Structure-Property Relationship In High Strength- High ...

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Structure-Property Relationship In High Strength- High Ductility Structure-Property Relationship In High Strength- High Ductility

Combination Austenitic Stainless Steels Combination Austenitic Stainless Steels

Chengyang Hu University of Texas at El Paso

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STRUCTURE-PROPERTY RELATIONSHIP IN HIGH STRENGTH-

HIGH DUCTILITY COMBINATION AUSTENITIC

STAINLESS STEELS

CHENGYANG HU

Doctoral Program in Materials Science and Engineering

APPROVED:

Devesh Misra, Ph.D., Chair

Guikuan Yue, Ph.D.

Singamaneni Srinivasa Rao, Ph.D.

Stephen L. Crites, Jr., Ph.D. Dean of the Graduate School

Copyright ©

by

Chengyang Hu

2021

STRUCTURE-PROPERTY RELATIONSHIP IN HIGH STRENGTH-

HIGH DUCTILITY COMBINATION AUSTENITIC

STAINLESS STEELS

by

CHENGYANG HU, M.E.

DISSERTATION

Presented to the Faculty of the Graduate School of

The University of Texas at El Paso

in Partial Fulfillment

of the Requirements

for the Degree of

DOCTOR OF PHILOSOPHY

Department of Metallurgical, Materials and Biomedical Engineering

THE UNIVERSITY OF TEXAS AT EL PASO

August 2021

iv

Acknowledgements

The research study described in this dissertation was carried out under the valuable

guidance of Professor Devesh Misra. The topic of dissertation, design of experimental

methodology, modeling, and preparing publications, were all possible because of the wisdom and

hard work of the Professor! The Professor's rigorous attitude, profound knowledge, keen insight,

strategic perspective, rich association, and broad mind greatly benefited me. Every time I discussed

with my Professor, it made me immediately enlightened. The teacher is not only my guide in

scientific research but also a great example in my life. Here, I would like to express my sincere

thanks and respect to my dear Professor Misra!

Thanks to Professor Kaiming Wu and Associate Professor Xiangliang Wan for their

guidance. Their scientific research attitude and numerous discussion provided new insights.

Thanks to Dr. M.C. Somani for the help with the processing of steel by Gleeble.

Grateful thanks are due to Yashwanth Injeti and Na Gong of my research group for their

guidance and useful discussion. Special thanks are to post-doctoral Dr. Kun Li and senior student

Bing Yu for their guidance and assistance. Thanks to Guanghui Wu, Lei Zhong, Hangyu Dong

and other students for their help in experimentation and many useful discussion. I would also like

to sincerely thank Dr. Guikuan Yue and Dr. Singamaneni Srinivasa Rao for serving on the

dissertation committee.

I also thank the department for providing access to experimental techniques.

Thanks are to my girlfriend Yang Yang's for encouragement, and parents and family for

their trust, support, and care!

v

Abstract

Austenitic stainless steels are widely used in our daily life, but their mechanical strength is

low. In order to improve their yield strength via grain refinement, an investigation was carried out

involving phase reversion annealing concept comprising of severe cold roll reduction followed by

annealing at different temperatures for short durations. During annealing reversion of deformation-

induced martensite to austenite occurred by shear mechanism, leading to fine-grained structure

and high strength-high ductility combination.

Nanoscale deformation studies suggested that the deformation mechanism of nanograined

structure was different from the coarse-grained counterpart. Post-mortem electron microscopy of

plastic zone surrounding the indent indicated that the active deformation mechanism was

nanoscale twinning with typical characteristics of a network of intersecting twins in the

nanograined structure, while strain-induced martensite transformation was the effective

deformation mechanism for the coarse-grained structure. The presence of ~3 wt % Cu in austenitic

stainless steel had a moderate effect on strain-rate sensitivity and activation volume at similar grain

size in relation to the Cu-free counterpart. The nanoscale twin density was noticeably higher in

Cu-bearing austenitic stainless steel as compared to Cu-free counterpart, a behavior that may be

related to the increase of stacking fault energy.

Furthermore, the synergistic effect of grain boundary and grain orientation on micro-

mechanical properties of austenitic stainless steel was studied. Micro/nano-scale deformation

behavior including hardness, elastic modulus, and pop-ins, was studied. Relatively higher hardness

and modulus was observed near {101} and more pop-ins occurred in this orientation at high

loading rate.

vi

From the perspective of engineering applications, the wear performance of fine-grained

austenitic stainless steel through three-body abrasive wear tests at room and high temperatures was

compared with the coarse-grained counterpart. The study demonstrated that fine austenite grains

with high yield strength and elongation exhibited superior wear resistance at high temperature

(250 °C), which was attributed to deformation twinning-induced plasticity in fine austenite grains.

The wear mechanisms were microploughing and microcutting.

vii

Table of Contents

Acknowledgements ................................................................................................................................. iv Abstract ....................................................................................................................................................... v Table of Contents .................................................................................................................................... vii List of Tables ............................................................................................................................................ xi List of Figures......................................................................................................................................... xiii Chapter 1: Introduction ........................................................................................................................... 1

1.1 WHAT ARE STAINLESS STEELS ........................................................................................ 1 1.2 DIFFERENT TYPES OF STAINLESS STEELS ...................................................................... 1 1.3 APPLICATIONS OF STAINLESS STEELS ............................................................................ 4 1.4 EFFECT OF ALLOYING ELEMENTS ON MICROSTRUCTURE ........................................... 6 1.5 MECHANICAL PROPERTIES OF STAINLESS STEELS ..................................................... 14 1.6 CORROSION RESISTANCE OF STAINLESS STEELS ........................................................ 21 1.7 DEFORMATION BEHAVIOR OF STAINLESS STEELS ...................................................... 30 1.8 STACKING FAULT ENERGY OF STAINLESS STEELS AND INFLUENCE OF STACKING

FAULT ENERGY ON DEFORMATION BEHAVIOR ........................................................... 34 1.9 SUMMARY ........................................................................................................................ 36

Chapter 2: Experimental Procedure ................................................................................................... 38 2.1 PHASE REVERSION .......................................................................................................... 38 2.2 METALLOGRAPHY ........................................................................................................... 45 2.3 X-RAY DIFFRACTION ...................................................................................................... 45 2.4 TENSILE TESTS ................................................................................................................. 46 2.5 FRACTURE SURFACE EXAMINATION BY SEM ............................................................ 46 2.6 NANOINDENTATION ........................................................................................................ 47 2.7 TEM FOIL PREPARATION AND TEM ............................................................................ 48 2.8 EBSD SAMPLE PREPARATION AND EBSD .................................................................. 51

Chapter 3: Improving the yield strength of an antibacterial 304Cu austenitic stainless steel by the reversion treatment ......................................................................................................................... 56

3.1 MATERIAL AND EXPERIMENTAL PROCEDURE............................................................. 56 3.2 RESULTS ........................................................................................................................... 58

3.2.1 Cold rolling ............................................................................................................. 58 3.2.2 Reversed microstructures ..................................................................................... 58 3.2.3 Grain size ................................................................................................................. 63 3.2.4 Precipitation structure ........................................................................................... 64 3.2.5 Tensile properties and strain-induced martensite ............................................ 67 3.2.6 Hardness ................................................................................................................... 71

3.3 DISCUSSION ...................................................................................................................... 72 3.3.1 Reversion behavior ................................................................................................ 73 3.3.2 Precipitation kinetics ............................................................................................. 79

viii

3.3.3 Enhanced strength .................................................................................................. 81 3.4 CONCLUSIONS .................................................................................................................. 84 3.5 SUMMARY ........................................................................................................................ 85

Chapter 4: On the mechanical behavior of austenitic stainless steel with nano/ultrafine grains and comparison with micrometer austenitic grains counterpart .................................................. 86

4.1 MATERIALS AND EXPERIMENTAL PROCEDURE........................................................... 86 4.1.1 Materials .................................................................................................................. 86 4.1.2 Nanoscale deformation ......................................................................................... 87

4.2 RESULTS AND DISCUSSIONS .......................................................................................... 87 4.2.1 Load-controlled nanoscale deformation experiments: load-displacement plots 88 4.2.2 Nanoscale deformation ......................................................................................... 89

4.3 CONCLUSIONS .................................................................................................................. 93 4.4 SUMMARY ........................................................................................................................ 93

Chapter 5: The significance of phase reversion-induced nanograined/ultrafine-grained structure on the load-controlled deformation response and related mechanism in copper-bearing austenitic stainless steel ......................................................................................................... 95

5.1 MATERIALS AND EXPERIMENTAL PROCEDURE........................................................... 95 5.2 RESULTS ........................................................................................................................... 96

5.2.1 Microstructure of CG and NG/UFG austenitic stainless steels .................... 96 5.2.2 Mechanical properties ........................................................................................... 97 5.2.3 The tensile fracture surface .................................................................................. 97 5.2.4 Nanoindentation experiments .............................................................................. 98

5.2.4.1 Load-controlled nanoindentation experiments ........................................ 98 5.2.4.2 Strain rate controlled nanoindentation experiments ............................ 100

5.2.5 Deformation structure ......................................................................................... 101 5.3 DISCUSSION .................................................................................................................... 103

5.3.1 Strain-rate sensitivity and activation volume ................................................. 103 5.3.2 Deformation mechanism in NG/UFG and CG structure.............................. 104 5.3.3 Fracture behavior of NG/UFG and CG ........................................................... 107 5.3.4 The relationship between austenite stability and strain energy .................. 107 5.3.5 The effect of Cu addition on 304 stainless steel ............................................ 109

5.4 CONCLUSIONS ................................................................................................................ 110 5.5 SUMMARY ...................................................................................................................... 111

Chapter 6: The synergistic effect of grain boundary and grain orientation on micro-mechanical properties of austenitic stainless steel .............................................................................................. 112

6.1 MATERIALS AND EXPERIMENTAL PROCEDURE......................................................... 112 6.1.1 Material .................................................................................................................. 112 6.1.2 Nanoindentation and post-mortem characterization ..................................... 113

6.2 RESULTS ......................................................................................................................... 113

ix

6.2.1 Microstructure ...................................................................................................... 113 6.2.2 Nanoindentation behavior .................................................................................. 114

6.3 DISCUSSION .................................................................................................................... 119 6.3.1 Effect of grain orientation on nanoindentation behavior ............................. 119 6.3.2 Effect of grain boundaries on nanoindentation behavior ............................. 123


Chapter 7: On the impacts of grain refinement and strain-induced deformation on three-body abrasive wear responses of 18Cr–8Ni austenitic stainless steel ................................................. 128

7.1 EXPERIMENTAL METHODS ........................................................................................... 128 7.1.1 Materials ................................................................................................................ 128 7.1.2 Microstructural characterization ....................................................................... 128 7.1.3 Mechanical property tests................................................................................... 129 7.1.4 Three-body abrasive wear tests ......................................................................... 130

7.2 RESULTS ......................................................................................................................... 131 7.2.1 Microstructure ...................................................................................................... 131 7.2.2 Mechanical properties ......................................................................................... 133 7.2.3 Three-body abrasive wear performance .......................................................... 133

7.3 DISCUSSION .................................................................................................................... 139 7.3.1 Effects of grain refinement on mechanical properties in austenitic stainless steel 139 7.3.2 Effects of grain refinement and test temperature on wear resistance in austenitic stainless steel ...................................................................................................... 140 7.3.3 Wear mechanisms of austenitic stainless steel ............................................... 142


Chapter 8: Conclusions and future work ........................................................................................ 146 8.1 CONCLUSIONS ................................................................................................................ 146

8.1.1 Improving the yield strength of an antibacterial 304Cu austenitic stainless steel by the reversion treatment ........................................................................................ 146 8.1.2On the mechanical behavior of austenitic stainless steel with nano/ultrafine grains and comparison with micrometer austenitic grains counterpart and their biological functions ...................................................................................................... ............................................................................................................................... 147 8.1.3 The significance of phase reversion-induced nanograined/ultrafine-grained structure on the load-controlled deformation response and related mechanism in copper-bearing austenitic stainless steel ......................................................................... 148 8.1.4 The synergistic effect of grain boundary and grain orientation on micro-mechanical properties of austenitic stainless steel ........................................................ 149 8.1.5 On the impacts of grain refinement and strain-induced deformation on three-body abrasive wear responses of 18Cr–8Ni austenitic stainless steel ...................... 149

x

8.2 FUTURE WORK ............................................................................................................... 150 References ............................................................................................................................................. 151 Vita ......................................................................................................................................................... 185

xi

List of Tables

Table 1.1: Uniform corrosion resistance of different grades stainless steels ................................ 25

Table 1.2: Critical pitting corrosion potential of super ferritic stainless steels at 3.5%NaCl, pH6.5

....................................................................................................................................................................... 27

Table 1.3: Pitting corrosion potential of different materials ............................................................. 27

Table 1.4: Critical pitting temperature of different materials ........................................................... 27

Table 1.5: Critical crevice corrosion temperature of different materials ....................................... 28

Table 1.6: Stress corrosion cracking resistance of different materials ........................................... 29

Table 1.7: Stress corrosion cracking resistance of different materials at 100 °C 40% CaCl2 ... 29

Table 3.1: Chemical composition (wt. %) of the experimental Cu-bearing austenitic stainless

steel ............................................................................................................................................................... 56

Table 3.2: Tensile properties of the 304Cu steel after reversion annealing treatments compared

to those of as-received (hot-rolled) and cold-rolled conditions. ...................................................... 68

Table 3.3: Martensite content before and after tensile testing and formed during tensile test of

samples annealed at different conditions. ............................................................................................. 71

Table 3.4: Number weighted average GS and corresponding calculated and measured YS after

reversion annealing at different conditions (°C-s). ............................................................................. 82

Table 5.1: Chemical composition (wt. %) of experimental Cu-bearing austenitic stainless steel.

....................................................................................................................................................................... 95

Table 6.1: Chemical composition (wt. %) of the investigated medical austenitic stainless steel.

..................................................................................................................................................................... 112

xii

Table 6.2: Elastic modulus and hardness for each orientation based on the data of 208 indents.

..................................................................................................................................................................... 116

Table 6.3: The strain rate sensitivity index (m), activation volume (v) calculated from the data

of the loading stage for nanoindentation tests for nine indents near {111}, {101} and {001}

grains. ......................................................................................................................................................... 122

Table 7.1: Chemical composition (wt. %) of the investigated 18Cr-8Ni stainless steel. .......... 128

Table 7.2: The experimental parameters for the stirring wear test. ............................................... 130

Table 7.3: The measured mechanical properties of the investigated steels. ................................ 133

Table 7.4: The hardness of the worn surface for investigated steels (HV0.5). ........................... 137

Table 7.5: The average martensite volume percentage of FG annealed sample and as-received

CG sample before and after wear tests (vol. %). ............................................................................... 138

xiii

List of Figures

Figure 2.1: Illustration of phase reversion process for a metastable austenitic stainless-steel that

includes cold rolling and annealing process. [134, 137] ..................................................................... 40

Figure 2.2: Reversion in 304Cu ASS occurred by the shear reversion (a), where dislocation free

grains are formed by continuous recrystallization (white arrows marked part) [136], diffusional

reversion (b) [136]. ..................................................................................................................................... 42

Figure 2.3: Time-Temperature- Reversion (TTR) diagram and an example of the reversion

treatment at 700 °C for the studied 304Cu steel. [136] ....................................................................... 44

Figure 2.4: Schematic of electron beam in TEM [208,209]. ............................................................. 50

Figure 2.5: An EBSD system. (a) Principle components of an EBSD system, (b) a photograph

showing the EBSD system integrated with an EDS system [210]. .................................................. 52

Figure 2.6: The formation of the electron backscattered diffraction pattern (EBSP). (a) Cones

(green and blue) generated by electrons from a divergent source which satisfy the Bragg equation

on a single lattice plane. These cones project onto the phosphor screen, and form the Kikuchi

bands which are visible in the EBSP. (b) Generated EBSP [210]. ................................................... 53

Figure 2.7: The spherical diffraction patterns generated by different orientations of a cubic

structure. [210]. ........................................................................................................................................... 55

Figure 3.1: Formation of αʹ-martensite during cold rolling. .............................................................. 58

Figure 3.2: Austenitic grain structure after annealing at 900 °C for 1 s. EBSD grain boundary

map (a) and the orientation image map (b)............................................................................................ 59

Figure 3.3: Reversed grain structure after annealing at 850 °C-1 s (a) and 800 °C-10 s (b and c).

Grains containing low angle grain boundaries pointed by arrows in (a), presence of irregular grain

xiv

(b) and a non-recrystallized deformed austenite grain in (c). (Austenite gray, martensite red in

color).............................................................................................................................................................. 60

Figure 3.4: Microstructure obtained after annealing at 700 °C for 10 s. Phase map (a) and OIM

map (b). Martensite red-colored in (a). .................................................................................................. 61

Figure 3.5: Microstructure obtained after annealing at 700 °C for 1800 s at two different

magnifications (OIM maps). ..................................................................................................................... 61

Figure 3.6: Microstructure obtained after annealing at 650 °C for 3600 s. Martensite red in the

phase map (left). .......................................................................................................................................... 62


phase map. .................................................................................................................................................... 62

Figure 3.8: Fraction of martensite retained after annealing at 750, 700 and 650 °C for various

annealing durations. .................................................................................................................................... 63

Figure 3.9: Grain size distribution after reversion annealing at different conditions based on high

angle grain boundaries (HAGBs) (a) or both HAGBs and low angle grain boundaries (LAGBs;

misorientation 2–15°) (b). ......................................................................................................................... 64

Figure 3.10: STEM micrograph after annealing at 700 °C for 1.5 h (a), the corresponding X-ray

map (b) and electron diffraction patterns of austenite (c) and martensite (d and e) taken from

areas marked in (a) by dashed circles. .................................................................................................... 66

Figure 3.11: A local view of dislocation-free austenite grains in a sample annealed at 700 °C for

1.5 h. Bright field (a) and dark field (b) images revealing nano-size particles. A magnified view

(c) of the square area marked with red line in (b) and corresponding X-ray map of Cu distribution

in this area (d). Black spots in (c) are holes (i.e. lost precipitates) and are not seen in (d). ........ 66

xv

Figure 3.12: A TEM 2-beam BF image revealing the coherence contrast of Cu precipitates in

austenite (a) and an HR-STEM image of a Cu particle (b). Annealing at 700 °C for 1.5 h. ....... 66

Figure 3.13: A STEM micrograph of the sample annealed at 650 °C for 1.5 h showing small

reversed dislocation-free austenite grains surrounded by deformed structure. Coherent Cu

precipitates in grains 1 and 2. ................................................................................................................... 67

Figure 3.14: Stress-strain curves of a cold rolled specimen and some reversion annealed ones in

different conditions. .................................................................................................................................... 69

Figure 3.15: Effect of annealing duration at 750, 700 and 650 °C on yield strength. .................. 69

Figure 3.16: Strain hardening rate as a function of true strain for the specimens annealed at

different conditions: (a) 750–900 °C with varying holding times 10–100 s, (b) 700 °C/100–

5400 s and (c) 650 °C/1800–5400 s. ....................................................................................................... 70

Figure 3.17: The amount of new DIM formed during tensile straining of the samples annealed at

650, 700 and 750 °C for different durations. ......................................................................................... 71

Figure 3.18: Hardness variation after annealing at different temperatures for 1, 10 and 100 s.

Some data from Mészáros and Prohászka [219] for 1 h and Martins et al. [220] for 0.5 h are

included. The shaded area highlights the temperature range, where the influence of annealing

duration is significant. ................................................................................................................................ 72

Figure 3.19: Formation of defect-free austenite grains during annealing at 700 °C for 10 s (a) and

600 s (b) indicating the shear reversion mechanism followed by continuous recrystallization.

Low angle grain boundaries are white lines in the orientation image map (a), and martensite is

red in the phase map (b). ........................................................................................................................... 75

Figure 3.20: Examples of big difference in the grain size in reversed dislocation-free grains after

xvi

annealing at 700 °C for 10 s (a,b) and 600 s (c,d). DA is retained deformed austenite grain (a).

Martensite is in red in the phase map (b,d)............................................................................................ 76


treatment at 700 °C for the studied 304Cu steel. .................................................................................. 78

Figure 3.22: Yield strength versus total elongation after different reversion conditions compared

to reversion treated 3XX grade austenitic stainless grades (data from Ref. [243]). ..................... 82

Figure 4.1: Light and TEM micrographs illustrating the microstructure of coarse-grained (CG)

and nanogrianed/ultrafine-grianed (NG/UFG) austenitic stainless steels with an average grain

size of ~55 ± 20 μm and ~200–400 nm, respectively. ......................................................................... 88

Figure 4.2: Load-displacement plots at constant load rate of 2 uNs−1 for CG and NG/UFG steel,

respectively. ................................................................................................................................................. 89

Figure 4.3: Hardness versus strain rate plots for CG and NG/UFG austenitic stainless steels at

different strain rates. ................................................................................................................................... 90

Figure 4.4: Post-mortem transmission electron microscopy of the plastically deformed region

surrounding the indented region illustrating twinning as the actual deformation mechanism in

NG/UFG austenitic stainless steel. (a) bright field micrograph and (b) dark field micrograph. The

inset in (a) is the electron diffraction pattern from the twinned region. .......................................... 92


surrounding the indented region illustrating strain-induced martensite as the actual deformation

mechanism in CG austenitic stainless steel. The inset is the electron diffraction pattern from the

martensite region. ........................................................................................................................................ 92

Figure 5.1: (a) Light and (b) transmission electron micrographs of CG and NG/UFG structure,

xvii

respectively in Cu-bearing austenitic stainless steel. ........................................................................... 97

Figure 5.2: Typical engineering stress-strain curves for CG and NG/UFG Cu-bearing austenitic

stainless steels. ............................................................................................................................................. 97

Figure 5.3: SEM fractographs at identical magnifications illustrating microvoid coalescence type

of fracture in CG (a and b) and line-up of voids along the striations in NG/UFG (c and d) in Cu-

bearing austenitic stainless steels. Figures (b) and (d) are processed images with Image Pro

software to clearly illustrate striations observed in NG/UFG Cu-bearing austenitic stainless steel

(c). .................................................................................................................................................................. 98

Figure 5.4: Load-displacement plots at fixed loading rate of 2 μN s−1 for NG/UFG and CG Cu-

bearing austenitic stainless steels obtained via load controlled nanoindentation experiments. .. 99

Figure 5.5: Hardness versus strain rate plots for CG and NG/UFG Cu-bearing stainless steels

obtained via strain rate controlled nanoindentation experiments. Please note that the hardness is

in GPa. Thus, there is significant difference in the hardness of NG/UFG and CG Cu-bearing

austenitic stainless steel. .......................................................................................................................... 101

Figure 5.6: Post-mortem electron microscopy of the plastic zone surrounding the indented region

in Cu-bearing CG austenitic stainless steel illustrates stain-induced martensite. ........................ 102


in Cu-bearing NG/UFG austenitic stainless steel. .............................................................................. 102

Figure 6.1: The SEM micrograph (a), EBSD grain boundary map (b) and TEM micrographs (c,

d) for the original microstructure of the investigated steel. The blue lines in (b) implying grain

boundary misorientation greater than 15°. ........................................................................................... 114

Figure 6.2: Representative post-mortem EBSD orientation map (a), load-displacement plots for

xviii

indents in group I in 2a (b) for the sample........................................................................................... 115

Figure 6.3: (a–c) Load-displacement plots from loading to unloading for nine samples

representing indentations in grains near {111}, {001}, and {101}, respectively and (d) load-

induced displacement as a function of loading time. ......................................................................... 118

Figure 6.4: (a–c) Stress - strain rate curves during the loading stage for nine samples representing

indentations on grains near {111}, {001}, and {101}, respectively. ............................................. 121

Figure 6.5: The distribution of the first pop-in displacement (a) and load (b) as a function of

distance to grain boundary of the indents located in grains with orientation close to {001}, {101},

and {111}, symbolized with squares, triangles and cross, respectively. ....................................... 123

Figure 6.6: Schematic illustration for the plastic zone radius (c), where point A is the dislocation

source in the neighboring grain [297]. .................................................................................................. 124

Figure 6.7: Distributions of ratio (c/d) for the indents located in grains with orientation close to

{001}, {101}, and {111}, symbolized with triangles, circles and squares, respectively,

superimposed with amplitude version of Gaussian peak function. ................................................ 125

Figure 7.1: (a) Schematic illustration of the three-body abrasive wear test and dimensions of the

specimens and (b) the shape and size of quartzite stones used in the experiment. ..................... 130

Figure 7.2: The microstructure of the as-received CG (a) and FG annealed (b) samples. ........ 132

Figure 7.3: TEM bright field micrographs of (a, b) as-received CG and (c, d) FG annealed

samples, respectively................................................................................................................................ 132

Figure 7.4: EBSD results for grain boundary reconstruction maps of austenite in as-received CG

(a) and FG annealed (b) samples combined with grain size distribution fraction in as-received

CG (c) and FG annealed (d) samples. ................................................................................................... 133

xix

Figure 7.5: The average accumulated weight loss (a, c) combined with their weight loss rate (b,

d) of the investigated steels in room temperature (a, b) and high temperature (c, d) stirring wear

test. ............................................................................................................................................................... 135

Figure 7.6: The SEM pictures for worn surface morphology of edge part (left and/or right view

of wear part) and center part (front and/or back view of wear part) of investigated samples in

both the room and high temperature work condition stirring wear test. ........................................ 136

Figure 7.7: The harness versus depth plots of subsurface deformation layer of FG annealed

sample and as-received CG sample before (a), after the wear tests at room temperature (b) and

high temperature (c). ................................................................................................................................ 138

Figure 7.8: Schematic illustrations for wear mechanisms in wear process. (a) Microploughing;

(b) Microcutting. ....................................................................................................................................... 143

1

Chapter 1: Introduction

1.1 WHAT ARE STAINLESS STEELS

Stainless steels have alloying elements (mainly Cr and Ni) to improve its corrosion

resistance and toughness [1]. Since the invention and production of world stainless steel by Krupp

in the late 1920s, they have been widely used all over the world. In Europe and America, stainless

steel industry developed rapidly in 1950s and 1960s, while the output of stainless steel in Japan

ranked first in the world in the 1970s. Since the 1980s, stainless steel production in Asia has rapidly

progressed [2].

At present, there is continued interest in the deformation behavior of low stacking fault

energy non-stationary austenitic stainless steels, mainly because of the uncertainty of stress-strain

behavior during deformation of stainless steel. There are many reasons for this uncertainty, such

as chemical composition, temperature, strain conditions (strain rate, strain path, etc.), grain size,

etc.

1.2 DIFFERENT TYPES OF STAINLESS STEELS

There are many kinds of stainless steel with different properties. The classification methods

are as follows: (i) chemical composition of stainless steel or some characteristic elements of steel,

such as Cr stainless steel and Cr-Ni stainless steel, (ii) properties and uses of stainless steel, such

as non-magnetic stainless steel, plasticity of stainless steel, and low-temperature stainless steel,

(iii) general grade of stainless steel, such as 300 series, 400 series and (iv) microstructure of

stainless steel.

They are usually classified according to the microstructure of steel: austenitic stainless steel,

ferritic stainless steel, martensitic stainless steel, duplex stainless steel, and precipitation hardening

stainless steel.

2

The characteristics of stainless steel are introduced based on the microstructural

classification of stainless steels [3,4]:

(1) Austenitic stainless steel: Austenitic stainless steels are most widely used stainless

steels. Austenitic stainless steel refers to the addition of a variety of elements, mainly Cr and Ni,

as well as a small amount of Mn, N, C, etc. during the process of ferroalloy smelting, when the

combined action of these elements renders them to have austenite structure at room temperature

[5].

Austenitic stainless steel is a face-centered cubic structure, and the representative steel

types are 301, 304, 321, and 316. The main characteristics are as follows:

a) Under normal heat treatment conditions, the matrix structure of steel is austenite. Under

improper heat treatment or different heating conditions, there may be a small amount of carbide

and ferrite in the austenite matrix.

b) The mechanical properties of austenitic stainless steel cannot be changed by heat

treatment, but can only be strengthened by cold deformation.

c) By adding alloying elements such as Mo, Cu, and Si, different steel grades, such as 316L

and 304Cu, are obtained.

d) Nonmagnetic, good low-temperature performance, easy forming, and weldability are the

important characteristics of steel.

Austenitic stainless steel not only have excellent corrosion resistance but also have good

plasticity, low-temperature toughness, work hardening ability, and weldability. They are widely

used to store nitric acid, organic acid, salt, alkali, and in other industries.

(2) Ferritic stainless steel: The corrosion resistance, plasticity, and weldability of ferritic

stainless steels are better than martensitic stainless steel, but their strength is lower. This kind of

3

steel is mainly used for mechanical parts and structural parts with low mechanical properties but

have requirements for corrosion resistance, such as in the nitric acid absorption tower, heat

exchanger, phosphoric acid tank, etc. and can also be used as anti-oxidation material at high

temperature.

(3) Martensitic stainless steel: Martensitic stainless steel has good hardenability, high

hardness, and martensitic structure at room temperature, representing 410 and 420 steel grades.

The main characteristics are:

a) Martensitic stainless steel has strong magnetic properties at room temperature. Generally

speaking, its corrosion resistance is not outstanding, but its strength is high. It is used as high

strength structural steel.

b) It has a stable austenite structure at high temperature, martensite phase under air cooling

or oil cooling, and full martensite structure at room temperature.

Martensitic stainless steel has good corrosion resistance in oxidizing medium (steam,

atmosphere, seawater, oxidizing acid), but poor in nonoxidizing medium (alkali solution,

hydrochloric acid).

(4) Duplex stainless steel: It has high Cr and N composition, austenite and ferrite mixed

phase at room temperature. The representative steel grades are 2304, 2205, and 2507. Main

characteristics are:

a) The matrix is ferrite at high temperature, and has a 30-50% ferrite + austenite dual-phase

structure when cooled to room temperature.

b) High yield strength, strong pitting corrosion resistance, stress corrosion resistance, easy

forming, and welding.

