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NOVEL TRIBOLOGICAL SYSTEMS USING SHAPE-
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NOVEL TRIBOLOGICAL SYSTEMS USING SHAPE-MEMORY ALLOYS AND
THIN FILMS

By

Yijun Zhang

A DISSERTATION

Submitted to
Michigan State University
in partial fulfillment of the requirements
for the degree of

DOCTOR OF PHILOSOPHY
Department of Chemical Engineering and Materials Science

2006

ABSTRACT

NOVEL TRIBOLOGICAL SYSTEMS USING SHAPE-MEMORY ALLOYS AND
THIN FILMS

By
Yijun Zhang

Shape memory alloys and thin films are shown to have robust indentation-induced
shape memory and superelastic effects. Loading conditions that are similar to
indentations are very common in tribological systems. Therefore novel tribological
systems that have better wear resistance and stronger coating to substrate adhesion can be
engineered using indentation-induced shape memory and superelastic effects.

By incorporating superelastic NiTi thin films as interlayers between chromium nitride
(CrN) and diamond-like carbon (DLC) hard coatings and aluminum substrates, it is
shown that the superelasticity can improve tribological performance and increase
interfacial adhesion. The NiTi interlayers were sputter deposited onto 6061 T6 aluminum
and M2 steel substrates. CrN and DLC coatings were deposited by unbalanced magnetron
sputter deposition. Temperature scanning X-ray diffraction and nanoindentation were
used to characterize NiTi interlayers. Temperature scanning wear and scratch tests
showed that superelastic NiTi interlayers improved tribological performance on
aluminum substrates significantly.

The two-way shape memory effect under contact loading conditions is demonstrated

for the first time, which could be used to make novel tribological systems. Spherical

for the first time, which could be used to make novel tribological systems. Spherical
indents in NiTi shape memory alloys and thin fihns had reversible depth changes that
were driven by temperature cycling, afier thermomechanical cycling, or one-cycle
slip-plasticity deformation training. Reversible surface topography was realized after the
indents were planarized. Micro- and nano- scale circular surface protrusions arose from
planarized spherical indents in bulk and thin film NiTi alloy; line surface protrusions
appeared from planarized scratch tracks. Functional surfaces with reversible surface
topography can potentially result in novel tribological systems with reversible friction
coefficient.

A three dimensional constitutive model was developed to describe shape memory
effects with slip-plasticity. After implementing the model in a commercial finite element
program (ABAQUS), spherical indentation-induced superelasticity in NiTi was simulated
and the results were compared to the results of experiments, including measurement of
load-displacement curves and recovery ratios. Shallow indents were shown to have
complete recovery upon unloading; deep indents had partial recovery. Stress-strain
distributions, and the extent of the indentation-induced phase transformation region were

also analyzed.

Acknowledgement

I am much obliged to Prof. David.S.Grummon at Michigan State University and Dr.
Yang-Tse Cheng at General Motors Corp. for their advice, help, encouragement and
inspiration in the art and practice of scientific research, which are truly invaluable to me.
It would have been impossible for me to complete even a minute piece of this work
without their support.

I give sincere appreciation to my committee members: Dr. Carl Boehlert, Dr. Martin
A. Crimp, Dr. John Shaw and Dr. Thomas Pence for their teachings, expertise,
suggestions and generous help.

I am also gratefiil for the help received fiom : Anita Weiner, Mike Luktich, Erkan
Konca, Jihui Yang, and Yue Qi at General Motors Corp.; Dr. Thomas R. Bieler, Dr.
Yuchen Lan, Hairong J iang, F ei Ren, John Foltz, Kebin Low, and J iaming Zhang at

Michigan State University

Table of Contents

List of Tables .................................................................................................................... viii
List of Figures ..................................................................................................................... ix
Chapter 1 Introduction ......................................................................................................... 1
Chapter 2 Literature Review .............................................................................................. 6
2.1 NiTi Shape Memory Alloys ............................................................................... 6
2.1.1 Historical Development of Shape Memory Alloys .................................... 6

2.1.2 Phases in NiTi Alloy .................................................................................. 7

2.1.3 The Shape Memory Effect ......................................................................... 9

2.1.4 Phase Transformation Temperatures ........................................................ 11

2.2 Sputter-Deposited NiTi Thin Films .................................................................. 13
2.2.1 Fabrication of NiTi Thin Films ................................................................ 13

2.2.2 Effect of Stress on Transformation: the Clausius-Clapeyron Equation 14

2.3 Superelastic Effect of NiTi Shape Memory Alloy ........................................... 17
2.3.1 Residual Stress in NiTi Thin Films .......................................................... 17
2.3.2 Temperature Dependence of Superelasticity ............................................ 19
2.3.3 Superelasticity and Phase Transformations .............................................. 20
2.3.4 Mechanical and Microstructural Behaviors in Superelasticity ................. 24
2.3.5 Effects of Thermo-Mechanical Treatments on Superelasticity ................ 25

2.4 The Two-way Shape memory Effect ................................................................ 27
2.4.1. Introduction .............................................................................................. 27
2.4.2. Training Methods and Mechanisms of TWSME ..................................... 28
2.4.3 Phase Transformation Temperatures of TWSME .................................... 33

2.5 Contact Mechanics and Tribological Study ..................................................... 36
2.5.1 Elastic Contact Problems .......................................................................... 36
2.5.2 Elastic-Plastic Contact Problems and Instrumented Indentation ............. 38
2.5.3 Indentation-induced Shape Memory Effect ............................................. 41
2.5.4 Wear and Scratch Damage ....................................................................... 42
2.5.5 Tribological Study on Bulk NiTi Shape Memory Alloy .......................... 46
2.5.6 Hard Coatings and Interlayers .................................................................. 48

2.6 Tribological Study of Hard Coatings on Soft Substrates ................................. 50

2.6.1 Experimental Research ............................................................................. 50
2.6.2 Theoretical Research ................................................................................ 53
2.6.3 Preliminary Study of CerNiTizAl Tribological System ......................... 57
2.7 Constitutive Modeling of Shape Memory Effect ............................................. 59
2.7.1 One dimensional Constitutive Model of SME ......................................... 59
2.7.2 One—dimensional Ivshin-Pence Model ..................................................... 61
2.7.3 Three Dimensional Constitutive Models of SME .................................... 63
2.8 Summary .......................................................................................................... 65
Chapter 3 Methods and Materials .................................................................................... 67
3.1 Tribological Study ............................................................................................ 67
3.1.1 Substrate Materials ................................................................................... 67
3.1.2 NiTi Interlayer and Hard Coating Deposition .......................................... 68
3.1.3 Special Treatments of DLCzNiTi Interfaces ............................................ 70
3.1.4 Characterization of NiTi Interlayer Properties ......................................... 70
3.1.5 Tribological Testing Methods .................................................................. 74
3.2 Training and Characterization of The Two-way Shape Memory Effect .......... 75
3.3 Finite Element Modeling of Spherical Indentation on Superelastic NiTi ........ 78
3.3.1 Formulation of Three-dimensional Constitutive Model ........................... 78
3.3.2 Formulation of Slip Plasticity ................................................................... 81
3.3.3 Finite Element Program Implementation ................................................. 83
3.3.4 Validation and Discussion ........................................................................ 85
Chapter 4 Results and Discussion .................................................................................... 90
4.1 Tribological System with NiTi Interlayers ....................................................... 90
4.1.1 Residual and Phase Transformational Stress in NiTi Interlayers ............. 91
4.1.2 Wear Resistance of CrN Coatings on 6061 Al Substrates with NiTi
Interlayers ............................................................................................................. 94
4.1.3 Contact Stress Induced Phase Transformation in NiTi Interlayer ............ 99
4.1.4 Wear Resistance of DLC Coatings on Soft Substrates with NiTi
Interlayers ........................................................................................................... 102
4.1.5 Wear Resistance of CrN Coatings on Hard Substrates with NiTi
Interlayers ........................................................................................................... 104
4.1.6 Interfacial Adhesion between CrN and NiTi .......................................... 105

vi

4.1.7 Interfacial Adhesion between DLC:NiTi ............................................... 107
4.1.8 Finite Element Modeling of Indentation on CrN:NiTi:A1 and CrN:Al.. 110

