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University of Texas at El Paso University of Texas at El Paso
ScholarWorks@UTEP ScholarWorks@UTEP
Open Access Theses & Dissertations
2021-08-01
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
Follow this and additional works at: https://scholarworks.utep.edu/open_etd
Part of the Mechanics of Materials Commons
Recommended Citation Recommended Citation Hu, Chengyang, "Structure-Property Relationship In High Strength- High Ductility Combination Austenitic Stainless Steels" (2021). Open Access Theses & Dissertations. 3270. https://scholarworks.utep.edu/open_etd/3270
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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
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
6.4 CONCLUSIONS ................................................................................................................ 126 6.5 SUMMARY ...................................................................................................................... 127
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
7.4 CONCLUSIONS ................................................................................................................ 144 7.5 SUMMARY ...................................................................................................................... 145
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
Figure 3.7: Microstructure obtained after annealing at 650 °C for 5400 s. Martensite red in the
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
Figure 3.21: Time-Temperature- Reversion (TTR) diagram and an example of the reversion
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
Figure 4.5: Post-mortem transmission electron microscopy of the plastically deformed region
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
Figure 5.7: Post-mortem electron microscopy of the plastic zone surrounding the indented region
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
Figure 2.3: Time-Temperature- Reversion (TTR) diagram and an example of the reversion
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
61
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).
Figure 3.7: Microstructure obtained after annealing at 650 °C for 5400 s. Martensite red in the
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.
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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
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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
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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.
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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.
Figure 3.21: Time-Temperature- Reversion (TTR) diagram and an example of the 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
90
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]
^ = √3k`abc 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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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.
Fig. 4.5 shows representative TEM micrograph illustrating strain-induced martensite in the
plastic zone of CG steel. Martensite contributes to strain-hardening.
Figure 4.5: Post-mortem transmission electron microscopy of the plastically deformed region
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.
5.1 MATERIALS AND EXPERIMENTAL PROCEDURE
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
99
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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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.
Figure 5.7: Post-mortem electron microscopy of the plastic zone surrounding the indented region
in Cu-bearing NG/UFG austenitic stainless steel.
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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]:
^ = √35`a 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].
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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 MATERIALS AND EXPERIMENTAL PROCEDURE
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
^ = √35`a 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.
123
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
124
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
129
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.
130
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.
131
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).
133
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.
135
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
structure on the load-controlled deformation response and related mechanism in copper-
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
References
[1] Lu, S. 2013. Introduction to stainless steel. Beijing: Chemical Industry Press. (in Chinese)
[2] Hu, B., and S. Tu. Iron Steel Technology, 6(2008): 24-2 6. (in Chinese)
[3] Ji, W. Science Technology Information, 4(2012): 455. (in Chinese)
[4] Zhang, W. 2010. Stainless steel and its heat treatment.Shenyang: Liaoning Science and
Technology Press.(in Chinese)
[5] Lu, S., and T. Zhang. 1995. Stainless steel. Beijing: Atomic Energy Press. (in Chinese)
[6] Zhao, L., B. Wang, Y. Ma, J. Cao, G. Ding, J. Zhao, J. Zhang, and H. Wang. Iron Steel,
38.2(2003): 22-23. (in Chinese)
[7] Yang, J., L. Wang, and J. Sheng. Iron Steel, 43.3(2008): 29-32. (in Chinese)
[8] Wang, W., and H. Mao. Baosteel Technology, 4(2005): 56-60. (in Chinese)
[9] Bi, H., X. Li, X. Ou, and Y. Wu. Baosteel Technology, 4(2007): 1-4. (in Chinese)
[10] Wang, W., H. Yu, and H. Mao. Steel Rolling, 23.6(2006): 8-11. (in Chinese)
[11] Zhang, W., and C. Yang. Mechanic Management Development, 4(1999): 46-48. (in Chinese)
[12] Rong, H. China Stainless Market and Information, 24(2004): 19-20. (in Chinese)
[13] Yang, J., L. Zhou, J. Zhang, and H. Su. Journal of Inner Mongolia University of Science and
Technology, 33.4(2014): 353-356. (in Chinese)
[14] Ryoo, D., N. Kang, and C. Kang. "Effect of Ni content on the tensile properties and strain-
induced martensite transformation for 304 stainless steel." Materials Science and
Engineering: A 528.6 (2011): 2277-2281.
[15] Xiang, H., and X. Gu. Metal World, 1(2013): 30-36. (in Chinese)
[16] Wang, W., S. Chen, W. Yan, L. Zhao, Y. Shan, and K. Yang. Transaction Materialrs Heat
Treatment, 7( 2010): 59-65. (in Chinese)
152
[17] Ma, S., Z. Zhang, and S. Chu. Journal of Iron Steel Research, 20.12(2008): 10-13. (in
Chinese)
[18] Gao, Y., G. Chen, and W. Jin. Special Steel, 26.2(2005): 51-53. (in Chinese)
[19] Kang, X. Special Steel Technology, 17.3(1994):1-14. (in Chinese)
[20] Uggowitzer, P. J., et al. "Nickel-free high nitrogen austenitic stainless steels produced by
metal injection moulding." Materials science forum. Vol. 318. Trans Tech Publications Ltd,
1999.
[21] Chen, D., J. Huang, Z. Cai. Shanghai Metals, 35.6(2013): 6-10. (in Chinese)
[22] Nakazawa, T., et al. "Effects of silicon, molybdenum, and nitrogen on elevated temperature
properties of low carbon austenitic stainless steels." Tetsu-to-Hagane 91.8 (2005): 670-675.
[23] Xue, R., Z. Song, W. Zheng, Z. Du, and J. Ren. Journal of Iron Steel Research,25.10(2013):
36-41. (in Chinese)
[24] Pu, J., T. Chen, and B. Wang. Science Technology Information, 35(2013): 54. (in Chinese)
[25] Chen, S., and Z. Qin. Valve, 1(2005): 20-25. (in Chinese)
[26] Jung, S., et al. "Effects of Mn and Mo addition on high-temperature tensile properties in
high-Ni-containing austenitic cast steels used for turbo-charger application." Materials
Science and Engineering: A 682 (2017): 147-155.
[27] Zandrahimi, M., J. Vatandoost, and H. Ebrahimifar. "Al, Si, and Al–Si coatings to improve
the high-temperature oxidation resistance of AISI 304 stainless steel." Oxidation of
metals 76.3 (2011): 347-358.
[28] Bei, H., et al. "Aging effects on the mechanical properties of alumina-forming austenitic
stainless steels." Materials Science and Engineering: A 527.7-8 (2010): 2079-2086.
[29] La, P., Y. Li, S. Liu, D. Shen, and H. Wang. Iron Steel, 5(2010): 71-75. (in Chinese)
153
[30] La, P., Y. Li, and S. Liu. Materials Protection, 43.12(2010): 62-64. (in Chinese)
[31] Zhao, X., Z. Yang, W. Song., and J. Liang. 2010. Functionalization of steel structural
materials. Beijing: Metallurgical Industry Press. (in Chinese)
[32] Wang, J., F. Hua, Z. Liu, and G. Wang. Journal of Iron Steel Research, 24.4(2012): 42-46.
(in Chinese)
[33] Chen, T. H., K. L. Weng, and J. R. Yang. "The effect of high-temperature exposure on the
microstructural stability and toughness property in a 2205 duplex stainless steel." Materials
Science and Engineering: A 338.1-2 (2002): 259-270.
[34] Park, C., H. Kwon, and M. M. Lohrengel. "Micro-electrochemical polarization study on 25%
Cr duplex stainless steel." Materials Science and Engineering: A 372.1-2 (2004): 180-185.
[35] Schober, M., et al. "Precipitation behavior of intermetallic NiAl particles in Fe-6 at.% Al-4
at.% Ni analyzed by SANS and 3DAP." Intermetallics 18.8 (2010): 1553-1559.
[36] Dunning, J. S., D. E. A., and J. C. Rawers. "Influence of silicon and aluminum additions on
the oxidation resistance of a lean-chromium stainless steel." Oxidation of Metals 57.5 (2002):
409-425.
[37] Simmons, J. W. "Overview: high-nitrogen alloying of stainless steels." Materials Science
and Engineering: A 207.2 (1996): 159-169.