It plays an important role in fertilizer plant equipment, petroleum refining industry, marine

4

condenser, etc.

(5) Precipitation hardening stainless steel: Precipitation hardening stainless steel can be

divided into martensitic precipitation hardening stainless steel (represented by 0Cr17Ni4Cu4Nb),

semi austenitic precipitation hardening stainless steel (represented by 0Cr17Ni7Al and

0Cr15Ni25Ti2MoVB), and austenitic plus ferrite precipitation hardening stainless steel

(represented by ph55A, B and C). These kinds of steels improve the strength of the material by

precipitation of intermetallic compounds such as Cu, Al, Ti, Nb after heat treatment. The main

features are as follows:

a) This type of stainless steel can be easily processed and formed. Semi austenitic

precipitation hardening stainless steel has high strength and good toughness through martensitic

transformation and precipitation hardening, while austenite and martensite precipitation hardening

stainless steel has high strength and good toughness through precipitation hardening treatment.

b) The content of Cr is about 17%, in addition to Ni, Mo, and other elements, such that the

corrosion resistance of 18-8 type austenitic stainless steel is close to 18-8 type austenitic stainless

steel.

They are mainly used in some pressure vessels, pipes, springs, diaphragms, etc.

1.3 APPLICATIONS OF STAINLESS STEELS

Stainless steel is widely used in the chemical industry, biology, aerospace, nuclear energy,

medical equipment, and bioengineering because of their superior corrosion resistance, welding

performance, and non-magnetic character and excellent processing performance.

Among them, 304 austenitic stainless steel is common. Each element in the product

standard must meet the requirements, otherwise, it cannot be called 304 stainless steel. 304

stainless steel is also called 18/8 stainless steel, Cr content is more than 18%, Ni content is more

5

than 8%. At present, based on the wide application prospect of 304 stainless steel, trace alloying

elements are added to enhance their performance, to further expand the market application range.

409L steel is ferritic stainless steel with good corrosion resistance and a small thermal

expansion coefficient [6-11]. Compared with stainless steel containing a large amount of Ni,

stainless steel has been widely used in automobile exhaust systems due to its low cost. According

to the literature [12], the stainless steel used for the exhaust system of each vehicle is about 24 kg,

of which 409L ferritic stainless steel accounts for 80%. 409L ferritic stainless steel contains 12

wt.% Cr, which makes the ferrite phase zone expand and the austenite phase zone shrink. There is

almost no γ phase during the solidification process, such that there is no solid phase transformation

between austenite and ferrite.

The stainless steel used for rail metro body should have t advantages of good strength,

reduced thickness, good weldability, and easy cold working. The first two requirements are based

on the lightweight of the car body, and the latter is applied to the processing formability of the car

body structure. At present, austenitic stainless steel containing Cr and Ni is commonly used to

produce a rail car body.

Since the 1930s, Bard company in the United States has produced the first stainless steel

rail car. France also followed suit in manufacturing stainless steel rail vehicles. Bombardier

Canada manufactured more than 1500 rail cars in the years after 1982, of which nearly 90% were

stainless steel buses. During the same period, the company also manufactured 252 stainless steel

passenger cars of 6 different types for the British French subsea tunnel. At present, the production

scale, production process, and raw material research and development of stainless steel vehicles

are at the world's leading level. Japan started to develop stainless steel car body later than the

United States. The R & D process of the Japanese stainless steel car body has gone through four

6

stages, namely skin stainless steel, semi-stainless steel, all stainless steel, and lightweight stainless

steel. The stainless steel used has developed from SUS201, SUS304, SUS301 to SUS301L. In

recent years, a large number of Japanese lightweight stainless steel car bodies have been

manufactured, which greatly reduces the maintenance work and cost of car body structure, and

further realizes the goal of green energy saving. Since SUS301L was launched, many countries

and companies have focused on this material. In order to control the body weight of a high-speed

tilting train, AISI301L/ 1.4318 is selected as the material for body structure.

1.4 EFFECT OF ALLOYING ELEMENTS ON MICROSTRUCTURE

Alloying element refers to a certain amount of one or more kinds of metal or nonmetal

added during the process of smelting. The addition of alloying elements can optimize the physical

and chemical properties of the materials, such as increasing strength, improving oxidation

resistance, improving plasticity, and processing properties. Therefore, the addition of alloying

elements is also an important means to improve the properties of stainless steel in industrial

production. Most of the chemical elements that make up the alloying additions are metallic

elements, such as Cu, Mn, Cr, Mo, Ni, and rare metals, and a few are non-metallic elements, such

as C, N, S, etc.

Alloying elements can be classified according to the following three characteristics:

First, according to the characteristics of interaction with Fe, they can be divided into

austenite forming elements, such as C, N, Cu, Mn, Ni, etc., and ferrite forming elements, such as

Cr, Si, Al, Mo, etc. The two forming elements are represented by Cr and Ni respectively. Therefore,

the sum of the ability of each element to form austenite or ferrite is usually called Ni equivalent

Nieq and Cr equivalent Creq [4]. It can be expressed as follows:

�� = �� + 30 + �� + 0.5�� (1.1)

7

�� = � + ��1.5�� + 0.5�� (1.2)

Generally, austenite forming elements are preferentially distributed in austenite, while

ferrite forming elements are preferentially distributed in ferrite. However, the actual distribution

of alloying elements in the alloy is also related to the heat treatment condition.

Second, according to the characteristics of interaction with C, it can be divided into non-

carbide forming elements, such as Ni, Cu, Si, Al, etc., and carbide forming elements, such as Cr,

Mo, V, etc. Non-carbide forming elements can be easily dissolved in ferrite or austenite, while

carbide forming elements usually exist in carbides. However, when the amount of carbide forming

elements is small, these elements will also dissolve into solid solution or cementite. Only when

the amount of carbide forming elements is more, special carbides are formed.

Third, according to the classification of the influence on the stacking fault energy of

austenite, it can be divided into the elements that can improve the stacking fault energy of austenite,

such as Ni, Cu, C, etc., and the elements that reduce the stacking fault energy of austenite, such as

Mn, Cr, Ru, Ir, etc.

The alloying elements in austenitic stainless steel mainly play the following role: first, tune

the structure of steel, reduce or eliminate its non-uniformity, so as to enhance its stability; second,

strengthen the steel matrix, optimize the mechanical properties, improve cold and hot working

properties; third, improve the corrosion resistance of steel. The following are the specific roles of

various alloying elements in austenitic stainless steel:

(1) The role of Ni: Ni is an excellent corrosion-resistant material, which is usually added

to Fe-C alloy to improve its corrosion resistance. The main role of Ni in stainless steel is to form

and stabilize austenite. However, in order to obtain completely austenite structure in low-C Ni

steel, the Ni content must exceed 24%, and the Ni content must reach 27% to improve the corrosion

8

resistance of steel in some media [4]. When Ni and Cr coexist in steel, the effect of Ni will change

greatly. For example, in ferritic stainless steel with 17% Cr and about 2% Ni, the ferritic steel will

become martensitic steel, and the properties of the steel itself will change greatly. In addition,

when the content of Ni in the steel reaches about 8%, the single-phase austenite structure can be

obtained, that is, the widely used 18-8 austenite stainless steel. Compared with ferritic steel and

martensitic steel, this kind of steel has better corrosion resistance, processability, weldability, low-

temperature plasticity, and impact toughness. Of course, Ni can play a good role in stainless steel

because of the coordination with Cr. Therefore, in order to ensure that Ni plays a better role in

stainless steel, it is necessary to design a reasonable ratio of Cr and Ni. However, due to the high

price and other factors, the current research on Ni in stainless steel is often focused on using other

elements instead of Ni or using other effective processing methods to improve the properties of

stainless steel, so as to reduce the content of Ni in steel.

Yang et al. [13] carried out high-temperature tensile tests on 07Cr17Ni12Mo2N austenitic

stainless steel with different Ni content at four different temperatures from 950 °C to 10 °C using

a Gleeble thermal simulator to study the effect of Ni on the microstructure and high-temperature

tensile properties. The results showed that when the Ni content is reduced from 10.23% to 8.14%,

the microstructure of the test steel is still fully austenite, but the hot plasticity of the steel decreases.

With the increase of temperature, the tensile strength of low Ni steel is greater than high Ni steel

in the temperature range of 950-1050 °C. When the temperature reaches 1100 °C, the tensile

strength of low Ni steel changes in the opposite direction because the high temperature weakens

the pinning effect of N atoms.

Ryo et al. [14] studied the change of tensile properties and strain hardening behavior of 304

stainless steel with a Ni content of 8.3-12%. The tensile test temperature decreased from 25 °C to

9

- 196 °C. At room temperature, hardening and ductility (tensile strength, strain hardening rate, and

elongation) increases with the decrease of Ni content. In the case of steels with 8.3-9.0% Ni, a

lower yield point was observed at temperatures below - 60 °C. The reason is that dynamic strain-

softening or strain-induced plasticity (TRIP) is accompanied by a rapid increase in the amount of

strain-induced martensite (α′) at low strain. For 12% Ni, no dynamic strain softening and TRIP

were observed because martensitic transformation occurred only at low strain.

(2) The role of N: N is a strong alloying element to expand and stabilize austenite structure,

and its effect is 25-30 times that of Ni. The role of N in stainless steel includes the following: (1)

the addition of N can significantly improve the strength and local corrosion resistance of steel,

reduce the precipitation of σ phase, prevent high-temperature brittleness, and render austenite with

good anti-sensitization ability; (2) N can replace part of Ni with Mn to reduce production cost; (3)

N atom can reduce the production cost in austenitic stainless steel, the majority of N is dissolved

in austenite and plays a solid solution strengthening role; (4) when the C content in stainless steel

decreases, the volume fraction of ferrite increases, and the addition of N can reduce the adverse

effect of C reduction on the microstructure. The content and morphology of ferrite in austenitic

stainless steel are also affected by the increase of N content. The increase of N content can reduce

the volume fraction of ferrite and change the ferrite from network and strip to short rod and island,

thus reducing the adverse effect of network ferrite; (5) N can also block the interstitial impurity

clusters by reducing dislocation density in austenite. The strength of austenitic stainless steel is

improved by preventing movement of dislocations [15, 16]. High N content in stainless steel may

lead to porosity and other defects in castings. Therefore, N content should be controlled within a

reasonable range.

The difficulty with the production of N-containing stainless steel is to prevent the overflow

10

of N during the process of cooling and solidification and improve the solubility of N in steel. In

order to solve this problem, researchers tried to use different processes to prepare N bearing

stainless steel. At present, N bearing stainless steel is often produced by adding a N bearing alloy

smelting method, injecting N-containing gas smelting method, pressurized smelting method, and

powder metallurgy method [15]. Ma et al. [17] obtained Cr18Mn18N steel with N content as high

as 1%. When this method is used, bubbles and inhomogeneous microstructure are easy to occur in

ingot due to insufficient pressure. The equipment required for the preparation method is simple

and easy to operate. Gao et al. [18] used AOD Process to blow N into molten steel many times and

smelt 1Cr2Mn15N high N austenitic stainless steel with a N content of 0.56% by adding Cr-nitride

in the later stage. In addition, the N content in the steel can reach 1.2%. This process is more

suitable for large-scale industrial production of N bearing stainless steel than adding N-containing

alloy. A large number of studies have proved that powder metallurgy is a process with great

economic potential for the preparation of N bearing stainless steel, so this method has attracted

more and more attention of researchers and producers [19-21].

In order to improve the high-temperature mechanical properties of low-C medium N 316

stainless steel, Nakazawa et al. [22] studied the effects of Si, Mo, and N on the high-temperature

mechanical properties of 316 stainless steel. The tensile strength of 316 stainless steel at high

temperature was enhanced by all the elements. The effect of N was most obvious, followed by Mo

and Si, and the effect of Si was 1/2 of Mo. The addition of 3% Si, 2% Mo, and 0.8% N can improve

the tensile strength and elongation at the same time. Among the three, N has the most obvious

effect in improving the creep fracture strength. Adding Si with or without Mo and N can slightly

increase the creep fracture strength, while adding Mo or N can significantly improve the fracture

strength. Xue et al. [23] studied the effect of N on the microstructure and properties of nuclear

11

grade 316LN stainless steel. It was found that N can significantly refine the grain size of 316LN

stainless steel. With the increase of N content, the strength and hardness of 316LN stainless steel

was increased linearly, and the strength at - 196 °C was much higher than at room temperature;

the elongation of 316L stainless steel decreased gradually at room temperature, but increased first

and then decreased at - 196 °C. According to the microstructure, the formation and growth of

ferrite phase and other precipitates in steel were strongly hindered by N, which significantly

delayed the occurrence of severe plastic deformation of austenite at - 196 °C, resulting in

deformation-induced martensitic transformation.

(3) The role of Mn: Mn is an element that expands the austenite phase zone and stabilizes

austenite structure, but the effect is not strong, and is only equivalent to 1/2 of Ni [24]. The main

function of Mn in stainless steel is to replace part of Ni with N to save Ni and reduce cost. In

addition, the addition of Mn can improve the solubility of N in stainless steel. However, as far as

Mn itself is concerned, its addition has no beneficial effect in improving the corrosion resistance

of stainless steel. According to the pitting corrosion equivalent (PRE = Cr + 30N + 3.3Mo - Mn),

the addition of Mn reduces the pitting resistance equivalent, that is, the pitting corrosion resistance

of stainless steel [4]. This is because Mn will combine with sulfur in steel to form harmful MnS

inclusions, and MnS is the source of crevice corrosion and pitting corrosion. Moreover because of

good reducibility, Mn can act as a deoxidizer in steel [1]. In addition, the addition of Mn can

promote the precipitation of γ phase and reduce the low-temperature toughness and weldability of

the steel [25]. Jung et al. [26] studied the effect of Mn and Mo on the high-temperature tensile

properties of high Ni austenitic cast steel. 6% Ni in N20 (0.4C-1.2Si-1.0Mn-20Ni-25Cr) austenitic

stainless steel was replaced by 6.9% Mn and 2-4% Mo was added. Generally, when Mn is used

instead of Ni, the stability of austenite decreases, and a small amount of ferrite appears at high

12

temperatures. However, there is no ferrite in N14 steel, and it has similar or superior high-

temperature tensile properties as N20 steel.

(4) The role of C: Carbon is the main component of steel. Since C strongly expands the

austenite phase zone and stabilizes austenite structure, its effect is about 30 times that of Ni. Its

content and distribution in steel determines the structure and properties of steel [4]. For example,

in stainless steel with 17% Cr, when the C content is less than 0.12%, it is ferritic stainless steel

without phase transformation and cannot be strengthened by heat treatment, and the annealing

hardness is less than 20 HB; but when the C content is greater than 0.7%, it is martensitic stainless

steel, and the hardness after quenching and tempering can reach more than 50 HRC. It can be seen

that C plays a key role in stainless steel. Of course, C also has an adverse effect on stainless steel.

Due to the strong affinity between C and Cr in stainless steel, it is very easy to form C-Cr

compounds. The more C content is, the more Cr is bound. This will inevitably reduce the solid

solubility of Cr in the matrix, which will adversely affect the corrosion resistance of the steel.

Especially, when the C-Cr compound precipitates along the grain boundary, it will lead to Cr poor

area in this area and cause intergranular corrosion. With the development of metallurgical

technology, more attention has been devoted to the research of low C, ultra-low C, and other new

steel grades.

(5) The role of Cr: The corrosion resistance of stainless steel is mainly due to the increase

of Cr content and the stability of stainless steel However, the existence of C will lead to the

formation of carbide, resulting in the formation of Cr depleted zone in the matrix, thus reducing

the corrosion resistance of the material.

(6) The role of Al: Given that Al is a ferrite forming element, it is inevitable to produce

some ferrite in Fe-based alloys, which reduced the creep strength. When the temperature is 500-

13

600 °C, the creep strength of ferrite is very low [27]. In some precipitation hardening stainless

steels, Al is often added as a precipitation forming element. According to the phase diagram of Fe-

Ni-Al alloy, with the increase in volume fraction of (Ni, Fe)Al phase, the solid solution temperature

of precipitated phase increases gradually, which makes Fe-Ni-Al alloy to be used as a structural

material at higher temperatures and may replace austenitic steel and wrought Ni-base superalloy.

Some significant findings have been made. For example, Pickering et al. studied NiAl precipitation

strengthened austenitic steel, and alloys containing a small amount of NiAl precipitated phase were

developed, such as 17-7 PH (1Cr17Ni7Al precipitation hardening semi austenitic stainless steel)

and 13-8Mo PH (0Cr13Ni8Mo2Al precipitation hardening martensitic stainless steel) have been

developed and widely used. However, the amount of precipitation phase is less because of low Al

content, and the Ni-Al-B2 Laves phase mainly precipitates in the form of the second phase in the

austenite matrix during aging. These precipitates increase the strength and decrease the toughness

of the alloy at room temperature. However, at 750 °C, the alloy does not exhibit strong

precipitation strengthening through these phases, and the elongation at break is not affected by

aging. The analysis of fracture morphology and cross-section microstructure after the tensile test

indicated that the difference in mechanical properties between room temperature and 750 °C is

because of ductile-brittle transition of B2 precipitate. B2 precipitates are hard and brittle at room

temperature, but they become flexible above the ductile-brittle transition temperature (DBTT).

Therefore, its service temperature is still limited to medium temperature, generally not exceeding

650 °C [28]. The addition of Al can significantly improve the uniform corrosion resistance and

intergranular corrosion resistance of as cast 316L stainless steel, and the mechanical properties do

not decrease [29, 30].

(7) The role of other elements: In addition to the above alloying elements, in order to

14

enhance other physical and chemical properties of stainless steel, other alloying elements are added

in the design and manufacture of stainless steel materials. The addition of Mo in stainless steel can

enhance the passivation effect of stainless steel and improve the corrosion resistance of steel [4];

the addition of Cu in stainless steel can improve the stability of the austenite phase and improve

the corrosion resistance of steel in sulfuric acid, especially when added together with Mo; the

addition of titanium and niobium in stainless steel can combine with C in steel to form carbides,

thus ensuring the presence of Cr in steel. The results show that the elements are stable in the solid

solution, which can effectively improve the intergranular corrosion resistance of steel; the addition

of Si in steel can improve the casting performance, improve the corrosion resistance, intergranular

corrosion resistance and pitting corrosion resistance of the steel in the oxidizing medium; the

hardness of the steel can be improved by adding Co, the thermal strength of the steel can be

improved by adding V, and the processing ability of steel can be improved by adding rare earth

elements [4].

The above is only the basic role of each alloy element in austenitic stainless steel. In order

to achieve the best effect, it is necessary to fully consider each alloy element and the interaction

among the elements when designing the composition of stainless steel.

1.5 MECHANICAL PROPERTIES OF STAINLESS STEELS

Hamda et al. [31] used laser recovery to treat 301LN metastable austenitic stainless steel

with the aim to refine grains and enhance mechanical properties. The results showed that laser

recovery annealing is an effective method to refine and homogenize the austenite grain structure

for 301LN. With the decrease of laser scanning speed from 10.5 m/s to 7.5 m/s, the temperature

increased from 590 °C to 820 °C, and the fully recovered fine-grained austenite structure with an

average grain size of 2 μm was obtained. When the laser scanning speed was greater than 8.5 m/s,

15

the yield strength and tensile strength were significantly increased.

Masyoshi et al. [32] studied the effects of V, Nb, and Ti addition and annealing temperature

on the microstructure and tensile properties of 301L stainless steel. The 301L stainless steel with

0.5% V was annealed at 850 °C for 30 s to obtain a smaller grain size, about 0.9 μm, compared

with other conditions. With the increase of V and Nb content to 0.5% and 0.1% respectively, the

grain size decreased. As the annealing temperature decreased from 1000 to 850 °C, the grain size

also decreased. With the increase of V and Nb content and the decrease of annealing temperature,

σ0.2 increased from 400 MPa to 750 MPa.

Li et al. [33] studied the effect of the cold rolling process on the microstructure and

properties of 301L stainless steel. Their study indicated that when the cold reduction was increased

from 20% to 40%, the content of strain-induced martensite in 301L stainless steel was gradually

increased, the yield strength of the material wase increased from 789 MPa to 1260 MPa, and the

tensile strength also increased from 977 MPa to 1317 MPa. The microhardness was increased by

120 HV. Grain refinement occurred in the material, resulting in fine-grain strengthening. At the

same time, due to martensitic transformation in 301L stainless steel, the tensile strain hardening

index of 301L stainless steel with 20% reduction was higher than 301L stainless steel with 30%

reduction.

Noriyuki et al. [34] studied the effects of tensile test temperature and strain rate on the

tensile properties of metastable 301L austenitic stainless steel. When the temperature was between

123 K and 373 K, the tensile strength increased from 622 MPa to 1560 MPa with the decrease in

temperature. The uniform elongation reached the maximum value of 5.3% at 323 K. Under the

same condition, the volume fraction of stress-induced martensite increased with the decrease of

temperature. The yield strength increased with the increase of strain rate. When the strain rate was

16

less than 100 s-1, the tensile strength gradually decreased, and when the strain was greater than this

value, the tensile strength started to increase. When the true strain exceeded 0.3, the excellent

combination of tensile strength and uniform elongation was obtained when the maximum

transformation rate is less than 3.

Anti et al. [35] studied the effect of austenite stability on the formation of α′- martensite

(DIM) induced by deformation in 301LN Cr-Ni austenitic stainless steel, and the cyclic

deformation behavior of grain refinement structure when the grain size was in the range of 13-0.6

μm under fatigue load. The transitions were recorded by magnetic saturator, electron backscatter

diffraction (EBSD), and X-ray diffraction (XRD) during the cyclic process with a constant total

strain amplitude of 0.4% and 0.6%. The cyclic deformation behavior was influenced by stress

amplitude. The results showed that the stability of austenite increased with the decrease of grain

size to 1 μm at 900 °C for 1 s. On the contrary, when annealed at a lower temperature of 800 -

700 °C, submicron grains were obtained, and the stability of the non-uniform grain structure was

dramatically decreased. In these structures, submicron grains were more stable, and CrN

precipitation reduced the stability of grains with several microns in size in submicron grains. The

volume fraction of martensite is 6% and 23% respectively in the two structures. Under cyclic

loading, the level of initial stress amplitude varied significantly with austenite grain size. At 0.6%

strain amplitude, the initial softening was followed by cyclic hardening. The level of the final stress

amplitude is related to the fraction of deformation-induced martensite formed during cyclic strain.

Eskandari et al. [36] studied the effect of continuous and discontinuous rolling processes

on martensite saturation strain value, austenite grain size, and mechanical properties of 301L

stainless steel. The results indicated that discontinuous rolling increases the volume fraction of

strain-induced martensite and decreases the saturation strain value of martensite. In addition, the

17

final austenite grain size obtained by discontinuous rolling was smaller. The hardness and yield

strength increased to the maximum value of 1970 MPa with the formation of the nanocrystalline

structure during annealing and then decreased to the minimum value of 1545 MPa with grain

growth.

Many studies reported that the mechanical properties of high N Ni-free austenitic stainless

steel are significantly better than ordinary Ni bearing austenitic stainless steel, especially the yield

strength and tensile strength. Some studies have shown that the tensile strength and yield strength

of high N Ni-free austenitic stainless steel are 2-4 times higher than AISI200 and 300 grade

austenitic stainless steels at room temperature [37]. On the other hand, N can decrease the grain

size of stainless steel [38]. Therefore, the addition of N significantly improves the strength of high

N austenite. In addition, studies by Stein et al. show that cold working can enhance the deformation

strengthening effect of high N austenite, and the increase of N content can significantly increase

the work hardening ability of austenitic stainless steel [39]. Another important role of N is to

maintain the stability of austenite during the deformation of high N austenitic stainless steel and

prevent the formation of magnetic martensite, which is of great significance for the application of

high N Ni-free austenitic stainless steel in the medical field [40]. Therefore, N in stainless steel

can improve the mechanical properties of stainless steel through solution strengthening, grain size

strengthening, and deformation hardening [41-43]. In addition, Schino [44] thinks that the fine

grain strengthening effect of high N austenitic stainless steel is significantly stronger than ordinary

304 stainless steel. Wang et al. have shown that N can significantly increase the deformation

hardening rate of austenitic stainless steel [41].

At present, the strengthening and toughening aspects of austenitic stainless steel focused

on the following aspects: (1) grain refinement, (2) using the grain size, phase or chemical

18

composition distribution, dislocation density to form a gradient structure, and (3) using multi-

phase or different single-phase structure to form a mixed structure.

(1) Grain refinement: When the grain boundary interacts with the dislocation, the

dislocation movement is blocked at the grain boundary, resulting in dislocation accumulation,

which improves the strength of the steel. The smaller the grain size is, the greater the blocking

effect of dislocations. The relationship between grain size and yield strength conforms to Hall-

Petch equation:

�� = �� + ��/� (1.3)

where, σ - yield strength, σ0 - lattice friction force, Ky constant, and d-grain diameter.

Cold rolling + annealing treatment is often used to refine austenitic stainless steel. This

method iinvolves cold rolling followed by annealing at high temperature such that the structure of

stainless steel will recrystallize and transform into the nanocrystalline or ultrafine grain. Eskandari

et al. [45,46] obtained homogeneous 301 austenite structure by cold rolling and annealing, and the

grain size was about 70 nm. The yield strength of the refined steel approached 1970 MPa, which

was 1.3 times of steel with grain size of 380 nm.

Although grain refinement can improve the yield strength of stainless steel, it will decrease

the plastic deformation ability. When the annealing time is same, the grain size of the sample

decreases with the annealing temperature from 900 °C to 700 °C. The stress-strain curves of

samples with different grain sizes showed that the yield strength increases with the decrease of

grain size, but the plasticity decreases.

Misra et al. [47] studied the evolution law of austenite microstructure in 301 stainless steel

during grain refinement, and the obtained grain size range was 200-500 nm. The results suggested

that austenite recovery consists of three stages: (a) strain-induced transformation of α'- martensite

19

into lath austenite, (b) formation of dislocation cells and new austenite subgrains, and (c) fine-

grain structure formed by the combination of subgrains. The deformation mechanism indicated

that the grain process belongs to shear type recovery, which is different from the diffusion type

recovery mechanism of 301LN stainless steel.

(2) Gradient structure: Wei et al. [48] introduced gradient twins into twinning induced

plastic steel by torsion, which significantly improved the strength of the sample and did not reduce

the plasticity. The experimental results confirmed that when the Fe-Mn-C steel was twisted to 360 °

the yield strength was doubled, but the elongation did not decrease. The microstructure analysis

showed that the twin density increases gradually from the inside to the surface of the sample. The

study indicated that the gradient structure of nano twin increase the strength. At the same time, the

interaction between the twist deformation twins and the twins nucleated in the next deformation,

and dislocation made the steel exhibit good plasticity.

The mechanical properties of high strength and high plasticity can be obtained by surface

nanocrystallization, which can improve the strength and maintain the plasticity of the matrix.

Surface mechanical attrition treatment (SMAT) and surface mechanical grinding treatment (SMGT)

are the two most commonly used methods to realize surface nanocrystallization. For SMAT, the

energy generated by the vibrator is transferred to the surface of the sample by a high-speed and

high-frequency moving ball in the cavity. Under the action of the high strain rate, a large number

of twins and dislocations are generated. When the defects on the surface of the material reach the

limit value, the defects interact with each other to form subgrain boundaries and new grain

boundaries. Finally, the grain size on the surface of the sample can be refined to nanometer level

by reciprocating cycle. Martensitic transformation occurs in austenitic stainless steel at a high

strain rate, and the refining mechanism is different. In addition, the strain rate also presents a

20

gradient change on the surface and the center, which eventually leads to the microstructure with a

gradient change.

Chen et al. [49] studied the microstructure and corresponding mechanical properties of 304

austenitic stainless steel after SMAT treatment at different strain rates. The research results showed

that twinning is the main deformation mechanism of stainless steel at high strain rate, resulting in

a large amount of ε-martensite and twins and a small amount of α-martensite. At low strain rate,

martensitic transformation and dislocation slip were the main deformation mechanisms. Thus, low

strain rate was more likely to induce martensitic transformation.

(3) Mixed structure: Milad et al. [50] studied the effect of different cold rolling processes

on the microstructure and mechanical properties of austenitic 304 stainless steel. After cold rolling,

α-martensite existed in the microstructure of AISI304 stainless steel. The results showed that with

the increase of rolling deformation, the yield strength increased about 5 times, from 258 MPa to

1260 MPa and the elongation at fracture decreased from 75% to 10%. The the strain-induced α-

martensite led to the increase of yield strength of cold-rolled samples.

Fahr et al. [51] carried out studies on 304 stainless steel myocardial infarction by warm

rolling technology. The results show that on rolling at 450 °C, with the increase of rolling

deformation, the yield strength was first increased and then decreased. When the deformation was

60%, the yield strength of the sample reached 1100 MPa, and the elongation was ~ 30%, showing

excellent combination of strength and plasticity. They believed that the rolling temperature and

deformation affected the mechanical properties by changing the stability of the austenite matrix.

On rolling at high temperature, the increase of deformation will make the sample stay at high

temperature for a longer time, and the amount of precipitated phase in the microstructure will also

increase, but the austenite is more stable under large deformation. The first increase and then

21

decrease in yield strength was the result of competition between the above two reasons.

In addition, martensite deformation can also be used to strengthen and toughen amorphous

alloys. Wu et al. [52] obtained amorphous/microcrystalline composite samples with excellent

plastic deformation ability using this method. The yield strength of the metallic glass was 1650

MPa at room temperature, and the plasticity was ~ 10%. In the amorphous alloy, martensitic

transformation occurs in a small number of microcrystalline particles.