4.2 Indentation-induced Two-way Shape Memory Effect ................................... 112
4.2.1 Indentation-induced TWSME fi'om Thermomechanical Training ......... 113

4.2.2 One-cycle Deformation Training and Reversible Surface Topography. 117

4.2.3 Layered Planarization ............................................................................. 121

4.2.4 Small Hysteresis and “Self Training” of Reversible Surface Protrusion 123

4.2.5 Mechanism of Indentation-induced TWSME ........................................ 127
Chapter 5 Conclusions and Future Work ....................................................................... 131
5.1 Conclusions: Hard Coating Systems with NiTi Interlayer ............................. 131

5.2 Conclusions: Indentation-Induced TWSME .................................................. 133

5.3 Suggestions to Future Work ........................................................................... 134
Appendices ....................................................................................................................... 136
References ........................................................................................................................ 229

vii

List of Tables

Table 1. Comparison of crystallographic data of NiTi martensite a, Otsuka et al [20]; b,
Hehmann and Sandrock [235]; c, Michal and Sinclair [21]; d, Kudoh et al [22]. .. 136

Table 2. Lattice structure and parameters of precipitates in NiTi alloys ................. 136

Table 3. Deposition parameters for NiTi interlayers on 6061 T6 substrates and M2
steel substrates, before CrN deposition ................................................................... 137

Table 4. Deposition parameters for NiTiCu interlayers on 6061 T6 substrates, before
CrN deposition ........................................................................................................ 137

Table 5. Deposition parameters for NiTi interlayers on 6061 T6 substrates, before
DLC deposition ....................................................................................................... 138

Table 6. Deposition parameters for NiTi interlayers on M2 tool steel substrates, before
CrN deposition ........................................................................................................ 138

Table 7. Phase transformation temperatures of thin film and bulk NiTi used in
TWSME study ........................................................................................................ 139

Table 8. Material parameters of NiTi and reference materials used in finite element
modeling. (The “Ref” stands for reference material) ........................................... 139

viii

List of Figures

Figure 1. Phase diagram of a NiTi alloy by Massalski[241], to which the phase
equilibrium between the B2 and Ti3Ni4 phases are added [242]. ........................... 140

Figure 2. Transmission electron micrographs of Ti3Ni4 precipitates in 51.3at%NiTi
thin films which were solution-treated at 973 K for 3.6 ks followed by age-treatment
at 773 K for (a) 3.6 ks; (b) 36 ks and (c) 360 ks, at 673 K for (d) 3.6 ks; (e) 36 ks and
(f) 360 ks, and at 573 K for (g) 3.6 ks; (h) 36 ks and (i) 360 ks (after reference [51]).
................................................................................................................................. 140

Figure 3. Schematic illustration of parent and martensite with invariant plane as a
habit plane. .............................................................................................................. 141

Figure 4. Stereographic projection of 24 habit plane variants from WLR theory, for
<01—l > Type II twinning. ..................................................................................... 142

Figure 5. Ms temperature as a function of Ni content for binary NiTi alloys. Different
data symbols represent data fi'om different authors. The solid line is given by
thermodynamic calculations (after reference [34]). ................................................ 143

Figure 6. Schematic stress-temperature diagram for isochronal annealing of sputtered
NiTi films on (100) silicon. Dark curves P1, P2 and P3 show the general intrinsic
stress—deposition temperature characteristics for different pressures of the working
gas. Trajectories a, b, c and d are for amorphous films deposited at different
pressures and at temperatures below Ted, the minimum in-situ crystallization
temperature. Isochronal crystallization occurs at Tc giving a transient tensile stress
spike at x. Curve d additionally shows the effect of an isothermal hold inserted in the
isochronal sequence. Trajectories e, f and g are for crystalline films deposited above
T64. Cooling from any point on a trajectory generally proceeds along a path with
slope 1.5 MPa/K, to points such as d‘, a‘ and g‘, allowing a wide range of stress
levels to be obtained at T a. (after reference [58]) ................................................... 144

Figure 7. Stress-strain curves as a function of temperature of 50.6 at.% NiTi alloy,
which was solution-treated at 1273 K for l h. The critical stress for inducing
martensites and that for reverse transformation are indicated by an arrow (—-)) and
by a double arrow (—-)—->) in (i), respectively. (after reference [68]) .................. 145

Figure 8. Critical stress as a function of temperature for inducing martensites (open
circles) and for reverse transformation (solid circles), which were plotted fi'om the

ix

data in Figure 7. (after reference [68]) ................................................................. 146

Figure 9. Schematic representation for the appearance of the shape memory effect
and superelasticity in temperature—stress space. (after reference [4]) .................. 147

Figure 10. Critical stress for inducing martensites and for reverse transformation of a
50.3at.%NiTi single crystal, which was solution-treated at 1273 K for 1 h followed
by aging at 673 K for 1 h. (after reference [243].) ............................................... 147

Figure 11. Stress—strain curves as a function of temperature of a Ti—50.5Ni single
crystal, which was aged at 673 K for l h after solution-treatment. The curves are
related with R-phase transformation or deformation of R-phase. T’m=Rs. (after
reference [72]) ......................................................................................................... 148

Figure 12. (a) Critical stresses for R-phase deformation and superelasticity, and (b)
the recoverable strain of R—phase superelasticity, which is a function of temperature.
(after reference [72]) ............................................................................................... 149

Figure 13. Elongation vs. temperature curves under 40 MPa load for Ti—Ni—Cu alloys
with various compositions indicated. (after reference [73]). ............................... 150

Figure 14. Stress—strain curves of 50.6 at.%NiTi alloy, which was deformed at 243 K
(>Aj). The alloy was solution-treated at 1273 K for 1 h. Dashed lines represent the
recovered strain upon heating to 373 K. The symbol (x) represents the fracture point.
(b) Plots of various strains in (a) as a function of total strain. (afier reference [68])
................................................................................................................................. 1 5 1

Figure 15. Deformation stage of martensitic NiTi alloy sample (after reference [78] ).
................................................................................................................................. 152

Figure 16. TEM images of in-situ tensile loading of superelastic NiTi. (a) martensite
plates 1, 2, 3, 4 formed at 118 um displacement. (b) High density of looped
dislocations (marked by ABCD) formed at 150 um displacement. (c) Magnified
photo of the area between martensite plates 1, 2 and 3, in (b). (d) Diffraction pattern
corresponding to the area indicated by single-headed arrow in (c). (After reference
[79].) The martensite plate 2 shows some detwinning at the lower part; but the plate
3 does not seems to detwin, even though dislocations have been generated. ......... 153

Figure 17. (above) Tensile cyclic stress—strain response of single crystal NiTi oriented
in the [112] direction. The test temperature is above the A, transformation

X

temperature. (below) Compressive cyclic stress-strain response of single crystal NiTi
oriented in the [112] direction. The test temperature is above the A f transformation
temperature. (after reference [76]) ........................................................................ 154