[38] Speidel, M. O. "New nitrogen-bearing austenitic stainless steels with high strength and
ductility." Metal Science & Heat Treatment 47 (2005): 489-493.
[39] Stein, G., J. Menzel, and M. Wagner. N-Alloyed steels for retaining rings and other
applications, in: HNS 90. Verlag Stahleisen GmbH Dusseldorf, 1990: 399.
[40] Yi, B., and Y. Hu. Iron Steel, 33.3(1998): 43-45. (in Chinese)
[41] Wang, S. 2008, Ph.D Thesis, Mechanical behavior of high nitrogen austenitic stainless steel
154
and mechanism of nitrogen. Shenyang: Institute of metal research, Chinese Academy of
Sciences. (in Chinese)
[42] Wu, X., Y. Fu, W. Ke, S. Xu, B. Feng, B. Hu, and J. Lu. Journal of Chinese Society Corrosion
Protection, 36.3(2016): 197-204. (in Chinese)
[43] Li, K., X. Qu, and D. Cui. Powd. Metal. Ind., 15.2(2005): 20-25. (in Chinese)
[44] Di Schino, A., and J. M. Kenny. "Grain refinement strengthening of a micro-crystalline high
nitrogen austenitic stainless steel." Materials Letters 57.12 (2003): 1830-1834.
[45] Eskandari, M., et al. "Potential application of nanocrystalline 301 austenitic stainless steel
in lightweight vehicle structures." Materials & Design 30.9 (2009): 3869-3872.
[46] Eskandari, M., A. Kermanpur, and A. Najafizadeh. "Formation of nanocrystalline structure
in 301 stainless steel produced by martensite treatment." Metallurgical and Materials
Transactions A 40.9 (2009): 2241-2249.
[47] Misra, R. D. K., et al. "Nanograined/ultrafine-grained structure and tensile deformation
behavior of shear phase reversion-induced 301 austenitic stainless steel." Metallurgical and
Materials Transactions A 41.8 (2010): 2162-2174.
[48] Wei, Y., et al. "Evading the strength–ductility trade-off dilemma in steel through gradient
hierarchical nanotwins." Nature communications 5.1 (2014): 1-8.
[49] Chen, A. Y., et al. "The influence of strain rate on the microstructure transition of 304
stainless steel." Acta Materialia 59.9 (2011): 3697-3709.
[50] Milad, M., et al. "The effect of cold work on structure and properties of AISI 304 stainless
steel." Journal of materials processing technology 203.1-3 (2008): 80-85.
[51] Fahr, D.. "Stress-and strain-induced formation of martensite and its effects on strength and
ductility of metastable austenitic stainless steels." Metallurgical Transactions 2.7 (1971):
155
1883-1892.
[52] Wu, Y., et al. "Bulk metallic glass composites with transformation‐mediated work‐hardening
and ductility." Advanced materials 22.25 (2010): 2770-2773.
[53] Ramakrishnan, V., J. A. McGurty, and N. Jayaraman. "Oxidation of high-aluminum
austenitic stainless steels." Oxidation of metals 30.3 (1988): 185-200.
[54] Brady, M. P., et al. Alumina-forming austenitics: a new class of heat-resistant stainless steels.
Oak Ridge National Lab.(ORNL), Oak Ridge, TN (United States); Shared Research
Equipment Collaborative Research Center, 2008.
[55] Xu, X., et al. "Improvement of high-temperature oxidation resistance and strength in
alumina-forming austenitic stainless steels." Materials Letters 65.21-22 (2011): 3285-3288.
[56] Wang, H., et al. "Effect of aluminium and silicon on high temperature oxidation resistance
of Fe-Cr-Ni heat resistant steel." Transactions of Tianjin University 15.6 (2009): 457.
[57] Wang, S., J. Zhang, and Y. Wang. Journal of University Jinan(Science Technology),
16.2(2002): 177-179. (in Chinese)
[58] Liu, J., J. Guo, C. Ji, and Y. Liu. Ferro-Alloys, 1(2010): 26-28. (in Chinese)
[59] Lu, Y. C., M. B. Ives, and C. R. Clayton. "Synergism of alloying elements and pitting
corrosion resistance of stainless steels." Corrosion Science 35.1-4 (1993): 89-96..
[60] Bandy, R. "Measurement of corrosion rate using two electrodes." Electrochimica Acta 26.1
(1981): 149-159.
[61] Lu, Y. C., and M. B. Ives. "Chemical treatment with cerium to improve the crevice corrosion
resistance of austenitic stainless steels." Corrosion Science 37.1 (1995): 145-155.
[62] Bandy, R., and D. A. Jones. "Analysis of errors in measuring corrosion rates by linear
polarization." Corrosion 32.4 (1976): 126-134.
156
[63] Grabke, H. J. "The role of nitrogen in the corrosion of iron and steels." Isij International 36.7
(1996): 777-786.
[64] Shi, H., et al. "Corrosion behavior of high-nitrogen stainless steel in NaCl
solution." International Journal of Electrochemical Science 12.12 (2017): 11298-11308.
[65] Bandy, R., and D. Van Rooyen. "Pitting-resistant alloys in highly concentrated chloride
media." Corrosion 39.6 (1983): 227-236.
[66] Lee, J., and S. Yoon. "Effect of nitrogen alloying on the semiconducting properties of passive
films and metastable pitting susceptibility of 316L and 316LN stainless steels." Materials
Chemistry and Physics 122.1 (2010): 194-199.
[67] Wan, P., et al. "Effect of nitrogen on biocorrosion behavior of high nitrogen nickel-free
stainless steel in different simulated body fluids." Materials Science and Engineering: C 32.3
(2012): 510-516.
[68] Malik, A. U., et al. "Corrosion behavior of steels in Gulf seawater
environment." Desalination 123.2-3 (1999): 205-213.
[69] Zhu, B., and G. Lindbergh. "Corrosion behaviour of high-chromium ferritic steels in molten
carbonate in cathode environment." Electrochimica acta 46.17 (2001): 2593-2604.
[70] Olubambi, P. A., J. H. Potgieter, and L. Cornish. "Corrosion behaviour of superferritic
stainless steels cathodically modified with minor additions of ruthenium in sulphuric and
hydrochloric acids." Materials & Design 30.5 (2009): 1451-1457.
[71] Palcut, M., et al. "Corrosion stability of ferritic stainless steels for solid oxide electrolyser
cell interconnects." Corrosion Science 52.10 (2010): 3309-3320.
[72] Chasse, K. R., and P. M. Singh. "Corrosion study of super ferritic stainless steel UNS S44660
(26Cr-3Ni-3Mo) and several other stainless steel grades (UNS S31603, S32101, and S32205)
157
in caustic solution containing sodium sulfide." Metallurgical and Materials Transactions
A 44.11 (2013): 5039-5053.
[73] Li, R., et al. "Localized corrosion performance of laser surface cladded UNS S44700
superferritic stainless steel on mild steel." Surface and Coatings Technology 88.1-3 (1997):
96-102.
[74] Malik, A. U., et al. "The effect of dominant alloy additions on the corrosion behavior of some
conventional and high alloy stainless steels in seawater." Corrosion science 37.10 (1995):
1521-1535.
[75] Kang, X. Super ferritic stainless steel, in: The 4th modern ferritic martensitic stainless steel
Conference, Beijing: Stainless steel branch of China Special Steel Enterprise Association,
2011. (in Chinese)
[76] Qiu, D., H. Liu, and C. Zhao. Sulphur Phosphorus & Bulk Materials Handling Related
Engineering, 4(2007): 22-25. (in Chinese)
[77] Liu, H., and J. Ye. Sulphur Phosphorus & Bulk Materials Handling Related Engineering,
5(2006): 32-36. (in Chinese)
[78] Qin, L., M. Sun, W. Fan, and Y. Zou. Taigang Science & Technology, 4(2009): 25-29. (in
Chinese)
[79] Mangonon, P. L., and G. Thomas. "Structure and properties of thermal-mechanically treated
304 stainless steel." Metallurgical transactions 1.6 (1970): 1587-1594.
[80] Tamura, I. "Deformation-induced martensitic transformation and transformation-induced
plasticity in steels." Metal Science 16.5 (1982): 245-253.
[81] Eckstein, C. B., and J. R. C. Guimarães. "Microstructure-property correlation in martensite-
austenite mixtures." Journal of materials science 19.9 (1984): 3043-3048.