1.6 CORROSION RESISTANCE OF STAINLESS STEELS

In the 1980s, Ramakrishnan et al. studied the oxidation properties of high Al austenitic

stainless steel. The cyclic oxidation behavior of austenitic stainless steel (24% Ni, 10% Cr, 5% Al)

at 800-1300 °C was studied. The results showed that Fe-Ni-Cr-Al stainless steel exhibited excellent

oxidation resistance at up to 1300 °C, which was mainly due to the formation of α-Al2O3 film [53].

At the same time, Takashi et al. from Nippon Steel also studied the oxidation resistance and heat

corrosion resistance of high Al austenitic stainless steel.

Compared with conventional Al-800 stainless steel, a new type of austenitic stainless steel

(AFA) was developed. In austenitic stainless steels, they found that only 2.5-3wt.% Al was added

to form a protective film of Al2O3. In the laboratory, no cracking phenomenon was found in the

AFA alloy produced by arc casting in a small batch during cold rolling with a reduction of 50-70%,

indicating that the alloy has good machinability. The main difficulty facing AFA alloy was to

further enhance the forming ability of Al protective film without reducing the mechanical

properties and welding properties of the materials. Up to now, AFA alloys have only been prepared

in a small amount by arc casting in the laboratory [54].

The State Key Laboratory of new metal materials, University of Science and Technology

Beijing studied the effect of Al formation on high-temperature oxidation resistance and strength

22

of austenitic stainless steel. A new type of austenitic stainless steel with Al formation was

developed by adding 3.0wt.% Al into the Fe-25Ni-18Cr alloy, which greatly improved the high-

temperature oxidation resistance and strength of the alloy. At 800 °C, continuous, stable, and

unique Al film was formed either in dry air or in the air mixed with 10% water vapor. The long-

term high-temperature oxidation performance was enhanced, which is related to the high-density

B2-NiAl precipitates in the surface layer of Al2O3. In addition, when tested in dry air at 750 °C,

the new steel exhibited tensile yield strength and fracture strength of 310-335 MPa and 480-500

MPa, respectively [55].

The high-temperature oxidation resistance of Fe-Cr-Ni-Al alloy containing Al and Si was

studied. When the content of Al was less than 5wt. %, a duplex heat-resistant steel with austenite

(about 95%) and ferrite (no more than 8%) was formed. The thermodynamic properties,

microstructure, phase transformation, high-temperature performance, and oxidation resistance of

the alloy were studied. The high-temperature oxidation resistance of Al was better than Cr (the

oxidation resistance temperature of Al2O3 was close to 1200 °C), and the high-temperature strength

and castability of the alloy were also improved. Weldability and machinability were improved,

even at 1200-1300 °C, and good oxidation resistance and mechanical properties were obtained

[56]. Wang et al. developed Fe-Cr-Ni-Al (composition: 0.08-0.12C, 20-27Cr, 8-12Ni, 3-4Al) heat-

resistant alloy through mechanical properties, high-temperature oxidation resistance, and saving

Ni resources. The alloy has excellent castability and weldability and has good oxidation resistance

and mechanical properties at 1200-1300 °C, which has certain development and application

meaning [57].

It has been shown that N in stainless steel can improve the ability of uniform corrosion

resistance in a specific solution system. This is mainly because N can be enriched at the

23

metal/passive film interface to form Cr nitride, which avoids the dissolution of Cr, improves the

performance of the passive film, and makes it more compact [58].

Compared with total corrosion, N can improve the local corrosion ability of stainless steel

more significantly, especially pitting corrosion and intergranular corrosion. There is no unified

view on the mechanism that N can improve the local corrosion resistance of stainless steel, but

there are several viewpoints [59-65]: (a) acid consumption theory, (b) surface enrichment theory,

and (c) synergistic effect with Cr and Mo.

(a) Acid consumption theory: acid consumption theory is also called ammonia formation

theory. In the early stage of pitting corrosion, the hydrolysis of metal ions in the pitting pit will

form H+, which will reduce the pH value in the pitting pit, and accelerates the dissolution of metal

ions and renders the pitting process "autocatalytic". Lu [59] and Bandy [60] believe that N in the

alloy will react with H+ and increase the pH in the pitting pit:

[N]+4H++3e-→NH4+ (1.4)

This process can effectively alleviate the local acidification in the pitting pit, inhibits the

anodic dissolution process in the pitting process, promotes the re-passivation of stainless steel, and

improve its pitting resistance.

(2) Surface enrichment theory: Lu et al. [61] analyzed the surface of N-containing

stainless steel by Auger electron and photoelectron and found that N was enriched at the

metal/passive film interface. Bandy et al. [60, 62] considered that N can be adsorbed on the oxide

layer of metal and enhance its passivation ability. The research results of Grabk et al. [63] showed

that N can affect the kinetics of re-passivation process and accelerate the re-passivation process,

thus inhibiting the further growth of pits.

(3) Synergistic effect with Cr and Mo: Lu et al. [59, 64,65] noted that N on the subsurface

24

of the passivation film can form nitrides with Cr and Mo to achieve enrichment, inhibit the

dissolution of these two elements, strengthen the corrosion resistance in austenitic stainless steel,

and make the passive film more compact and stable.

Lee et al. [66] studied the corrosion resistance of passive film of 316L and 316LN alloy

with different N-content. In 0.1 mol/L NaCl solution, N in the alloy can significantly increase the

pitting potential, and the fluctuation of current in the metastable pitting stage is significantly

reduced. With the increase of N-content, the pitting sensitivity of the alloy decreases gradually.

Peng et al. [67] studied the biological corrosion behavior of high N steel and 317L stainless steel

with different N-content in phosphate buffer solution. With the increase of N-content, the

impedance value of alloy increases. Moreover, the author also found that N can be alloyed with

superior biocompatibility, which is more suitable for implant devices.

Super ferritic stainless steel has excellent corrosion resistance, and its corrosion resistance

is much better than conventional ferritic stainless steel. The corrosion resistance of high-grade

ferritic stainless steel is equivalent to super austenitic stainless steel and Ni-based corrosion

resistant alloy [68-74].

(1) Uniform corrosion: Table 1.1 shows uniform corrosion resistance of super ferritic

stainless steel. Compared with austenitic stainless steel, Ni-based corrosion resistant alloy, Ti plate,

and other materials, the super ferrite stainless steel shows excellent corrosion resistance in various

acid media, and the annual corrosion rate reached 0.2 mm/yr and below [75]. It can be seen that

the corrosion rate of stainless steel increases with the increase of sulfuric acid concentration and

temperature in medium concentration sulfuric acid solution. Under identical test conditions, 446

stainless steel has excellent uniform corrosion resistance to sulfuric acid, and its corrosion

resistance becomes more and more superior with the change of concentration, followed by 904L

25

stainless steel. Many researches studied super ferritic stainless steel, including isocorrosion curves

of several materials in high temperature concentrated sulfuric acid [76], the isocorrosion curve [5]

of high purity Cr30Mo2 steel in sulfuric acid, with 0.1g / (M2×h) as the boundary, the corrosion

curve of 25-4-4 in sulfuric acid [5], and the corrosion rate of ultra-low C super ferritic stainless

steel (Cr28-30%, Mo3.6-4.2%, C ≤ 0.03%) in high temperature concentrated sulfuric acid [77].

Hydrochloric acid is one of the most corrosive acid. Stainless steel has corrosion resistance in

dilute hydrochloric acid solution with low concentration at room temperature. Based on the test

results of 446 super ferritic stainless steel and 904L super austenitic stainless steel in 4%

hydrochloric acid solution at different temperatures for 24 h [78], it is not difficult to see that in 4%

dilute hydrochloric acid solution, the corrosion rate of 904L stainless steel gradually increased

with the increase of temperature, while the corrosion rate of 446 stainless steel had no obvious

change with the increase of temperature. The results showed that in 4% dilute hydrochloric acid

solution, the corrosion resistance of 446 stainless steel was the best, while that of 904L stainless

steel was poor.

Table 1.1: Uniform corrosion resistance of different grades stainless steels

Materials

Corrosion rate in boiling medium

65%

HNO3

50%H2SO4+

Fe2(SO4)3

45%

HCOOH

20%

CH3COOH

10%

H2SO4

1%

HCl

AISI304

0Cr18Ni9 0.2 0.6 44 0.1 400 81

AISI316

0Cr17Ni12Mo2 0.3 0.6 13 0.1 22 71

Carpenter 0.2 0.2 0.2 0.1 1.1 0

26

20Nb3Cr29Ni34Mo2Cu3Nb

Hastelloy C

0Cr16Ni60Mo16W4 11.4 6.1 0.1 0 0.4 0.3

Ti 0.3 5.9 22 0 160 5.6

DIN 1.4575

00Cr28Ni4Mo2Nb 0.2 0.3 0.1 0 0.2 0

High cleanness

Cr29Mo4 0.1 0.2 0.1 0 - 0.2

S44800

00Cr29Mo4Ni2 0.1 0.2 0.1 0 0.2 0.2

S44660

00Cr27Mo4Ni2TiNb 0.1 - - - <0.1 <0.1

(2) Pitting: The excellent pitting resistance is the outstanding advantage of ferritic stainless

steel, which can be confirmed by pitting equivalent of PRE = Cr + 3.3Mo. Generally, the pitting

index of super ferritic stainless steel is 35 or above. Pitting is characterized by pitting potential and

critical pitting temperature CPT. Table 1.2 shows the critical pitting potential of super ferritic

stainless steel in 3.5% NaCl solution [75]. It can be seen that at 80 °C, the pitting potential of

S44660 steel was still greater than 600 mV, and that of S44635 steel was still greater than 800 mV,

reaching or exceeding the level of 254 alloys. Tables 1.3 and 1.4 show the pitting corrosion

performance comparison of super ferritic stainless steel, super austenitic stainless steel, and

corrosion-resistant alloy. Table 1.3 shows the pitting potential in 5% NaCl at different temperatures.

The higher the pitting potential, the better the pitting resistance. Table 1.4 shows the critical pitting

temperature in 10% FeCl3-6H2O solution. The higher the critical pitting temperature, the superior

27

is the pitting resistance.

Table 1.2: Critical pitting corrosion potential of super ferritic stainless steels at 3.5%NaCl, pH6.5

Chemical

composition Grade

Critical pitting corrosion potential/mV (SHE)

60 °C 80 °C 100 °C

00Cr28Ni4Mo2Nb Remanit4575 - 625 400

00Cr26Ni2Mo3Ti S44660 965 640 380

00Cr25Ni4Mo4Ti monitS44635 - 820 480

00Cr20Ni25Mo45Cu 254SLX 920 750 506

Table 1.3: Pitting corrosion potential of different materials

Material Pitting corrosion potential/mV (SCE)

60 °C 80 °C

AISI316 125 35

Alloy 825 320 190

904L 515 290

25-4-4 950 685

Table 1.4: Critical pitting temperature of different materials

Material Critical pitting temperature/°C

AISI316 15

Alloy 825 29

904L 42

25-4-4 55

(3) Crevice corrosion: The crevice corrosion of stainless steel is mainly caused by the

solution acidification and anodic reaction in the crevice, which results in the destruction of the

passive film on the surface. Therefore, improving the stability of stainless steel passivation film

28

and passivation, re-passivation ability is also an important measure to improve the ability of

stainless steel to resist crevice corrosion. Therefore, some measures to select pitting corrosion-

resistant materials are also applicable to the selection of crevice corrosion-resistant materials.

Table 1.5 shows the critical crevice corrosion temperature of super ferritic stainless steel [75]. It

can be seen that the critical crevice corrosion temperature of super ferritic stainless steel is

generally above 45 °C, which is greater than other materials.

(4) Stress corrosion: Table 1.6 shows the stress corrosion resistance of super ferritic

stainless steel in various stress corrosion tests [75]. It can be seen that the super ferrite stainless

steel has good stress corrosion resistance except boiling 45% MgCl2 solution. It can be seen from

Table 1.7 that S44635 stainless steel does not undergo stress corrosion cracking in 40% CaCl2

solution at 100 ° C after 5000 h.

Table 1.5: Critical crevice corrosion temperature of different materials

Material Critical crevice corrosion temperature / °C

SANTRON FeCl3

29(00Cr29Mo4Ti) 90 >55

Monit(00Cr25Mo4Ni4Ti) 67.5 47

254SMO(00Cr20Ni18Mo6N) 62.5 46

0Cr20Ni15Mo5Mn5CuN 62.5 33

Ferraliccm255

(00Cr25Ni6Mo3Cu2N) 60 37

SC-

1(S44660)(00Cr25Mo4Ni2NbTi) 60 45

AL-6X(00Cr20Ni24Mo6) 57.5 37

29

JS-700

(00Cr20Ni25Mo4.5CuNb) 45 31

00Cr20Ni25Mo5Cu 42.5 22

AISI316 - <15

Alloy825 (0Cr21Ni42Mo3Cu21) - <15

904L (00Cr20Ni25Mo4Cu) - 20

Table 1.6: Stress corrosion cracking resistance of different materials

Material Boiling 45 %

MgCl2

Boiling 26 % NaCl

pH7 Boiling LiCl2

304(0Cr18Ni9) <3h cracking 72 h cracking cracking

316(0Cr17Ni12Mo2) - cracking -

-6X(00Cr20Ni06) cracking uncracking -

254SMO(00Cr20Ni06CuN) cracking uncracking -

DIN 1.4575

(00Cr28Ni4Mo2Ti) cracking uncracking -

S44660 (00Cr27Mo3Ni2Ti) cracking uncracking uncracking

Al29- (00Cr29Mo4Ti) uncracking uncracking -

S44800 (00Cr29Ni4Mo2) cracking uncracking -

Monit S44635

(00Cr25Mo4Ni4Ti) cracking uncracking uncracking

Table 1.7: Stress corrosion cracking resistance of different materials at 100 °C 40% CaCl2

Material Cracking time / h

AISI316 200-500

30

904L (00Cr20Ni25Mo4Cu) >2000 uncracking

Monit S44635 (00Cr25Mo4Ni4Ti) >5000 uncracking

1.7 DEFORMATION BEHAVIOR OF STAINLESS STEELS

It is generally believed that the work hardening of α-martensite is the main reason to

improve the strength of austenitic stainless steel. α-martensite contains high-density of dislocations

after deformation. Stainless steel is usually regarded as austenite with high plasticity after

deformation, and dispersed hard phase α-martensite is distributed on the matrix [79,80]. This

structure is very beneficial to improve the strength and plasticity of materials.

Eckstein et al. [81] carried out morphological studies via microscopy and corresponding

mechanical property tests of γ-martensite composite structure. It is shown that the main

strengthening mechanism is to prevent the movement of dislocations, the dispersed martensite

phase mainly plays a strengthening role, and the plastic deformation is mainly concentrated in

austenite.

Narutani [82] studies showed that the flow stress had a linear relationship with the square

root of dislocation density during deformation of 301 stainless steel. If there is α-martensite in the

material, the above relationship is also true as long as the content is less than 20%. It was believed

that α-martensite is produced during transformation, resulting in the expansion of martensite. In

order to control the two-phase structure and promote more slip systems to operate in austenite, a

large number of dislocations slide, which increases the dislocation density in stainless steel and

strengthens the material.

It can be seen that the transformation temperature, time, and degree in austenitic stainless

steel can be controlled by various process methods or changing process parameters, such as

changing deformation temperature, deformation rate, heat treatment temperature, or time and

31

deformation degree, so as to obtain varied microstructure and/or manage size, in order to obtain

high strength and high plastic properties of austenitic stainless steel.

The matrix structure of austenitic steel is usually metastable at room temperature. Therefore,

there are many deformation mechanisms. The main influencing factors are composition, lath size,

and plastic deformation conditions. Austenitic stainless steel is a kind of austenitic steel, and its

deformation mechanism determines its strain hardening ability and plastic deformation ability.

Therefore, in order to obtain austenitic stainless steel with excellent mechanical properties, it is

very important to study its deformation mechanism. The most important deformation mechanisms

of austenitic stainless steel are dislocation slip, twin-induced plasticity, and phase transformation

induction plasticity mechanisms [83].

(1) Dislocation slip mechanism: Dislocation slip is the main deformation mechanism of a

crystal. When the shear stress on the slip surface in the crystal is greater than the critical slip stress,

one part of the crystal moves relative to the other part along a certain slip plane and direction,

resulting in plastic deformation of the crystal. With the increase of stress, a large number of

dislocations will move until they encounter obstacles, such as other dislocations, second phase

particles, or grain boundaries, so as to strengthen the material. On the other hand, the yield strength

of materials is also closely related to the movement of dislocations. The yield strength of materials

mainly depends on whether the stress concentration generated by the dislocations near the grain

boundary of the sliding grains can be energized, and the dislocation sources in the sliding system

of adjacent grains can be activated, so as to promote coordinated multiple slips. When the applied

stress and other conditions are same, the number of dislocations is directly proportional to the

distance from the grain boundary to the fault source. The larger is the grain size, the greater is the

distance. Therefore, more dislocations are accumulated, and once the stress concentration exceed

32

the yield strength of the material, plastic deformation of adjacent grains occurs. For fine grains,

the number of dislocations near the grain boundary of the sliding grains is small, so the stress

concentration is small. In order to make the adjacent grains plastic deformation, more external

stress should be applied. This is the reason why grain refinement improves yield strength.

(2) Twining induced plasticity (TWIP): Twinning is another mode of plastic deformation.

When twin deformation occurs, a part of the crystal will shear uniformly with respect to another

part of the crystal along a certain mirror surface (twin plane) and a certain crystal direction (twin

direction) under the action of shear stress. This shear will not change the lattice structure of the

crystal but will change the relationship between the deformed part and the matrix. In twinning,

a part of atomic lattice is deformed and forms mirror image of lattice next to it. Generally speaking,

two parts of symmetrical grains are called twins, and the process of forming twins is twinning. For

the face-centered cubic metal of austenitic stainless steel, the twin plane is {111}, and the twin

direction is "112". The displacement of atoms in each layer (Bi, CJ, DK, etc.) is directly

proportional to the distance from the twin plane. Similar to slip, twinning can occur only when the

shear stress in the twin direction reaches the critical value.

Twinning induced plasticity (TWIP) mechanism can significantly improve the strength and

plasticity of materials. Twinning induced plasticity steel (TWIP), a typical representative of the

second generation advanced high strength steel, is a kind of high strength steel developed by the

TWIP mechanism. In recent years, the research on the role of TWIP mechanism in austenitic

stainless steel has been given strong attention. Wu et al. [84] studied 316L austenite micro/nano

composite structure, it was found that in the process of tensile deformation, dislocations in the

micron grain first slipped and then piled up at the grain boundary, then deformation twins were

produced. However, the stacking fault energy (SFE) increased rapidly in the nanograins, and the

33

twinning does not occur, which is beneficial to the structural strength. Wittig et al. [85] studied the

effect of temperature on the deformation mechanism of Fe-16.5Cr-8Mn-3Ni austenitic stainless

steel. It was found that the main deformation structure was α'-martensite, accounting for about

90%, and the deformation mechanism was transformation induced plasticity; at room temperature,

the deformation mechanism was composed of TRIP effect/TWIP effect and dislocation slip; when

the tensile temperature was increased to 200 °C, the deformation mechanism was composed of

TRIP effect / TWIP effect and dislocation slip, dislocation slip is the main deformation mechanism.

Misra et al. [86,87] found that when the grain size is refined from coarse grained to ultra-fine

grained/nanocrystalline, the deformation mechanism of stainless steel changes from TRIP to TWIP,

which is the reason why ultra-fine grains/nanocrystalline austenitic stainless steel can maintain

good plasticity while improving strength.

(3) Transformation induced plasticity (TRIP): Phase transformation induced plasticity

is the strengthening and toughening mechanism of high strength and high toughness steel plate

produced by the automobile industry in recent years. The principle is that when the austenite is

deformed under stress, the austenite will be transformed into martensite at the place where the

strain is concentrated. Due to the high hardness of martensite, the local hardness is improved, and

it becomes difficult to continue deformation. The deformation transfers to the surrounding

structure and the occurrence of necking is delayed. With the continuous development of the strain,

the material obtains higher plasticity [88]. Generally speaking, there are two nucleation modes of

the austenite phase to α'- martensite phase in austenitic stainless steel, one is an austenitic phase

(γ) → ε martensite → α' - martensite; the other is γ phase → α' phase. The nucleation of α'-

martensite is generally believed to increase through repeated nucleation and coalescence at the

junction of shear zones. The structure of the shear band depends on the overlapping process. If the

34

stacking faults of γ parent phase regularly overlap on each {111} plane, ε-martensite with dense

hexagonal crystal structure (HCP) is formed. In fact, two different nucleation methods can occur

independently or simultaneously. Huang et al. [89] studied the nucleation mechanism of martensite

in Fe-18Cr-12Ni austenitic stainless steel during plastic deformation. It was found that during

plastic deformation, γ → ε, γ → ε → α' and γ → α' can occur, and two martensite formation modes,

namely, stress- assisted and strain-induced, can occur. Chen et al. [90] studied the effect of

temperature on martensite transformation in SU304 austenitic stainless steel during the tensile

process, and found a similar behavior. The results show that there is little martensite at room

temperature, and the increase of martensite volume was not obvious with the increase of tensile

strain; however, at - 60 °C, martensite was formed even under very small strain, and with the

increase of strain, martensite was formed at room temperature, martensite mainly formed near the

triple node of the grain boundary or inside the twin. At - 60 °C, the martensite mainly formed near

the austenite grain boundary.

The deformation mechanisms of austenitic stainless steel, dislocation slip, TWIP and TRIP

play an important role in improving the strength and plasticity of austenitic stainless steel, and

many external factors can affect it. Therefore, in order to obtain high strength and high plasticity

austenitic stainless steel, it is particularly important to study these factors and their internal

relationship.

1.8 STACKING FAULT ENERGY OF STAINLESS STEELS AND INFLUE NCE OF STACKING FAULT

ENERGY ON DEFORMATION BEHAVIOR

The study of work hardening behavior of austenitic stainless steel is an important research

direction. During cold deformation, the plastic deformation of austenitic stainless steel occurs in

different ways, which mainly depends on chemical composition and deformation temperature of

35

stainless steel, but also related to stacking fault energy and other factors. It was not until 1940 that

strain-induced martensite transformation occurred in 301 austenitic stainless steel at room

temperature, and the formation of deformation-induced martensite was related to dislocation,

which made the deformation mode more complicated. During the process of plastic deformation,

the deformation-induced martensitic transformation had a significant impact on the mechanical

properties of 301L metastable austenitic stainless steel, such as improving strength, toughness, and

formability, which rendered 301L stainless steel having wide applications. As mentioned above,

the decisive influence of alloying elements on stacking fault energy and austenite stability is

mentioned. Ludwigson et al. [91] believed that C and N have a significant influence on the

plasticity of 301 stainless steel, and were added as interstitial elements to austenite and played a

strengthening role when austenite transformed to austenite. They considered austenite hardening,

the content and strength of α′-martensite, and established a relationship between tensile flow stress

and strain and suggested that the increased strength after plastic deformation was directly related

to the formation of high strength α′-martensite. In other words, the formation of martensite leads

to increase of dislocation density and strengthens austenite. Llewlyn et al. [92] showed that when

the stacking fault energy was less than or equal to 20 MJ/m2, stainless steel is prone to deformation-

induced martensitic transformation, that is, transformation induced plasticity effect. Olson [93]

found that a large number of dislocations were introduced during the transformation of γ austenite

to α′-martensite induced by deformation, and α′-martensite hindered dislocation slip, resulting in

a large number of dislocations piling up in austenite, increasing dislocation density and producing

work hardening.

The good plasticity of austenitic stainless steel is related to the FC structure of austenite

and low stacking fault energy. There is almost no lattice distortion when a stacking fault is formed

36

in materials, but it will destroy the normal integrity and periodicity of the crystal, and leads to an

abnormal diffraction effect of electrons, which increases the energy of the crystal. This part of the

energy is usually called stacking fault energy. The greater is the stacking fault energy, the lower is

the probability of stacking faults. Stacking fault energy is an important parameter to measure the

plastic deformation process. It determines the difficulty of material to experience slip, affects the

nucleation, movement, and decomposition of dislocations, and affects the strength and creep of

materials. The deformation stability of austenite in metastable austenitic stainless steel plays an

important role in the plasticity and formability of stainless steel. The martensitic transformation

temperature Ms and the strain-induced martensitic transformation temperature Md30 are important

parameters affecting the stability of austenite. The lower the two temperature, the higher the

stability of austenite. The deformation-induced martensitic transformation temperature Md30 is the

temperature at which 30% plastic strain occurs, resulting in 50% deformation-induced martensite

phase transformation [24]. Specifically, when the material is in a certain temperature range above

the Ms point, plastic deformation of stainless steel will cause martensitic transformation of

austenite at this temperature, that is to say, plastic deformation causes Ms point to increase. This

transformation of martensite due to deformation is called deformation-induced martensite

transformation, and the obtained martensite is called deformation-induced martensite. The strain

in the forming process provides the mechanical driving force for martensitic transformation and

provided the required chemical driving force.

1.9 SUMMARY

In this chapter, we introduced stainless steels and their recent development. Based on the

above discussion, we proposed a strategy in my research of obtaining high strength and high

ductility in austenitic stainless steel (a) by utilizing different cold reduction and annealing

37

parameters to obtain nano-grained or ultra-fine grained structure to achieve good combination of

mechanical properties and (b) establish the relationship between phase reversion parameters,

mechanical properties and microstructures and understand the strengthening mechanisms and

deformation behavior of austenitic stainless steel. Furthermore, as the aim of engineering

application, the wear performance of austenitic stainless steel was studied.

38

Chapter 2: Experimental Procedure

2.1 PHASE REVERSION

Grain refinement is an effective method to improve the strength (Hall-Petch relationship)

and fatigue properties of metals and alloys, especially in the high cycle fatigue (HCF) region. The

grain size (GS) of commercial austenitic stainless steels (ASSs) is usually greater than 10 µm.

Although the effect of GS on the strength of ASSs is not as great as that of ferritic steel, a large

number of studies indicated that the refinement of GS can significantly improve the yield strength

(YS) of austenite [94–111]. The conventional hot rolling process or cold rolling together with

recrystallization annealing cannot effectively refine the GS of austenite phase, although the GS of

2 µm can be obtained by warm rolling and annealing [112]. However, in the past 30 years, a large

number of studies have found that the phase reversion treatment of deformation induced α′-

martensite (DIM) and transformation back to austenite can refine GS to submicron size, thus

making the perfect combination of YS and tensile elongation (TE).

The martensite-austenite reversion process in austenitic Cr-Ni steels was studied in 1970s

and 1980s [113–118], and a comprehensive study was carried out in Japan in the early 1990s [119–

121] and later in many countries [45, 47, 86, 94–105,112,122–176]. In laboratory scale research,

the process has been applied to several commercial Cr-Mn and Cr-Ni steels such as 201, 201L,

204Cu, 301, 301LN, 304 and 304L.

Since the phase reversion process is carried out simply by cold rolling and annealing, it

seems to be more suitable for bulk production of large-size sheets and has practical application

potential than many other severe plastic deformation techniques used for GS refinement.

However, despite extensive academic research work, there only exist a couple of industrial

companies utilizing the reversion-treatment for ASSs. In Japan, Nano grains Co. Ltd

39

(Komatsuseiki Kosakusho Co., Ltd.) produces grain-refined 304, 316 and 301 foils (thickness

range 80–300 µm), with GS finer than 1 µm, using repeated reversion treatment [177]. Especially

the enhanced properties of micro-scale cutting and hole piercing were utilized in the manufacturing

of orifices for electronic fuel injection [178–180]. Nippon Steel & Sumitomo Metal company lists

in its product catalogue fine-grained 304 (SUS304 BA19) and 301L (NSSMC-NAR-301L BA1)

grades: thin sheets, strips and foils (0.08–0.6 mm) having both high strength, ductility and

formability and smooth formed surface due to refinement of GS [181,182]. The feasibility of the

process for an industrial manufacturing using a continuous annealing line has been demonstrated

in one laboratory study [96]. In recent studies by Järvenpää et al. [166], a pilot induction heating

line has been employed to simulate industrial conditions in reversion annealing of the 301LN grade.

The first stage of reversion treatment is cold rolling of ASS sheet to obtain DIM, which can

be reverted to fine-grained austenite during subsequent annealing stage. In the following

paragraphs, we point out the different stability of the austenite in different ASSs towards

transformation to martensite, depending mainly on the chemical composition of steel. Some

commercial ASS grades are compared. The degree of cold rolling reduction is a factor affecting

the fraction of DIM, and in industrial rolling, it cannot be very high. Therefore, only partial

transformation of austenite to DIM can happen affecting the microstructure obtained in the

annealing stage. Further, the DIM forms gradually to different degrees, which has an influence on

microstructure heterogeneity and thereby on mechanical properties.

The reversion heat treatment for ASSs consisting of cold rolling and annealing stages is

schematically presented in Fig. 2.1. A metastable ASS must first be cold worked to transform

austenite to DIM. There is some times hexagonal ε-martensite formed as well, but its fraction and

contribution are minor in the reversion process, so that much attention has not been paid on that

40

phase. Highly-deformed cell-type DIM is preferred for large number of nucleation sites for new

austenite grains to attain the desired highly-refined GS in the subsequent reversion annealing [119–

121]. If the total cold-rolling reduction is small, the transformation of austenite to DIM tends to

remain partial and coarse-grained deformed austenite (DA) grains are retained in the structure.

Furthermore, the lath-type, slightly deformed DIM reverts to austenite with coarser GS.

Accordingly, very high cold rolling reduction of 90% to 95% were recommended originally and

applied in numerous studies, e.g., [98,99,101,103,104,120,121], which can, however, be

impractical in the industry.

