ferrite. Magneti
saturation measurements were made using a vibrating sample magnetometer (VSM)
This machine creates a large, variable and uniform magnetic eld, and the sampl
being tested is vibrated perpendicular to the applied eld. The applied magnetic
eld and the resulting magnetic moment created by the magnetized sample are mea
sured. The magnetic response of the material under a high magnetic eld is used t
determine the magnetic saturation.
3.2.2 Magnetic Sample Preparation
The samples tested with the VSM were discs with a ve millimeter diameter, cu
out of one millimeter thick heat-treated sheet samples. The discs were machine
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Experimental Methods 44
using a water-jet cutter and were left in place in the heat-treated samples with tabs
attaching the disc to the surrounding material. The discs were then broken out
of the samples and the tabs ground down using sandpaper. The top and bottom
surfaces of the material were also ground down to remove any oxide incurred durithe heat treatment. This grinding was done so that the exact mass of the magnetizing
steel could be determined. The samples were weighed using a mass balance with
uncertainty of 0.05 mg.
3.2.3 Magnetic Testing Conditions
For all magnetic testing, a VSM facility at the Royal Military College of Canad(RMC) was used. The system consists of ve pieces of equipment including the VS
itself. The VSM uses two large, water-cooled electromagnets to produce a consta
magnetic eld between +6500 and -6500 Oer. A Hall probe placed between th
magnets at the height of the sample measures the applied magnetic eld, while fou
pick-up coils measure the magnetic moment produced by the sample. The sample
placed on the end of a long, non-magnetizing rod which is oriented perpendicular
the magnetic eld. The rod is vibrated along its length, moving the sample perpen
dicular to the magnetic eld. The end of the rod can be raised, lowered and angle
so that the sample vibrates exactly in the centre of the pick-up coils. A schematic o
the VSM is shown in Figure 3.5. The VSM is powered by a large transformer which
converts AC electricity into a high-current DC output, and is capable of producing 26 A. A current control box determines the level of current supplied to the magnet
and can either be controlled manually or through the VSM controller. The VSMcontroller displays the measurements made by the pick-up coils and the Hall probe
This controller can also be used to step the applied magnetic eld up and down t
determine the hysteresis properties of the sample. The measurement conditions can
also be set using the VSM controller. A computer interface allows the user to chan
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Experimental Methods 45
the settin
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