Figure 18. Observation of Luders deformation by photography for Ti-50.5Ni alloy.
Black regions represent the parent phase, and lighter colored regions martensite. The
strain increases with increasing numbers flour (1) to (13). (afier reference [83]).155

Figure 19. Effect of thermomechanical treatment on the stress—strain curves as a
function of deformation temperature for 49.8at.%NiTi alloy. Solution-treatment at
1273 K for 1 h followed by annealing at 673 K for 1 h. (after reference [244]) 155

Figure 20. Effect of thermomechanical treatment on the stress—strain curves as a
function of deformation temperature for 49.8at.%NiTi alloy. Annealed at 673 K for
1 h immediately after cold-working. (after reference [244]) .................................. 156

Figure 21. TEM images of 49.8at.% NiTi alloy, which were annealed for l h at
various temperatures indicated after cold-working; (a) as cold-worked, (b) at 67 3 K,
(c) at 773 K and (d) at 873 K. (after reference [245]) .......................................... 157

Figure 22. Stress—strain curves as a function of testing temperature of 50.6at.%NiTi
alloy, which was aged at 773 K for 1 h after solution-treatment at 1273 K for l h.
................................................................................................................................. 1 5 8

Figure 23. Effect of aging temperature on the critical stresses for inducing
martensites and for reverse transformation in a 50.6at.% NiTi alloy, which was aged
at various temperatures for 1 h after solution treatment. (after reference [244]). 159

Figure 24. TEM images and corresponding diffraction patterns of 50.6at.% NiTi
alloy specimens, which was aged at various temperatures after solution-treatment at
1273 K for 1 h; (a and a’) at 673 K, (b and b’) at 773 K. (after reference [245]) 160

Figure 25. Stress—strain curves as a function of deformation temperature of
50.6at.%NiTi alloy, which was aged at 673 K immediately after cold-working. (after
reference [245]) ....................................................................................................... 161

Figure 26. The procedure of processing TWSME TiNi SMAs spring (afier reference
[87]) ......................................................................................................................... 162

Figure 27. The two-way shape memory effect obtained after a single deformation by

xi

6.2% of equiatornic NiTi alloy (after reference [102]). .......................................... 162

Figure 28. Optical microscopy verification of the predicted microstructure memory for
an aged Auled 49 martensite and its disappearance after aging in the parent phase.
(a) Domain pattern of the aged martensite (aged at room temperature for six months).
The sample is heated up to (b) parent phase (60 °C), and immediately cooled down to
(c) martensite. Then the sample is heated up again to (d) parent phase (60 °C) and
aged for 30 minutes, followed by cooling down to (e) martensite. (after reference
[104]) ....................................................................................................................... 163

Figure 29. TWSME due to internal tensile residual stress: elongation during cooling
and shrinkage during heating. ................................................................................. 164

Figure 30. External Stress and Internal Compressive Stress Field. The tensile external
stress increases the internal tensile stress, elevating phase transformation
temperatures but compressive external stress decreases the internal tensile stress. 164

Figure 31. Strain-temperature curves during (a) cooling and (b) heating through the
transformation region of a trained specimen at different external stresses. The
external stresses are indicated in the rectangulars (in MPa). The pure elastic
deformation which is caused by the loading in the austenitic condition has been
subtracted to obtain a common curve in the austenitic condition. .......................... 165

Figure 32. DSC curves of bulk NiTi samples used in TWSME study .................... 166

Figure 33. Stress distributions at the surface and along the axis of symmetry caused by
(left) uniform pressure and (right) Hertz pressure acting on a contacting radius a (after
K.L. Johnson, Contact Mechanics, Cambridge University Press, 1985, pp.94, [108]).
................................................................................................................................. 167

Figure 34. Two-dimensional photo-elastic fringe patterns (contours of shear stress).
(a) point load; (b) uniform pressure; (c)rigid flat punch; ((1) contact of cylinders.
(after reference [108], pp. 103). ............................................................................... 168

Figure 35. The “pile-up” (above) where contact radius is larger than linear case and,
“sink-in” (below) where contact radius is smaller. ................................................. 169

Figure 36. Representative strain vs. hardness in spherical indentation, and its
correlation with true stress-true strain in uniaxial tension. Metal A is steel and metal B
is copper. (after reference [I I I, I 12]) .................................................................... 170

xii

Figure 37. Finite element modeling of indentation of an elastic-plastic half space by a
rigid sphere: development of the plastic zone. Broken line: contour of .12. Py is the
normal load needed to first initiate plastic flow. (after reference [114]) .............. 171

Figure 38. Recovery of spherical indents in NiTi alloy by the shape memory effect.
(after reference [231] ) ............................................................................................ 172

Figure 39. Load-displacement curves of (a) Berkovich and (b) spherical indentations
in SE NiTi. (after reference [1 1 7] ) ......................................................................... 172

Figure 40. TEM images from the indent-affected zone in 50.9at%NiTi single crystals.

Images (a) to (c) were taken adjacent to the same indent at different magnifications.
(afier reference [I 18]). ............................................................................................ 173

Figure 41. Images of wear marks on NiTi after 200 cycles of scanning scratch under
different normal loads and temperatures. A f is 16 °C. (afier reference [144] ) ....... 174

Figure 42. Calculated stresses due to an isolated unit normal line load applied at the
origin as a function of substrate modulus E2. Variation of bearing stress 03, and shear
stress try at the film-substrate interface are shown for a 1000 nm film with a reference
elastic modulus E1=l . (after reference [198]) ....................................................... 175

Figure 43. Calculated tensile stress a, at the film-substrate interface for a 1000 mm film
with a reference elastic modulus E1=1 for a unit normal line load at the origin.
Substrate modulus E2 is varied in the upper figure. The effect of varying the film
thickness on the flexure stress is shown in the lower figure for a substrate of modulus
E2=0.2. (after reference [198]) ............................................................................ 176

Figure 44. Wear test results of samples with] um CrN coating and with: (from left to
right) no interlayer, with 4 um SE NiTi interlayer (S1), 4 um Cr interlayer and 4pm
SMA NiTi interlayer (82). (after reference [3] ) .................................................. 177

Figure 45. SEM Images of scratch track end with (a) lum CrN coating, no interlayer,
(b) 5 um CrN coating no interlayer, (0) 1pm CrN coating 4pm SE NiTi interlayer, (d)
lum CrN coating and 4pm SMA NiTi interlayer and (e) 1pm CrN coating and 4m Cr
interlayer. (after reference [3]) ............................................................................. 178

Figure 46. Vogit (parallel) model for stress induced martensite in austenite phase.
(after reference [204] ) ............................................................................................ 179

xiii

Figure 47. Series model for stress induced martensite in austenite phase. (afier
reference [204] ) ...................................................................................................... 179

Figure 48. X ray diffraction patterns of NiTi thin films (sample #1) at various
temperatures ............................................................................................................ 180

Figure 49. X ray diffraction patterns of NiTi thin films (sample #2) at various
temperatures sample ................................................................................................ 180

Figure 50. Temperature scanning X-ray diffraction of NiTi thin film on 6061T6
substrates, from which phase transformation temperatures were estimated and listed
in TABLE. CrN coating was later deposited onto the NiTi thin films. ................ 181

Figure 51. Temperature scanning X-ray diffraction of NiTiCu thin film on 6061T6
substrates. CrN coatings were later deposited on the NiTi thin films .................. 181

Figure 52. Temperature scanning X-ray diffraction of NiTi thin film on 6061 T6
aluminum substrates. The NiTi thin films were interlayers of DLC coatings ...... 182