Figure 2.1: Illustration of phase reversion process for a metastable austenitic stainless-steel that

includes cold rolling and annealing process. [134, 137] During annealing stage of cold-rolled ASS sheet containing DIM, the reversion of DIM

back to austenite can take place, refining the GS and enhancing the mechanical properties. There

are two reversion mechanisms, shear and diffusional, the type depends on the chemical

composition of the steel, heating rate and annealing temperature but hardly on the degree of cold

rolling reduction. In this part, we first discuss the type and kinetics of the reversion in various Cr-

41

Ni type ASSs and the factors affecting them. This section is followed by discussion on the

temperature range suitable for the reversion treatment, accounting for the different reversion

mechanism. It can be noticed that certain typical temperatures (600–1000 °C) have been applied

in experiments for commercial Cr-Ni and Cr-Mn ASSs, and the duration of annealing can be

selected to be very short (less than 1 s) or even hours, highlighting the flexibility of the process.

As regards the reversion mechanism in metastable ASSs, DIM reverts back to austenite

during continuous heating or isothermal annealing at an appropriate temperature. Reversion of

martensite to austenite is a phenomenon also taking place during intercritical annealing in low

carbon martensitic stainless steels [183] and medium-Mn steels, e.g., [184]. A variety of

experimental methods have been employed to study both martensitic transformation and DIM

reversion in ASSs, e.g., microscopic methods, dilatometry, calorimetry, X-ray diffraction (both

postmortem and in situ), internal friction, various magnetic, positron annihilation or

hardness/mechanical properties measurements [185–190]. A small amount of ε-martensite can

form in some ASSs at small degree of deformation, and it reverts at much lower temperatures

compared to α′-martensite (see e.g., [143]). Singh [118] reported that the ε-martensite was stable

up to 200 °C, and according to Santos and Andrade [186], it reverts in the temperature range 50–

200 °C and between 150–400 °C according to Dryzek et al. [187]. Very recently a latent

strengthening mechanism, bake hardening without interstitials, due to the reversion of ε-martensite,

has been reported in a metastable FCC high entropy alloy by Wei et al. [191]. Annealing for 20

min at 200 °C was adequate for complete reversion accomplished by a shear-assisted displacive

mechanism. The reversion of α′-martensite (i.e., DIM) to austenite can occur by two different

mechanisms, diffusionless shear or diffusion-controlled one, as reported already in 1967 by Guy

et al. [117]; see also [120,126]. Guy et al. observed that austenite with mechanical twins formed

42

first from martensite in 18Cr-8Ni and 18Cr-12Ni steels, which then recovered to a sub-grain

structure. As regards the GS refinement, both reversion mechanisms can readily lead to a sub-

micron scale GS, though in principle the diffusional reversion is more efficient [119]. In Fe-Cr-Ni

ternary alloys, in the first stage, the shear phase reversion results in austenite which contains traces

of prior α′-martensite morphology, the same grain boundaries as those of original austenite and a

high density of defects. After the fast transformation, defect-free austenite subgrains are formed

which coalesce to a structure resembling recrystallized structure [119,120]. An example of the

formation of dislocation free grains from subgrains is shown in Fig. 2.2a. On the contrary, the

diffusional reversion is characterized by nucleation and growth of randomly oriented equiaxed

austenite grains and the result is shown in Fig. 2.2b. The nucleation occurs at cell or lath boundaries

of deformed DIM, and austenitic grains grow in size with time but stay in a nanometer or

submicron range. Secondary phase precipitates can also form in the course of the reversion, for

instance nano-size chromium nitrides in 301LN [106,110] and carbides in 301 [188,192].

Figure 2.2: Reversion in 304Cu ASS occurred by the shear reversion (a), where dislocation free

grains are formed by continuous recrystallization (white arrows marked part) [136], diffusional reversion (b) [136].

With respect to the temperature range of reversion, the difference in the reversion

mechanisms can be illustrated by reversion-temperature-time diagram where the start and finish

temperatures of the martensite reversion to austenite are drawn [120]. Moreover, from the diagram,

it is possible to judge how the reversion process makes progress under certain conditions. In Fig.

43

2.3, the respective temperatures are As′ and Af′ for the shear reversion and As and Af for the

diffusional one. The martensitic shear reversion proceeds during heating in a narrow temperature

range As′ – Af′. These temperatures depend on the chemical composition of steels and are lowered

by increasing the Ni/Cr ratio, but they are independent of the heating rate. The shear reversion rate

is fast and independent of prior cold rolling reduction, and the reversed fraction is independent of

isothermal holding time between these temperatures. On the other hand, the As and Af temperatures

for the diffusional reversion depend on the heating rate in addition to the chemical composition of

the steel (Fig. 2.3). The effect of heating rate was already investigated in Fe-Ni-C alloys by Apple

and Krauss [193] in 1972. In isothermal annealing, the diffusional reversion proceeds more rapidly

with increasing annealing temperature, and the Af temperature depends on soaking time, as seen

from the As and Af curves.

44


treatment at 700 °C for the studied 304Cu steel. [136] There are two methods to use the phase reversion approach. One way is to achieve through

Gleeble simulator. For cold rolling, the steel was received in the form of a hot rolled sheet, ~3.2

mm in thickness. The as-received steel sheet was cold rolled in a laboratory rolling mill to 0.93

mm thickness (~71% reduction) and subsequently annealed in a Gleeble 3800 simulator at various

temperatures in the range 650-950 °C using isothermal holding times between 1 and 5400 s. For

reversion treatment, the samples were heated at 200 °C/s to the annealing temperature, held for

desired duration and then cooled at the same rate at least down to 300 °C.

45

The other way is to achieve through tubular resistance furnace. Cold rolling was performed

up to 30-90% reduction at room temperature. Subsequently, the strips were annealed at various

temperatures in the range 800-1000 °C using isothermal holding times between 1 and 36000 s in

a tubular resistance furnace filled with argon, followed by quenching in ice-water.

2.2 METALLOGRAPHY

Standard metallographic techniques were used to ground and polish the specimens to

mirror finish and then electrochemically etched with 60% nitric acid solution. Microstructure was

observed by optical microscopy (OM) and field emission scanning electron microscopy (FE-SEM).

2.3 X-RAY DIFFRACTION

X-ray diffraction (XRD) is mainly used for phase identification of phase and can provide

information on cell size. The analyzed material was finely ground and homogenized, and its

average volume fraction of bulk scale was determined.

X-ray diffraction is based on the phase diagram interference between monochromatic X-

rays and crystal samples. These X-rays are generated by the cathode ray tube, filtered to produce

monochromatic radiation, collimated, and focused on the sample. When Bragg's law (nλ=2d sin θ)

is satisfied, the interaction between incident light and sample will produce constructive

interference (and diffraction light). This law relates the wavelength of electromagnetic radiation

to the diffraction angle and lattice spacing of the crystal sample. Then the X-ray diffraction is

detected, processed and counted. By scanning the sample in the 2θ angle range, all possible

diffraction directions of the lattice should be obtained due to the random orientation of the bulk

material. The transformation of diffraction peak to d-spacing can identify minerals, because each

mineral has a unique set of d-spacing. Usually, this is achieved by comparing the d-spacing with

the standard reference pattern.

46

All diffraction methods are based on the generation of X-rays in X-ray tubes. These X-rays

irradiate the sample directly and collect the diffraction rays. A key component of all diffraction is

the angle between the incident ray and the diffracted ray. In addition, the diffractometer of powder

and single crystal are also different.

The contents of martensite and austenite were measured by X-ray diffraction (XRD) using

Cu Kα radiation (PANslytical, Netherlands, 40 kV, 40 mA). The obtained data were analyzed in

Jade software. The volume fractions of austenite and martensite were calculated by the integrated

intensities of (110)α, (211)α, (200)α, and (202)α martensite peaks and (111)γ, (220)γ, (200)γ, and

(311)γ austenite peaks by Eqs. (2.1) and (2.2) [194,195].

! = 1.4 #!/#$ + 1.4#!� (2.1)

$ = 1 − ! (2.2)

where Vγ and Vα are the volume fractions of austenite and martensite, respectively, Iγ and

Iα are the integrated intensities of austenite and martensite peaks, respectively.

2.4 TENSILE TESTS

Mechanical properties of the cold-rolled and reversion annealed specimens were

determined by tensile testing. Uniaxial tensile tests were conducted at room temperature using a

Zwick Z100 machine on specimens, taken along rolling direction, with the gage dimensions of 15

× 5 × 1 mm at an initial strain rate of 0.008 s-1 (according to standard EN ISO-10002-1) or 65 ×

20 × 1 mm at an initial strain rate of 5×10-4 s-1 (according to standard ISO 6892). Generally, tests

were repeated twice. The hardness tests were carried out by use of Vickers method with a 5 kg or

0.5 kg load.

2.5 FRACTURE SURFACE EXAMINATION BY SEM

Scanning electron microscopy (SEM) uses a focused electron beam onto a surface to

47

produce an image. The electrons in the electron beam interact with the sample to generate various

signals, which can be used to obtain information about the surface morphology and composition.

Electron microscope was developed when wavelength became the limiting factor of optical

microscope. The wavelength of the electron is much shorter, so the resolution is higher.

Tensile samples tested until fracture were examined in a FE-SEM to study the mode of

fracture. The SEM micrographs of the fracture surface were processed using Image Pro software

to clearly delineate the fracture morphology.

2.6 NANOINDENTATION

In order to evaluate the mechanical properties of finite size structures, such as

nanostructured materials, films and ion irradiated damage areas, small-scale deformation is usually

required [196-200]. Nanoindentation is a robust technique to study the local deformation behavior

in nano/micro scale by continuously controlling and recording the load and displacement depth of

the indenter on the specimen surface. It provides an economical and effective method to understand

the deformation mechanism in multi-scale modeling. The accuracy of load is 1nn and the accuracy

of displacement depth is 0.1nm, which can effectively eliminate the influence of surrounding

structure and base on experimental data [201-207].

Two types of nanoscale deformation experiments were conducted.

The first type was conducted in load-controlled mode at a loading rate of 2 μN·s-1 with the

maximum load set to 0.5 mN or at a loading rate of 6 mN∙min-1 with the maximum load set to 1000

μN, dwell time of 10 s, followed by unloading. Here the objective was to observe any differences

in load-displacement plots that may provide an insight on the deformation mechanism. Load-

controlled nanoindentation experiments at a constant loading rate can elucidate the indentation-

induced deformation phenomenon as a function of displacement (or strain) that is difficult to

48

achieve from the strain rate-controlled experiments. This is because the minimum strain rate

available with the instrument is relatively quite high, and any discrete bursts in the load-

displacement plots associated with dislocation nucleation or phase transformation cannot be

recorded.

The second type of experiment was conducted in displacement-controlled mode, which

involved indentation at various constant strain rates in the range 0.01–1 s-1. The maximum

displacement was fixed at 500 nm or 2100 nm. Here the aim was to study the strain-rate sensitivity

at low strain rate and the hardness distribution beneath the worn subsurface.

The nanoindentation test system (Keysight Nanoindenter G200) consisted of a Berkovich

three-sided pyramidal diamond indenter with a nominal angle of 65.3° and indenter tip diameter

of 20 nm.

2.7 TEM FOIL PREPARATION AND TEM

Transmission electron microscope (TEM) is a powerful tool in materials science. The

interaction between electrons and atoms can be used to observe the crystal structure and structural

features, such as dislocations and grain boundaries. Chemical analysis can also be carried out.

TEM can be used to study the growth, composition and defects of layers in semiconductors. High

resolution can be used to analyze the mass, shape, size and density of quantum wells, quantum

wires and quantum dots.

TEM works on the same principle as optical microscope, but uses electrons instead of light.

Because the wavelength of electron is much smaller than light, the optimal resolution of TEM

image is many orders of magnitude higher than that of optical microscope. Thus, TEM can reveal

the finest details of the internal structure - in some cases, even as small as a single atom.

The electron beam from the electron gun is focused onto a small and thin coherent beam

49

through the focusing lens. This kind of beam is limited by the focusing hole, which excludes high

angle electrons. Then, the light beam irradiates on the sample, and part of the sample is transmitted

according to the thickness and electronic transparency of the sample. The transmission part is

focused by the objective lens into the image on the fluorescent screen or charge coupled device

(CCD) camera. The optional objective aperture can be used to enhance contrast by blocking high

angle diffraction electrons. Then the image is transmitted downward through the middle lens and

the projector lens, and is magnified all the time.

The image hits the screen and generates light, allowing the user to see the image. In the

image, the darker region represents the sample region with less electron transmission, while the

brighter region represents the sample region with more electron transmission.

Fig. 2.4 [208,209] shows a simple sketch of the path of the electron beam from the sample

to the screen in TEM. As electrons pass through the sample, they are scattered by the electrostatic

potential generated by the constituent elements in the sample. After passing through the sample,

they focus all the electrons scattered from one point of the sample to a point on the image plane

through the electromagnetic objective. In addition, the dotted line shown in Fig. 2.4 shows that the

electrons scattered by the sample in the same direction are collected to a point. This is the back

focal plane of the objective, where the diffraction pattern is formed.

50

Figure 2.4: Schematic of electron beam in TEM [208,209].

TEM samples must be thin enough to transmit enough electrons to form image with

minimal energy loss. Therefore, sample preparation is an important aspect of TEM analysis. For

electronic materials, a common preparation technology is ultrasonic disc cutting, indentation and

ion milling. Indentation method is a kind of preparation technology, which can make the central

area of the sample thinner and the outer edge of the sample have enough thickness to facilitate

handling. Traditionally, ion grinding is the final form of sample preparation. In this process, the

charged argon ions are accelerated to the sample surface by high pressure. Due to momentum

transfer, the material will be removed from the sample surface by ion impact.

Post-mortem TEM study of indented NG/UFG and CG samples was carried out to explore

the deformation mechanisms in the plastic zone surrounding the indented region. This involved

removal of indented 3 mm punched disks from the mount and electropolishing from the side

opposite to the indented surface, whereas the side with the indentations was masked with an

aluminum foil. Thin foils were prepared by twin-jet electropolishing of 3 mm disks using a solution

of 10% perchloric acid in acetic acid as electrolyte at 0 °C. Using this approach, the area

surrounding the indents present around the jet-polished hole, was electron transparent thus

51

enabling study of the deformation behavior by TEM. During TEM studies, the focus was in the

center of the deformation zone. The data presented here had excellent reproducibility, as confirmed

by a number of experiments for each set of conditions.

2.8 EBSD SAMPLE PREPARATION AND EBSD

In general, EBSD system (Fig. 2.5) [210] consists of a crystal sample tilted from the

horizontal direction to 70 ° using a SEM stage or a pre-filter bracket, a fluorescent screen emitting

fluorescence from electrons scattered by the sample, a sensitive camera, an optical element for

observing the pattern formed on the screen, an insertion mechanism, etc. It precisely controls the

position of the detector when in use, and retracts the detector to a safe position when not in use. In

order to prevent interference with the operation of the scanning electron microscope, control the

electronic equipment of the scanning electron microscope, including the movement of the beam

and the workbench, control the computer and software of the EBSD experiment, collect and

analyze the diffraction patterns, and display the results. The forward scattering diode (FSD)

installed around the fluorescent screen is used to generate the microstructure image of the sample

before collecting the EBSD data. The EBSD system can be selected Integration with EDS system.

52

Figure 2.5: An EBSD system. (a) Principle components of an EBSD system, (b) a photograph

showing the EBSD system integrated with an EDS system [210]. The following model describes the main features of pattern formation and collection for

EBSD analysis. The electron beam points to a point of interest on the tilted crystal sample. The

atoms in the material scatter some electrons inelastically, and the energy loss is very small, forming

a divergent electron source near the sample surface. Some of the electrons are incident on the

atomic surface at an angle satisfying the Bragg’s law (nλ=2d sin θ).

These electrons are diffracted to form a pair of large angle cones corresponding to each

diffraction plane. The image generated on the screen contains a characteristic Kikuchi band formed

53

at the intersection of the enhanced electron intensity region and the screen (Fig. 2.6) [210]. The

pattern we see is the projection of the diffractive cone, which makes the band edge hyperbolic.

Figure 2.6: The formation of the electron backscattered diffraction pattern (EBSP). (a) Cones

(green and blue) generated by electrons from a divergent source which satisfy the Bragg equation on a single lattice plane. These cones project onto the phosphor

screen, and form the Kikuchi bands which are visible in the EBSP. (b) Generated EBSP [210].

The mechanism causing the intensity and shape of Kikuchi band is complex. As an

approximation, the strength of planar Kikuchi band (hkl) is given by the following formula:

#&'( = )∑ +,-� cos 22ℎ4, + 56, + 78,�, 9� + )∑ +,-� sin 22ℎ4, + 56, + 78,�, 9�

(2.3)

where fi(θ) is the atomic scattering factor of the electron and (xi, yi, zi) is the fractional coordinate

in the unit cell of atom i. This ensures that only the plane that produces the visible Kikuchi band

is used to solve the diffraction pattern.

The center line of Kikuchi band corresponds to the position where the diffraction plane

54

intersects the fluorescent screen. Therefore, each Kikuchi band can be expressed by Miller index

of diffraction plane. The intersection of the Kikuchi band corresponds to the region axis in the

crystal.

This pattern is a kind of Gnostic projection of the electron diffraction cone on the screen.

The half angle of the electron diffraction cone is (90 - θ)°. For EBSD, this is a large angle, so the

Kikuchi band is approximately a straight line. For example, the wavelength of 20kV electron is

0.00859nm, and the spacing of (111) plane in aluminum is 0.233nm, so that the half angle of cone

is 88.9°.

The width W of Kikuchi band near the center of the pattern is given by the following

formula:

< ≈ 27- ≈ >(?@ (2.4)

where l is the distance from the sample to the screen. The plane with wide d-spacing gives

a thinner Kikuchi band than the plane with narrow d-spacing. Because the diffraction pattern is

related to the crystal structure of the sample, the diffraction pattern changes with the change of

crystal orientation. Therefore, the position of the Kikuchi band can be used to calculate the

orientation of the diffractive crystal (Fig. 2.7) [210].

55

Figure 2.7: The spherical diffraction patterns generated by different orientations of a cubic structure. [210].

Samples for EBSD scans were first ground gently to 600 grit and then to avoid any DIM

formation during the mechanical preparation electropolished with perchloric acid solution (20%

perchloric acid-80% ethanol solution operated at 25 °C at an applied potential of 15 V). The EBSD

scans were performed using the accelerating voltage of 15 kV, the working distance of 11 mm;

step size was varied according to magnification being between 0.2 and 0.05 μm. All EBSD scans

were made in RD-TD plane, i.e. ¾ thickness from the top surface. The analyses of EBSD

measurements were carried out using an Oxford HKL acquisition and analysis software in order

to characterize the grain structure, grain and sub-boundaries and various phases. The boundary

with a misorientation larger than 2° was regarded as the boundary of two crystallographic grains.

56

Chapter 3: Improving the yield strength of an antibacterial 304Cu austenitic stainless steel

by the reversion treatment

In this chapter, we have conducted an in-depth understanding of the effect of annealing

temperature-time on the mechanical properties of 304Cu stainless steel. The potential of

optimizing the combination of annealing temperature-time in obtaining high strength-high ductile

steel is elucidated in the study presented here. The microstructure and nanoscale Cu precipitates

were also analyzed to understand their contribution toward strengthening.

3.1 MATERIAL AND EXPERIMENTAL PROCEDURE

The chemical composition of the experimental 304L austenitic stainless steel containing

3.15% Cu is listed in Table 3.1. The typical features of this steel composition are its low C content

(0.023% C) and high stability as estimated using the stability index, Md30 = -11.2 °C (for GS ASTM

#7), as given below [211]:

MBC�°C� = 552 − 462C + N� − 9.2Si − 8.1Mn − 13.7Cr − 29Ni + Cu� −18.5Mo − 68Nb − 1.42GS − 8� (3.1)

Table 3.1: Chemical composition (wt. %) of the experimental Cu-bearing austenitic stainless steel

C Si Mn Cr Ni Cu S P Fe Md30(°C)

0.023 0.55 0.85 17.40 7.32 3.15 0.011 0.025 Balance -11.16

The steel was made as a laboratory casting using standard melting practice. For cold rolling,

the steel was received in the form of a hot rolled sheet, ~3.2 mm in thickness. The as-received steel

sheet was cold rolled in a laboratory rolling mill to 0.93 mm thickness (~71% reduction) and

subsequently annealed in a Gleeble 3800 simulator at various temperatures in the range 650–

950 °C using isothermal holding times between 1 and 5400 s. For reversion treatment, the samples

were heated at 200 °C/s to the annealing temperature, held for desired duration and then cooled at

57

the same rate at least down to 300 °C. The initial grain size of the sample was estimated as 31 ± 4

μm.

Specimens for the metallographic characterization were ground and polished according to

standard metallographic practice. Microstructural characterization was performed using a Zeiss

Ultra Plus field emission gun scanning electron microscope (FEG-SEM) equipped with an electron

backscatter diffraction (EBSD) device. Samples for EBSD scans were first ground gently to 600

grit and then to avoid any DIM formation during the mechanical preparation electropolished with

perchloric acid solution. The EBSD scans were performed using the accelerating voltage of 15 kV,

the working distance of 11 mm; step size was varied according to magnification being between

0.2 and 0.08 μm. All EBSD scans were made in RD-TD plane, i.e. ¾ thickness from the top

surface. The analyses of EBSD measurements were carried out using an Oxford HKL acquisition

and analysis software in order to characterize the grain structure, grain and sub-boundaries and

various phases.

Examination of Cu precipitation was performed on a JEOL JEM-2200FS

scanning/transmission electron microscope (STEM/TEM) operated at 200 kV. Specimens for

TEM/STEM were first ground to a thickness of 80 μm and then prepared using twin-jet

electropolishing at - 10 °C using 23 V DC in an electrolyte consisting of perchloric acid, ethanol,

butyl cellosolve and distilled water (Struers A2).

DIM fractions were determined by a Ferritescope (Helmut Fisher FMP 30) instrument. The

readings obtained were multiplied by a factor 1.7 for α′-martensite fractions.

Mechanical properties of the cold-rolled and reversion annealed specimens were

determined by tensile testing. Uniaxial tensile tests were conducted at room temperature using a

Zwick Z100 machine on specimens, taken along rolling direction, with the gage dimensions of 15

58

× 5 × 1 mm at an initial strain rate of 0.008 s-1 (according to standard EN ISO-10002-1). Generally,

tests were repeated twice. The hardness tests were carried out by use of Vickers method with a 5

kg load.

3.2 RESULTS

3.2.1 Cold rolling

Cold rolling in 13 passes (about 15% each) to the total reduction of 71% resulted in the

DIM fraction of 80%, so that still a significant amount of deformed austenite (DA) was retained.

Fig. 3.1 shows the evolution of the DIM fraction during cold rolling. A gradual increase of

martensite also means that inevitably some of DIM becomes deformed only slightly, for instance

the half of martensite was deformed to a 30% reduction at the maximum.

Figure 3.1: Formation of αʹ-martensite during cold rolling.

3.2.2 Reversed microstructures

As the microstructure after cold rolling consisted of two phases, viz., DIM, i.e. αʹ-

martensite with varying degree of deformation, and deformed DA with ~71% reduction, these two

phases ought to behave differently during the reversion annealing treatment and inevitably results

in variable microstructural features, which are visibly discernible especially for samples annealed

at low temperatures.

59

We begin the description of microstructures from more simple structures obtained at high

annealing temperatures and continue towards low-temperature annealing structures, the latter

being more complex but providing better strength properties. Fig. 3.2 displays the microstructure

after annealing for 1 s at 900 °C (annealing at 950 °C resulted in more or less similar but still

coarser structure, not shown here), as observed by EBSD. The structure is fully austenitic, and the

GS is slightly non-homogeneous consisting essentially of fine grains of few microns in size and

larger grains up to 10 μm (Fig. 3.2a). (The GS distribution is shown later in section 3.2.4) This

inhomogeneity is due to a mixture of reversion-refined fine grains formed from DIM and

recrystallized grains formed from DA, as reported in many papers, e.g. Ref. [163,212]. The

orientation image microscopy (OIM) map (Fig. 3.2b) reveals that the grains appear in different

colors, i.e. they have random orientation, although green colored {110} ⟨hkl⟩ grains seem to be

most prominent.

Figure 3.2: Austenitic grain structure after annealing at 900 °C for 1 s. EBSD grain boundary

map (a) and the orientation image map (b). At 850 °C-1 s hold, the local inhomogeneity in GS was even more pronounced due to lesser

grain growth than that occurred at 900 °C (Fig. 3.3a). At this temperature, few larger grains

containing low-angle grain boundaries (LAGBs; the misorientation between 2 and 15°) were found

to exist (marked by arrows in Fig. 3.3a), i.e. shear reversed austenite displaying substructure. A

very small fraction of unreversed martensite (red grains) was also detected. Similar features were

present in the structure of the sample annealed at 800 °C within 10 s, though the large irregular-

60

shaped grains with LAGBs were far more numerous and the amount of unreversed martensite also

increased (Fig. 3.3b). After annealing at 800 °C for 10 s, few non-recrystallized DA grains were

also observed, an example is given in Fig. 3.3c.

Figure 3.3: Reversed grain structure after annealing at 850 °C-1 s (a) and 800 °C-10 s (b and c).

Grains containing low angle grain boundaries pointed by arrows in (a), presence of irregular grain (b) and a non-recrystallized deformed austenite grain in (c).

(Austenite gray, martensite red in color). At lower annealing temperatures of 750–650 °C, the microstructures were strikingly

different from those created at higher temperatures; some examples are shown in Figs. 3.4–3.7.

The EBSD phase maps and Ferritscope measurements indicated clearly that even after a short

annealing duration, the major phase was austenite with only a minor amount of unreversed

martensite (red colored grains in phase maps), for instance 81% austenite after 1 s hold at 700 °C

(not shown) and 86% after 10 s (Fig. 3.4a). The austenite consisted of refined grains, though with

different sizes and various colors, as highlighted in the figures, but also green-colored, coarse

elongated grains were present with the shape of the original cold-rolled grains, as can be seen in

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Figs. 3.4b, 3.5 and 3.6. They also contained a large number of LAGBs (white-colored boundaries).

The fraction of fine grains became smaller in structures annealed at lower temperatures and

correspondingly the fraction of large austenite grains with LAGBs increased (compare Figs. 3.4

and 3.7).

Figure 3.4: Microstructure obtained after annealing at 700 °C for 10 s. Phase map (a) and OIM

map (b). Martensite red-colored in (a).

Figure 3.5: Microstructure obtained after annealing at 700 °C for 1800 s at two different

magnifications (OIM maps).

62

Figure 3.6: Microstructure obtained after annealing at 650 °C for 3600 s. Martensite red in the phase map (left).


phase map. In addition to austenite, some retained DIM existed in the structures annealed at 800 °C

and lower temperatures (Figs. 3.4–3.7). The DIM fraction existing after 1 s holding depended on

the annealing temperature, being 8%, 12% and 19% after annealing at 800, 750 and 700 °C

respectively.

The unreversed DIM fractions after annealing at 750, 700 and 650 °C are plotted as a

function of annealing time in Fig. 3.8. A complete list of DIM fractions at different annealing

temperatures and/or times including the data plotted in Fig. 3.8 will be presented later in Table 3.3.

It is seen that the DIM fraction decreased with prolonged soaking time, being lower after higher

annealing temperatures at a given annealing duration. Thus, the reversion phenomenon continued

and was dependent on temperature and time.

63

Figure 3.8: Fraction of martensite retained after annealing at 750, 700 and 650 °C for various

annealing durations. 3.2.3 Grain size

As noticed from previous figures (Figs. 3.4–3.7), the GS after reversion treatment is not

uniform and the non-homogeneity increases with decreasing the annealing temperature. To

illustrate this, the area weighted GS distributions after various reversion conditions are plotted in

Fig. 3.9, based on high angle grain boundaries (HAGBs; misorientation >15°) or both LAGBs and

HAGBs (H&LAGBs) (Fig. 3.9 a and b, respectively). It is seen that after annealing at 800–900 °C,

the peak in the area frequency is between 1 and 3 μm, though few much larger grains do also exist.

GS distributions of structures obtained by reversion at 800–900 °C are more uniform (specially

for HAGBs) compared to those at lower temperatures and the average GS is almost constant

meaning that reversion is completed at these temperatures. A comparison of HAGBs and

H&LAGBs also shows that the area fraction of LAGBs is quiet low for high temperatures

indicating again the completion of reversion. The peak shifted to slightly smaller GSs while

annealing was performed at 700 or 650 °C.

64

Figure 3.9: Grain size distribution after reversion annealing at different conditions based on high

angle grain boundaries (HAGBs) (a) or both HAGBs and low angle grain boundaries (LAGBs; misorientation 2–15°) (b).

A distinct feature in the distributions is the appearance of large (30–50 μm) grains after

annealing at low temperatures of 700 and 650 °C resulting from the existence of non-recrystallized

DA grains. Another feature after the same conditions is a high fraction of LAGBs (Fig. 3.9b)

highlighting the presence of subgrains in the shear reversed austenite and recovery in DA grains.

The area fraction of LAGBs decreases with prolonged holding, especially at 700 °C as an obvious

consequence of the progress of recovery and subgrain coalescence.