158
[82] Liu, Y. H., et al. "Super plastic bulk metallic glasses at room temperature." science 315.5817
(2007): 1385-1388.
[83] Wang, Z. 2003. Metalworking in plasticity. Beijing: Metallurgy Industry Press. (in Chinese)
[84] Wu, H., F. Wu, S. Yang, and D. Tang. Acta Metallurgica Sinica (Chinese edition), 50.3(2014):
269-274. (in Chinese)
[85] Wittig, J. E., et al. "Temperature Dependent Deformation Mechanisms of a High Nitrogen‐
Manganese Austenitic Stainless Steel." steel research international 80.1 (2009): 66-70.
[86] Misra, R. D. K., et al. "On the significance of nature of strain-induced martensite on phase-
reversion-induced nanograined/ultrafine-grained austenitic stainless steel." Metallurgical
and Materials Transactions A 41.1 (2010): 3.
[87] Misra, R. D. K., et al. "Relationship of grain size and deformation mechanism to the fracture
behavior in high strength–high ductility nanostructured austenitic stainless steel." Materials
Science and Engineering: A 626 (2015): 41-50.
[88] De Cooman, B. C. "Structure–properties relationship in TRIP steels containing carbide-free
bainite." Current Opinion in Solid State and Materials Science 8.3-4 (2004): 285-303.
[89] Huang, C. X., et al. "Investigation on the nucleation mechanism of deformation-induced
martensite in an austenitic stainless steel under severe plastic deformation." Journal of
Materials Research 22.3 (2007): 724-729.
[90] Chen, M., et al. "Identical area observations of deformation-induced martensitic
transformation in SUS304 austenitic stainless steel." Materials Transactions 54.3 (2013):
308-313.
[91] Frommeyer, G., U. Brüx, and P. Neumann. "Supra-ductile and high-strength manganese-
TRIP/TWIP steels for high energy absorption purposes." ISIJ international 43.3 (2003): 438-
159
446.
[92] Olson, G. B., and M. Cohen. "Stress-assisted isothermal martensitic transformation:
application to TRIP steels." Metallurgical Transactions A 13.11 (1982): 1907-1914.
[93] Hamada, A. S., et al. "Enhancement in grain-structure and mechanical properties of laser
reversion treated metastable austenitic stainless steel." Materials & Design 94 (2016): 345-
352.
[94] Di Schino, A., M. Barteri, and J. M. Kenny. "Development of ultra fine grain structure by
martensitic reversion in stainless steel." Journal of materials science letters 21.9 (2002): 751-
753.
[95] Schino, A. D., I. Salvatori, and J. M. Kenny. "Effects of martensite formation and austenite
reversion on grain refining of AISI 304 stainless steel." Journal of Materials Science 37.21
(2002): 4561-4565.
[96] Srikanth, S., et al. "Property Enhancement in metastable 301LN austenitic stainless steel
through strain-induced martensitic transformation and its reversion (SIMTR) for metro
coach manufacture." Int. J. Metall. Eng 2 (2013): 203-213.
[97] Kumar, B. R., et al. "Ultrafine grained microstructure tailoring in austenitic stainless steel
for enhanced plasticity." Materials & Design 68 (2015): 63-71.
[98] Forouzan, F., et al. "Production of nano/submicron grained AISI 304L stainless steel through
the martensite reversion process." Materials Science and Engineering: A 527.27-28 (2010):
7334-7339.
[99] Moallemi, M., et al. "Effect of reversion annealing on the formation of nano/ultrafine grained
structure in 201 austenitic stainless steel." Materials Science and Engineering: A 530 (2011):
378-381.
160
[100] Rezaee, A., et al. "Production of nano/ultrafine grained AISI 201L stainless steel through
advanced thermo-mechanical treatment." Materials Science and Engineering: A 528.15
(2011): 5025-5029.
[101] Sabooni, S., F. K., and M. H. Enayati. "Thermal stability study of ultrafine grained 304L
stainless steel produced by martensitic process." Journal of materials engineering and
performance 23.5 (2014): 1665-1672.
[102] Behjati, P., et al. "Effect of annealing temperature on nano/ultrafine grain of Ni-free
austenitic stainless steel." Materials Science and Engineering: A 592 (2014): 77-82.
[103] Baghbadorani, H. S., et al. "An investigation on microstructure and mechanical propertiesof
a Nb-microalloyed nano/ultrafine grained 201 austenitic stainless steel." Materials Science
and Engineering: A 636 (2015): 593-599.
[104] Xu, D. M., et al. "Deformation behavior of high yield strength–high ductility ultrafine-
grained 316LN austenitic stainless steel." Materials Science and Engineering: A 688 (2017):
407-415.
[105] Ma, Y., J. Jin, and Y. Lee. "A repetitive thermomechanical process to produce nano-
crystalline in a metastable austenitic steel." Scripta Materialia 52.12 (2005): 1311-1315.
[106] Rajasekhara, S., et al. "Hall–Petch behavior in ultra-fine-grained AISI 301LN stainless
steel." Metallurgical and Materials Transactions A 38.6 (2007): 1202-1210.
[107] Somani, M. C., et al. "Enhanced mechanical properties through reversion in metastable
austenitic stainless steels." Metallurgical and Materials transactions A 40.3 (2009): 729-744.
[108] Misra, R. D. K., et al. "Microstructure and deformation behavior of phase-reversion-induced
nanograined/ultrafine-grained austenitic stainless steel." Metallurgical and Materials
Transactions A 40.10 (2009): 2498-2509.
161
[109] Misra, R. D. K., et al. "Probing deformation processes in near-defect free volume in high
strength–high ductility nanograined/ultrafine-grained (NG/UFG) metastable austenitic
stainless steels." Scripta Materialia 63.11 (2010): 1057-1060.
[110] Rajasekhara, S., et al. "Microstructure evolution in nano/submicron grained AISI 301LN
stainless steel." Materials Science and Engineering: A 527.7-8 (2010): 1986-1996.
[111] Järvenpää, A., M. Jaskari, and L. P. Karjalainen. "Reversed microstructures and tensile
properties after various cold rolling reductions in AISI 301LN steel." Metals 8.2 (2018): 109.
[112] Poulon, A., et al. "Fine grained austenitic stainless steels: the role of strain induced α′
martensite and the reversion mechanism limitations." ISIJ international 49.2 (2009): 293-
301.
[113] Smith, H., and D. R. F. West. "The reversion of martensite to austenite in certain stainless
steels." Journal of materials science 8.10 (1973): 1413-1420.
[114] Smith, H., and D. R. F. West. "Annealing of austenite formed by reversion from martensite
in an Fe–16Cr–12Ni alloy." Metals Technology 1.1 (1974): 37-40.
[115] Coleman, T. H., and D. R. F. West. "Deformation-induced martensite and its reversion to
austenite in an Fe–16Cr–12Ni alloy." Metals Technology 3.1 (1976): 49-53.
[116] Guy, K. B., E. P. Butler, and D. R. F. West. "Reversion of bcc α′ martensite in Fe–Cr–Ni
austenitic stainless steels." Metal science 17.4 (1983): 167-176.
[117] Guy, K., E. P. Butler, and D. R. F. West. "ε and α'martensite formation and reversion in
austenitic stainless steels." Le Journal de Physique Colloques 43.C4 (1982): C4-575.
[118] Singh, J.. "Influence of deformation on the transformation of austenitic stainless
steels." Journal of Materials Science 20.9 (1985): 3157-3166.
[119] Tomimura, K., et al. "Optimal chemical composition in Fe-Cr-Ni alloys for ultra grain
162
refining by reversion from deformation induced martensite." ISIJ international 31.7 (1991):
721-727.
[120] Tomimura, K., S. Takaki, and Y. Tokunaga. "Reversion mechanism from deformation
induced martensite to austenite in metastable austenitic stainless steels." ISIJ
international 31.12 (1991): 1431-1437.
[121] Takaki, S., K. Tomimura, and S. Ueda. "Effect of pre-cold-working on diffusional reversion
of deformation induced martensite in metastable austenitic stainless steel." ISIJ
international 34.6 (1994): 522-527.
[122] Di Schino, A., M. Barteri, and J. M. Kenny. "Grain size dependence of mechanical, corrosion
and tribological properties of high nitrogen stainless steels." Journal of Materials
Science 38.15 (2003): 3257-3262.