Figure 53. The normalized load-depth curves from Berkovich indentation of the NiTi
thin films at different temperatures. ........................................................................ 182

Figure 54. Diagram of thermomechanical training method to create two-way indent
depth change. .......................................................................................................... 183

Figure 55. Stress-strain curves of simple loading conditions from the 3D constitutive
model for superelastic effect with plasticity. For the biaxial loading curve, the strain
rate is set the same in two loading directions. The dashed curve represents a high

uniaxial load which causes plastic hardening and permanent plastic strain. .......... 183
Figure 56. Finite element mesh of indentation model. The refined region is the
contact region. Total node number is around 15000. ........................................... 184

Figure 57. Absolute (a) and normalized (b) load—displacement curves of SE NiTi and
reference materials from finite element modeling. Indentation-induced
superelasticity enables large or 100% indent depth recovery. ................................ 185

Figure 58. Recovery ratios vs. representative strains at different indentation depths,
and the size of indentation-induced phase transformation region at different

xiv

indentation depths, measured in contact radius. ..................................................... 186

Figure 59. Comparison of experimental and calculated load-displacement curves.
The experimental results are adapted from Ni et 01.6 Modeling SE NiTi has a phase
transformation hysteresis A (A= (A f- A,)/2-(Ms - Mf)/2 ) 15K and 30K for the two
computed curves. .................................................................................................... 186

Figure 60. Contours of the austenite phase volume fiaction at indentation depth of
0.018um (a), and complete indent recovery after unloading (b), where few residual
martensites are present after unloading ................................................................... 187

Figure 61. Contours of the austenite phase volume fraction (a) and plastic strains (b) at
indentation depth of 0.10am. (c) Contours of the austenite volume fiaction after
unloading. For the partially recovered indent (recovery ratio 71%), there were much
more residual martensites than in FIG. 5(b). .......................................................... 188

Figure 62. Residual stress evolution of NiTi thin film on Al substrates with
temperature, for three heating-cooling cycles ......................................................... 189

Figure 63. First cycle residual stress evolution of NiTi thin film on Al substrates with
temperature. ............................................................................................................ 189

Figure 64. The stress/temperature rate of NiTi thin film on Al substrates during
martensite phase transformation. ............................................................................ 190

Figure 65. The stress/temperature rate of NiTi thin film on Al substrates from CTE
mismatch between thin film and substrates, without the influence of phase
transformation, ........................................................................................................ 190

Figure 66. Curvature change of 2‘SDLCaniTigégAl sample between room
temperature around Mfand elevated temperature above A f. The curvature change
mainly resulted from residual stress change in NiTi interlayer. ............................. 191

Figure 67. Temperature scanning wear deformation of CrN:NiTi:Al and CrN:Al
samples, under 0.49 and 0.98 N load, using a 200 um tip radius spherical indenter.
Superelastic NiTi improves the wear resistance by limiting plastic deformation to the
soft aluminum substrates ......................................................................................... 192

Figure 68. SEM images of temperature wear tracks of 3"’CerA1 samples (top row) and
3“S'CrN:'5NiTi353:Al sample (bottom row), under 0.98N load. The wear data was

XV

presented in Figure 67 ............................................................................................. 193

Figure 69. SEM images of temperature scanning wear tracks of 3‘9CerAl samples
(first row) and 3'9CrN: lsNiTi353,:Al sample (second row), under 0.49N. The wear data
was presented in Figure 67. .................................................................................... 193

Figure 70. SEM images of temperature scanning wear tracks of l'7CrN:9NiTiCu343:Al

samples. The wear data was presented in Figure 71 . ............................................ 194
Figure 71. Temperature scanning wear deformation of "7CermiTiCu343zAl and
l'7CrN2Al samples, under 0.49 and 0.98 N load. ..................................................... 194

Figure 72. The effect of interlayer and CrN coating thickness on wear resistance at
room temperature. The wear test was done using 6.35mm diameter WC under 5N
load. Thicker coating and interlayer gave better wear resistance. ........................ 195

Figure 73. Wear loss vs. temperature using a different loading condition (6.3 5mm ball,
0.98N load). Wear resistance was not improved at temperature above .................. 196

Figure 74. Elastic contact stress in elastic halfspace from 3.175mm radius and 200nm
radius sphere under different loads. ........................................................................ 196

Figure 75. Temperature scanning wear tests of 2'SDLCzNiTimzAl samples, using
200nm tip radius diamond indenter under 0.49N load. Thick NiTi interlayer and SE
NiTi interlayer at temperature above A f showed better wear. ................................. 197

Figure 76. SEM images of temperature scanning wear tracks of 2'5DLC :NiTi363zAl and
2‘SDLC:Al samples. The wear data was presented in Figure 88. .......................... 198

Figure 77. The wear resistance of CrN:NiTi:M2 samples at room temperature as a
function of CrN coating thickness. ......................................................................... 199

Figure 78. The wear resistance of CrN:NiTi:M2 samples at room temperature as a
function of interlayer thickness ............................................................................... 199

Figure 79. SEM images of temperature scanning scratch tracks on 3'S’CergNiTi353:Al.
Images of the whole scratch tracks (1 mm in length )are placed at the upper right
comer. ..................................................................................................................... 200

xvi

Figure 80. SEM images of temperature scanning scratches tracks on 3'S’CerAl.
Images of the whole scratch tracks (1 mm in length )are placed at the upper right
corner. ..................................................................................................................... 200

Figure 81. Acoustic emission of scratch tests on 3'9Cer9NiTi353:Al. It showed strong
temperature dependence that matched phase transformation temperatures of NiTi
interlayers ................................................................................................................ 201

Figure 82. Acoustic emission of scratch tests on 3'9CrN:Al. There is no significant
temperature dependence .......................................................................................... 201

Figure 83. EDS spectrum of the delarrrinated area in scratch tracks on
3'9Cer9NiTi353zAl, showing that delamination occurred at the interface between CrN
and NiTi interlayer. ................................................................................................. 202

Figure 84. EDS spectrum of the delaminated area in scratch tracks on 3‘9CerAl,
showing the delamination occurred at the interface between CrN and the Cr interlayer.

The Cr interlayer (<um) was deposited before CrN as a precursor layer. It is an
indispensable process of making CrN ..................................................................... 202

Figure 85. Critical load of CrN coating in scratch tests as a function of temperature.

The superelasticity of NiTi interlayer improved the scratch resistance of CrN coating.
................................................................................................................................. 203

Figure 86. SEM images of scratch track on 2'SDLCzAl sample, showing severe DLC
coating failure and delamination. 7N progressive load was used. ........................ 203

Figure 87. SEM images of scratch track on 2‘SDLC:'5NiTi3¢;3:Al sample, showing
much less coating failure compared to the previous figure without NiTi interlayer.
7N progressive load. ............................................................................................... 204

Figure 88. SEM images of scratch track on 2'SDLC:'5NiTi3¢;g:Al sample, under high
3ON progressive load. Although there was coating failure, no dramatic delamination
occurred, proving the strong interface between DLC coating and NiTi interlayer. 205

Figure 89. EDS of delarninated region in scratch tracks on 2'5DLC:'5NiTi363:Al sample,
shows the delarrrination occurred at DLC and (Cr)NiTi interface. The extra Cr layer
on NiTi is represented by the high Cr peak here, compared to the weak Cr peak in
Figure 83. ................................................................................................................ 206

xvii

Figure 90. Computational load-displacement curves from finite element modeling of
spherical indentation on 3'9CrN:15NiTi:Al and 3'9CerAl, using 200 um tip radius
spherical diamond indenter. .................................................................................... 206