3.2.4 Precipitation structure

Precipitation of Cu particles at surface region is required for the antibacterial property [213-

217]. According to Luo et al. [217], the optimal aging is 1.5 h at 650 °C. Therefore, TEM, STEM

and X-ray mapping were employed to check the presence of Cu after 1.5 h holding at 700 and

650 °C. Examples of the structure recorded on a sample annealed at 700 °C for 1.5 h are shown in

Figs. 3.10–3.12. Also, an X-ray map revealing the distribution of Cu in the examined field is

65

displayed in Figs. 3.10b and 3.11d. Selected area electron diffraction (SAED) patterns as presented

in Fig. 3.10c–e were used to identify the austenite and martensite grains in the reversion treated

microstructure. In Fig. 3.10a, a reversed austenite grain with low dislocation density (upper part),

unreversed martensite and sheared austenite with subgrains are found to co-exist based on SAED

patterns. The X-ray Cu map in Fig. 3.10b indicates that Cu precipitates did form and here they

seem to be mainly distributed along specific zones (black channels) at phase, grain and subgrain

boundaries. A local view of dislocation-free austenite grains is shown in Fig. 3.11, where Cu

precipitates are distributed quite uniformly. White spots in bright field (BF) image in Fig. 3.11a

are seen as black spots in dark field (DF) image in Fig. 3.11b, which appear as empty holes, where

particles have fallen off during the foil preparation. White spots in Fig. 3.11c are particles rich in

Cu, as also confirmed by X-ray map in Fig. 3.11d. The coherent character of the precipitates is

shown in a TEM two-beam BF image in Fig. 3.12a. The coherence is further verified by high-

resolution image of one particle (Fig. 3.12b).

66

Figure 3.10: STEM micrograph after annealing at 700 °C for 1.5 h (a), the corresponding X-ray map (b) and electron diffraction patterns of austenite (c) and martensite (d and e)

taken from areas marked in (a) by dashed circles.

Figure 3.11: A local view of dislocation-free austenite grains in a sample annealed at 700 °C for

1.5 h. Bright field (a) and dark field (b) images revealing nano-size particles. A magnified view (c) of the square area marked with red line in (b) and corresponding X-ray map of Cu distribution in this area (d). Black spots in (c) are holes (i.e. lost

precipitates) and are not seen in (d).

Figure 3.12: A TEM 2-beam BF image revealing the coherence contrast of Cu precipitates in

austenite (a) and an HR-STEM image of a Cu particle (b). Annealing at 700 °C for 1.5 h.

Fig. 3.13 displays a local area in a sample annealed at 650 °C for 1.5 h, revealing

dislocation-free austenite grains (or subgrains) which are surrounded by deformed structure.

Coherent Cu precipitates could be detected in at least two grains, as pointed out. However, the

67

presence or absence of precipitates in other grains has not been resolved. This would require further

studies.

Figure 3.13: A STEM micrograph of the sample annealed at 650 °C for 1.5 h showing small

reversed dislocation-free austenite grains surrounded by deformed structure. Coherent Cu precipitates in grains 1 and 2.

3.2.5 Tensile properties and strain-induced martensite

The objective of the applied reversion treatment applied was to improve the strength

properties of the steel. The hot rolled sheet before the cold rolling had the YS of 286 MPa, UTS

553 MPa and TE 51%. The results from tensile tests of the reversion-treated samples are listed in

Table 3.2 and corresponding engineering tensile stress-strain curves are plotted in Fig. 3.14. In

reversion experiments, the lowest YS was 256 MPa after annealing at 950 °C for 100 s

(microstructures or stress-strain curves are not shown here). It is seen that after 1h annealing at

700 °C, the YS value of the reversion-treated structure is about twice (≈524 MPa) higher than the

lowest YS and jumps to a level about thrice (≈790–830 MPa) higher after annealing at 650 °C. In

Fig. 3.15, the influence of the annealing time at 750, 700 and 650 °C on the YS is shown, revealing

fast drop of YS corresponding to annealing at 750 and 700 °C, but much less at 650 °C. The fracture

elongation decreases slightly with increasing strength, but it stays around 36% even after annealing

68

at 650 °C. Yielding however, appears to be the Lüders type after annealing at 650 °C for long times,

although not so distinctly as obtained by Sun et al. [218], who reported a long Lüders strain of 10%

for a reversion-treated structure of a 17Cr–6Ni–2Cu steel, annealed at 700 °C.

Table 3.2: Tensile properties of the 304Cu steel after reversion annealing treatments compared to those of as-received (hot-rolled) and cold-rolled conditions.

Temperature

(°C)-Time (s)

Yield strength

(MPa)

Tensile strength

(MPa)

Uniform

elongation (%)

Total elongation

(%)

As received 285 553 46.4 51.0

Cold-rolled 1228 1260 0.5 10.4

950-1 330 652 51.0 65.1

900-1 351 672 50.9 68.4

900-10 336 664 48.4 62.7

850-1 421 719 46.7 59.9

850-10 397 704 47.9 61.2

800-100 403 735 44.9 57.8

750-100 588 873 32.2 43.5

700-10 824 995 22.7 36.6

700-100 776 972 23.9 35.1

700-600 653 924 29.1 40.0

700-1800 602 898 29.2 40.8

700-3600 524 870 31.6 43.3

700-5400 507 876 31.2 43.1

650-1800 831 1007 24.6 35.8

650-3600 812 997 25.0 36.3

69

650-5400 791 1008 25.2 36.0

Figure 3.14: Stress-strain curves of a cold rolled specimen and some reversion annealed ones in

different conditions.

Figure 3.15: Effect of annealing duration at 750, 700 and 650 °C on yield strength. It is also seen that the stress-stress curves are convex in shape after annealing at

temperatures of 800–900 °C, but they become concave soon after the start of yielding for samples

annealed at 750–650 °C for 100 s or longer. The difference in tensile behavior becomes more

evident in strain hardening rate (SHR) vs. true strain curves predicted from the stress-strain data,

which are plotted in Fig. 3.16. The curves corresponding to reversion treatments at 750 °C and

lower temperatures reveal peaks in the incremental strain hardening rate vs. true strain plots, whose

height increases with prolonged duration at 700 and 650 °C. The high peak suggests more

significant αʹ-martensite formation during straining in the structures formed under certain

70

conditions with close dependence on annealing temperature and holding time.

Figure 3.16: Strain hardening rate as a function of true strain for the specimens annealed at

different conditions: (a) 750–900 °C with varying holding times 10–100 s, (b) 700 °C/100–5400 s and (c) 650 °C/1800–5400 s.

DIM fractions in some samples, reversion annealed for short times, were measured before

and after tensile testing and the values are listed in Table 3.3. In the table, the values of DIM formed

during tensile testing are also given, based on the difference between the values after and before

the tests. They indicate that the amount of DIM formed during straining to fracture does not depend

on annealing duration and it only increases slightly with increasing the annealing temperature in

the range of 850–950 °C, where minor change can be related to the coarsening of GS with

increasing annealing temperature [163]. Instead, after annealing at 750 and 700 °C, the fraction

increases with annealing time. This dependence is more readily seen in Fig. 3.17, where the amount

of new DIM formed during tensile straining to fracture is plotted as a function of annealing

71

duration at temperatures of 750, 700 and 650 °C. The figure reveals that the amount of new DIM

increases very significantly from about 50 to 90%, slightly faster corresponding to a higher

annealing temperature. This means that the stability of austenite decreases as a consequence of

annealing at these temperatures, and obviously due to the precipitation of Cu.

Table 3.3: Martensite content before and after tensile testing and formed during tensile test of samples annealed at different conditions.

Temperature/Time Before tensile testing After tensile testing During tensile testing

1 s 10 s 100 s 1 s 10 s 100 s 1 s 10 s 100 s

700 18.9 14.0 10.8 68.2 60.2 75.0 49.2 46.2 64.2

750 11.6 7.7 6.5 61.4 63.2 77.7 49.7 55.5 71.2

800 7.8 6.6 2.8 68.3 71.6 66.1 60.6 65.0 63.3

850 3.7 2.6 2.1 62.6 64.1 61.0 58.8 61.5 59.0

900 2.8 2.8 2.1 64.3 64.4 64.4 61.5 61.7 62.4

950 2.8 2.7 2.3 63.6 65.8 65.6 60.8 63.1 63.3

Figure 3.17: The amount of new DIM formed during tensile straining of the samples annealed at

650, 700 and 750 °C for different durations. 3.2.6 Hardness

The hardness values of the reversion-treated samples after annealing at various

72

temperatures for a short holding time of 1–100 s are plotted in Fig. 3.18. The corresponding

hardness data following reversion annealing of 0.5 and 1 h durations for 304L steels, taken from

Mészáros and Prohászka [219] and Martins et al. [220] respectively, are also included for

comparison. The drop in hardness is steep in a temperature interval between 700 and 800 °C as

shown by a green highlight in Fig. 3.18. In this temperature range, the amount of retained DIM

decreases as seen in Table 3.3, but the softening continues at 850–950 °C, where no retained DIM

existed. Excellent agreement can be noticed between the present and literature data also revealing

an influence of prolonged annealing times on hardness, particularly at 700–800 °C.

Figure 3.18: Hardness variation after annealing at different temperatures for 1, 10 and 100 s.

Some data from Mészáros and Prohászka [219] for 1 h and Martins et al. [220] for 0.5 h are included. The shaded area highlights the temperature range, where the

influence of annealing duration is significant. 3.3 DISCUSSION

The results clearly indicated that the grain structure of the 304Cu steel could be modified

by cold rolling and reversion annealing sequence resulting in an increase in the YS, while the

elongation remained reasonably high. We discuss below the reversion process in general and the

complex microstructures created in the temperature regime relevant to obtain the antibacterial

property in this steel (≤750 °C) and assess the improvement of the strength achieved.

73

3.3.1 Reversion behavior

Different opinions have been presented so far on the characteristics of the reversion process

in 304/304L steels with and without Cu alloying, so it is worthy to discuss shortly the mechanism

and kinetics of the reversion. One practical variable in the reversion treatment is the thickness

reduction in the cold rolling stage before the annealing. In relatively stable alloys such as the 304

grade, severe deformation is required at RT to obtain a structure consisting of 100% DIM;

reductions of 90% or beyond being applied in some studies [98,221]. The Cu alloying further

increases the stability of the steel by increasing its SFE and decreasing the Md temperature

[222,223]. However, very high cold rolling reductions are not quite practical in industry. In this

study, the steel was rolled to provide about 71% reduction, which resulted in the martensite fraction

of about 80% (Fig. 3.1). This means that two deformed phases, DIM and DA, were present, which

must be accounted for microstructure analysis. DA cannot, however, be refined to the same extent

as the highly deformed DIM. Also, after gradual formation of DIM during cold rolling without a

saturation stage, as is evident in the present case, a certain fraction of martensite remains as slightly

deformed lath martensite and does not reverse in a manner similar to the highly deformed cell

martensite [107, 111, 163, 224]. The behavior of DA among DIM has been discussed in several

papers, e.g. Ref. [212,225]. Järvenpää et al. [111, 163] have investigated in detail the evolution of

grain structure after different cold rolling reductions (32–63%), i.e., containing various retained

DA fractions in a 301LN steel, showing formation of complex structures after annealing, where

GS can be classified in four classes, the fractions dependent on the annealing temperature, in

particular.

The reversion mechanism in the 304 type stainless steel has been investigated in numerous

studies, e.g. Ref. [141-143, 149, 218, 225, 226], but somewhat different opinions exist. Some

74

researchers have reported that the diffusional reversion can occur at low annealing temperatures

of 550–650 °C [149, 226] and the shear reversion occurs at higher temperatures, e.g. above 750 °C

[149]. The effect of heating rate has been observed, suggesting diffusional reversion at low heating

rates (<10 °C/s) and shear reversion at higher ones (>40 °C/s) [218]. Consistently, Sun et al. [225]

envisaged that fine austenite grains formed via diffusional reversion of martensite as using the

heating rate of ≈10 °C/s. However, Cios et al. [143] and Ondobokova et al. [141] reported shear

type reversion mechanism in the temperature range 400–700 °C and 600–800 °C, respectively.

Even the concurrent operation of both the diffusion and shear mechanisms have also been claimed

[142].

Tomimura et al. [120] have shown that an increase in the Ni/Cr ratio causes an increase in

the Gibbs free energy change between the fcc and bcc structures and thereby lowers the martensitic

shear reversion temperature. The ratio Cr/Ni ≈ 0.63 was found to favor the shear reversion

mechanism in their experiments. In the present steel, Ni/Cr is low, about 0.42, so the diffusional

reversion would be preferred. However, in addition to Ni, Cu too is an austenite stabilizing element

(with the same power as Ni in Md [211]), and therefore we can expect it to favor the shear reversion

mechanism, but even the (Ni þ Cu)/Cr ≈ 0.54 is not very high in the present instance. However, a

high heating rate of 200 °C/s was used in the annealing experiments, so this may be the main

reason for the occurrence of the shear mechanism, in agreement with the observations of Sun et al.

[9]. A firm evidence for the shear reversion, as seen in Fig. 3.19, is that the structure is almost fully

coarse-grained austenite even after very short annealing time of 10 s at 700 °C (a decrease of the

DIM fraction from 80% to 19%, Table 3.3, also seen in Fig. 3.4). This implicates that the reversion

has been very fast, i.e., faster than those employed in the experiments of Sun et al. [218], who

found that only 25% of DIM reversed within 2 min at 700 °C. The shear mechanism is evident

75

also from the elongated shape of austenite grains after annealing at 650 and 700 °C (Figs. 3.4–3.7),

containing traces of prior αʹ-martensite morphology, as a clarifying feature of the shear reversion

[120]. After this fast reversion transformation, subgrains are formed in reversed austenite which

coalesce into a structure resembling the defect-free recrystallized structure with time. The

occurrence of continuous recrystallization mechanism is demonstrated in local views in Fig. 3.19,

depicting two examples.

Figure 3.19: Formation of defect-free austenite grains during annealing at 700 °C for 10 s (a) and

600 s (b) indicating the shear reversion mechanism followed by continuous recrystallization. Low angle grain boundaries are white lines in the orientation

image map (a), and martensite is red in the phase map (b). After low-temperature annealing, very pronounced size differences appear in dislocation-

free grains (concluded from a high image quality (IQ) in EBSD images), as shown in Fig. 3.20.

There are fine grains but also large grains (few microns), often in groups, and one might speculate

that the large grains have formed by the local occurrence of diffusional reversion. A low

temperature of 650–700 °C would favor diffusional reversion in 304 type steel [149, 226]. Takaki

et al. [224] demonstrated that in the diffusional reversion from lath-type martensite, i.e. from

slightly deformed martensite, austenite nucleates on lath boundaries with a shape of thin plate and

that the same kind of austenite gathers in a group forming blocks with certain orientation. One

kind of austenite grains nucleate within one martensite block and the reversed austenite inherits

the morphological characteristics of lath-martensite even in a diffusional reversion. However, in

the present instance, these grains do not have a lath shape, but they are largely equiaxed or have a

76

typical subgrain shape, though large in size. These grains are surrounded by areas of subgrains,

where new strain-free grains are forming by continuous recrystallization, obviously as a follow-

up of the shear reversion. Therefore, it seems that even these large dislocation-free grains are

formed very quickly from the shear reversed austenite, which in turn has formed from slightly

deformed martensite. Some of these grains still contain LAGBs inside (see Fig. 3.20a).

Figure 3.20: Examples of big difference in the grain size in reversed dislocation-free grains after

annealing at 700 °C for 10 s (a,b) and 600 s (c,d). DA is retained deformed austenite grain (a). Martensite is in red in the phase map (b,d).

The DA grains are always green in color, i.e. Brass oriented, as in Fig. 3.20a (a DA grain

marked), showing the highest stability [163, 227]. They can be distinguished from reversed coarse

grains on the basis that hardly any new grains would have formed in them after short time

annealing (DA in Fig. 3.20a) and often long parallel shear bands are also seen in them (see Figs.

3.4 and 3.6). Recrystallization of DA is a slow process, starting in 1 h at 700 °C [220, 228]. Larger

grains seen after annealing at 850–900 °C are formed from the DA by recrystallization (Figs. 3.2

77

and 3.3), which can also be concluded from the GS distribution in Fig. 3.9. However, it has been

found that the continuous recrystallization by formation of subgrains and their evolution to grains

can also happen in DA grains [163].

A special feature in the present microstructures is that some αʹ-martensite exists even after

annealing for 1.5 h at 650–700 °C. However, it is clearly seen that the DIM content decreased

during isothermal holding at temperatures below 850 °C (Fig. 3.8, Table 3.3). This can be

explained by the occurrence of diffusional reversion following shear reversion. Tomimura et al.

[120] have presented the time-temperature-reversion (TTR) diagram to describe the reversion

under different heating-annealing conditions. In Fig. 3.21, the reversion start and finish

temperatures are shown for the shear and diffusional reversion mechanisms (Asʹ, Afʹ and As, Af

respectively). Thus, under certain conditions: at a high heating rate to the regime between Asʹ and

Af ʹ results in partial reversion by the shear mechanism. In the two-phase regime, during isothermal

holding, the diffusional reversion starts at time corresponding to As and becomes completed at

time corresponding to Af. Shakhova et al. [142] predicted and observed Af around 800 °C for

S304H. Also for the present steel, Afʹ can be evaluated as or slightly above 800 °C from the data

in Table 3.3. Recently Sohrabi et al. [229] have investigated the remaining DIM in 304L, 301LN,

316L type steels showing it to be thermodynamically stable at temperatures below 700 °C. In

continuous heating at 15 °C/min the DIM disappeared at 750 °C in 304L [230], in fair consistency

with the present observations. Martins et al. [220] and Shakhova et al. [142] reported reversion

starting between 500 and 550 °C for a 304L and S304H steels, respectively, and Shakhova et al.

[142] measured about 8% and 20% DIM (ferrite) at 700 and 650 °C for 30 min, respectively. In

the present experiments, some DIM (ferrite) still remains in the structure at 700 and 650 °C,

although the decrease seems to continue at 700 °C after 1.5 h. Based on this information, the

78

approximate scales are drawn in the TTR diagram in Fig. 3.21 and a processing route at 700 °C

for 5400 s is shown comprising two successive reversion mechanisms leading to partial reversion.


treatment at 700 °C for the studied 304Cu steel. Luo et al. [217] mentioned that martensite is formed during an annealing treatment for the

antibacterial property. Of course, martensite cannot form at such a high temperature, but a bcc

phase is retained during annealing below a certain temperature. Also in the experiments of Cios et

al. [143], Shen et al. [231], Mészáros and Prohászka [219], Odnobokova et al. [141], and Shakhova

et al. [142], up to 10% martensite was left unreversed at 700 °C (after 0.5–1 h), in agreement with

the present observation (Figs. 3.6 and 3.7).

According to Luo et al. [217], the presence of a small amount of martensite has no effect on

the precipitation of the Cu-rich phase, so it does not affect the antibacterial properties of the sample.

However, we may need to address a concern of the pitting corrosion resistance of the structure,

which may be detrimentally affected by the presence of the martensite phase [232], further

debilitated by the Cu precipitation [214,216,233,234]. This needs further studies.

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3.3.2 Precipitation kinetics

It is known that in supersaturated ferrite, Cu-rich coherent clusters, initially also rich in Fe,

nucleate fast. During growth, these bcc-Cu preprecipitates run through two structural

transformations, viz. twinned 9R and untwinned 3R, until they ultimately transform to the

equilibrium fcc structure of pure Cu with an incoherent interphase boundary [235–237].

Segregation of Cu on grain boundaries is accompanied with the cluster formation, too. However,

in austenite, the precipitation of Cu-rich phase is just a gradual chemical composition change

without any transformation of crystallographic structure, because both the Cu-rich phase and

austenitic matrix have the same fcc crystallographic structure and close lattice parameters.

As regards the duration of aging needed for the precipitation, Hong and Koo [214] showed

that an ageing time of 4 h at 700 °C is long enough for a 304-2.5Cu steel to generate sufficient Cu-

precipitates for the antibacterial property. Chi et al. [238] detected Cu segregated areas after 1 h

aging at 650 °C in a 304-Nb-N-3Cu alloy which developed into coherent particles keeping a fine

size of 34 nm even until 10000 h. Also, Luo et al. [217] observed Cu-rich precipitates of 15 nm in

size in a 304L-3Cu alloy after 1 h annealing at 650 °C and found that the optimal heat treatment

process comprised aging at 650 °C for 1.5 h, following solid solution at 1050 °C for about 30 min.

Recently, Luo et al. [239] reported the existence of coherent Cu particles and good antibacterial

property in a 304L-3Cu alloy after aging for 30 min at 750 °C.

All the same, the above aging experiments have been performed for annealed austenite. In

the present instance, the initial structure is mainly deformed DIM (about 80%), from which the

austenite is shear reversed that contain a higher density of dislocations and subboundaries, which

later annihilate or coalesce to form dislocation-free austenite grains. In addition to shear reversed

austenite, there is about 20% DA, containing recovered structure, and a small fraction of retained

80

DIM. The precipitation kinetics can be much faster in martensite and dislocated austenite

compared to that in the annealed austenite. The Cu precipitation kinetics in martensitic 17-4 PH

stainless steel has been analyzed by Mirzadeh and Najafizadeh [240] showing that the activation

energy of precipitation was close to that of Cu diffusion in ferrite. Precipitation took place at

temperatures much lower than tested here. Stechauser and Kozeschnik [236] have presented a

TTP-diagram based on simulation and also experimental data for the Cu precipitation in bcc α-

iron, and accordingly, the precipitation would start in 100 s at 600–700 °C and be completed within

an hour. Soylo and Honeycombe [241] found coherent Cu-rich bcc zones in martensite/ferrite in a

30Cr–8Ni–3Cu steel quenched from 1300 °C, followed by reversion annealing at 700 °C for 30 s,

and within 1 min incoherent fcc particles developed from them. During the reversion annealing of

cold rolled 301LN steel, precipitation of CrN has been found to occur within few seconds at 700 °C

[106, 110]. Thus, the precipitation can be very fast, if it occurs in bcc structure and a high

dislocation density further accelerates it [242].

On the other hand, it is noteworthy that the shear reversion was very fast, so that the

reversion was almost completed on heating at 200 °C/s and holding for 1 s at 700 °C, for instance.

Therefore, this duration does not provide much time for the Cu precipitation in the DIM. The

coherence of the particles in the austenite grain (as seen in Figs. 3.12 and 3.13) means that they

have the fcc structure, but it is difficult to know if they formed in DIM or austenite. Hence, we

have to conclude that the present circumstances for the precipitation are complex and this study

cannot reveal and explain them, but only demonstrates that the Cu precipitation has occurred

during the reversion treatment at 650–700 °C for 1.5 h, as is evident from Figs. 3.10–3.13. A shorter

time might be enough for the purpose, even beneficial for the strength, but this needs to be

ascertained in a future study.

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Tensile stress-strain and SHR curves (Figs. 3.14 and 3.16, respectively) might also augment

further information regarding the precipitation kinetics. The evolution of the YS does not give such

information, as it is affected by changes in dislocation structure in austenite and also the decrease

of the retained DIM fraction. However, a pronounced change in the stability of austenite occurred,

appearing as a peak in SHR curves (Fig. 3.16), after annealing at 750–650 °C, being dependent on

aging time. Also, the martensite fraction formed during tensile straining increased (Table 3.3; Fig.

3.17). The dependence of the stability on the previous annealing treatment can be connected with

the precipitation of Cu out of the solid solution, while the austenite stability decreases. From Fig.

3.16, it can be seen that at 700 °C, an annealing duration of 600 s resulted in the maximum SHR

of 2 GPa, though even 100 s annealing at 750 °C caused a similar peak of 2 GPa, but at 650 °C

about 1 h was required to reach that peak level. Thus, at 700–750 °C, time even shorter than 0.5 h

seems to be adequate to result in pronounced precipitation of Cu, but the process continues at least

until 1.5 h, as is evident from Figs. 3.16 and 3.17.

3.3.3 Enhanced strength

As regards the targeted strength, the results of the reversion annealing experiments carried

out for a 304L-3.15Cu steel indicate that the YS can be improved significantly without impairing

its ductility considerably. Hence, excellent YS-TE combinations are possible to achieve, similarly

as reported in numerous studies for various austenitic stainless steels earlier. In Fig. 3.22, some

YS-TE combinations, taken from Table 3.2, are included in the data shown in Ref. [243] for

reversion treated Cr–Ni steels. It can be realized that the present results are quite typical for

reversion treated structures, although at a lower regime. Thus, from the figure it is possible to

estimate and conclude which mechanical properties can be achieved for 304Cu steel, if other

properties are not taken into account.

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Figure 3.22: Yield strength versus total elongation after different reversion conditions compared

to reversion treated 3XX grade austenitic stainless grades (data from Ref. [243]). Järvenpää et al. [243] have recently presented on overview of the Hall-Petch type

relationships presented for reversion-treated austenitic stainless steels and pointed out a broad

scatter between proposed relationships. Shakhova et al. [142] suggested a relationship between the

YS (in MPa) and GS (D in μm) for grain-refined Cr–Ni/Cr–Ni–Cu steels (Eq. (3.2)):

YS = 205 + 395D��.R (3.2)

To check briefly the relevance of that equation, the number weighted average GS and

corresponding calculated and measured YS for different reversion conditions are listed in Table

3.4. For structures created at high annealing temperatures (at 800 °C and above), obviously the GS

has an important contribution, although it seems that the present YS values are relatively low

compared to those reported for 304 [95], 301LN [107, 111, 243] and 204Cu [162] steel at a given

GS. After annealing at lower temperatures, in addition to GS, the retained phases and dislocations

contribute to the YS in addition to grain boundaries so that any simple relationship cannot be

expected. From the present results, it is seen that the measured YS is distinctly higher than the

predicted one.

Table 3.4: Number weighted average GS and corresponding calculated and measured YS after reversion annealing at different conditions (°C-s).

Specimen 650- 650- 700-10 700- 700- 700- 800-10 850-1 900-1

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3600 5400 600 1800 3600

AGS

(Intercept) 0.6 0.9 0.6 0.9 0.7 0.9 1.3 1.24 1.72

Calculated

YS 715 621 715 617 677 621 551 560 506

Measured

YS 812 791 824 653 602 524 416 421 351

It can be pointed out that the GS obtained at low reversion annealing temperatures is neither

very fine, nor uniform. However, recent studies have indicated that complex structures with non-

uniform GS and non-recrystallized retained austenite among the reversed fine grains can provide

good mechanical properties [226, 244–246]. Bimodal GS has been shown to enhance ductility

[246]. Anyhow, it is important to note that an enhanced strength, for instance, YS to the tune of

812 or 791 MPa can be achieved on reversion annealing at 650 °C for 1 and 1.5 h respectively, i.

e. under the conditions where the Cu precipitation has been found to be sufficient for the

antibacterial property even in annealed austenite [217]. At 700 °C, with the same aging treatment,

the YS of about 507–524 MPa can be achieved, which is still almost double compared to that of

the annealed state.

According to a review of Bauer et al. [247], the 316L grade itself can be used as 30% cold

rolled, while its YS is around 790 MPa. However, cold rolling then must be carried out as a separate

operation after the precipitation aging, if the antibacterial property is desired.

Finally, we again highlight that in addition to improved static and fatigue strength, ultrafine

grain-refined austenitic stainless steels have greater open lattice in the position of high angle grain

boundaries and high hydrophilicity [248–252]. Therefore, they have been found to possess

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favorable enhancement of osteoblast functions, protein adsorption on surface and consequently

improved cell attachment, proliferation, and expression level of actin, vinculin and fibronectin.

This may be a factor to favor the adoption of grain-refined austenitic stainless steel as an implant

material.

3.4 CONCLUSIONS

Various reversion treatments were applied to a 304L-3.15Cu austenitic stainless steel with

the objective to improve significantly its yield strength without considerably impairing ductility,

under conditions suitable for the antibacterial property. Microstructures were characterized and

hardness and tensile properties determined. The main observations and conclusions are as follows:

The 71% cold rolling reduction results in the structure containing about 80% deformation-

induced martensite and 20% retained deformed austenite.

Short reversion annealing (1–100 s holding) at 800–900 °C results in fully austenitic grain

structure with the average grain size of few microns, but also larger grains inherited from retained

austenite grains exist.

At lower annealing temperatures of 700–650 °C, the reversion occurred very fast by the

shear mechanism, further followed by the diffusional mechanism. Depending on the annealing

duration (1 s up to 1.5 h), the complex structure consisted of reversed grains with different sizes

(below one micron and few microns), large grains with subgrains (which coalesce and recrystallize

with the continuous recrystallization mechanism), large retained austenite grains and a small

amount of retained martensite (ferrite).

Cu-precipitation occurred during annealing at temperatures of 750–650 °C, concluded from

the decrease in the stability of austenite and increase of strain hardening rate in tensile tests and

the observations made by transmission electron microscopy.

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This study reveals that following the grain size refinement and retained phases obtained by

the reversion annealing treatment at 700–650 °C for 1–1.5 h, the yield strength of the present 304L-

3.15Cu steel increases by 2–3 times that of the annealed structure, while the ductility remains high.

Based on the occurrence of Cu-precipitation, it can be concluded that the antibacterial property is

obtained under these conditions.

3.5 SUMMARY

In this chapter we have fundamentally elucidated here the concept of phase reversion

annealing to obtain a series of revised 304Cu stainless steel. The cold rolled steel was characterized

by majority of deformation-induced martensite and minority of retained deformed austenite.

HRTEM was used to understand the nanoscale crystal structure and precipitation behavior of

nanoscale precipitates. Short reversion annealing (less than100 s holding) at 800–900 °C results in

fully austenitic grain structure with the average grain size of few microns, but also larger grains

inherited from retained austenite grains exist. At lower annealing temperatures of 700–650 °C, the

reversion occurred very fast by the shear mechanism, further followed by the diffusional

mechanism. Depending on the annealing duration (1 s up to 1.5 h), the complex structure consisted

of reversed grains with different sizes (below one micron and few microns), large grains with

subgrains (which coalesce and recrystallize with the continuous recrystallization mechanism),

large retained austenite grains and a small amount of retained martensite (ferrite). Cu-precipitation

occurred during annealing at temperatures of 750–650 °C, concluded from the decrease in the

stability of austenite and increase of strain hardening rate in tensile tests and the observations made

by transmission electron microscopy.

