Lawrence Benjamin R 201001 MSc

8/13/2019 Lawrence Benjamin R 201001 MSc

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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

iv


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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

xii


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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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Literature Review 5

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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Literature Review 6

Figure 2.1The aff ect of alloying elements on the thermal processing of TRIPsteels [10]


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Literature Review 7

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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Literature Review 8

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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Literature Review 9

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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Literature Review 10

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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Figure 2.13Volume fraction of retained austenite with increasing major strain forTRIP steel tested in several strain paths [17]


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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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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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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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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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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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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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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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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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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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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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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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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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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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the settin

Lawrence Benjamin R 201001 MSc

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