Figure 91. Horizontal strain component 8,, in 3'SiCerlsNiTizAl and 3'9CrN:Al fi'om
FEA ......................................................................................................................... 207

Figure 92. Vertical strain component 3,, in 3'9CrN:15NiTi:Al and 3‘9CerAl from FEA
................................................................................................................................. 208

Figure 93. Shear strain component 8,, in 3'9CrN: 15NiTizAl and 3"”CerAl from F EA
................................................................................................................................. 209

Figure 94. Martensite volume fiaction in NiTi interlayer from FEA. The blank region
represent the coating and substrate because there is no martensitic phase. ............ 209

Figure 95. Horizontal stress component on in 3 ‘9CrN:15NiTi:Al and 3'9CerAl fi'om
FEA ......................................................................................................................... 210

Figure 96. Vertical stress component oyy in 3‘9CerlsNiTizAl and 3'9CrN:Al from F EA
................................................................................................................................. 211

Figure 97. Shear stress component 0,, in CrN:NiTi:Al and CerAl from F EA ...... 212

Figure 98. (a) 3D profile of a heated and cooled indent, (b) Sectional profile of the

heated and cooled indent ......................................................................................... 213
Figure 99. Two—way indent depth change over several thermal cycles. .................. 214
Figure 100. Indent depth change ratio and absolute indent depth change. .............. 214

Figure 101 . Indent depth change ratio (DM-DA)/DM vs. horizontal direction, referred to
center. The center of indent had smallest ratio of depth change. Further away from
the indent, both DM and DA become small and dominant by the surface roughness,
causing the fluctuation in the ratio. Ideally the (DM-DA)/DM should approach zero
for surfaces far away from indents because (DM'DA) should equal zero. But Did-DA is
not zero because: 1) small surface roughness change in NiTi between martensitic and
austenitic phase, 2) measurement error between two measurements. .................... 215

xviii

Figure 102. Schematic diagrams of experimental procedures and mechanisms for
achieving RSP. (a), indentation is accommodated by the deformed martensite and
slip—plasticity. (b), When heated to the austenite phase, the indent is partially
recovered. (c), when cooled down to martensite phase, the residual indent became
deeper. The indent depth change is reversible with heating and cooling temperatures
(Figure 2). The dislocation structure and its associated stress field can facilitate the
growth of oriented martensite variants during austenite-to-martensite transformation,
leading to two-way indent depth change. (d), The residual indents are removed by
polishing. (e), The oriented martensite region and dislocations can still exist after the
indent is removed. (f), A reversible surface protrusion is formed upon heating. . 216

Figure 103. Reversible indent depth change. A spherical indent made with a steel ball
pressed relatively deeply into a martensitic sample of NiTi not only recovers a
substantial percentage of the indent depth when heated above the Af temperature, but
also subsequently displays cyclic displacement behavior as the temperature is cycled
fi'om below M/to above Af. .................................................................................... 217

Figure 104. Thermally-reversible surface protrusions. A 3x3 matrix of
thermally-reversible protrusions produced in a NisogTiae-r alloy by first creating deep
spherical indentations and then planarizing the specimen while it was still in the
martensitic condition. The protrusions appear and disappear as the specimen is
heated from below M; to above the A f temperature. .............................................. 218

Figure 105. Sectional profiles of reversible circular surface protrusions and peak
height over five thermal cycles. .............................................................................. 219

Figure 106. Actual shape of the circular surface protrusion .................................... 219
Figure 107. Schematic of different kind of surface fluctuations. Image from ...... 220

Figure 108. Reversible GM logo made from adjoining reversible surface protrusions.
Imaged in normal bright field optical microscopy .................................................. 220

Figure 109. 3D profiles of reversible line surface protrusions from scratch. .......... 221

Figure 110. 3D profiles of nanoscale circular surface protrusions on NiTi thin
film/Silicon substrate. ............................................................................................. 221

Figure 111. 3D profiles of reversible nano scale line protrusion on NiTi thin
film/Silicon substrate .............................................................................................. 222

xix

Figure 112. Sectional profiles of the nano circular surface protrusions and peak height
over five thermal cycles. ......................................................................................... 222

Figure 113. Decay of protrusion amplitude with surface-removal. When successive
surface layers of thickness dp; are removed after initial planarization to a depth ho, the
amplitude of protrusion appearance and disappearance, hp, diminishes and approaches
zero when material has been removed to a depth roughly equal to the radius of the
initial indent contact ................................................................................................ 223

Figure 114. Plastic strain distribution under an indent. Theoretical and experimental
results , adapted from Samuel and Mulhearn [23 2] showing the expected magnitude
of plastic strain under an indent in a ductile metal. ................................................ 224

Figure 115. Height of reversible circular protrusions vs. temperature. The phase
transformation temperature of NiTi sample is also marked. The hysteresis of the
reversible protrusion is much narrower. ................................................................. 225

Figure 116. Schematic of Residual Stress in Reversible Surface Protrusion. The
residual stress is in the opposite direction of TWSME during heating from martensite

to austenite phase, and not in the opposite direction during cooling, narrowing
hysteresis. The residual stress also stabilizes the TWSME. ...... . ............................ 226

Figure 117. Geometrically necessary dislocations. It is assumed that the indentation is
accommodated by circular loops of geometrically necessary dislocations with
Burgers vectors normal to the plane of the surface. As the indenter is forced into the
surface of a single crystal, geometrically necessary dislocations are required to
account for the permanent shape change at the surface. ......................................... 227

Figure 118. Residual stress in y direction (03,.) after spherical indentation in austenitic
NiTi alloy, by finite element analysis. The negative sign is compressive direction.
................................................................................................................................. 228

Images in this dissertation are presented in color.

Chapter 1

Introduction

Shape memory alloys and thin films have been shown to have robust indentation-
induced shape memory effect and superelasticity. Experimental and theoretical work
was undertaken to explore potential novel tribological systems utilizing
indentation-induced superelasticity and shape memory effects:

(1) Using superelastic NiTi thin films as interlayers between hard coatings and
soft substrates as a method to enhance tribological performance of hard
coatings/soft substrates systems;

(2) Demonstration of the indentation-induced two-way shape memory effect,
which can lead to future tribological systems with fiiction coefficient that
changes with temperature.

(3) Finite element modeling of spherical indentation on superelastic NiTi alloy to
understand stress-strain and phase transformation evolution in indentation-
induced superelasticity.

Wear of machine components made from the lightweight alloys such as aluminum

and magnesium is an important issue in the automobile and aerospace industries. The
past two decades have seen rapid development of hard coatings produced by various

physical and chemical vapor deposition methods. The successful application of hard

coatings to hard substrates like machine tools and cast dies has stimulated research efforts
towards deploying hard coatings on soft metal substrates. However, one of the most
difficult obstacles is that soft substrates are prone to plastic deformation before the failure
of the hard coating. The following question is addressed: What is the best ‘foundation’
(interlayer) to support a stiff, hard, wear resistant coating for substrate materials that have
arbitrarily poor elastic-plastic support characteristics and a high CTE mismatch? Such an
interlayer must be capable of shielding the substrate from plastic damage while
accommodating large displacements in the hard coating, all without allowing cyclic
damage accumulation (fatigue failure) in either the interlayer or the substrate. This
interlayer must also manage thermal mismatch strains to allow use of thicker coatings to
limit 0', (the stress component parallel to the coating surface ). What is needed is an
elastically compliant material with very high effective resilience - one in which fully
reversible strains beyond l-2% can occur at moderate stress without causing
unacceptable lattice damage. An interlayer with such properties must be made
sufficiently thick to disperse the strain field associated with contact loading, to a degree
sufficient to shield the substrate material from cyclic plastic strain.