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Chapter 4: On the mechanical behavior of austenitic stainless steel with nano/ultrafine

grains and comparison with micrometer austenitic grains counterpart

In the last chapter, we conducted an in-depth understanding of the effect of phase reversion

on the mechanical behavior of austenitic stainless steel. In sequel to the previous chapter, we

present here a systematic study of grain size effect on the nanoscale mechanical properties of the

austenitic stainless steel. In order to neglected the effect of copper, a copper-free austenitic

stainless steel was selected. The microstructural evolution and deformation mechanism were

critically analyzed using a combination of nanoindenter, optical microscopy (OM), scanning

electron microscopy (SEM) and transmission electron microscopy (TEM) to study deformation

mechanism in this austenitic stainless steel under nanoscale deformation.

4.1 MATERIALS AND EXPERIMENTAL PROCEDURE

4.1.1 Materials

The starting material was a commercial biomedical grade of austenitic stainless steel 18Cr–

8Ni with nominal chemical composition of (wt%) Fe–0.04C–1.52Mn–17.8Cr–8.1Ni–0.005P–

0.005S. To obtain the NG/UFG structure, the solution-treated (1050 °C for 10–15 min) steel was

subjected to severe cold reduction (90% reduction in thickness to ~0.8 mm) via multiple passes.

Subsequently, the strips were cut to dimensions of 0.8 mm × 70 mm × 210 mm and annealed at a

temperature of 800 °C for 10 s, when the cold rolled martensite reverts to NG/UFG austenite via

diffusional reversion mechanism [108, 169, 171, 253, 254]. The grain structure of NG/UFG and

CG austenite was studied by TEM and light microscopy, respectively. For TEM of NG/UFG steel,

the steel was metallographically ground to 50 μm thickness and 3 mm diameter disks were punched.

These disks referred as foils were electropolished in an electrolyte with 10% perchloric acid and

90% ethanol at 25 V for 30 s. The tensile properties of NG/UFG and CG steels were determined

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by tensile test using a CMT5605 tensile machine at room temperature.

4.1.2 Nanoscale deformation

Nanoscale deformation experiments were carried out using a nanoindenter (Keysight

Nanoindenter G200) consisting of a Berkovich three sided pyramid diamond indenter with a

nominal angle of 65.3° and indenter tip diameter of 20 nm under load control and loading rate

(strain rate) conditions. Given that post-mortem electron microscopy was to be carried out to the

study deformation mechanisms, NG/UFG and CG austenitic stainless steels were first ground to

50 μm thickness and 3 mm diameter disks were punched from the foil. As stated above, two sets

of nanoscale deformation experiments were designed with an array of indents of matrix 10✕10

with spacing of 10 μm between the two indents in the center of the NG/UFG and CG austenitic

stainless steels. The first set of nanoindentation experiments were carried out at a fixed loading

rate of 2 μNs−1 and a maximum load of 0.5 mN. Here the objective was to observe any differences

in load-displacement plots that may provide insights on the deformation mechanisms between

NG/UFG and CG austenitic stainless steels. The second set of experiments were carried out at

different loading rates (0.01–1 s−1) to determine strain rate sensitivity. The maximum displacement

was 500 nm. The aim here was to study strain rate sensitivity at low strain rates for the two steels.

After the nanoindentation experiments, the 3 mm disks were electropolished in an

electrolyte containing 10% perchloric acid and 90% ethanol at 25 V for 30 s to obtain an electron

transparent region in the center of the 3 mm disks for the study of deformation mechanism in the

plastically deformed region surrounding the indent. The focus of the TEM study was in the center

of the deformed zone.

4.2 RESULTS AND DISCUSSIONS

The microstructure of CG and NG/UFG austenitic steels is presented in Fig. 4.1. The

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average grain size of solution-treated CG steel was ~55 ± 20 μm, while the NG/UFG steel (cold

rolled to 90% and annealed at 800 °C for 10 s) was in the range of ~200–400 nm. The yield strength

and tensile elongation of CG and NG/UFG steels were as follows: CG (yield strength: 277 ± 41

MPa, elongation: 70 ± 0.8%), NG/ UFG (yield strength: 557 ± 30 MPa, elongation: 44 ± 1%).

Thus, through the phase reversion annealing concept there was significance increase in the strength

of steel because of refinement of grain size.

Figure 4.1: Light and TEM micrographs illustrating the microstructure of coarse-grained (CG)

and nanogrianed/ultrafine-grianed (NG/UFG) austenitic stainless steels with an average grain size of ~55 ± 20 μm and ~200–400 nm, respectively.

4.2.1 Load-controlled nanoscale deformation experiments: load-displacement plots

Load-controlled nanoindentation experiments at a constant loading rate can provide an

insight on the deformation processes as a function of displacement (or strain), which is difficult

from the strain rate-controlled experiments. The underlying reason is that the minimum strain rate

available with the technique is high and any discrete bursts in the load-displacement associated

with nucleation of dislocations or strain-induced transformation cannot be recorded [108, 253,

255].

Fig. 4.2 shows load-displacement plots at a constant loading rate of 2 uNs−1 for CG and

NG/UFG steel, respectively. There is clear and distinct difference between NG/UFG and CG

austenitic stainless steels. Beyond the elastic region (point 1), CG steel shows pop-ins, while no

such behavior is observed for NG/UFG steel at an applied loading rate of 2 uNs−1. The appearance

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of first pop-in is related to nucleation of dislocations during deformation of CG steel, while the

subsequent pop-ins (points 2, 3, 4) represent austensite-to-martensite phase transformation. The

horizontal arrest or pop-in represents geometrical softening caused by martensite variant selection,

minimizing the total energy change during austensite-to-martensite transformation [255]. The

initial region of NG/UFG steel and the region prior to the first pop-in in CG steel follow the power

law relationship (L × h1.5) consistent with the Hertzian contact solution, where L is the applied load

and h is the displacement.

Figure 4.2: Load-displacement plots at constant load rate of 2 uNs−1 for CG and NG/UFG steel,

respectively. 4.2.2 Nanoscale deformation

To study the impact of loading rate, four different stain rates (0.01, 0.1, 0.5 and 1s−1) were

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studied for CG and NG/UFG steels. The indentation strain rate is defined as the displacement rate

divided by the displacement and is given by [256]:

ST = �&� B&

BU (4.1)

where ST is the indentation strain rate, h is the displacement, t is the loading time and dh/dt is the

displacement rate. Here the displacement rate depends on the maximum displacement depth (set

at 500 nm) and the loading time. The data presented is an average of at least 10 experiments with

95% confidence level.

The indentation hardness-strain rate plots for CG and NG/UFG steels at different strain

rates are presented in Fig. 4.3. The hardness data is directly obtained from the instrument. From

the plots in Fig. 4.3, it may be noted that hardness increased with strain rate for both the steels, but

hardness of NG/UFG was greater than CG steel at a constant strain rate. Therefore, we can deduce

that the NG/UFG steel exhibits higher strain rate sensitivity than the CG steel.

Figure 4.3: Hardness versus strain rate plots for CG and NG/UFG austenitic stainless steels at

different strain rates. The strain rate sensitivity, m, is given by [257-259]:

V = √CXYZ[\

= C√CXYZ] (4.2)

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where k is the Boltzmann constant, T is the absolute temperature, σf is the flow stress, H is the

hardness and is generally assumed to be three times of the flow stress, v is the activation volume

and is the rate of decrease of activation enthalpy with respect to the flow stress at a fixed

temperature. v is given by: [257]

^ = √3kàbc dT �a[ � (4.3)

where ST is the strain rate. The strain rate sensitivity parameter, m, and the activation volume, v,

provide insight into the sensitivity of flow stress to strain rate and may provide similarity or

differences in the deformation mechanism [260].

The strain rate sensitivity, m, is 0.147 and 0.086 and for NG/UFG and CG steels, respectively.

The m value of NG/UFG steel is almost twice that of CG steel. If we consider grain size, based on

previous studies [256, 257], the strain rate sensitivity m was 0.022 at grain size of 100 nm and

0.012 at grain size of 1000 nm for Cu. Thus, our value of m is high, such that the grain size effect

is small, implying that the differences in strain hardening between NG/UFG and CG austenitic

stainless steel are small.

Eq. (4.3) was used to calculate the activation volume (v) of the two steels. The calculation

shows that the v value for CG steel is ~19b3, where b is the magnitude of Burgers vector. However,

the v value for NG/UFG steel is ~6b3. These differences point to the differences in the

deformation mechanism between NG/UFG and CG steels (see below).

TEM micrographs depicting representative illustration of deformation-induced processes

in the plastic zone for NG/UFG steel are presented in Fig. 4.4. Nanoscale twinning was the

effective deformation mechanism. Twin boundaries are like grain boundaries and act as obstacles

to the movement of dislocations. The stain hardening effect of twin boundaries acting as strong

boundaries to dislocations is known.

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surrounding the indented region illustrating twinning as the actual deformation mechanism in NG/UFG austenitic stainless steel. (a) bright field micrograph and (b)

dark field micrograph. The inset in (a) is the electron diffraction pattern from the twinned region.

Fig. 4.5 shows representative TEM micrograph illustrating strain-induced martensite in the

plastic zone of CG steel. Martensite contributes to strain-hardening.


surrounding the indented region illustrating strain-induced martensite as the actual deformation mechanism in CG austenitic stainless steel. The inset is the electron

diffraction pattern from the martensite region. The difference in ‘m’ observed between NG/UFG and CG structures is envisaged to

represent differences in deformation mechanism (mechanical twinning versus strain-induced

transformation), such that the NG/UFG structure stabilized austenite and promoted twinning, while

in the CG structure, strain-induced martensitic transformation occurred. Both these mechanisms

are effective strain hardening mechanisms and prevent strain localization and thereby enhance

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ductility (inclusive of uniform elongation). Thus, twinning substituted for martensite nucleation

with a decrease in grain size from CG to NG regime. This is definitely a case of grain size effect

(and strength) and is related to increased stability of austenite with decrease in grain size.

It is pertinent to emphasize here that the nature of deformation mechanism is expected to

alter the surface during nano/microscale motion and impact cell attachment and proliferation. It is

in this context that nano/micromotion is of significance.

4.3 CONCLUSIONS

1) Severe cold deformation of conventional coarse-grained biomedical austenitic

stainless steel followed by annealing for short durations enabled NG/UFG stainless steel to be

obtained with high strength-high ductility combination.

2) There was a distinct difference in the mechanical behavior of load-displacement plots.

In the CG steel, pop-ins reflecting austenite-to-martensite phase transformation were observed,

while they were absent in the case of NG/UFG steel. NG/UFG steel had higher strain rate

sensitivity and lower activation volume than CG steel. Post-mortem electron microscopy of plastic

zone associated with the nano/microscale deformed regions indicated twinning as an active

deformation mechanism in NG/UFG steel. In contrast, strain-induced martensite was the

deformation mechanism in CG steel. Twinning contributed to the ductility of high strength

NG/UFG steel, while strain-induced martensite was responsible for the high ductility of low

strength CG steel.

4.4 SUMMARY

In this chapter, we studied the dependence of grain size on the deformation mechanism in

nanoscale deformation in copper-free austenitic stainless steel. We elucidate here the impact of

grain size on deformation mechanism on copper-free austenitic stainless steel. There was a distinct

difference in the mechanical behavior of load-displacement plots. In the CG steel, pop-ins

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reflecting austenite-to-martensite phase transformation were observed, while they were absent in

the case of NG/UFG steel. NG/UFG steel had higher strain rate sensitivity and lower activation

volume than CG steel. Post-mortem electron microscopy of plastic zone associated with the

nano/microscale deformed regions indicated twinning as an active deformation mechanism in

NG/UFG steel. In contrast, strain-induced martensite was the deformation mechanism in CG steel.

Twinning contributed to the ductility of high strength NG/UFG steel, while strain-induced

martensite was responsible for the high ductility of low strength CG steel.

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Chapter 5: The significance of phase reversion-induced nanograined/ultrafine-grained

structure on the load-controlled deformation response and related mechanism in copper-

bearing austenitic stainless steel

Based on the aforementioned research, the effect of grain size on the deformation

mechanism during nanoscale deformation process in a copper-free austenitic stainless steel has

been achieved. Meanwhile the deformation mechanism that contribute to high strength-high

ductility of copper-bearing austenitic stainless steels has not been explored to the best of our

understanding. The objective of this chapter is to elucidate the deformation behavior of copper-

bearing austenitic stainless steel via post-mortem electron microscopy of nanoindented samples.


The chemical composition of experimental austenitic stainless steel containing Cu (3.15

wt%) is listed in Table 5.1. The steel was made inhouse in a laboratory using standard melting

practice. For cold rolling, the steel was received in the form of a hot rolled sheet, about 3 mm in

thickness. The as-received steel sheet was cold rolled in a laboratory rolling mill to 1 mm thickness

(66.7% reduction) and subsequently annealed at 800 °C for 10 s to obtain NG/UFG structure and

950 °C for 100 s to obtain the CG counterpart. The annealing was carried out in a Gleeble 3800

thermo-mechanical simulator. At 950 °C for 100 s, the final grain size of the experimental steel

was similar to the as-received steel. The microstructure of NG/UFG and CG steel in terms of grain

size was examined by transmission electron microscopy (TEM) and light optical microscopy,

respectively. The annealed steels were subsequently tensile tested according to the ASTM standard

E8. The fracture surface after the tensile tests was studied by scanning electron microscopy (SEM).

Table 5.1: Chemical composition (wt. %) of experimental Cu-bearing austenitic stainless steel.

C Si Mn Cr Ni Cu S P Fe

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0.023 0.55 0.85 17.40 7.32 3.15 0.011 0.025 Balance

Two types of nanoscale deformation experiments were conducted. The first type was

conducted in load-controlled mode at a loading rate of 2 μN·s-1 with the maximum load set to 0.5

mN. Here the objective was to observe any differences in load-displacement plots that may provide

an insight on the deformation mechanism. The second type of experiment was conducted in

displacement-controlled mode, which involved indentation at various constant strain rates in the

range 0.01–1 s-1. The maximum displacement was fixed at 500 nm. Here the aim was to study the

strain-rate sensitivity at low strain rate and compare it with that of Cu-free steel. The

nanoindentation test system (Keysight Nanoindenter G200) consisted of a Berkovich three-sided

pyramidal diamond indenter with a nominal angle of 65.3° and indenter tip diameter of 20 nm. An

array of indents (10 × 10) were made with the indent gap of 10 μm. Post-mortem TEM study of

indented NG/UFG and CG samples was carried out to explore the deformation mechanisms in the

plastic zone surrounding the indented region. This involved removal of indented 3 mm punched

disks from the mount and electropolishing from the side opposite to the indented surface, whereas

the side with the indentations was masked with an aluminum foil. Using this approach, the area

surrounding the indents present around the jet-polished hole, was electron transparent thus

enabling study of the deformation behavior by TEM. During TEM studies, the focus was in the

center of the deformation zone. The data presented here had excellent reproducibility, as confirmed

by a number of experiments for each set of conditions.

5.2 RESULTS

5.2.1 Microstructure of CG and NG/UFG austenitic stainless steels

Fig. 5.1 illustrates light and TEM micrographs of CG and NG/UFG structure, respectively.

The average grain size of CG steel (cold rolled to 66.7% reduction and reversion annealed at

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950 °C for 100 s) was 22 ± 5 μm, while that of NG/UFG steel (cold rolled to 66.7% reduction and

reversion annealed at 800 °C for 10 s) mainly consisted of nanograins of size less than 100 nm and

a few ultrafine grains of size ~100–500 nm.

Figure 5.1: (a) Light and (b) transmission electron micrographs of CG and NG/UFG structure,

respectively in Cu-bearing austenitic stainless steel. 5.2.2 Mechanical properties

Tensile stress-strain plot depicting yield strength and elongation of CG and NG/UFG steels

are presented in Fig. 5.2. The yield strength and elongation for CG steel are 297 MPa and 68%,

respectively, while for the NG/UFG steel are 769 MPa and 38%, respectively. NG/UFG steel has

shown ~2.5 times higher yield strength than CG steel at high level of elongation of 40%.

Figure 5.2: Typical engineering stress-strain curves for CG and NG/UFG Cu-bearing austenitic

stainless steels. 5.2.3 The tensile fracture surface

The fracture surface for NG/UFG and CG structures are presented in Fig. 5.3. In NG/UFG

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austenitic stainless steel, fine striations similar to those observed in fatigue fracture were observed,

except that there is a line-up of voids along the striations. The striations on NG/UFG steel appear

distinctly clear following processing the SEM micrographs with Image Pro software. It appears

that tearing occurred along the striations. In the CG steel, microvoid coalescence leading to cup-

and-cone type fracture was observed, which is commonly observed in ductile metals and alloys.

Interestingly, the microvoids in CG austenitic steel are similar to the line-up of voids observed in

NG/UFG steel along the striations. The microvoids corresponding to the coalescence in CG steel

were only slightly larger in size in comparison to the line-up of voids along the striations in

NG/UFG steel. The difference in the behavior of fracture surface is discussed in section 5.3.3.

Figure 5.3: SEM fractographs at identical magnifications illustrating microvoid coalescence type

of fracture in CG (a and b) and line-up of voids along the striations in NG/UFG (c and d) in Cu-bearing austenitic stainless steels. Figures (b) and (d) are processed

images with Image Pro software to clearly illustrate striations observed in NG/UFG Cu-bearing austenitic stainless steel (c).

5.2.4 Nanoindentation experiments

5.2.4.1 Load-controlled nanoindentation experiments

Load-controlled nanoindentation experiments at a constant loading rate can elucidate the

indentation-induced deformation phenomenon as a function of displacement (or strain) that is

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difficult to achieve from the strain rate-controlled experiments. This is because the minimum strain

rate available with the instrument is relatively quite high, and any discrete bursts in the load-

displacement plots associated with dislocation nucleation or phase transformation cannot be

recorded. Fig. 5.4 presents representative load–displacement plots at constant loading rate of 2

μN·s-1. It may be noted that if the indentation hardness is assumed to be directly related to the

strength, the loading rate can then be regarded to be equivalent to the indentation strain rate.

Figure 5.4: Load-displacement plots at fixed loading rate of 2 μN s−1 for NG/UFG and CG Cu-

bearing austenitic stainless steels obtained via load controlled nanoindentation experiments.

At loading rate of 2 μN·s-1, there was no discontinuity in any of the load-displacement plots

recorded for NG/UFG steel. A representative example is presented in Fig. 5.4a. On the contrary,

the CG steel showed two discontinuous horizontal displacement bursts or arrests (referred as pop-

ins) with the first one at a displacement of ~20 nm and at an applied load of 0.05 mN. The first

pop-in in CG steel corresponds to dislocation nucleation [261, 262] and the subsequent pop-in in

CG steel results from the strain-induced transformation of austenite-to-martensite (see section

5.3.2). The horizontal arrest represents the geometric softening caused by martensite variant

selection, thereby minimizing the total energy change during austenite-to-martensite

transformation process [206]. The initial region of NG/UFG steel and the region prior to the first

pop-in in CG steel follows the power law relationship (P ∝ h1.5) and is in line with the Hertzian

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contact solution.

The similar penetration depth in nanoindentation experiments has been previously

observed when the differences in yield strength are approximately two times. This is attributed to

the indentation effect (IE) after 40–50 μm displacement for both the coarse-grained (CG) and

nanograined/ultrafine-grained (NG/UFG) materials. The IE is expected to be less in NG steel as

compared to CG steel, which depends on the microstructural aspects such as grain size and

dislocation density present within the material [263-266]. The size effects have been attributed to

geometrically necessary dislocations that are introduced by strain gradients [267] and have

recently been discussed for coarse-grained (CG) and ultrafine-fine grained (UFG) materials [268].

5.2.4.2 Strain rate controlled nanoindentation experiments

In addition to load-controlled nanoindentation experiments, CG and NG/UFG steels were

also subjected to depth-sensing nanoindentation experiments at strain rates in the range 0.01–1 s-1

(0.01, 0.1, 0.5, and 1 s-1) to study the strain-rate sensitivity. The indentation strain rate is derived

using the following equation [256]:

ST = �&� B&

BU (5.1)

where, ST is the indentation strain rate, h is the displacement, t is the loading time, dh/dt is the

displacement rate. Here the displacement rate depends on the maximum displacement depth (set

as 500 nm) and the required loading time. The data presented is an average of at least 10

experimental measurements with 95% confidence interval.

The hardness data of the samples that is directly obtained from the nanoindentation

experiments was utilized to determine the strain-rate sensitivity. Hardness vs. strain rate plots for

both the samples are presented in Fig. 5.5 at various strain rates. It can be seen that hardness

increased with increase in strain rate and was greater for NG/UFG steel in comparison to that of

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the CG steel at an identical strain rate. Please note that the unit of hardness is GPa, and hence there

is an appreciable difference in the hardness values of NG/UFG and CG structures at a given strain

rate.

Figure 5.5: Hardness versus strain rate plots for CG and NG/UFG Cu-bearing stainless steels

obtained via strain rate controlled nanoindentation experiments. Please note that the hardness is in GPa. Thus, there is significant difference in the hardness of NG/UFG

and CG Cu-bearing austenitic stainless steel. 5.2.5 Deformation structure

The results of post-mortem electron microscopy study of nanoindented samples are

presented in Figs. 5.6 and 5.7 for CG and NG/UFG steels, respectively. In the CG structure, only

strain-induced martensite was observed (Fig. 5.6). In comparison, a number of representative

electron micrographs are presented for NG/UFG structure, because this essentially constitutes the

focus of the present study. Referring to the NG/UFG austenitic stainless steel, in general, a number

of intersecting nanoscale twins were present in a number of regions (Fig. 5.7a). Furthermore, there

were regions where high dislocation density was observed in the vicinity of nanoscale twins and

the twin boundaries appeared extremely blurred because of the significant dislocation pile-ups

(Figs. 5.7b–d). Thus, we can conclude that there was a clear and distinct transition in the

mechanism of deformation from CG to the NG/ UFG structure.

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in Cu-bearing CG austenitic stainless steel illustrates stain-induced martensite.


in Cu-bearing NG/UFG austenitic stainless steel.

103

5.3 DISCUSSION

5.3.1 Strain-rate sensitivity and activation volume

It is evident from Fig. 5.5 that the hardness of NG/UFG steel was higher than that of CG

steel. For instance, the average indentation hardness of CG and NG/UFG steel at the lowest strain

rate (0.01 s-1) was 0.9 GPa and 1.1 GPa, respectively.

The strain-rate sensitivity is calculated by using the following equation [201, 257]:

V = √3k` ^�⁄ = 3√3k` ^f⁄ (5.2)

where, m is a non-dimensional strain rate sensitivity index, k is the Boltzmann constant, T is the

absolute temperature, σ is the flow stress, H is the hardness (which is generally assumed to be three

times the flow stress) and v is the activation volume, which is the rate of decrease of the activation

enthalpy with respect to the flow stress at a fixed temperature [201, 257]:

^ = √35à bc dTa[ � (5.3)

Strain-rate sensitivity parameter m and activation volume v provide insight into the

sensitivity of flow stress to strain rate, and point out the similarity or difference in the deformation

mechanism between NG/UFG and CG structures. According to the data in Fig. 5.5, the estimated

strain-rate sensitivity values are 0.14 and 0.21 for the CG and NG/UFG structures, respectively,

suggesting that the strain-rate sensitivity (m) of NG/UFG steel is 1.5 times of the CG counterpart.

If we consider grain size based on a previous study [257], the strain-rate sensitivity m was only

~0.022 at a grain size of 100 nm and dropped further to ~0.012 at a grain size of 1000 nm for Cu

and Ni. Thus, the m values of our experimental steel, both for CG and NG/UFG structures, are

comparatively very high, suggesting that the effect of grain size is very small. According to the

definitions of strain-rate sensitivity and activation volume stated in Eq. (5.3), activation volume

(v) for both the CG and NG/UFG steels was calculated using Fig. 5.5. In order to simplify the

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expression, Burgers vector (b) was used. The activation volume of CG steel is ~13b3, and the

corresponding value for NG/UFG steel is ~3b3. Although the activation volume value of CG is ~4

times larger than the NG/UFG structure, the difference is not very large for our stainless steel in

absolute terms. These estimates point to the fact that the effective contribution of deformation

mechanism to strain-hardening behavior in NG/UFG and CG structures is quite similar.

Nevertheless, it is obvious that the indentation response of NG/UFG stainless steel to strain-rate

sensitivity is greater than that of the CG material. Therefore, the strain-rate sensitivity of NG/UFG

stainless steel is larger than that of CG steel. All the same, the differences in strain-rate sensitivity

and activation volume values of CG and NG/UFG structures are quite obvious and must be related

to differences in their deformation mechanisms.

5.3.2 Deformation mechanism in NG/UFG and CG structure

Based on the results in Figs. 5.6 and 5.7, it is concluded that nanoscale twinning is an active

deformation mechanism in NG/UFG steel, whereas strain-induced martensite formation is the

effective deformation mechanism in CG steel. Both mechanisms, however, are responsible for the

high ductility of the steel [269-271]. We know that twin nucleation is promoted by emitted multiple

partial dislocations without dislocation rearrangement. In these situations, dissociation produces a

fixed part and a twinning part; twin growth includes the twinning part experiencing double cross-

slip [272]. At the same time, Figs. 5.7b–d implied that the existing twin boundaries (TBs) behaved

similar to grain boundaries in acting as obstacles to strain propagation [273, 274]. The

strengthening effect of twin boundaries acting as strong barriers to dislocation motion has also

been demonstrated in an in situ TEM observation of the deformation process in a nanocrystalline

Cu specimen [275]. Considering Fig. 5.7d, interestingly, the thickness of twin at the right top is

evidently less than that at the left bottom. This can be attributed to the interaction between slip

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dislocations and the twin boundary. A perfect dislocation in the matrix (a primary plane) can

dissociate into Frank sessile dislocations and Shockley partials, stopped by an obstacle such as

twin boundary [274, 276]:

�� )1g019 → �

C )1g1g19 + �i )1g219 (5.4)

The partial dislocation glides on the conjugate twinning plane, while the sessile dislocation

is stopped at the intersection of the primary and conjugate planes. The pronounced dislocation

accumulation at the TBs leads to multiple consecutive interaction events between dislocations and

the TB, which consequently decreases the thickness of twin lamellae [271] as marked in Fig. 5.7d.

Therefore, twinning is an effective method to improve the strength and ductility of metallic

materials with remarkable strain-hardening ability.

As regards the CG structure, the orientation of martensite transformed from austenite was

determined using the electron diffraction pattern (Fig. 5.6). The orientation relationship between

austenite and indentation-induced martensite followed the Kurdjumov-Sachs (K-S) orientation

relationship, i.e., {111}<110>γ//{011}<111>α′. Considering that each of the 24 K-S variants has a

compression axis and two tensile axes for martensite transformation, termed as Bain distortion,

the variants whose compression axis is almost parallel to the indentation direction have a high

probability of selection during nanoindentation [206].

It is worth noting that TEM of all the NG/UFG and CG samples showed the above

observations, while the areas far from the deformation zone did not exhibit the aforementioned

observations.

Fig. 5.7 illustrates that mechanical twinning was an effective active deformation

mechanism in NG/UFG structure. On the contrary, Fig. 5.6 illustrates that strain-induced α’-

martensite was an active deformation mechanism in the CG structure. Both mechanisms have

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positive effect on preventing strain localization and hence, improve the ductility. Therefore,

twinning replaced the nucleation of strain-induced martensite, when the grain size was

substantially reduced from CG to NG/UFG. This certainly is a consequential effect of the grain

size refinement, thus leading to a noticeable increase in hardness and strength and presumably

enhanced the austenite stability too with decrease in grain size. Therefore, twinning becomes the

preferred mechanism when the weighted average grain size is ~340 nm. Twinning promoted good

ductility in “high strength” NG/UFG steel, but for the “low strength” CG steel, strain-induced

martensite contributed to the high ductility. It is emphasized that twinning is a main factor leading

to the high ductility of “high strength” NG/UFG structure and is an active and governing

deformation mechanism, while for the “low strength” CG structure, the ductility expectedly was

also very high, but without the occurrence of twinning.

Both deformation twinning and strain-induced α’-martensite formation are essentially

strain hardening mechanisms that inhibit local strain and contribute to ductility. In addition, both

mechanisms involve diffusionless shear of a constrained plate-like region of parent crystals.

Twinning must be related to the enhanced contribution of grain boundaries that increase the

stability of NG/UFG austenite, which limits the occurrence of strain-induced martensite, both of

which effectively control the deformation mechanism and ultimately lead to fracture [255, 277]. It

is generally believed that when fcc austenite is transformed into bcc martensite, anisotropic strain

is introduced into the adjacent untransformed austenite to reduce the total strain energy [278, 279].

However, when the austenite grain size is smaller than the martensite lath, such as in NG/UFG

structure, the number of martensite variants participating in an austenite grain is significantly

reduced because of the high strain energy (~850 MJ/m3), thereby reducing the ability to potentially

nucleate the martensite [280].

107

5.3.3 Fracture behavior of NG/UFG and CG

There are interesting differences in the fracture mode of NG/UFG and CG steels (Fig. 5.3).

In NG/UFG steel, where nanoscale twinning was obtained, the striations with line-up of voids are

observed (Figs. 5.3c and d). While in CG steel, when parent austenite transforms into martensite,

the fracture is characterized by microvoid coalescence or dimple fracture (Figs. 5.3a and b).

While this aspect is currently being further studied via fracture toughness tests, the

difference in fracture surface between the two steels merits a preliminary interpretation. We

currently envisage that the fracture process in NG/UFG is step-wise or quasi-static in nature that

produces a striated fracture. Striations had a spacing of ~5 μm. It is likely that the voids grow in

front of the asserted crack. When the crack advances, the tearing of the intervoid area forms a ridge,

which defines a new crack front. This process is repeated as a quasi-static crack growth process,

such that a number of striations are observed.

5.3.4 The relationship between austenite stability and strain energy

The deformation mechanism in NG/UFG structure is related to the high density of grain

boundaries, which led to the strength enhancement of NG/UFG austenite and prevented strain-

induced martensite formation. This must be because of higher austenite stability with the decrease

of grain size that led to the change in deformation mechanism.