[123] Di Schino, A., M. Barteri, and J. M. Kenny. "Effects of grain size on the properties of a low
nickel austenitic stainless steel." Journal of Materials Science 38.23 (2003): 4725-4733.
[124] Di Schino, A., and J. M. Kenny. "Grain size dependence of the fatigue behaviour of a
ultrafine-grained AISI 304 stainless steel." Materials Letters 57.21 (2003): 3182-3185.
[125] Choi, J., and W. Jin. "Strain induced martensite formation and its effect on strain hardening
behavior in the cold drawn 304 austenitic stainless steels." Scripta Materialia 36.1 (1997):
99-104.
[126] Lee, S., Y. Park, and Y. Lee. "Reverse transformation mechanism of martensite to austenite
in a metastable austenitic alloy." Materials Science and Engineering: A 515.1-2 (2009): 32-
37.
[127] Jung, Y., et al. "Effect of grain size on strain-induced martensitic transformation start
temperature in an ultrafine grained metastable austenitic steel." Metals and Materials
163
International 17.4 (2011): 553-556.
[128] Huang, J., et al. "Enhanced mechanical properties of type AISI301LN austenitic stainless
steel through advanced thermo mechanical process." Materials Science and Engineering:
A 532 (2012): 190-195.
[129] Shen, Y. F., et al. "Suppression of twinning and phase transformation in an ultrafine grained
2 GPa strong metastable austenitic steel: Experiment and simulation." Acta Materialia 97
(2015): 305-315.
[130] Gong, N., et al. "Effects of annealing temperature on nano/ultrafine-grained structure in
austenite stainless steel." Materials Science and Technology 33.14 (2017): 1667-1672.
[131] Gong, N., et al. "Studying mechanical properties and micro deformation of ultrafine-grained
structures in austenitic stainless steel." Metals 7.6 (2017): 188.
[132] Lei, C., et al. "Deformation mechanism and ductile fracture behavior in high strength high
ductility nano/ultrafine grained Fe-17Cr-6Ni austenitic steel." Materials Science and
Engineering: A 709 (2018): 72-81.
[133] Xu, D. M., et al. "Effect of grain refinement on strain hardening and fracture in austenitic
stainless steel." Materials Science and Technology 34.11 (2018): 1344-1352.
[134] Xu, D. M., et al. "The significant impact of cold deformation on structure-property
relationship in phase reversion-induced stainless steels." Materials Characterization 145
(2018): 157-171.
[135] Kumar, B. R., and S. Sharma. "Recrystallization behavior of a heavily deformed austenitic
stainless steel during iterative type annealing." Metallurgical and Materials Transactions
A 45.13 (2014): 6027-6038.
[136] Somani, M. C., et al. "Improving the yield strength of an antibacterial 304Cu austenitic
164
stainless steel by the reversion treatment." Materials Science and Engineering: A 793 (2020):
139885.
[137] Hu, C. Y., et al. "On the impacts of grain refinement and strain-induced deformation on three-
body abrasive wear responses of 18Cr–8Ni austenitic stainless steel." Wear 446 (2020):
203181.
[138] Mallick, P., et al. "Microstructure-tensile property correlation in 304 stainless steel after cold
deformation and austenite reversion." Materials Science and Engineering: A 707 (2017):
488-500.
[139] Mallick, P., et al. "Effect of TMCP on microstructure and mechanical properties of 304
stainless Steel." steel research international 89.8 (2018): 1800103.
[140] Mateo, A., A. Zapata, and G. Fargas. "Improvement of mechanical properties on metastable
stainless steels by reversion heat treatments." IOP Conference Series: Materials Science and
Engineering. Vol. 48. No. 1. IOP Publishing, 2013.
[141] Odnobokova, M., et al. "Annealing behavior of a 304L stainless steel processed by large
strain cold and warm rolling." Materials Science and Engineering: A 689 (2017): 370-383.
[142] Shakhova, I., et al. "Effect of large strain cold rolling and subsequent annealing on
microstructure and mechanical properties of an austenitic stainless steel." Materials Science
and Engineering: A 545 (2012): 176-186.
[143] Cios, G., et al. "The investigation of strain-induced martensite reverse transformation in AISI
304 austenitic stainless steel." Metallurgical and Materials Transactions A 48.10 (2017):
4999-5008.
[144] Fargas, G., et al. "Correlation between microstructure and mechanical properties before and
after reversion of metastable austenitic stainless steels." Metallurgical and Materials
165
Transactions A 46.12 (2015): 5697-5707.
[145] Hedayati, A., et al. "The effect of cold rolling regime on microstructure and mechanical
properties of AISI 304L stainless steel." Journal of Materials Processing Technology 210.8
(2010): 1017-1022.
[146] Rezaee, A., et al. "The influence of reversion annealing behavior on the formation of
nanograined structure in AISI 201L austenitic stainless steel through martensite
treatment." Materials & Design 32.8-9 (2011): 4437-4442.
[147] Sadeghpour, S., A. Kermanpur, and A. Najafizadeh. "Influence of Ti microalloying on the
formation of nanocrystalline structure in the 201L austenitic stainless steel during martensite
thermomechanical treatment." Materials Science and Engineering: A 584 (2013): 177-183.
[148] Sadeghpour, S., A. Kermanpur, and A. Najafizadeh. "Formation of nano/ultrafine grain
structure in a Ti-modified 201L stainless steel through martensite thermomechanical
treatment." ISIJ international 54.4 (2014): 920-925.
[149] Shirdel, M., H. Mirzadeh, and M. H. Parsa. "Nano/ultrafine grained austenitic stainless steel
through the formation and reversion of deformation-induced martensite: Mechanisms,
microstructures, mechanical properties, and TRIP effect." Materials Characterization 103
(2015): 150-161.
[150] Baghbadorani, H. S., et al. "Influence of Nb-microalloying on the formation of
nano/ultrafine-grained microstructure and mechanical properties during martensite reversion
process in a 201-type austenitic stainless steel." Metallurgical and Materials Transactions
A 46.8 (2015): 3406-3413.
[151] Rasouli, D., et al. "On the reversion and recrystallization of austenite in the interstitially
alloyed Ni-free nano/ultrafine grained austenitic stainless steels." Metals and Materials
166
International 25.4 (2019): 846-859.
[152] Kheiri, S., H. Mirzadeh, and M. Naghizadeh. "Tailoring the microstructure and mechanical
properties of AISI 316L austenitic stainless steel via cold rolling and reversion
annealing." Materials Science and Engineering: A 759 (2019): 90-96.
[153] Droste, M., et al. "Fatigue behavior of an ultrafine-grained metastable CrMnNi steel tested
under total strain control." International Journal of Fatigue 106 (2018): 143-152.
[154] Maréchal, D.. Linkage between mechanical properties and phase transformations in a 301LN
austenitic stainless steel. Ph.D. Thesis, Diss. University of British Columbia, BC, Canada,
2011.
[155] Fava, J., et al. "Characterization of reverse martensitic transformation in cold-rolled
austenitic 316 stainless steel." Matéria (Rio de Janeiro) 23.2 (2018).
[156] Souza Filho, I. R., et al. "Strain partitioning and texture evolution during cold rolling of AISI
201 austenitic stainless steel." Materials Science and Engineering: A 702 (2017): 161-172.
[157] Souza Filho, I. R., et al. "Effects of strain-induced martensite and its reversion on the
magnetic properties of AISI 201 austenitic stainless steel." Journal of Magnetism and
Magnetic Materials 419 (2016): 156-165.
[158] Souza Filho, I. R., et al. "Austenite reversion in AISI 201 austenitic stainless steel evaluated
via in situ synchrotron X-ray diffraction during slow continuous annealing." Materials
Science and Engineering: A 755 (2019): 267-277.
[159] Hamada, A. S., et al. "Enhancement of mechanical properties of a TRIP-aided austenitic
stainless steel by controlled reversion annealing." Materials Science and Engineering: A 628
(2015): 154-159.
[160] Kisko, A., et al. "The influence of grain size on the strain-induced martensite formation in
167
tensile straining of an austenitic 15Cr–9Mn–Ni–Cu stainless steel." Materials Science and
Engineering: A 578 (2013): 408-416.