A superelastic NiTi interlayer can potentially meet the above requirements.
Preliminary studies have shown that a superelastic NiTi interlayer can improve
tribological performance [1-3], but there is no systematic information on how the
evolution of martensite-austenite phase transformations affects tribological performance.

This knowledge can only be obtained fiom wear and scratch tests carried out over the

2

complete transformation temperature range. Another key problem is how the
coating:interlayer thickness ratio affects tribological performance. The residual stresses
in a CrN coating and a NiTi interlayer interact with each other, and also with the contact
stress during tribological testing. Residual stresses also increase with coating and
interlayer thickness. Detailed knowledge of this problem can be obtained from
tribological tests over a matrix of coating:interlayer thickness ratios and over a
temperature range that covers the phase transformation temperatures.

The two-way shape memory effect (T WSME) under contact loading conditions,
which can be utilized to make novel tribological systems, has been demonstrated for the
first time in this work. Spherical indents in NiTi shape memory alloys and thin films,
after thermomechanical cycling or one-cycle slip-plasticity deformation training,
displayed a reversible depth change that was driven by temperature cycling.

Polymer thin films that exhibit photonic-controlled reversible surface relief gratings
can be utilized in optical data storage and diffractive optical elements. Metals with
similar ability can find optical and mechanical applications for their greater load bearing
ability and often better durability than polymers. This idea has been realized in a NiTi
SMA by indentation-induced two-way shape memory effect. Micro and nano scale
circular surface protrusions arose from planarized spherical indents in bulk and thin film
NiTi alloys. Line surface protrusions appeared fi'om planarized scratch tracks. This
reversible surface topography was driven by martensiteHaustenite phase transformations.

This interesting phenomenon raises many intriguing questions as to the basic mechanism

3

driving the indentation-induced two-way effect. The results point to novel applications,
from self-healing of wear damage, to creation of surfaces with thermally-variable
coefficient of friction. Furthermore, the ease with which large numbers of regularly
spaced features can be produced by this method may lead to highly robust
optically-fimctional surfaces, and the ability to create the effect at very small length
scales may enable new kinds of high-energy actuators for micro-electro-mechanical
systems.

Finite element modeling of indentation in superelastic NiTi alloy is helpful to better
understand the novel tribological systems using shape memory alloys and thin films. A
three—dimensional constitutive model was developed to describe shape memory effects
with plasticity. After implementing the model in a commercial finite element program
(ABAQUS), spherical indentation-induced superelasticity in NiTi was simulated and
compared to results fi'om experiments, including measurement of load-displacement
curves and recovery ratios. Shallow indents were shown to have complete recovery
upon unloading; deep indents had partial recovery. The stress-strain distribution and the
extent of the indentation-induced phase transformation region were also modeled.

Chapter 2 briefly reviews recent work in the various research areas involved: shape
memory alloys and thin films, superelastic effects, the two-way shape memory effect,
contact mechanics, hard coatings on soft substrates, and finite element modeling.

Chapter 3 contains the details of materials and methods, including descriptions of

experimental materials, equipment, and test parameters. The formulation of

4

three-dimensional constitutive model of shape memory material and its finite element
implementation is also included in this chapter, along with the validation of the model by
comparison with experimental results.

Chapter 4 discusses the experimental and computational results. The first section
addresses the novel tribological systems with NiTi interlayers. The phase
transformational stress in NiTi thin films on aluminum substrates, and these covered by a
hard coating is strong evidence of phase transformations occur in the NiTi interlayer.
Tribological test results on hard coating on aluminum and steel substrates with NiTi
interlayers are presented and discussed, along with finite element modeling of the
tribological systems. The second part concerns the indentation-induced two-way shape
memory effect. Two-way shape memory effects fi'om two different training methods are
presented. Reversible surface topography can be realized from the two-way shape
memory effect by planarization. Possible mechanisms of this two-way shape memory
effect are then discussed.

Chapter 5 summarizes the dissertation and gives suggestions for future work.

Chapter 2

Literature Review

2.1 NiTi Shape Memory Alloys

2.1.1 Historical Development of Shape Memory Alloys

The discovery of thermoelastic martensite phase transformation preceded the
discovery of the shape memory effect. According to Otsuka and Wayman [4] , A.
Olander [5] discovered the pseudoelastic behavior of the Au-Cd alloy in 1932.
Greninger & Mooradian [6] observed the formation and disappearance of a martensitic
phase by decreasing and increasing the temperature of a Cu-Zn alloy in 1938. The first
attempts to clarify the basic mechanisms of thermoelastic martensite phase
transformations was made a decade later by Kurdjumov & Khandros (1949) [7] and also

by Chang & Read [8].

In the early 19605, Buehler and his co-workers at the US. Naval Ordnance
Laboratory discovered the functionality of the thermoelastic martensite phase
transformation in equiatomic NiTi and termed it “the shape memory effect”, which has
attracted academic and industrial interest ever since. This alloy was named Nitinol
(Nickel-Titanium Naval Ordnance Laboratory). Intensive investigations have been
made to elucidate the metallurgy and mechanics of various kinds of shape memory alloys,

which also have been used in a wide range of structures for actuation purposes [9-11].

6

Due to NiTi’s biocompatibility and excellent mechanical properties, shape memory
alloys have recently been used to make implantable medical devices, like cardiovascular

stents [12-18].

2.1.2 Phases in NiTi Alloy

The near-equiatomic NiTi alloy system has a high temperature parent phase
(austenite phase) that transforms displacively to a daughter phase (martensite phase) at
lower temperatures. The hi gh-temperature austenite phase has a B2 CsCl type ordered
structure, with lattice parameter of 0.301 5nm [19]. The low temperature martensite
phase was determined to be monoclirric B19’ type lattice after a series of electron
microscopy studies [20-22]. The ultra-fine size of NiTi martensite twins and variants
made its characterization difficult. The lattice parameters are slightly different in
different reports (Table l). The austenite phase has a higher symmetry than the
martensite phase.

There also exists an R-phase transformation at intermediate temperatures between
austenite-martensite phase transformations, characterized by a sharp increase of electrical
resistivity and a small hysteresis (1-2K). The R-phase has a trigonal P3 type lattice [23,
24]. It was initially considered a precursor effect, or pre-martensitic behavior, but is

now established as martensitic transformation itself, which precedes the B19’ martensite

transformation.

NiTiCu with Cu content above 8-10 at.% has a B19 orthorhombic type martensite
with lattice parameters a0=0.2881 nm, b0=0.4279 nm, c0=0.4514 nm [25-27]

The phase diagram for the binary NiTi system is shown in Figure 1. The 50-50
at.% NiTi matrix has almost no solubility of extra Ti at temperatures below 773K;
solubility of extra Ni decreases rapidly as temperature cools below 1373K.

Various precipitate phases, Ti3Ni4, T12N13, TiNi3, form under certain heat-treatment
conditions. Their lattice structures are listed in Table 2. The size and distribution of
the precipitates are determined by composition and heat treatment. Figure 2 shows how
aging temperature and time affects the size and distribution of Ti3Ni4 precipitates in
51 .3at%NiTi thin films. At aging temperatures around 773K to 1073K, the precipitates
form by diffusion controlled phase transformations. Generally, at lower aging
temperature and for a shorter aging time, Ti3Ni4 appears; at higher aging temperature and
longer aging time, TiNi3 appears; intermediate aging time and temperature, TizNi3
appears. Prolonged aging also decomposes Ti3Ni4, and TizNi3 into TiNi3. These
transformations are summarized in the TTT (isothermal time-temperature-transformation)
diagram [28]. The precipitates are important to the phase transformation temperatures,

mechanical properties and phase transformation characteristics.