Although the thermal stability of austenite is considered to be governed by grain size [277,

281, 282], the effect on mechanical stability is still not clear. As recently reported for TRIP steels

[282], the stability of austenite grains is controlled by local carbon concentration. It is widely

accepted that the transformation of austenite-to-martensite results in an anisotropic strain in

adjacent untransformed austenite. The approximate equidistribution of transformation strain

demands that a number of multivariant transformations coincide in an austenite grain to minimize

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the total strain energy [277]. However, if the austenite grain size is equal to or smaller than the

martensite lath, the possibility of several martensite laths appearing simultaneously in an austenite

grain decreases with the decrease of space. Therefore, it is impossible to minimize the strain energy

by martensitic transformation in NG/UFG steel. In summary, it is impossible to reduce strain

energy when the NG/UFG austenite is transformed into martensite under single deformation mode

due to space constraint effect. The following is a brief explanation on the effect of grain size.

Based on the physical energy and transformation from austenite-to-martensite [281], the

mechanism of austenite stabilization induced by grain refinement is as follows. If austenite is

transformed into martensite by single variant mode, the increase of elastic strain energy is defined

by Eq. (5.5) [281]:

∆kZ = 1 2⁄ �k�S�� + 1 2⁄ �k�S�� + 1 2⁄ �kCSC� (5.5)

where E and ε are Young’s modulus and elastic strain in each lattice plane, respectively. According

to previous studies [280, 281], there are two methods of atomic motion during the fcc-bcc

transformation lattice displacement: the first one is shear deformation of 36% to [110] direction

and the second one is anisotropic deformation accompanying the volume expansion of 4.5%,

which includes 13.9% (ε1) expansion to [001] direction (1st direction), 7.0% (ε2) contraction to

[110] direction (2nd direction) and 1.4% (ε3) contraction to [110] direction (3rd direction); then

E1, E2 and E3 in the three directions are 132.1 GPa, 220.8 GPa and 220.8 GPa, respectively.

According to Eq. (5.5), the evaluated increment of elastic strain energy is ~1840 MJ/m3.

By modifying Eq. (5.5), the increase of elastic strain energy can be obtained according to

the following equation:

∆kZ = 1 2⁄ �k�S��4 �⁄ �� + {1 2⁄ �k�S�� + 1 2⁄ �kCSC�}4 �⁄ � (5.6)

where x is the thickness of the martensite plate and lattice strain is elastically accommodated over

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the space of austenitic grain (grain size: d). On substituting the value of Young’s modulus and

strain into Eq. (5.6), the increased elastic strain energy is:

∆kZ = 1276.14 �⁄ �� + 562.64 �⁄ � (5.7)

For CG steel (average grain size of ~22 μm), ΔEv is ~6 MJ/m3 and for the NG/UFG steel

(average size is ~340 nm), ΔEv increases significantly to ~860 MJ/m3. Therefore, the nucleation

ability of martensite reduces with finer grain size. Considering that multivariant transformation is

difficult in the NG/UFG structure, the transformation from austenite to strain-induced martensite

is inhibited. Therefore, the lattice displacement related to strain can be adjusted by dislocation

sliding and twinning, suggesting that the twinning tendency increases with the decrease of grain

size.

5.3.5 The effect of Cu addition on 304 stainless steel

We now discuss the effect of copper by comparing the results reported here with that of the

Cu-free 304 stainless steel. The following are the differences between Cu-bearing and Cu-free 304

steels: the differences listed below are deduced from the work carried out by our group [283] in a

manner identical to that described here. The strain-rate sensitivity for Cu-free steel (NG/UFG:

0.147; CG: 0.086) calculated by the nanoindentation experiments are not very different from the

values calculated here for both the CG structure and NG/UFG structure in Cu-bearing austenitic

stainless steel (NG/UFG: 0.21; CG: 0.14). The similarity in activation volume of Cu-bearing

(NG/UFG: 3b3; CG: 13b3) and Cu-free (NG/UFG: 6b3; CG: 19b3) is consistent with the

deformation mechanism observed through the post-mortem electron microscopy of nanoindented

Cu-bearing steel in this study and Cu-free NG/UFG and CG steels; i.e., twinning was observed in

NG/UFG structure and strain-induced martensite in the CG structure. There was, however, an

important difference in the degree of twinning (twin density) in NG/UFG structure of Cu-bearing

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(% area fraction: ~25%) and Cu-free (% area fraction: ~12%) 304 steels. The twin density was

significantly greater in Cu-bearing 304 steel. The higher twin density in Cu-bearing steel is related

to the fact that Cu promotes twinning because of its high stacking fault energy at 55 mJ/m2, which

is well above the value facilitating martensitic transformation in steels (<30 mJ/m2) [284].

5.4 CONCLUSIONS

The load-controlled deformation response and strain-rate sensitivity of copper-bearing

austenitic stainless steels (NG/UFG and CG) were studied using a nanoindenter (nanoscale

deformation experiments) and post-mortem electron microscopy. The behavior was compared with

that of the Cu-free austenitic stainless steel. The following are the conclusions:

(1) The strain-rate sensitivity of NG/UFG structure was about 1.5 times (0.21) that of its

CG counterpart (0.14). Using strain-rate sensitivity data, the activation volume of NG/UFG

structure is about one-fourth (3b3) of that of the CG structure (13b3).

(2) Post-mortem TEM studies indicated that the deformation mechanism of NG/UFG and

CG stainless steel was dramatically different. Deformation twinning resulted in high ductility of

“high strength” NG/UFG steel, while in “low strength” CG steel, ductility was also very good but

as a result of strain-induced martensitic transformation.

(3) In NG/UFG structure, the twinning was the active deformation mechanism and the

fracture morphology was characterized by striations (river markings) with line-ups of voids just

along the striations. In contrast, in the CG structure, microvoid coalescence occurred leading to

dimple type fracture with strain-induced martensite as the governing deformation mechanism.

(4) The shift of deformation mechanism from strain-induced martensite in CG structure to

nanoscale twinning in NG/UFG structure is related to the austenite stability that increased with the

finer grain size.

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(5) The addition of Cu had moderate effect on the strain-rate sensitivity and activation

volume of the austenitic stainless steel. However, there was noticeable difference in twin density,

which was significantly greater in Cu-bearing steel compared to the Cu-free steel.

5.5 SUMMARY

In this chapter we elucidate here the deformation mechanism of copper bearing austenitic

stainless steel is involved multiple deformation processes such as strain-induced martensite, and

deformation twins depend on their grain size. The high ductility of low strength CG copper-bearing

austenitic stainless steel is attributed to strain-induced martensitic transformation. While in

NG/UFG structure, the twinning was the active deformation mechanism, and the fracture

morphology was characterized by striations (river markings) with line-ups of voids just along the

striations. In contrast, in the CG structure, microvoid coalescence occurred leading to dimple type

fracture with strain-induced martensite as the governing deformation mechanism. The shift of

deformation mechanism from strain-induced martensite in CG structure to nanoscale twinning in

NG/UFG structure is related to the austenite stability that increased with the finer grain size. The

addition of Cu had moderate effect on the strain-rate sensitivity and activation volume of the

austenitic stainless steel. However, there was noticeable difference in twin density, which was

significantly greater in copper-bearing steel compared to the copper-free steel.

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Chapter 6: The synergistic effect of grain boundary and grain orientation on micro-

mechanical properties of austenitic stainless steel

In previous chapter when we investigate phase reversion annealing, the reversion of

martensite (α’) to austenite (γ) is an important constituent that is believed to control the final

structure and influence the mechanical properties. For the nanoscale deformation, namely,

nano/micro mechanical properties, who will govern them? The synergistic effect of grain boundary

and grain orientation on micro-mechanical properties continues to be unclear. In the study

described here, we have used a combination of nanoindentation and electron back-scattered

diffraction (EBSD), to explore the deformed reverted austenite in copper-free austenitic stainless

steel and elucidate the different group of grain orientation and grain boundary, to understand the

deformation mechanism of reverted austenite and their synergistic effect on micro mechanical

properties.


6.1.1 Material

The chemical composition of experimental medical austenitic SS is listed in Table 6.1. The

as-received steel sheet was cold rolled in a laboratory rolling mill to 1 mm thickness via 66.7%

reduction and subsequently annealed at 1000 °C for 600 s in a tubular resistance furnace under

argon atmosphere, to obtain nearly defect-free equiaxed microstructure sample with yield strength

of ~251 MPa and elongation of ~88%. The microstructure of experimental steel was observed by

scanning electron microscope (SEM; JEOL JSM-7001F) and transmission electron microscope

(TEM, JEM-2100).

Table 6.1: Chemical composition (wt. %) of the investigated medical austenitic stainless steel.

C Si Mn Cr Ni S P Mo N Fe

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0.04 0.34 1.15 18.06 8.33 0.03 0.04 0.051 0.0008 Balance

6.1.2 Nanoindentation and post-mortem characterization

After electro-polishing the samples with a 20% perchloric acid ethanol solution at 25 °C

with an applied voltage of 15 V, the nanoscale deformation experiments were conducted in load-

controlled mode at a loading rate of 6 mN·min-1 with the maximum load set to 1000 μN, dwell

time of 10 s, followed by unloading. The objective was to observe the difference in load-

displacement curves that may provide an insight on the deformation mechanism. The

nanoindentation test system (Hysitron, TI950) consisted of a Berkovich-type three-sided

pyramidal diamond indenter. 30 random arrays containing 25 indents (5 × 5) were made with the

indent gap of 4 μm (to avoid the influence of stress field between each indent [285]). The hardness

and modulus values for each indent were obtained using the method proposed by Oliver and Pharr

[286]. Post-mortem EBSD study of indented samples was carried out at a step size of 50 nm to

explore the grain orientation. The indentation results obtained in the grain interior were statistically

analyzed. The indents were divided in different groups based on their local grain orientation. In

order to study the influence of grain boundary on nanoindentation behavior, the closest distance

from grain boundaries to the indent in 2D surface served as the distance between grain boundary

and indent. Indentations on sample surface close to {001}, {101}, and {111} were selected for

detailed characterization and analysis.

6.2 RESULTS

6.2.1 Microstructure

Fig. 6.1 illustrates SEM (EBSD) and TEM micrographs of initial microstructure of the

experimental steel. It is characterized by a coarse-grained austenite structure (Fig. 6.1a) and some

annealing twins (Fig. 6.1b) with average grain size of ~14 μm measured by EBSD with twins taken

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into consideration. The microstructure at high magnification reflects that nearly defect-free

equiaxed austenite grains with present a few annealing twins were present in the annealed sample

(Fig. 6.1c and d).

Figure 6.1: The SEM micrograph (a), EBSD grain boundary map (b) and TEM micrographs (c,

d) for the original microstructure of the investigated steel. The blue lines in (b) implying grain boundary misorientation greater than 15°.

6.2.2 Nanoindentation behavior

Fig. 6.2 gives two representative EBSD orientation maps and load-displacement curves of

indents in the indented region. As presented in Fig. 6.2a, the indents can be divided into two groups

based on the grain orientation: group I (00, 01, 05, 06, 10, 11, 15, 16, 20, 21) with Miller index (2,

2, 3), and group II (02–04, 07–09, 12–14, 17–19, 22–24) with Miller index (1, 1, 2). Only indents

located in grains with orientations close to {001}, {101}, and {111} were selected based on the

IPF map as shown in Fig. 6.2a, and subjected to statistical analysis. In this case, only indents

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located in group I with the above selected Miller indices {111} were considered, and their load-

displacement curves are presented in Fig. 6.2b. The average modulus and average hardness were

measured to be 167 ± 4 GPa and 4.17 ± 0.09 GPa, respectively.

Figure 6.2: Representative post-mortem EBSD orientation map (a), load-displacement plots for

indents in group I in 2a (b) for the sample. Through the same approach, 208 indents located on the grains with orientations close to

{001}, {101}, and {111} were selected and their elastic modulus and hardness results are

summarized in Table 6.2. The elastic moduli in Table 6.2 varied from 162 to 194 GPa, and are in

the range for austenitic 316L SSs (~150–210 GPa) [287]. The grains close to {101} have the

highest moduli of average ~181 GPa, followed by the grains close to {111} (~179 GPa) and then

{001} (~175 GPa). The component of elastic tensor for the whole sample was calculated as C11 =

197 GPa, C12 = 125 GPa, C44 = 120 GPa, respectively, according to the method in Ref. [288], and

the values were in good agreement with those of single crystal in either 304SS (C11 = 209 GPa,

C12 = 133 GPa, and C44 = 121 GPa) or 316 SS (C11 = 198 GPa, C12 = 125 GPa, and C44 = 122 GPa)

[288, 289]. The values of C12 and C44 are not equal, implying the anisotropy phenomenon in moduli

for our SS. The dependence of hardness on grain orientation in Table 6.2 is relatively weak such

that the grains near {101} (average ~3.94 GPa) and {111} (average ~3.95 GPa) have similar

hardness, and marginally greater than the hardness near {001} (average ~3.88 GPa). The hardness

measured in the range of 3.41–4.53 GPa in our study is greater than the austenitic 316L SSs (~2–

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3.5 GPa) [290], and must be related to relatively higher loading rate (6 mN min-1 in this study is

greater than the commonly used (2.6 mN min-1) [290]). The anisotropy in elastic modulus and

hardness is in qualitatively in agreement with the results of 316L steels reported earlier [290, 291].

Table 6.2: Elastic modulus and hardness for each orientation based on the data of 208 indents. Miller indices

(h, k, l) Quantity of

indents E

(GPa) H (GPa) Number of indents that exhibited more than one pop-in

Near {111}

2, 2, 3� 10 167±4 4.17±0.09 5 3, 2, 2� 15 173±12 3.41±0.14 3 3, 2g, 3� 8 174±10 3.69±0.13 2 2g, 3, 3g� 7 175±6 4.19±0.11 6 2, 2, 3� 4 181±4 3.51±0.16 0 3, 2, 3� 7 182±10 4.31±0.33 7 2g, 3g, 2g� 1 187 3.61 0 3, 2, 3� 1 189 4.28 1 2g, 2g, 3g� 1 177 4.53 1 3, 3, 2� 14 186±5 3.78±0.07 8

Near {101}

0,1, 1g� 4 172±11 3.87±0.09 3 1g, 1g, 0� 1 192 4.38 1 7g, 6g, 1g� 2 194 3.94±0.14 1 2g, 3, 0� 3 172±1 3.90±0.15 3 0, 2, 3� 2 173±6 3.76±0.21 1 0, 1g, 1� 7 185±8 3.66±0.09 0 1, 1, 0� 8 175±6 4.01±0.10 8 3g, 2g, 0� 24 183±5 4.03±0.16 22 6g, 7, 1� 4 180±3 3.95±0.03 4 0, 3, 2� 4 188±9 3.76±0.07 1 1, 5, 6� 2 175±3 4.17±0.04 2 0, 2g, 3� 2 169±10 3.73±0.09 1 1g, 1, 0� 8 186±5 4.00±0.13 7 0, 1g, 1� 6 179±5 4.13±0.11 6 0, 2, 3g� 7 186±8 3.98±0.11 7 0, 1g, 1� 1 180 3.74 0 2, 0, 3g� 1 184 3.69 0 3g, 2g, 0� 5 179±7 3.76±0.09 3 1g, 7, 6g� 21 192±9 3.94±0.12 19 1g, 6, 5� 1 183 4.37 1

Near {001}

5g, 0, 1� 3 177±7 3.80±0.12 2 4g, 0, 1� 3 167±4 3.73±0.19 1 1g, 0, 4g� 3 171±12 3.62±0.23 0 1, 0, 6� 5 167±4 4.02±0.20 4 2g, 8, 1g� 1 176 3.96 1 4g, 1g, 0� 3 176±6 3.95±0.22 3

2g, 3g, 14� 1 182 3.64 0

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3, 12, 2� 3 174±7 3.83±0.06 3 1g, 0, 4� 1 172 3.91 1

3, 2g, 14� 2 190±9 4.02±0.03 2 1g, 4, 0� 1 176 3.92 1

3, 2, 16� 1 172 4.15 1 Note: Thirty regions (25 indents per region) were selected randomly in this experiment. The same

orientation from different region is regard as two individual data source.

Fig. 6.3a–c presents the load (P) - displacement (h) plots for nine representative indentation

tests on individual grains selected based on the highest and average number of pop-ins having the

surface normal close to <111>, <101>, and <001>, denoted by dotted, dashed and solid lines,

respectively. The minimum and maximum number of pop-ins indicates the minimum and

maximum value in this orientation, while the average number of pop-ins reflects the common value

in this orientation. The displacement as a function of loading time (t) is presented in Fig. 6.3d. The

displacement increases with the progress in loading time, while the strain rates were observed to

be similar for indentations close to grains with orientation among {111}, {101}, and {001}.

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Figure 6.3: (a–c) Load-displacement plots from loading to unloading for nine samples

representing indentations in grains near {111}, {001}, and {101}, respectively and (d) load-induced displacement as a function of loading time.

As presented in Fig. 6.3 and Table 6.2, the pop-in effect occurred for almost all the selected

grains. Besides, 48.5% in the group of {111} expressed more than one pop-in, while for groups

{101} and {001}, the percentages are 79.7% and 70.4%, respectively. Even in a grain with more

indents, the percentage did not change (grains with more than 3 indents, the percentages are ~40%,

~80% and ~68% for groups {111}, {101} and {001}, respectively.)

As shown in Fig. 6.3a–c, the circled part implies discontinuous horizontal displacement bursts

or arrests (referred as pop-ins) for the three orientations. The first pop-in occurred from ~7 nm to

~20 nm for different grain orientations, and even in the same group, the pop-in occurred differently

with specific indents. This implies the relationship between pop-in effects and grain orientation.

Voyiadjis et al. [292] has reported that grain boundary also influences the hardening phenomenon

119

during nanoindentation in FCC metals. Thus, besides grain orientation, other factors such as nature

of grain boundary may also influence this behavior, which is discussed in detail in section 6.3.2.

6.3 DISCUSSION

6.3.1 Effect of grain orientation on nanoindentation behavior

The anisotropy in moduli has been observed in uniaxial mechanical tests with single

crystals. The anisotropy in hardness can be qualitatively correlated with the resolved shear stresses

on slip systems estimated from Schmid’s law for uniaxial compression, and (111) expressed the as

lowest absolute value of Schmid factor in perfect FCC crystals, indicating the largest resistance to

deformation [291]. However, either modulus or hardness is not maximum for (111) orientation in

our study, which might be caused by the pop-in effect occurred in our experiments at relatively

higher loading rate.

The pop-in effect is a typical phenomenon observed in SSs, where the first pop-in corresponds

to the nucleation of glissile dislocation loops [261, 262] as the transition from pure elastic to

elastic-plastic deformation. The subsequent pop-ins represent the geometric softening caused by

martensite variant selection, minimizing the total energy change during austenite-to-martensite

transformation process [206]. As reported in previous studies [255, 260, 293, 294] on

nanoindentation behavior for austenitic SS, the martensite was obviously observed through post-

mortem TEM technique. Considering the higher frequency of the second and subsequent pop-ins

observed in {101} (79.7%) and {001} (70.4%) and larger Schmid factor in {101} (0.41) and {001}

(0.41), which reflects the more easily deform and apparently softer in such orientation in perfect

FCC crystal [291], the second and subsequent pop-ins should occur more easily in the initial softer

orientation under large loading rate.

In a perfect FCC crystal, the Schmid’s factor for <101> is greater than <111>, leading to

120

easier slip of dislocations and lower hardness and modulus in {101}. In our case, {101} group was

characterized by more pop-ins, which reflect strain-induced martensite formation during the

loading process, and contributed to hardening in this orientation. Dislocations promote the

nucleation of martensite at high loading rate, resulting in stronger hardening effect in {101}. Thus,

the hardness of group {101} is similar to group {111}.

Thermally activated mechanisms contributing to plastic deformation processes in metals and

alloys are generally quantitatively interpreted by examining the rate sensitivity index, m (a non-

dimensional rate-sensitivity index), and activation volume, v (the rate of decrease of the activation

enthalpy with respect to flow stress at a fixed temperature). The m is defined by [201, 257]:

V = √C'YZ[ (6.1)

where k is the Boltzmann constant, T is the absolute temperature, σ is the uniaxial flow stress [201,

257], and

^ = √35à bc dTa[ � (6.2)

where ST is the instantaneous strain rate and is deduced from equation (6.3) [256]:

ST = �& @&

@U� (6.3)

and the relationship between projecting area A and displacement h for the Berkovich

indenter is listed in equation (6.4) [295]:

o = 3√3ℎ� tan� θ (6.4)

where θ is the half angle (65.3°) of the Berkovich tip. Using equation (6.4), we derive equation

(6.5) to define the relationship amongst project area (A), load (P) and stress (σ) near the surface as

follows:

121

� = st = s

C√C&u vwcu x = s�y.Ri&u (6.5)

The stress (σ) is plotted as a function of instantaneous strain rate (ST) in Fig. 6.4a–c for the

nine representative indentation tests during the loading process.

Figure 6.4: (a–c) Stress - strain rate curves during the loading stage for nine samples representing

indentations on grains near {111}, {001}, and {101}, respectively. The ln σ - lnε˙ data plotted in Fig. 6.4a–c shows a transition between the two linear fitting

segments. Depending on the indentation orientation, the transition point corresponds to the time

point of 0.75–0.95 s during the loading stage, which separates the curves by the part circled in Fig.

6.3d. When we move from a higher strain rate (elastic regime) to a lower strain rate (plastic regime),

the slopes of the three groups of data decrease, implying that the transition in slope is common for

all the indentation orientations. Table 6.3 lists the strain rate sensitivity m fitted from the plastic

regime based on equations (6.1) and (6.2), and Fig. 6.4.

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Table 6.3: The strain rate sensitivity index (m), activation volume (v) calculated from the data of the loading stage for nanoindentation tests for nine indents near {111}, {101} and

{001} grains. Indented plane m υ, b3

Near {111} (2, 2, 3) 0.20±0.03 9.21 (3, 3, 2) 0.13±0.03 13.98 3, 2g, 3� 0.12±0.02 14.57

Near {101} 0, 1g, 1� 0.17±0.04 10.77 1g, 7, 6g� 0.19±0.04 9.49 3g, 2g, 0� 0.19±0.04 9.31

Near {001} 5g, 0, 1� 0.18±0.03 10.08 (3, 12, 2) 0.10±0.03 17.72 (1, 0, 6) 0.12±0.03 14.81

As mentioned above, strain-rate sensitivity (m) of flow stress is an important parameter for

identifying deformation mechanism in materials. Definition of m is based on incremental changes

in strain rate during tests performed at a fixed temperature and fixed microstructure with

corresponding changes in flow stress. Activation volume (v) is the rate of decrease of the activation

enthalpy with respect to flow stress at a fixed temperature, which reflects the dislocation

mechanism controlling the deformation process. In other words, it expresses the volume which is

physically swept by a dislocation during thermally activated process. υ shows a small value (tends

to be atomic volume or less than b3) for the diffusion mechanism, including grain boundary sliding,

Nabarro-Herring and Coble creep, while very large value (the order of 1000 b3) for the forest

mechanism (where a long dislocation segment moves forward by a few Burgers vectors to cut

through a forest dislocation) [296].

As presented in Table 6.3, although the activation volume value of indent (3, 12, 2) is ~2

times larger than for the indent (2, 2, 3), their absolute values in the range of ~10–20 b3 are not

very different in our study, and similar to the results obtained in earlier studies [255, 260, 294].

This indicated that neither conventional dislocation segments passing through dislocation forests

nor diffusional creep processes controls plastic deformation in our case. Thus, the key to the

difference in nanoindentation behavior may lie somewhere else, such as the grain boundary.

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6.3.2 Effect of grain boundaries on nanoindentation behavior

As mentioned above, the properties of the first pop-in and the values of m and v may have

a relationship with the nature of grain boundary. The distance from each indent to the closest grain

boundary in 2D surface was measured based on EBSD orientation maps. The displacement and

load of the first pop-in is plotted in Fig. 6.5. The distributions appear to be random for both

displacement and load for all the indents. Thus, there should be an underlying reason to explain

this phenomenon.

Figure 6.5: The distribution of the first pop-in displacement (a) and load (b) as a function of

distance to grain boundary of the indents located in grains with orientation close to {001}, {101}, and {111}, symbolized with squares, triangles and cross,

respectively. Given that indentation is a plastic deformation process, the plastic zone radius (c) was taken

into account to obtain further insights. Fig. 6.6 is a schematic illustration of plastic zone radius

given by equation (6.6) [297]:

z = { Cs�|[}~ (6.6)

where P is the load when the first pop-in occurred during indentation for specific indent and σYS is

the uniaxial yield strength. Furthermore, the ratio of c/d, is relevant to the properties of the grain

boundary rather than the load applied and has a relationship with the slope of Hall-Petch equation

for shear stress in a given material [298]. With d as the distance of the indent from the grain

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boundary, the ratio of c/d was estimated plotted in Fig. 6.7. To better visualize the distribution,

amplitude version of Gaussian peak function fitted curves were superimposed on the statistical

data. As presented in Fig. 6.7, the distribution of c/d ratio for all the three {001}, {101}, and {111}

groups followed amplitude version of Gaussian peak function distribution. Both {111} and {101}

had peak in distribution at ~1.33, whereas the {001} grains only had an unobvious peak. Thus, this

phenomenon indicates that although the pop-in load/displacement can vary depending on the

distance from the grain boundary as well as the grain boundary concerned, the highest frequency

of ratio c/d was observed nearly 1.33 when the indent is made in an identical grain orientation and

made near a given grain boundary segment. This shows that the ratio c/d is relevant to the

properties of the grain boundary rather than the load applied. In order to have a good compare with

previously study, we attempt to find some reports for this ratio, however, only Wang et al. [298]

reported that the peak of ratio c/d was ~2 and varied from 1.5 to 5 for BCC Nb, which is slightly

greater than the ratio obtained in our case (~1.33). This is because the Hall–Petch slope is steeper

for BCC metals as compared to FCC metals, providing greater resistance to intergranular slip

transmission [298].

Figure 6.6: Schematic illustration for the plastic zone radius (c), where point A is the dislocation

source in the neighboring grain [297].

125

Figure 6.7: Distributions of ratio (c/d) for the indents located in grains with orientation close to

{001}, {101}, and {111}, symbolized with triangles, circles and squares, respectively, superimposed with amplitude version of Gaussian peak function.

For the situation illustrated in Fig. 6.6, at point A (neighbor to grain boundary in another

grain), the maximum shear stress is [298]

τ ≈ [}~� �

@�C (6.7)

and the emitted dislocation from point A led to the emission of high density of dislocations,

which can be explained by [298]:

�� = �� (6.8)

where, r0 is the distance of the source at point A from the grain boundary, and Kc is a critical stress

intensity factor for the emission. Hence, the critical condition for the emission of high density of

dislocations is:

�@�C ≈ ��

[}~�� (6.9)

A rough estimate of the source distance r0 was ~0.1 μm [298], and c/d is ~1.33 for different

grain boundaries, as observed in Fig. 6.7, σYS (= 251 MPa) is the uniaxial yield strength. Kc is

estimated to be ~93.0 MPa·μm1/2, where Kc is a factor relates to shear stress [see Eq. (8)], and is

~2.7 times [299] smaller than the macroscopic Hall–Petch slope in FCC metal. The macroscopic

126

Hall–Petch slope because of lower Kc values here is ~251.1 MPa·μm1/2, and this compares

reasonably well with the experimental value of ~214.8 MPa·μm1/2 in our previous study [137] for

phase reversion SS.

6.4 CONCLUSIONS

To study the nanoscale deformation behavior of a medical austenitic SS, systematic

nanoindentation tests were carried out together with post-mortem EBSD studies. The following

are the conclusions:

(1) The average modulus was calculated for each grain orientation under a large loading rate

condition as: {001} (175 GPa), {111} (179 GPa) and {101} (181 GPa), expressing a similar result.

Similar behavior was observed for hardness, which was 3.88 GPa, 3.94 GPa and 3.95 GPa for

{001}, {111} and {101} grains, respectively.

(2) This phenomenon had a relationship with the number of pop-ins during the loading stage. The

number density and percentage were different for the three orientations, which occurred at {101}

group (79.4%), followed by {001} group (70.4%) and {111} group (48.5%), respectively. As an

initial softer orientation in perfect FCC crystal, group {101} expressed the highest pop-ins

percentage, which contributes to a stronger hardening effect, leading to a similar hardness to {111}

under a large loading rate.

(3) The strain rate sensitivity (m) and activation volume (v) obtained from nanoindentation

had weak dependence on grain orientation and v was ~10–20 b3, indicating that neither diffusional

creep processes nor conventional dislocation segments passing through dislocation forests controls

plastic deformation in our study.

(4) The highest frequency of ratio of c/d was observed as ~1.33 no matter which orientation

the indents located, implying that this ratio is a property related to the grain boundary.

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6.5 SUMMARY

Micro/nano-scale deformation behavior including hardness, elastic modulus, and pop-ins,

was studied in a medical austenitic stainless steel followed by post-mortem EBSD characterization.

Relatively higher hardness and modulus was observed near {101} and more pop-ins occurred in

this orientation at high loading rate. The activation volume (v) obtained from nanoindentation had

weak dependence on grain orientation and was ~10–20 b3, indicating that neither diffusional creep

processes nor conventional dislocation segments passing through dislocation forests controls

plastic deformation in our study. The plastic zone radius (c) and the distance of the indent from the

grain boundary (d) were used to describe the effect of grain boundary on the pop-in effect. The

ratio of c/d meets amplitude version of Gaussian peak function distribution for a given orientation,

whose peak value remains nearly constant for all the orientations.