[161] Kisko, A., et al. "Effect of Nb microalloying on reversion and grain growth in a high-Mn
204Cu austenitic stainless steel." ISIJ International 55.10 (2015): 2217-2224.
[162] Kisko, A., et al. "Effects of reversion and recrystallization on microstructure and mechanical
properties of Nb-alloyed low-Ni high-Mn austenitic stainless steels." Materials Science and
Engineering: A 657 (2016): 359-370.
[163] Järvenpää, A., et al. "Austenite stability in reversion-treated structures of a 301LN steel
under tensile loading." Materials Characterization 127 (2017): 12-26.
[164] Järvenpää, A., et al. "Stability of grain-refined reversed structures in a 301LN austenitic
stainless steel under cyclic loading." Materials Science and Engineering: A 703 (2017): 280-
292.
[165] Järvenpää, A., et al. "Demonstrating the effect of precipitation on the mechanical stability of
fine-grained austenite in reversion-treated 301LN stainless steel." Metals 7.9 (2017): 344.
[166] Järvenpää, A., M. Jaskari, and L. P. Karjalainen. "Properties of induction reversion-refined
microstructures of AISI 301LN under monotonic, cyclic and rolling deformation." Materials
Science Forum. Vol. 941. Trans Tech Publications Ltd, 2018: 601-607.
[167] Misra, R. D. K., et al. "Deformation processes during tensile straining of
ultrafine/nanograined structures formed by reversion in metastable austenitic steels." Scripta
Materialia 59.1 (2008): 79-82.
[168] Misra, R. D. K., et al. "Martensite shear phase reversion-induced nanograined/ultrafine-
grained Fe–16Cr–10Ni alloy: The effect of interstitial alloying elements and degree of
austenite stability on phase reversion." Materials Science and Engineering: A 527.29-30
168
(2010): 7779-7792.
[169] Misra, R. D. K., et al. "Nanomechanical insights into the deformation behavior of austenitic
alloys with different stacking fault energies and austenitic stability." Materials Science and
Engineering: A 528.22-23 (2011): 6958-6963.
[170] Challa, V. S. A., et al. "Strain hardening behavior of phase reversion-induced
nanograined/ultrafine-grained (NG/UFG) austenitic stainless steel and relationship with
grain size and deformation mechanism." Materials Science and Engineering: A 613 (2014):
60-70.
[171] Challa, V. S. A., et al. "Significance of interplay between austenite stability and deformation
mechanisms in governing three-stage work hardening behavior of phase-reversion induced
nanograined/ultrafine-grained (NG/UFG) stainless steels with high strength-high ductility
combination." Scripta Materialia 86 (2014): 60-63.
[172] Challa, V. S. A., et al. "Strain hardening behavior of nanograined/ultrafine-grained (NG/UFG)
austenitic 16Cr–10Ni stainless steel and its relationship to austenite stability and deformation
behavior." Materials Science and Engineering: A 649 (2016): 153-157.
[173] Misra, R. D. K., V. S. Y. Injeti, and M. C. Somani. "The significance of deformation
mechanisms on the fracture behavior of phase reversion-induced nanostructured austenitic
stainless steel." Scientific reports 8.1 (2018): 1-13.
[174] Chlupová, A., et al. "Microstructural investigation and mechanical testing of an ultrafine-
grained austenitic stainless steel." Proc. of 5th Int. Conf. NANOCON. Tanger Ltd.:
Ostrava, Czech Republic, 2013: 733–738.
[175] Chlupová, A., et al. "LCF behaviour of ultrafine grained 301LN stainless steel." Procedia
Engineering 74 (2014): 147-150.
169
[176] Man, J., et al. "Microstructural changes during deformation of AISI 300 grade austenitic
stainless steels: Impact of chemical heterogeneity." Procedia Structural Integrity 2 (2016):
2299-2306.
[177] Komatsuseiki Kosakusho Co., Ltd. Available online:
https://www.komatsuseiki.co.jp/english/future/03.php (accessed on 9 August 2016).
[178] Komatsu, T., T. Matsumura, and S. Torizuka. "Effect of Grain Size in Stainless Steel on
Cutting Performance in Micro-Scale Cutting." IJAT 5.3 (2011): 334-341.
[179] Komatsu, T., et al. "Micro hole piercing for ultra fine grained steel." Materials Science
Forum. Vol. 783. Trans Tech Publications Ltd, Stafa-Zurich, Switzerland, 2014: 2653-2658.
[180] Komatsu, T., et al. "Effect of crystal grain size in stainless steel on cutting process in
micromilling." Procedia Cirp 1 (2012): 150-155.
[181] Nippon Steel & Sumitomo Metal Product Catalog: SUS304 BA1. Available online:
https://stainless.nipponsteel.com/product/grade/nssmc_series/product/sus304_ba1.php
(accessed on 6 December 2019).
[182] Nippon Steel & Sumitomo Metal Product Catalog: NSSMC-NAR-301L BA1. Available
online: https://stainless.nipponsteel.com/product/grade/nssmc_series/product/nssmc-nar-
301l_ba1.php (accessed on 6 December 2019).
[183] Nießen, F.. "Phase Transformations in Supermartensitic Stainless Steels." Ph.D. Thesis,
Technical University of Denmark, Lyngby, Denmark, (2018).
[184] Yang, D. P., D. Wu, and H. L. Yi. "Reverse transformation from martensite into austenite in
a medium-Mn steel." Scripta Materialia 161 (2019): 1-5.
[185] Santos, T. F. de Abreu, and M. S. Andrade. "Avaliação dilatométrica da reversão das
martensitas induzidas por deformação em um aço inoxidável austenítico do tipo ABNT
170
304." Matéria (Rio de Janeiro) 13.4 (2008): 587-596. (in Spanish)
[186] Santos, T. F. de Abreu, and M. S. Andrade. "Internal Friction on AISI 304 Stainless Steels
with Low Tensile Deformations at Temperatures between 50 and 20." Advances in Materials
science and Engineering 2010 (2010).
[187] Dryzek, E., M. Sarnek, and M. Wróbel. "Reverse transformation of deformation-induced
martensite in austenitic stainless steel studied by positron annihilation." Journal of Materials
Science 49.24 (2014): 8449-8458.
[188] Knutsson, A., Peter H., and M. Odén. "Reverse martensitic transformation and resulting
microstructure in a cold rolled metastable austenitic stainless steel." steel research
international 79.6 (2008): 433-439.
[189] Talonen, J., P. Aspegren, and H. Hänninen. "Comparison of different methods for measuring
strain induced α-martensite content in austenitic steels." Materials Science and
Technology 20.12 (2004): 1506-1512.
[190] Cios, G., T. Tokarski, and P. Bała. "Strain-induced martensite reversion in 18Cr–8Ni steel–
transmission Kikuchi diffraction study." Materials Science and Technology 34.5 (2018):
580-583.
[191] Wei, S., M. Jiang, and C. C. Tasan. "Interstitial-free bake hardening realized by epsilon
martensite reverse transformation." Metallurgical and Materials Transactions A 50.9 (2019):
3985-3991.
[192] Johannsen, D. L., A. Kyrolainen, and P. J. Ferreira. "Influence of annealing treatment on the
formation of nano/submicron grain size AISI 301 austenitic stainless steels." Metallurgical
and Materials Transactions A 37.8 (2006): 2325-2338.
[193] Apple, C. A., and G. Krauss. "The effect of heating rate on the martensite to austenite
171
transformation in Fe-Ni-C alloys." Acta metallurgica 20.7 (1972): 849-856.
[194] Lindström, A.. "Austempered high silicon steel: Investigation of wear resistance in a carbide
free microstructure.", Master’s thesis, Luleå University of Technology, Sweden, (2006).
[195] Li, Z., and D. Wu. "Effects of hot deformation and subsequent austempering on the
mechanical properties of Si–Mn TRIP steels." ISIJ international 46.1 (2006): 121-128.
[196] Marques, V. M. F., et al. "Nanomechanical characterization of Sn–Ag–Cu/Cu joints—Part 2:
Nanoindentation creep and its relationship with uniaxial creep as a function of
temperature." Acta Materialia 61.7 (2013): 2471-2480.
[197] Ljungcrantz, H., et al. "Nanoindentation studies of single‐crystal (001)‐,(011)‐, and (111)‐
oriented TiN layers on MgO." Journal of applied physics 80.12 (1996): 6725-6733.