2.1.3 The Shape Memory Effect

Thermoelastic martensite-austenite phase transformations and the unique twinning
deformation mechanism of NiTi martensite phase are the two main pillars of the shape
memory and superelastic effects [29]. “Thermoelastic” means that the martensite-
austenite phase transformations are reversible with heating and cooling within a relatively
narrow hysteresis, and the boundary between the austenite phase and the martensite phase
is thermally reversible [30]. During the martensite phase transformation, each
higher-symmetry austenite grain is transformed into self-accommodating
lower-symmetry martensite twins with mobile twin boundaries. Subjected to external
stresses, twins that align with the shear components of the external stress become
energetically favorable and grow at the expenses of other twins. Finally the grain may
become a single crystal of the most favorable martensite variant under stress. In this
way the deformation strain is accommodated by the growth and alignment of martensite
twins. Slip plasticity is not involved until this deformation mechanism is exhausted.
Because of the lattice correspondence between two phases, the deformation strain is
cancelled when the martensite variants are transformed back to the original austenite
grains, and the shape memory alloys “remember” the original macroscopical shape.

The WLR theory is a phenomenological crystallographic theory of martensitic
transformation, named after its developers Wechsler, Lieberman and Read [31, 32]. The
deformation and rotation of a parent lattice is represented by matrices as operators, which

act on an arbitrary vector. The habit plane between parent and martensite is considered

9

to be an invariant plane, which is undistorted and un-rotated, since the invariant plane
minimizes the interfacial strain energy. This situation is shown schematically in Figure
3 for the case where the lattice invariant shear is twinning. The lattice invariant shear is
an inhomogeneous shear occurring without changing the martensite structure.

The martensite plate, which consists of martensite matrix and its twin, is called the
habit plane variant (h.p.v.). The twins in the martensite are the smallest scale martensite
variants. Each has a specific lattice correspondence with the parent phase, which is
called lattice correspondence variants. For < 011> type II twinning (a lattice invariant
shear), phenomenological crystallographic calculation gives 24 h.p.v. which are shown in
a stereographic projection in Figure 4.

The elastic strains associated with martensitic transformations are reduced by the
invariant plain strain condition across the austenite/martensite interface. However, this
condition cannot eliminate the transformation strains completely, because out-of-plane
shear components are not eliminated under this condition, but if two habit plane variants
with opposite shearing directions are formed side by side, the shear strains are eliminated
macroscopically. The hierarchy in reducing elastic strain energy upon transformation
by the combination of multiple variants is called self accommodation of martensite. The
mechanism of self-accommodation in most of the B-phase alloys such as Cu—Al—Ni,
Ag—Cd were successfully analyzed using phenomenological crystallographic theory [33].

According to the mechanism, four or two clusters of habit plane variants (h.p.v.) around

10

[011] poles self-accommodate, and the averaged shape strain matrix becomes close to

the identity matrix, which represents the strain free state.

2.1.4 Phase Transformation Temperatures

There are four phase transformation temperatures: Ms, Mf( martensite start and finish
temperatures ) and A s, A f( austenite start and finish temperatures). Ms and Marc the
temperatures that the martensitic phase starts to form (Ms) and completes the phase
transformation (Mf), respectively. The A s and A f temperatures, the austenite start and

finish temperatures, are defined in the same way. If MS and Marc close to A S and A f,

then the shape memory alloy has a narrow phase transformation hysteresis and vice versa.

One way to define hysteresis, A, is:

A: AS+Af _MS+Mf
2 2 '

 

Transformation temperatures determine the temperature range in which shape
memory effect (SME) or superelastic effect (SE) can be observed, which is important for
practical applications. Experimentally it is known that the martensite transformation
temperatures are strongly dependent on composition. For example, one percent change
in Ni content above 50 at.% can change transformation temperature by more than 100 K.
Figure 5 shows the phase transformation temperature dependency on the Ni content in

NiTi alloy [34]. Therefore, the most effective way of obtaining desired transformation

11

temperature is by tuning composition.

Ternary alloying elements are added to the binary NiTi system to alter
transformation temperatures. Most alloying elements lower transformation
temperatures, but Pd [35], Pt [36], Zr [3 7, 38] and Hf [39-41] increase transformation
temperatures. Pd and Hf are used in making shape memory alloys with high phase
transformation temperatures. The addition of Cu changes the lattice structure of
martensite fi'om B19’ to B19 when the Cu content exceeds 8-10 at.%. Compared to the
B2-B19’ phase transformation, the B2-B19 transformation has a narrower temperature
hysteresis. The size and distribution of precipitates in the ternary NiTiX system has
strong influence on mechanical properties and shape memory characteristics [40, 42].

Aging processes can also influence phase transformation temperatures, especially the
Ms temperature, in binary NiTi alloy. The solubility of Ni in the B2 matrix is shown in
the phase diagram in Figure 1. Quenching from solution temperature around 1173K can
prevent excess Ni fiom precipitating. It enables the study of transformation
temperatures vs. Ni content. Ni4Ti3 and other precipitates are formed in the aging
treatment, which lower the Ni content in matrix and elevate Ms temperatures [43-45].
Since excess Ni can dissolve into the matrix while Ti cannot, the transformation

temperatures of bulk Ti-rich NiTi are more stable [46].

2.2 Sputter-Deposited NiTi Thin Films
Development of NiTi thin films was driven by the prospect for microactuators in

micro-electrical-mechanical system (MEMS) [9]. Magnetron sputter deposition, among
other physical vapor deposition (PVD) methods, like evaporation and laser ablation [47],
is the most effective method of making NiTi thin films with strong SE and SME [48-50].
Sputter—deposited NiTi thin fihns have grain size ranging from 0.5m to several microns,
almost ten times smaller than that of bulk NiTi [51, 52]. The fine grain sizes give good
mechanical properties such as higher yield stress, which is another reason for using NiTi

thin films for actuation.

2.2.1 Fabrication of NiTi Thin Films

Magnetron sputter deposition is a complex process. During the magnetron sputter
deposition process, inert gas ions are accelerated by an electric field and bombard the
cathode surface. Atoms of the cathode are dislodged by this energetic process and are
deposited on the substrate.

Many factors influence the resulting thin films’ composition, mechanical properties
and SME characteristics. The most important parameters are target chemical
composition, deposition power and voltage, Ar pressure, and deposition temperature.
They are discussed briefly here:

(1) At steady state, deposition flux matches the target composition even though

13

different elements have different sputter yield, because the target surface composition is
altered by the differential sputter rate. Differential sputtering only affects the growing
film because it too suffers sputter ejection and doesn’t reach a steady state. As a result,
the sputter-deposited NiTi thin films usually have different chemical composition
compared to the target composition.

(2) Deposition power determines the deposition rate of thin films. Higher
deposition power gives more substrate heating and affects trapping times.

(3) Higher Ar pressure may produce thin films with a columnar structures while
lower Ar pressure gives a featureless structure [51]. Because the kinetic energy of
sputtered atoms may be reduced by collision with Ar ions and the absorbed Ar atoms
may impede the diffusion of Ni and Ti atoms by reducing the kinetic energy driving
surface diffusion [51].