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Chapter 7: On the impacts of grain refinement and strain-induced deformation on three-

body abrasive wear responses of 18Cr–8Ni austenitic stainless steel

Grain size and phase transformation play significant roles in governing wear resistance of

stainless steel. First, ultra-fine and nano-crystalline grains significantly enhance the wear

resistance of stainless steel by improving its hardness. Second, during the wear process, the

transformed hard martensite on the surface can be easily spalled from the soft untransformed

austenite such that the wear resistance is low at high loads. However, the wear mechanism of

stainless steel during the three-body wear process is still unclear.

In this chapter, the three-body abrasive wear mechanism in stainless steel with different

grain sizes was investigated at room and high temperatures to simultaneously elucidate the effects

of grain size and martensitic transformation on wear performance. The study aimed at determining

the optimal parameters to enhance the wear behavior of 18Cr–8Ni austenitic stainless steel.

7.1 EXPERIMENTAL METHODS

7.1.1 Materials

The chemical composition of 18Cr–8Ni austenitic stainless steel is listed in Table 7.1. The

thickness of the as-received material was 3 mm. Cold rolling was performed up to 30% reduction

at room temperature. Subsequently, the strips were annealed at 900 °C for 3 min in a tubular

resistance furnace filled with argon, followed by quenching in ice-water.

Table 7.1: Chemical composition (wt. %) of the investigated 18Cr-8Ni stainless steel. C Si Mn Cr Ni S P Mo N Fe

0.04 0.34 1.15 18.06 8.33 0.03 0.04 0.051 0.0008 Balance

7.1.2 Microstructural characterization

Standard metallographic techniques were used to ground and polish the specimens to

mirror finish and then electrochemically etched with 60% nitric acid solution. Microstructure was

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observed by scanning electron microscopy (SEM).

The grain structure was further examined by a transmission electron microscope (TEM,

JEM-2100) operating at 200 kV. Thin foils were prepared by twin-jet electropolishing of 3 mm

disks using a solution of 10% perchloric acid in acetic acid as electrolyte at 0 °C. Electron

backscattered diffraction (EBSD) analyses were carried out at a step size of 50 nm or 200 nm to

obtain crystallographic information of samples. The samples for EBSD were electrochemically

etched with 20% perchloric acid-80% ethanol solution operated at 25 °C at an applied potential of

15 V. The boundary with a misorientation larger than 2° was regarded as the boundary of two

crystallographic grains. The contents of martensite and austenite were measured by X-ray

diffraction (XRD) using Cu Kα radiation (PANslytical, Netherlands, 40 kV, 40 mA). The obtained

data were analyzed in Jade software. The volume fractions of austenite and martensite were

calculated by the integrated intensities of (110)α, (211)α, (200)α, and (202)α martensite peaks and

(111)γ, (220)γ, (200)γ, and (311)γ austenite peaks by Eqs. (7.1) and (7.2) [194, 195].

! = 1.4#! #$ + 1.4#!�⁄ (7.1)

$ = 1 − ! (7.2)

where Vγ and Vα are the volume fractions of austenite and martensite, respectively, Iγ and Iα are

the integrated intensities of austenite and martensite peaks, respectively.

7.1.3 Mechanical property tests

The as-received and annealed samples were machined to make tensile samples according

to ISO 6892 standard (length of 140 mm, width of 20 mm and gage length of 65 mm) and tested 3

times for each sample. The uniaxial tensile tests were conducted at room temperature at

engineering strain rate of (5 × 10-4 s-1).

Vickers hardness tests were conducted using a 0.5 kg load with pyramid hardness indenter.

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Hardness data reported (for different tests) is an average of at least ten tests. The nanoindentation

tests were conducted under displacement or depth-controlled mode, where an array of 40 indents

was made at depths in the range of 0–2100 nm to study the hardness distribution beneath the worn

subsurface. The nanoscale hardness was investigated by the situ nanoindentation test system

(Keysight Nano Indenter G200). A Berkovich tip (half angle 65.3°) was used for the hardness tests.

The strain rate was 0.01 s-1. All of the tests were conducted at room temperature.

7.1.4 Three-body abrasive wear tests

In order to simulate the working condition of a rotary drilling rig, three-body abrasive wear

tests were conducted at both 25 °C (room temperature) and 250 °C (assumed as the extreme upper

limit temperature and considered on the response of material). A digital stirrer stirred the specimens

in the abrasive medium (small quartzite stones). The surface of specimens was polished using 1000

mesh SiC grinding paper to ensure similar initial roughness of all the test samples. Fig. 7.1a

illustrates the arrangement and dimensions of specimens for the stirring wear test. Fig. 7.1b is an

image of abrasives used for the stirring wear test. Table 7.2 shows the experimental parameters

used in abrasive wear tests [300].

Figure 7.1: (a) Schematic illustration of the three-body abrasive wear test and dimensions of the

specimens and (b) the shape and size of quartzite stones used in the experiment. Note: t represents the thickness, the thicknesses of the FG and CG samples were ~2 mm and ~3 mm, respectively.

Table 7.2: The experimental parameters for the stirring wear test.

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Abrasive Size Hardness Density Rotating

speed Test duration

Quartzite stone with

quartz content >90

wt. %

φ5~15mm 1100HV 2.64~2.71

g/cm3 2150±20rpm 45min×4cycles

The stirring wear tests were carried out three times for each fine and coarse grained sample

and the average values were considered. Each specimen was tested at 4 different cycles and weight

loss was measured four times after each cycle by balance (Sartorius, SQP, 0.01 mg). The specimens

were cleaned before measuring the weight loss and the duration of each cycle was for 45 min. The

abrasives were changed after each cycle to ensure similar stirring wear conditions. The high

temperature condition was achieved by salt bath furnace. The test was conducted at both room

temperature and high temperature.

7.2 RESULTS

7.2.1 Microstructure

The microstructure of as-received and annealed samples is shown in Fig. 7.2. The SEM

image of Fig. 7.2a shows coarse-grained (CG) austenite structure of as-received steel. The SEM

image presented in Fig. 7.2b indicated that a number of fine grains were obtained in the annealed

sample and the grain size measured according to ASTM standard [301] was 11 (9.0 μm) and 15

(2.0 μm) for as-received CG and annealed FG steels, respectively. The microstructure at high

magnification of CG and annealed FG steels was characterized via TEM and is presented in Fig.

7.3. The TEM micrographs in Fig. 7.3a and b indicated the presence of a number of dislocations

and stacking faults (SF) in CG austenitic stainless steel, inherited from the production process.

When the stainless steel was cold rolled to 30% reduction and annealed at 900 °C, near defect-free

equiaxed austenite grains with some annealing twins were formed in the annealed sample (Fig.

7.3c and d). According to the detected XRD patterns, the martensitic volume fractions in the CG

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and FG annealed samples were below detection limit and 5%, respectively.

Figure 7.2: The microstructure of the as-received CG (a) and FG annealed (b) samples.

Figure 7.3: TEM bright field micrographs of (a, b) as-received CG and (c, d) FG annealed

samples, respectively. Crystallographic information of grain boundaries of CG and FG annealed samples was

analyzed by EBSD (Fig. 7.4a and b), which is of significance in studying grain refinement of

annealed samples. The densities of grain boundaries with misorientation angles 2°–5°, 5°–15°, and

15°–65° for the FG annealed sample were 0.01 μm-1, 0.05 μm-1 and 1.39 μm-1, respectively. The

average grain size of annealed FG steel was 2.0 μm and as-received CG was 9.0 μm, which is

similar to the results acquired by ASTM method (Fig. 7.4c and d).

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Figure 7.4: EBSD results for grain boundary reconstruction maps of austenite in as-received CG

(a) and FG annealed (b) samples combined with grain size distribution fraction in as-received CG (c) and FG annealed (d) samples.

7.2.2 Mechanical properties

Table 7.3 presents the mechanical properties of steels. The yield strength and elongation of

the CG sample were found as 281 MPa and ~52%, respectively, whereas the phase-reversion FG

annealed sample exhibited higher yield strength (380 MPa) and similar elongation (~57%) [134].

Table 7.3: The measured mechanical properties of the investigated steels. Steels σs, (MPa) σb, (MPa) A, (%) Hardness, (HV0.5)

As-received CG 281±32 644±19 52±2.7 174±8

FG annealed 380±37 813±31 57±5.2 206±9

7.2.3 Three-body abrasive wear performance

At room temperature (25 °C), the weight loss of both the CG and FG samples increased

gradually with the prolonged test time. The value of measured weight loss of FG annealed sample

was always higher than the as-received CG steel (Fig. 7.5a). However, there were subtle and

obvious differences between these two samples. The weight loss of CG sample initially increased

and then remained nearly constant (19–20 mg per cycle) with increase in the number of cycles.

134

The weight loss per cycle of FG annealed sample decreased sharply with increase in the number

of cycles and then remained nearly constant (~10 mg per cycle), as shown in Fig. 7.5b. The weight

loss for FG annealed sample was smaller than the CG sample when the test time was adequate.

When the test temperature was increased to 250 °C, the weight loss for both samples increased

gradually with increase in test time, and the value of weight loss of both the samples tested at

250 °C was larger compared to the test at room temperature. It was interesting to note that the

weight loss of FG annealed sample was smaller than the original CG sample when the test time

was larger than 90 min (Fig. 7.5c). This is related to the unequal weight loss rate of samples tested

at 250 °C. The weight loss rate of CG sample increased gradually with time. On the contrary, the

weight loss per cycle of FG annealed sample was nearly constant (20–22 mg per cycle) with

increase in the number of cycles (Fig. 7.5d). Hence, the wear resistance of FG annealed steel was

better than the CG steel, when tested at high temperature and/or for a longer time.

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Figure 7.5: The average accumulated weight loss (a, c) combined with their weight loss rate (b, d) of the investigated steels in room temperature (a, b) and high temperature (c, d)

stirring wear test. Note: The error bars here were the results from three tests.

The worn surfaces of both the steels after stirring wear test at both the temperatures were

similar. Fig. 7.6 displays the edge (left and/or right view of wear part) and center (front and/or

back view of wear part) morphologies of worn surfaces of both the samples after room or high

temperature wear test. Given that the specimens were rotated around the center axis, the line speed

increased from their center to the edge, therefore, the degree of wear from the edge to the center

decreased gradually. It was clear that the mode of wear was similar for both the samples after the

wear test. The mode may vary from mild to severe wear, and hence, the exact transition period

from mild to severe was very difficult to define [302]. From SEM micrographs of the worn surface

morphology, it was clear that microploughing was the main wear mechanism at the edge of the

specimens (Fig. 7.6a, c, e and g), while microcutting was the main wear mechanisms in the middle

of the specimens (Fig. 7.6b, d, f and h).

136

Figure 7.6: The SEM pictures for worn surface morphology of edge part (left and/or right view

of wear part) and center part (front and/or back view of wear part) of investigated samples in both the room and high temperature work condition stirring wear test.

137

The hardness values of samples after wear tests at both room and high temperature are

presented in Table 7.4. The surface hardness of the CG sample after room temperature stirring

wear test was increased from 174 to 291 HV0.5, whereas in the case of annealed FG sample, it

increased from 206 to 309 HV0.5. After performing wear test at 250 °C, the hardness values of

CG and FG annealed samples were found as 251 and 255 HV0.5, respectively.

Table 7.4: The hardness of the worn surface for investigated steels (HV0.5). Steels

Before

stirring

After room temperature

stirring

After high temperature

stirring

As-received

CG 174±8 291±9 251±12

FG annealed 206±9 309±7 255±10

The hardness distributions beneath the material sub-surface (200 nm–2100 nm) for all

samples were measured via nanoindentation. Fig. 7.7 reveals that hardness of all samples

decreased with increase of indentation depth. It was found that at similar indentation depth, the

hardness of FG annealed sample was higher than CG sample. Comparing with the samples before

the wear test, the hardness of both the samples after the wear test was greater and is related to

hardening induced by the transformation of austenite to strain-induced martensite [303]. The

hardening effect of the samples after wear test at high temperature was remarkably lower than that

at room temperature, and was a consequence of variation in martensite fraction in the worn surface

of steel samples.

138

Figure 7.7: The harness versus depth plots of subsurface deformation layer of FG annealed

sample and as-received CG sample before (a), after the wear tests at room temperature (b) and high temperature (c).

Table 7.5 shows the results of the average volume fractions of martensite in worn surfaces

of the samples before and after three-body abrasive wear tests. It is noticeable that at room

temperature, ~7.0 vol. % of martensite formed in the FG annealed sample after wear test, which is

remarkably higher than the amount of martensite of ~3.0 vol. % in CG samples (when the phase

volume fraction was less than 5 %, it can be considered that this value was no longer reliable).

However, when the wear test temperature was increased to 250 °C, nearly no martensite was

formed on the worn surfaces of the samples.

Table 7.5: The average martensite volume percentage of FG annealed sample and as-received CG sample before and after wear tests (vol. %).

Sample Before stirring After room temperature

(25 °C) stirring

After high temperature

(250 °C) stirring

As-received

CG Below detection limit 3.0 Below detection limit

FG annealed 5.0 7.0 Below detection limit

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7.3 DISCUSSION

7.3.1 Effects of grain refinement on mechanical properties in austenitic stainless steel

According to the mechanical properties and hardness results presented in Table 7.3 and Fig.

7.7, it is clear that the FG annealed sample had higher yield strength and hardness, and similar

elongation in comparison to CG steel specimen. The strengthening mechanism in FG annealed

steel is related to grain refinement, which can be described by the well-known Hall-Petch

relationship [304, 305].

σ = �� + 5�� ⁄ (7.3)

where, σ is the yield strength of experimental steel, σ0 is the lattice friction stress, which is constant

for an identical material, k is a constant, and d is the average grain size. In the present case, σ0 =

208.62 MPa and k = 214.75 MPa·μm1/2, which are different from the values proposed by

Karjalainen et al. [306] (σ0 = 150 MPa and k = 537.59 MPa·μm1/2 converted from k = 17

MPa·mm1/2) since the chemical compositions and grain size measured methods were different for

us, but still fine grain size would represent higher strength.

Generally, a material with higher yield strength is expected to exhibit higher hardness since

hardness is proportional to the yield strength [307]. Forouzan et al. [98] proposed a model for

hardness measurement based on the Hall-Petch relationship as shown in Eq. (7.4):

HV = HV� + k��.R (7.4)

where HV is the hardness of the material, HV0 is a constant and depends on the hardness of a single

crystal, d is the average austenite grain size, and k is a constant for a material. In the present case,

HV0 = 134.22 HV0.5 and k = 109.35 HV0.5∙μm1/2, which are slightly lower than the values

proposed by Forouzan et al. [98] (HV0 = 172.32 HV10 and k = 115.83 HV10∙μm1/2) since the

chemical compositions and grain size measured methods were different for us, but still fine grain

140

size would represent higher hardness.

Furthermore, a uniform grain structure with near defect-free equiaxed grain was obtained

in the annealed FG sample via phase reversion annealing method (Fig. 7.3c and d). The elongation

of FG annealed sample was similar to CG steel and is attributed to high dislocation storage capacity

of near defect-free austenite grains [134, 308]. Thus, cold reduction and annealing process

produced FG austenitic stainless steel with combination of higher yield strength, higher hardness

and higher elongation.

7.3.2 Effects of grain refinement and test temperature on wear resistance in austenitic

stainless steel

It is well known that the hardness of the material has a significant impact on wear properties

[309]. The hardening ability of a material is an important aspect because it influences the hardness

of the surface during wear [310]. It is evident from Table 7.3 and Fig. 7.7 that the FG annealed

sample had a higher initial hardness and worn sub-surface hardness than CG sample under

identical test conditions; thus manifesting better wear resistance.

However, there were several subtle changes during the wear process, such as the weight

loss rate of FG annealed sample was noticeably high in the first two cycles and then decreased to

a remarkably lower level with increase in the number of cycles at room temperature, which can be

ascribed to the presence of martensite in the sample. Although the BCC martensite structure was

not observed from the EBSD results of phase reversion annealed FG sample, ~5 vol% strain-

induced martensite formed during cold deformation did not transform into austenite during small

duration of annealing based on the XRD results. The block-like martensite, which is harder than

the FCC austenite, formed an interface with high stress concentration during wear process and can

be easily spalled, thus resulting in a larger weight loss rate (~30–35 mg/ cycle) at the beginning of

141

the wear test. In contrast, the CG sample was composed of typical CG austenite grains, which are

relatively softer than martensite and no obvious interface for stress concentration can be easily

ploughed without obvious spalling, resulting in a lower weight loss rate after the commencement

of wear test. With progress in time, the new lath-like strain-induced martensite formed inside both

the samples. However, the distribution of martensite formed in FG annealed sample was more

profound than in CG specimen as the hardness results presented in Table 7.4 (smaller data scatter),

because of high density of grain boundaries [87, 311, 312], which was beneficial in enhancing the

hardness and interface stress distribution would be more uniform would not be easily spalling as

compared to the CG sample. This phenomenon led to a sharp decrease in weight loss rate of the

annealed FG sample during subsequent cycles.

At high temperature (250 °C), the hardness of FG annealed sample after wear continued to

be higher than the CG sample, but there was no obvious strain-induced martensite formation at

this time, as represented in Table 7.5. Zhang’s study [313] indicated that the lower cold rolling

reduction and higher deformation temperature would inhibit the strain-induced martensite

formation. Furthermore, when the temperature was increased from 25 °C to 200 °C, the

deformation mechanisms in FG (~1.5 μm) samples varied from primary TRIP þ minor TWIP at

25 °C to primary TWIP + minor TRIP at 200 °C, in CG (~9 μm) samples varied from TRIP at

25 °C to TWIP at 200 °C in the other research work of our group [314]. This is most likely to be

because of higher temperature, which may increase stacking fault energy in austenitic stainless

steel and change the deformation mechanism from strain-induced martensite to deformation-

induced twinning [255]. Nanoscale deformation twins also enhanced the hardness of sample with

increase in wear resistance. The FG annealed sample had a constant weight loss rate (20–22

mg/cycle) during the wear test, because the initial retained strain-induced martensite was tempered

142

to a softer structure with hardness similar to the twinned structure. Once tempered martensite

experiences spalling during the first several cycles, the newly formed twins were able to form a

new surface with similar hardness. In the CG sample, the twins were more difficult to form and if

formed, they would be large and accumulate in certain part and formed interface with higher stress

concentration can be spalled easier, such that wear resistance was worse at high temperature.

7.3.3 Wear mechanisms of austenitic stainless steel

In comparison to traditional wear processes, namely, sliding wear [315, 316] and rolling-

sliding wear [317, 318], during three-body abrasive wear, abrasives interact with the material

surface under the action of external force. In view of uneven shape and direction of multi-angle

abrasive particles, different types of wear morphologies are formed on the surface.

It is clear from Fig. 7.6 that the major mechanism for the edge (Fig. 7.6a, c, e and g) and

center (Fig. 7.6b, d, f and h) parts were microploughing and microcutting, respectively. Fig. 7.8

[319] illustrated the wear mechanisms for microploughing and microcutting. The spalling of front

extruded accumulation in microploughing and debris in microcutting was the main reason for the

weight loss of each sample. Although the furrow part was similar in both mechanisms, the

differences between these two mechanisms in the view of morphology were the extruded

accumulation was less in microcutting as compared to microploughing and the debris formed in

microcutting was much easier to spalling as compared with the front extruded accumulation

formed in microploghing. Besides, the microploughing was formed mainly by the blunter abrasive

with a low contact angle with material worn the same place for several times, the microcutting was

formed mainly by the sharper abrasive with an identical contact angle with material worn the

surface as a knife worked. Furthermore, the identical region of both samples experienced a similar

wear mechanism. Thus, the shift in wear mechanism must be related to other reasons except the

143

grain size and test temperature.

Figure 7.8: Schematic illustrations for wear mechanisms in wear process. (a) Microploughing;

(b) Microcutting. According to the relationship between angular velocity and linear velocity, and impulse

144

theorem, the following equations can be given:

�̂ = �� ∙ � (7.5)

�̂ = �� ∙ � (7.6)

−�� ∙ � = V� ∙ �̂ � − V� ∙ �̂ (7.7)

−�� ∙ � = V� ∙ �̂ � − V� ∙ �̂ (7.8)

where ω is the angular velocity of the specimen, vi and ^,� are the linear velocities of the local

specimen before and after impulse with abrasives, respectively, mi is the mass of the local specimen

and similar for identical regions, t is the impulse time, and Fi is the external force induced by

impulse with abrasives. The angular velocity of the local specimen would change slightly (Δω)

after impulse with abrasives, thus, the external force can be solved by Eqs. (7.5)–(7.8) as follows:

�, = − ��∙��∙∆�U (7.9)

The only difference was ri in Eq. (7.9) for the identical regions in this test, thus, the F would

be differing for each region. This means that only the external force had an influence on the wear

mechanism, and played a more significant role than grain refinement and test temperature on wear

mechanism.

7.4 CONCLUSIONS

In the present study, FG 18Cr–8Ni austenitic stainless steel was obtained by phase

reversion annealing. The mechanical properties and wear resistance of FG annealed steel were

studied by tensile tests and three-body abrasive wear tests at room and high temperature. The

quartzite stones (quartz content over 90 wt. %) with diameter 5–15 mm and hardness of 1100HV

served as abrasive. The following are the conclusions:

(1) Through phase reversion annealing method, FG annealed austenitic stainless steel with

small near defect-free equiaxed grains were obtained with high yield strength-elongation

145

combination.

(2) Although the wear performance of annealed FG steel was worse in relation to CG steel

at the initial stages in the three-body abrasive wear test at room temperature, it was superior as

compared to CG steel at high temperature.

(3) Microploughing was the main wear mechanism at the edge part of sample in the three-

body abrasive wear tests, whereas microcutting was key wear mechanism in the center part of

sample. The shift in wear mechanism resulted from the external force, which played a more

significant role in governing the wear mechanism compared to grain refinement and test

temperature.

7.5 SUMMARY

In this chapter, the phase revised fine grained 18Cr–8Ni austenitic stainless steel was

obtained. The primary objective of the present study was to elucidate the wear performance of

fine-grained austenitic stainless steel through three-body abrasive wear tests at room and high

temperatures and compare with the coarse-grained counterpart. The quartzite stones (quartz

content over 90 wt. %) with diameter 5–15 mm and hardness of 1100HV were used as the abrasive

in three-body abrasive wear tests. The study demonstrated that the microstructure consisting of

near defect-free and equiaxed fine austenite grains with high yield strength and elongation

exhibited superior wear resistance at high temperature (250 °C), which is attributed twinning

induced plasticity deformation in fine austenite grains. The wear mechanism varied as a function

of distance from the center of the steel sample and was characterized by microploughing and

microcutting.

146

Chapter 8: Conclusions and future work

8.1 CONCLUSIONS

Austenitic stainless steel with and without Cu including nano/ultrafine grain and coarse

grain were used to find the interplay between grain structure and chemical elements on strength

mechanism and nanoscale deformation mecahnism.

Secondly, for the aim of engineering application, we apply a three-body abrasive wear tests

to study the wear performance of austenitic stainless steel with different grain size under different

work condition.

8.1.1 Improving the yield strength of an antibacterial 304Cu austenitic stainless steel by

the reversion treatment

The 71% cold rolling reduction results in the structure containing about 80% deformation-

induced martensite and 20% retained deformed austenite.

Short reversion annealing (1–100 s holding) at 800–900 °C results in fully austenitic grain

structure with the average grain size of few microns, but also larger grains inherited from

retained austenite grains exist.

At lower annealing temperatures of 700–650 °C, the reversion occurred very fast by the

shear mechanism, further followed by the diffusional mechanism. Depending on the

annealing duration (1 s up to 1.5 h), the complex structure consisted of reversed grains

with different sizes (below one micron and few microns), large grains with subgrains

(which coalesce and recrystallize with the continuous recrystallization mechanism), large

retained austenite grains and a small amount of retained martensite (ferrite).

147

Cu-precipitation occurred during annealing at temperatures of 750–650 °C, concluded

from the decrease in the stability of austenite and increase of strain hardening rate in

tensile tests and the observations made by transmission electron microscopy.

This study reveals that following the grain size refinement and retained phases obtained

by the reversion annealing treatment at 700–650 °C for 1–1.5 h, the yield strength of the

present 304L-3.15Cu steel increases by 2–3 times that of the annealed structure, while the

ductility remains high. Based on the occurrence of Cu-precipitation, it can be concluded

that the antibacterial property is obtained under these conditions.

8.1.2 On the mechanical behavior of austenitic stainless steel with nano/ultrafine grains

and comparison with micrometer austenitic grains counterpart and their biological

functions

Severe cold deformation of conventional coarse-grained biomedical austenitic stainless

steel followed by annealing for short durations enabled NG/UFG stainless steel to be

obtained with high strength-high ductility combination.

There was a distinct difference in the mechanical behavior of load-displacement plots. In

the CG steel, pop-ins reflecting austenite-to-martensite phase transformation were

observed, while they were absent in the case of NG/UFG steel. NG/UFG steel had higher

strain rate sensitivity and lower activation volume than CG steel. Post-mortem electron

microscopy of plastic zone associated with the nano/microscale deformed regions

indicated twinning as an active deformation mechanism in NG/UFG steel. In contrast,

strain-induced martensite was the deformation mechanism in CG steel. Twinning

contributed to the ductility of high strength NG/UFG steel, while strain-induced

martensite was responsible for the high ductility of low strength CG steel.

148

8.1.3 The significance of phase reversion-induced nanograined/ultrafine-grained


bearing austenitic stainless steel

The strain-rate sensitivity of NG/UFG structure was about 1.5 times (0.21) that of its CG

counterpart (0.14). Using strain-rate sensitivity data, the activation volume of NG/UFG

structure is about one-fourth (3b3) of that of the CG structure (13b3).

Post-mortem TEM studies indicated that the deformation mechanism of NG/UFG and CG

stainless steel was dramatically different. Deformation twinning resulted in high ductility

of “high strength” NG/UFG steel, while in “low strength” CG steel, ductility was also

very good but as a result of strain-induced martensitic transformation.

In NG/UFG structure, the twinning was the active deformation mechanism and the

fracture morphology was characterized by striations (river markings) with line-ups of

voids just along the striations. In contrast, in the CG structure, microvoid coalescence

occurred leading to dimple type fracture with strain-induced martensite as the governing

deformation mechanism.

The shift of deformation mechanism from strain-induced martensite in CG structure to

nanoscale twinning in NG/UFG structure is related to the austenite stability that increased

with the finer grain size.

The addition of Cu had moderate effect on the strain-rate sensitivity and activation

volume of the austenitic stainless steel. However, there was noticeable difference in twin

density, which was significantly greater in Cu-bearing steel compared to the Cu-free steel.

149

8.1.4 The synergistic effect of grain boundary and grain orientation on micro-

mechanical properties of austenitic stainless steel

The average modulus was calculated for each grain orientation under a large loading rate

condition as: {001} (175 GPa), {111} (179 GPa) and {101} (181 GPa), expressing a

similar result. Similar behavior was observed for hardness, which was 3.88 GPa, 3.94

GPa and 3.95 GPa for {001}, {111} and {101} grains, respectively.

This phenomenon had a relationship with the number of pop-ins during the loading stage.

The number density and percentage were different for the three orientations, which

occurred at {101} group (79.4%), followed by {001} group (70.4%) and {111} group

(48.5%), respectively. As an initial softer orientation in perfect FCC crystal, group {101}

expressed the highest pop-ins percentage, which contributes to a stronger hardening

effect, leading to a similar hardness to {111} under a large loading rate.

The strain rate sensitivity (m) and activation volume (v) obtained from nanoindentation

had weak dependence on grain orientation and v was ~10–20 b3, indicating that neither

diffusional creep processes nor conventional dislocation segments passing through

dislocation forests controls plastic deformation in our study.

The highest frequency of ratio of c/d was observed as ~1.33 no matter which orientation

the indents located, implying that this ratio is a property related to the grain boundary.

8.1.5 On the impacts of grain refinement and strain-induced deformation on three-body

abrasive wear responses of 18Cr–8Ni austenitic stainless steel

Through phase reversion annealing method, FG annealed austenitic stainless steel with

small near defect-free equiaxed grains were obtained with high yield strength-elongation

combination.

150

Although the wear performance of annealed FG steel was worse in relation to CG steel at

the initial stages in the three-body abrasive wear test at room temperature, it was superior

as compared to CG steel at high temperature.

Microploughing was the main wear mechanism at the edge part of sample in the three-

body abrasive wear tests, whereas microcutting was key wear mechanism in the center

part of sample. The shift in wear mechanism resulted from the external force, which

played a more significant role in governing the wear mechanism compared to grain

refinement and test temperature.

8.2 FUTURE WORK

In austenitic stainless steel research, the kinetic origin of nanoscale twins was observed,

and hypothesis was proposed meanwhile due to lack of direct evidence we cannot solidified it. In-

situ TEM or other technique which can provide direct evidence is required for further research.

Since experimental data has been sufficiently acquired, we will continue in thermodynamic

simulation of deformation behavior and interaction between precipitates and dislocation to further

facilitate the understanding in strengthening mechanism. Reverted austenite is essential for further

improvement of ductility of austenitic stainless steel.

For the aim of engineering application, future research will focus on more complex work

condition to achieve better guidance.

151

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Vita

Chengyang Hu earned his bachelor’s degree in Metallurgical Engineering from Wuhan

University of Science and Technology (WUST), China in 2015. In 2018 he earned his master’s

degree in Material Science and Engineering from Wuhan University of Science and Technology

(WUST). Later he joined the University of Texas at El Paso to pursue his doctoral degree in

Material Science and Engineering.

Chengyang Hu received graduate research assistantship to pursue doctoral research for the

entire duration of the doctoral program.

Chengyang Hu authored several peer-reviewed publications in international journals. The

publications related to his dissertation topic are the following:

1. Wear, 446(2020): 203181.

2. Journal of the Mechanical Behavior of Biomedical Materials, 118(2021): 104473.


4. Materials Science and Engineering: A, 793(2020): 139885.


Contact Information: [email protected]

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