[198] Hosemann, P.. "Studying radiation damage in structural materials by using ion
accelerators." Reviews Of Accelerator Science And Technology: Volume 4: Accelerator
Applications in Industry and the Environment. 2011. 161-182.
[199] Kiener, D., et al. "Application of small-scale testing for investigation of ion-beam-irradiated
materials." Journal of Materials Research 27.21 (2012): 2724.
[200] Lupinacci, A., et al. "Characterization of ion beam irradiated 304 stainless steel utilizing
nanoindentation and Laue microdiffraction." Journal of Nuclear Materials 458 (2015): 70-
76.
[201] Lu, L., et al. "Nano-sized twins induce high rate sensitivity of flow stress in pure
copper." Acta materialia 53.7 (2005): 2169-2179.
[202] Schwaiger, R., et al. "Some critical experiments on the strain-rate sensitivity of
nanocrystalline nickel." Acta materialia 51.17 (2003): 5159-5172.
[203] Roa, J. J., et al. "Dependence of nanoindentation hardness with crystallographic orientation
172
of austenite grains in metastable stainless steels." Materials Science and Engineering: A 645
(2015): 188-195.
[204] Roa, J. J., et al. "Influence of testing mode on the fatigue behavior of< 111> austenitic grain
at the nanometric length scale for TRIP steels." Materials Science and Engineering: A 713
(2018): 287-293.
[205] Roa, J. J., et al. "Reversible phase transformation in polycrystalline TRIP steels induced by
cyclic indentation performed at the nanometric length scale." steel research
international 89.11 (2018): 1800234.
[206] Ahn, T., et al. "Investigation of strain-induced martensitic transformation in metastable
austenite using nanoindentation." Scripta Materialia 63.5 (2010): 540-543.
[207] Sapezanskaia, I., et al. "Deformation mechanisms induced by nanoindentation tests on a
metastable austenitic stainless steel: A FIB/SIM investigation." Materials
Characterization 131 (2017): 253-260.
[208] Williams, D. B., and C. Barry Carter. Transmission Electron Microscopy: Spectrometry
David B. Williams and C. Barry Cart. IV. Plenum, 1996.
[209] Hirsch, P. B., et al. "Electron Microscopy of Thin Crystals, Butterworths." London,
UK (1965): 358.
[210] Principle Components of an EBSD System, Available online: http://www.ebsd.com/ebsd-
explained/principle-components-of-an-ebsd-system
[211] Nohara, K., Y. Ono, and N. Ohashi. "Composition and grain size dependencies of strain-
induced martensitic transformation in metastable austenitic stainless steels." Tetsu-to-
Hagané 63.5 (1977): 772-782.
[212] Poulon, A., et al. "Fine grained austenitic stainless steels: the role of strain induced α′
173
martensite and the reversion mechanism limitations." ISIJ international 49.2 (2009): 293-
301.
[213] Zhuang, Y., et al. "Antibacterial activity of copper‐bearing 316L stainless steel for the
prevention of implant‐related infection." Journal of Biomedical Materials Research Part B:
Applied Biomaterials 108.2 (2020): 484-495.
[214] Hong, I. T., and C. Hl Koo. "Antibacterial properties, corrosion resistance and mechanical
properties of Cu-modified SUS 304 stainless steel." Materials Science and Engineering:
A 393.1-2 (2005): 213-222.
[215] Chai, H., et al. "Antibacterial effect of 317L stainless steel contained copper in prevention
of implant-related infection in vitro and in vivo." Journal of Materials Science: Materials in
Medicine 22.11 (2011): 2525-2535.
[216] Ren, L., L. Nan, and K. Yang. "Study of copper precipitation behavior in a Cu-bearing
austenitic antibacterial stainless steel." Materials & Design 32.4 (2011): 2374-2379.
[217] Luo, F., et al. "Study on properties of copper-containing austenitic antibacterial stainless
steel." Materials Technology 34.9 (2019): 525-533.
[218] Sun, G., et al. "The significant role of heating rate on reverse transformation and coordinated
straining behavior in a cold-rolled austenitic stainless steel." Materials Science and
Engineering: A 732 (2018): 350-358.
[219] Mészáros, I., and J. Prohászka. "Magnetic investigation of the effect of α′-martensite on the
properties of austenitic stainless steel." Journal of Materials Processing Technology 161.1-2
(2005): 162-168.
[220] Martins, L. F. M., R. L. Plaut, and A. F. Padilha. "Effect of Carbon on the Cold-worked State
and Annealing Behavior of Two 18wt% Cr–8wt% Ni Austenitic Stainless Steels." ISIJ
174
international 38.6 (1998): 572-579.
[221] Sun, G. S., et al. "Low temperature superplastic-like deformation and fracture behavior of
nano/ultrafine-grained metastable austenitic stainless steel." Materials & Design 117 (2017):
223-231.
[222] Gonzalez, B. M., et al. "The influence of copper addition on the formability of AISI 304
stainless steel." Materials Science and Engineering: A 343.1-2 (2003): 51-56.
[223] Dumay, A., et al. "Influence of addition elements on the stacking-fault energy and
mechanical properties of an austenitic Fe–Mn–C steel." Materials Science and Engineering:
A 483 (2008): 184-187.
[224] Takaki, S., K. Tomimura, and S. Ueda. "Effect of pre-cold-working on diffusional reversion
of deformation induced martensite in metastable austenitic stainless steel." ISIJ
international 34.6 (1994): 522-527.
[225] Sun, G., et al. "On the influence of deformation mechanism during cold and warm rolling on
annealing behavior of a 304 stainless steel." Materials Science and Engineering: A 746
(2019): 341-355.
[226] Sun, G. S., et al. "Ultrahigh strength nano/ultrafine-grained 304 stainless steel through three-
stage cold rolling and annealing treatment." Materials characterization 110 (2015): 228-235.
[227] Poulon, A., et al. "Influence of texture and grain size on martensitic transformations
occurring during low‐cycle fatigue of a fine‐grained austenitic stainless steel." Advanced
Engineering Materials 12.10 (2010): 1041-1046.
[228] Padilha, A. F., R. L. Plaut, and P. R. Rios. "Annealing of cold-worked austenitic stainless
steels." ISIJ international 43.2 (2003): 135-143.
[229] Sohrabi, M. J., H. Mirzadeh, and C. Dehghanian. "Thermodynamics basis of saturation of
175
martensite content during reversion annealing of cold rolled metastable austenitic
steel." Vacuum 174 (2020): 109220.
[230] Sohrabi, M. J., H. Mirzadeh, and C. Dehghanian. "Significance of Martensite Reversion and
Austenite Stability to the Mechanical Properties and Transformation-Induced Plasticity
Effect of Austenitic Stainless Steels." Journal of Materials Engineering and
Performance 29.5 (2020): 3233-3242.
[231] Shen, Y. F., et al. "Twinning and martensite in a 304 austenitic stainless steel." Materials
Science and Engineering: A 552 (2012): 514-522.
[232] Hamada, A. S., L. P. Karjalainen, and M. C. Somani. "Electrochemical corrosion behaviour
of a novel submicron-grained austenitic stainless steel in an acidic NaCl solution." Materials
Science and Engineering: A 431.1-2 (2006): 211-217.
[233] Xi, T., et al. "Effect of copper addition on mechanical properties, corrosion resistance and
antibacterial property of 316L stainless steel." Materials Science and Engineering: C 71
(2017): 1079-1085.
[234] Jiang, J., et al. "Effects of aging time on intergranular and pitting corrosion behavior of Cu-
bearing 304L stainless steel in comparison with 304L stainless steel." Corrosion Science 113
(2016): 46-56.
[235] Gagliano, M. S., and M. E. Fine. "Characterization of the nucleation and growth behavior of
copper precipitates in low-carbon steels." Metallurgical and Materials Transactions A 35.8
(2004): 2323-2329.
[236] Stechauner, G., and E. Kozeschnik. "Thermo-kinetic modeling of Cu precipitation in α-
Fe." Acta Materialia 100 (2015): 135-146.
[237] Kapoor, M., et al. "Aging characteristics and mechanical properties of 1600 MPa body-
176
centered cubic Cu and B2-NiAl precipitation-strengthened ferritic steel." Acta materialia 73
(2014): 56-74.
[238] Chi, C., et al. "The precipitation strengthening behavior of Cu-rich phase in Nb contained
advanced Fe–Cr–Ni type austenitic heat resistant steel for USC power plant
application." Progress in Natural Science: Materials International 22.3 (2012): 175-185.
[239] Luo, F., et al. "Study on the performance of new process for treating copper-containing
antibacterial stainless steel." Materials Research Express 6.4 (2019): 046554.
[240] Mirzadeh, H., and A. Najafizadeh. "Aging kinetics of 17-4 PH stainless steel." Materials
chemistry and physics 116.1 (2009): 119-124.
[241] Soylu, B., and R. W. K. Honeycombe. "Microstructural refinement of duplex stainless
steels." Materials science and technology 7.2 (1991): 137-146.
[242] Zhang, Z. X., G. Lin, and Z. Xu. "Effects of light pre-deformation on pitting corrosion
resistance of copper-bearing ferrite antibacterial stainless steel." Journal of materials
processing technology 205.1-3 (2008): 419-424.
[243] Järvenpää, A., et al. "Processing and properties of reversion-treated austenitic stainless
steels." Metals 10.2 (2020): 281.
[244] Järvenpää, A., and L. P. Karjalainen. "An overview of mechanical properties of today’s grain-
refined austenitic stainless steels." Proceedings of the ESSC & DUPLEX, ASMET, Vienna,
Austria 31 (2019): 12-21.
[245] Kumar, B. R., et al. "Effect of cyclic thermal process on ultrafine grain formation in AISI
304L austenitic stainless steel." Metallurgical and Materials Transactions A 40.13 (2009):
3226-3234.
[246] Sharma, S., et al. "Effects of concurrent strain induced martensite formation on tensile and
177
texture properties of 304L stainless steel of varying grain size distribution." Materials
Science and Engineering: A 725 (2018): 215-227.
[247] Bauer, S., et al. "Engineering biocompatible implant surfaces: Part I: Materials and
surfaces." Progress in Materials Science 58.3 (2013): 261-326.
[248] Misra, R. D. K., et al. "Understanding the impact of grain structure in austenitic stainless
steel from a nanograined regime to a coarse-grained regime on osteoblast functions using a
novel metal deformation–annealing sequence." Acta biomaterialia 9.4 (2013): 6245-6258.
[249] Nune, C., et al. "Dependence of cellular activity at protein adsorbed biointerfaces with nano‐
to microscale dimensionality." Journal of Biomedical Materials Research Part A 102.6
(2014): 1663-1676.
[250] Gong, N., et al. "On the mechanical behavior of austenitic stainless steel with nano/ultrafine
grains and comparison with micrometer austenitic grains counterpart and their biological
functions." Journal of the mechanical behavior of biomedical materials 101 (2020): 103433.
[251] Rybalchenko, O. V., et al. "The influence of ultrafine‐grained structure on the mechanical
properties and biocompatibility of austenitic stainless steels." Journal of Biomedical
Materials Research Part B: Applied Biomaterials 108.4 (2020): 1460-1468.
[252] Misra, R. D. K., et al. "Cellular Mechanisms of Enhanced Osteoblasts Functions via Phase‐
Reversion Induced Nano/Submicron‐Grained Structure in a Low‐Ni Austenitic Stainless
Steel." Advanced Engineering Materials 13.12 (2011): B483-B492.
[253] Misra, R. D. K., et al. "Cellular response of preosteoblasts to nanograined/ultrafine-grained
structures." Acta biomaterialia 5.5 (2009): 1455-1467.
[254] Misra, R. D. K., et al. "Cellular activity of bioactive nanograined/ultrafine-grained
materials." Acta biomaterialia 6.7 (2010): 2826-2835.
178
[255] Misra, R. D. K., et al. "Interplay between grain structure, deformation mechanisms and
austenite stability in phase-reversion-induced nanograined/ultrafine-grained austenitic
ferrous alloy." Acta Materialia 84 (2015): 339-348.
[256] Schuh, C. A., T. G. Nieh, and Y. Kawamura. "Rate dependence of serrated flow during
nanoindentation of a bulk metallic glass." Journal of Materials Research 17.7 (2002): 1651-
1654.
[257] Asaro, R. J., and S. Suresh. "Mechanistic models for the activation volume and rate
sensitivity in metals with nanocrystalline grains and nano-scale twins." Acta
Materialia 53.12 (2005): 3369-3382.
[258] Maier, V., et al. "Nanoindentation strain-rate jump tests for determining the local strain-rate
sensitivity in nanocrystalline Ni and ultrafine-grained Al." Journal of materials
research 26.11 (2011).
[259] Wei, Q., et al. "Effect of nanocrystalline and ultrafine grain sizes on the strain rate sensitivity
and activation volume: fcc versus bcc metals." Materials Science and Engineering: A 381.1-
2 (2004): 71-79.
[260] Misra, R. D. K., et al. "Nanoscale deformation experiments on the strain rate sensitivity of
phase reversion induced nanograined/ultrafine-grained austenitic stainless steels and
comparison with the coarse-grained counterpart." Materials Science and Engineering: A 548
(2012): 161-174.
[261] Lorenz, D., et al. "Pop-in effect as homogeneous nucleation of dislocations during
nanoindentation." Physical review B 67.17 (2003): 172101.
[262] Mason, J. K., A. C. Lund, and C. A. Schuh. "Determining the activation energy and volume
for the onset of plasticity during nanoindentation." Physical review B 73.5 (2006): 054102.
179
[263] Yang, B., and H. Vehoff. "Dependence of nanohardness upon indentation size and grain size–
a local examination of the interaction between dislocations and grain boundaries." Acta
materialia 55.3 (2007): 849-856.
[264] Backes, B., K. Durst, and M. Goeken. "Determination of plastic properties of polycrystalline
metallic materials by nanoindentation: Experiments and finite element
simulations." Philosophical Magazine 86.33-35 (2006): 5541-5551.
[265] Hou, X. D., A. J. Bushby, and N. M. Jennett. "Study of the interaction between the
indentation size effect and Hall–Petch effect with spherical indenters on annealed
polycrystalline copper." Journal of Physics D: Applied Physics 41.7 (2008): 074006.
[266] Lodes, M. A., et al. "Influence of dislocation density on the pop-in behavior and indentation
size effect in CaF2 single crystals: Experiments and molecular dynamics simulations." Acta
Materialia 59.11 (2011): 4264-4273.
[267] Fleck, N. A., et al. "Strain gradient plasticity: theory and experiment." Acta Metallurgica et
materialia 42.2 (1994): 475-487.
[268] Backes, B., et al. "The correlation between the internal material length scale and the
microstructure in nanoindentation experiments and simulations using the conventional
mechanism-based strain gradient plasticity theory." journal of Materials Research 24.3
(2009): 1197-1207.
[269] Shen, Y. F., et al. "Twinning and martensite in a 304 austenitic stainless steel." Materials
Science and Engineering: A 552 (2012): 514-522.
[270] Shen, Y. F., et al. "Suppression of twinning and phase transformation in an ultrafine grained
2 GPa strong metastable austenitic steel: Experiment and simulation." Acta Materialia 97
(2015): 305-315.
180
[271] Shen, Y. F., et al. "Softening behavior by excessive twinning and adiabatic heating at high
strain rate in a Fe–20Mn–0.6 C TWIP steel." Acta Materialia 103 (2016): 229-242.
[272] Lagerlöf, K. P. D., et al. "Nucleation and growth of deformation twins: a perspective based
on the double-cross-slip mechanism of deformation twinning." Philosophical Magazine
A 82.15 (2002): 2841-2854.
[273] Christian, J. W., and S. Mahajan. "Deformation twinning." Progress in materials
science 39.1-2 (1995): 1-157.
[274] Meyers, M. A., and K. K. Chawla. 2008. Mechanical behavior of materials. Cambridge
university press.
[275] Youngdahl, C. J., et al. "Deformation behavior in nanocrystalline copper." Scripta
Materialia 44.8-9 (2001): 1475-1478.
[276] Shen, Y., et al. "Deformation mechanisms of a 20Mn TWIP steel investigated by in situ
neutron diffraction and TEM." Acta materialia 61.16 (2013): 6093-6106.
[277] Matsuoka, Y., et al. "Effect of grain size on thermal and mechanical stability of austenite in
metastable austenitic stainless steel." ISIJ international 53.7 (2013): 1224-1230.
[278] Koch, C. C. "Optimization of strength and ductility in nanocrystalline and ultrafine grained
metals." Scripta Materialia 49.7 (2003): 657-662.
[279] Misra, R. D. K., et al. "Microstructure and texture of hot-rolled Cb-Ti and V-Cb microalloyed
steels with differences in formability and toughness." Metallurgical and Materials
Transactions A 34.10 (2003): 2341-2351.
[280] Yamamoto, S., et al. "Recrystallization texture and Young's modulus of ceramic particle
dispersed ferrite steel bars." Tetsu-to-hagane 82.9 (1996): 771-776.
[281] Takaki, S., et al. "Effect of grain refinement on thermal stability of metastable austenitic
181
steel." Materials Transactions 45.7 (2004): 2245-2251.
[282] Jimenez-Melero, E., et al. "Martensitic transformation of individual grains in low-alloyed
TRIP steels." Scripta Materialia 56.5 (2007): 421-424.
[283] Gong, N., et al. "On the mechanical behavior of austenitic stainless steel with nano/ultrafine
grains and comparison with micrometer austenitic grains counterpart and their biological
functions." Journal of the mechanical behavior of biomedical materials 101 (2020): 103433.
[284] An, X. H., et al. "Significance of stacking fault energy in bulk nanostructured materials:
Insights from Cu and its binary alloys as model systems." Progress in Materials Science 101
(2019): 1-45.
[285] Huang, Y.L., 2006. Characterize the Relationship between Stress-Strain of Thin Films
through Nanoindentation Approach. Master thesis. Xiangtan University (in Chinese).
[286] Oliver, W. C., and G. M. Pharr. "An improved technique for determining hardness and elastic
modulus using load and displacement sensing indentation experiments." Journal of materials
research 7.6 (1992): 1564-1583.
[287] Tsuru, T., et al. "Incipient plasticity of twin and stable/unstable grain boundaries during
nanoindentation in copper." Physical Review B 82.2 (2010): 024101.
[288] Ledbetter, H. M. "Predicted monocrystal elastic constants of 304-type stainless
steel." Physica B+ C 128.1 (1985): 1-4.
[289] Ledbetter, H. M. Predicted single-crystal elastic constants of stainless-steel 316. Br. J. Non-
destr. Test. 23(1982): 286–287.
[290] Tromas, C., et al. "Hardness and elastic modulus gradients in plasma-nitrided 316L
polycrystalline stainless steel investigated by nanoindentation tomography." Acta
Materialia 60.5 (2012): 1965-1973.
182
[291] Chen, T., et al. "The effect of grain orientation on nanoindentation behavior of model
austenitic alloy Fe-20Cr-25Ni." Acta Materialia 138 (2017): 83-91.
[292] Voyiadjis, G. Z., and C. Zhang. "The mechanical behavior during nanoindentation near the
grain boundary in a bicrystal FCC metal." Materials Science and Engineering: A 621 (2015):
218-228.
[293] Misra, R. D. K., et al. "Nanoscale deformation behavior of phase-reversion induced
austenitic stainless steels: the interplay between grain size from nano-grain regime to coarse-
grain regime." Metallurgical and Materials Transactions A 43.13 (2012): 5286-5297.
[294] Hu, C. Y., et al. "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." Journal of the mechanical behavior of biomedical
materials 104 (2020): 103666.
[295] Fischer-Cripps, A. C. "Handbook of nanoindentation." Fischer-Cripps Laboratories Pty Ltd,
Forestville, Australia (2009).
[296] Morris, D. G. "Strengthening mechanisms in nanocrystalline metals." Nanostructured
Metals and Alloys. Woodhead Publishing, 2011. 299-328.
[297] Kramer, D., et al. "Yield strength predictions from the plastic zone around
nanocontacts." Acta Materialia 47.1 (1998): 333-343.
[298] Wang, M. G., and A. H. W. Ngan. "Indentation strain burst phenomenon induced by grain
boundaries in niobium." Journal of materials research 19.8 (2004): 2478-2486.
[299] Zhang, K., et al. "Assessment of advanced Taylor models, the Taylor factor and yield-surface
exponent for FCC metals." International Journal of Plasticity 114 (2019): 144-160.
[300] He, T. Q., et al. "Effect of feeding Ca–Mg–RE–Zr composite cored wire during refining of
183
liquid steel on abrasive wear resistance of high-strength steels." Wear 303.1-2 (2013): 524-
532.
[301] ASTM E112, standard test method for determining average grain size, in: ASTM Annual
Book of Standards, vol. 96, ASTM International, W., Conshohocken PA, 2004.
[302] Chattopadhyay, C., et al. "Improved wear resistance of medium carbon microalloyed bainitic
steels." Wear 289 (2012): 168-179.
[303] Lei, C., et al. "Mechanical properties and strain hardening behavior of phase reversion-
induced nano/ultrafine Fe-17Cr-6Ni austenitic structure steel." Journal of Alloys and
Compounds 689 (2016): 718-725.
[304] Hall, E. O. "The deformation and ageing of mild steel: III discussion of results." Proceedings
of the Physical Society. Section B 64.9 (1951): 747.
[305] Petch, N. J. "The cleavage strength of polycrystals." Journal of the Iron and Steel
Institute 174 (1953): 25-28.
[306] Karjalainen, L. P., et al. "Some strengthening methods for austenitic stainless steels." steel
research international 79.6 (2008): 404-412.
[307] Cahoon, J. R., W. H. Broughton, and A. R. Kutzak. "The determination of yield strength
from hardness measurements." Metallurgical transactions 2.7 (1971): 1979-1983.
[308] Sun, G. S., et al. "Microstructural evolution and recrystallization behavior of cold rolled
austenitic stainless steel with dual phase microstructure during isothermal
annealing." Materials Science and Engineering: A 709 (2018): 254-264.
[309] Bressan, J. D., et al. "Influence of hardness on the wear resistance of 17-4 PH stainless steel
evaluated by the pin-on-disc testing." Journal of materials processing technology 205.1-3
(2008): 353-359.
184
[310] Leiro, A., et al. "Wear of nano-structured carbide-free bainitic steels under dry rolling–
sliding conditions." Wear 298 (2013): 42-47.
[311] Matsuoka, Y., et al. "Effect of grain size on thermal and mechanical stability of austenite in
metastable austenitic stainless steel." ISIJ international 53.7 (2013): 1224-1230.
[312] Talonen, J., et al. "Effect of strain rate on the strain-induced γ→ α′-martensite transformation
and mechanical properties of austenitic stainless steels." Metallurgical and materials
transactions A 36.2 (2005): 421-432.
[313] Zhang, P. Strain induced martensite behavior of 316L stainless steel subjected to warm
deformation, Heat Treatment Metal 44.2 (2019): 44–49. (in Chinese)
[314] Xu, D. M., et al. "On the deformation mechanism of austenitic stainless steel at elevated
temperatures: A critical analysis of fine-grained versus coarse-grained structure." Materials
Science and Engineering: A 773 (2020): 138722.
[315] Shipway, P. H., S. J. Wood, and A. H. Dent. "The hardness and sliding wear behaviour of a
bainitic steel." Wear 203 (1997): 196-205.
[316] Yang, J., et al. "Sliding wear resistance and worn surface microstructure of nanostructured
bainitic steel." Wear 282 (2012): 81-84.
[317] Leiro, A., et al. "Tribological behaviour of carbide-free bainitic steel under dry rolling/sliding
conditions." Wear 273.1 (2011): 2-8.
[318] Kankanala, A.. "Unlubricated rolling/sliding wear behaviour of high silicon carbide-free
steels." Master’s Thesis, Luleå University of Technology, Sweden, (2010).
[319] Chen, H.H., J.D. Xing, and W. Li. 2006. Handbook for Wear Resistant Materials
Application, Beijing: China Machine Press. (in Chinese).
185
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.
3. Journal of the Mechanical Behavior of Biomedical Materials, 104(2020): 103666.
4. Materials Science and Engineering: A, 793(2020): 139885.
5. Journal of the Mechanical Behavior of Biomedical Materials, 101(2020): 103433.
Contact Information: [email protected]