(4) The deposition temperature determines whether NiTi thin films are amorphous
or crystalline. The thin fihns are heated by the kinetic energy of deposited atoms,
normally around 100K above from room temperature. Extra substrate heating is needed
to elevate the temperature of thin films above around 650K so that Ni and Ti atoms have

enough kinetic energy to crystallize [5 0, 53].

2.2.2 Effect of Stress on Transformation: the Clausius-Clapeyron Equation

The Clausius-Clapeyron equation is used widely to describe the pressure dependence

l4

of the equilibrium temperature in the reversible first order phase transformation between
gas-liquid, gas-solid, and solid-solid phases. Let the Gibbs fiee energy of phase A and
B be G .4 and G B . For first order phase transformation, the Gibbs free energy is
continuous but its derivative is not.
do, = d0, , (eq. 2.1)
610 = Vdp — SdT , (eq. 2.2)
For each phase, V is the molar volume, p is the hydrostatic (tri-axial) pressure, T is the
temperature and S is the entropy. Use (eq. 2.2) in (eq. 2.1) to give:
VAdp - SAdT = Vde — SBdT, (eq. 2.3).
If the phase transformation direction is from phase A to phase B, then

AS: 2S3 —SA

, . 2.4
AV,” = V, —V, (6‘1 )

where AH j is the enthalpy change from phase A to phase B

Rearrange (eq. 2.3) to give

 

dp _ 45,?
dTEQ - AVAB , (eq. 2.5)

(eq. 2.5) is the basic form of the Clausius-Clapeyron equation. The dTEQ is the
change in the equilibrium phase transformation temperature due to the pressure. Under
the conditions of constant pressure and constant temperature, and no external work
except fiom volume change from pressure, an entropy change can be expressed in terms

of an enthalpy change:

15

 

Thus (eq .25) can be expressed as

 

dp _ AH j
dTEQ TEQ . AVAB , (eq. 2.7).

For shape memory alloys, uniaxial tension and compression are the more common
loading conditions than hydrostatic pressure. The uniaxial loading also changes phase
transformation temperatures. This phenomenon can be explained by a modification of
the Clausius-Clapeyron equation. If the pressure is replaced by uniaxial stress:

(10 = Edd — SdT , (eq. 2.8).
(eq. 2.5) becomes

do _ASf_ AH:
dTEQ A8? TEQ-AaA ’ (eq' 2‘9)

 

 

The shift of phase transformation temperatures due to external stress can be
estimated for NiTi SMA. Consider uniaxially loading of a lmx lmx 1m NiTi sample
cooled from austenite phase to martensite phase. Typical material properties are:

transformation enthalpy: AH f = 20J / g ,

equilibrium temperature: T=333K,

density: 6.45 g/cm3,

Ni atomic number: 58.7, Ti atomic number: 47.8,

transformation strain: A3]? = 0.05

16

 

B B 6 3 3
d_a'= AS; : AHAB ___ 20J/gx6.45x10 g/m xlm =7.74x106 ”(57.74 MPa/K
dT AaA ToAaA 333Kx0.05

Experimental value shows % = (7 i 1)MPa/ K [4, 54-56]. It is close to the theoretical
prediction because enthalpy change and transformation strain may be different for

different materials. [5 7]

2.3 Superelastic Effect of NiTi Shape Memory Alloy
2.3.1 Residual Stress in NiTi Thin Films

The residual stress in NiTi thin films comes fi'om two origins: the intrinsic stress
from the deposition process, and the thermal stress from mismatch of the coefficient of
thermal expansion (CTE) between the substrate and the thin film [58-62]. The residual
stress in a thin film can be measured by the wafer-curvature method using the
well-known Stoney equation [63, 64], assuming the curvature change is elastic.

The intrinsic residual stress is the result of the non-equilibrium characteristics of the
sputter deposition process, which involves rapid thermal energy dissipation and
simultaneous energetic particle bombardment during film growth. It is natural that the
intrinsic residual stress is affected by deposition parameters such as Ar pressure,
deposition temperature, and whether the thin fihns are crystalline or amorphous. The
intrinsic residual stress can be relaxed by post annealing. Grummon er al. proposed a

general stress model (Figure 6) describing the stress relaxation processes of NiTi thin

films [58]. Stress relaxation path of isochronal and isothermal annealing of NiTi thin
films are given for different deposition conditions. The intrinsic stress in a NiTi thin
film is completely relaxed at crystallization temperature around 753K. Cooling from
that temperature, the coefficient of thermal expansion (CTE) mismatch between silicon
substrates and NiTi gives rise to thermal stress. For instance, the CTE mismatch
between NiTi and Silicon substrates gives -1.0 MPa/K thermal stress rate, which is

derived as:

am, =11.4x10_6,as,. =3x10‘6

EMT, = 83GPa, v = 0.33 ,

E 3
dflz—L-(as —-af) 2 83x10 (-8.4x10‘6) z —l.011lfl’a/K
dT l—v 1—0.33

 

where aNm , a s: are CTE of NiTi thin films and silicon substrates, E NiTi and v
are elastic modulus and Poisson’s ratio of NiTi thin fihns.

The thermal stress can be relaxed by martensite transformation partly or completely,
and restored by the reverse phase transformation. The transformational stress relaxation
rate is governed by the thermodynamics of martensite phase transformation. It can be
expressed in the form of Clausius-Clapeyron equation, which was discussed in 2.2.2.

This transformational stress relaxation and restoration cause a quite significant but
reversible wafer curvature /displacement change within a relatively small temperature
hysteresis of about 80K. This is the basis of some NiTi thin fihn micro-actuators for

MEMS applications [60, 65, 66].

18

2.3.2 Temperature Dependence of Superelasticity

Superelasticity originates from the stress-induced phase transformations upon loading,
and the subsequent reverse transformations upon unloading. The external stress changes
the free energy and induces phase transformations according to the Clausius-Clapeyron
equation discussed in section 2.2.2:

For instance, tensile loading an austenitic NiTi specimen at temperature 50K above
MS requires a = 7.74MPa/ K x 50K = 387MPa tensile stress to induce the martensitic
phase transformation at that temperature.

Superelasticity is therefore closely dependant on testing temperatures. The highest
temperature at which martensite can no longer be stress induced is called Md. Above
Md NiTi alloy deforms like ordinary metals by slip-plasticity. Below the A 3 temperature,
martensite can be stress induced but does not completely revert to austenite. Thus,
superelasticity appears in a temperature range from near A fand up to M4. The greatest
ability to recover occurs just above A f [4]. Figure 7 shows the stress-strain curves for
uniaxial loading of a NiTi specimen at different temperatures. The critical stress to
initiate the plateau region where stress-induced martensites begin to form, increases with
increasing temperature, which is seen in Figure 8, where open circles represent the
critical stresses upon loading, and closed ones are critical stresses upon unloading.

The dependence of the shape memory effect and superelasticity on test temperatures
can be further presented in the schematic diagram of Figure 9. The line with a positive

slope represents the critical stress to induce martensite, which is similar to the one with

19

open circles in Figure 8. The lines with negative slopes represent the critical stress for
slip-plasticity. If the critical stress to introduce plasticity is very low (line B), complete
superelasticity can not be expected, because slip is an irreversible process and has a
detrimental effect on the thermoelastic phase transformations. If the critical stress for
slip is high enough (line (A)), superelasticity can appear at temperatures above A fwithin
the shaded region. Reverse transformation occurs upon unloading, since the martensite
is unstable at temperatures above A fin the absence of stress. Similarly, at temperatures
below A s the deformed martensites are stable, and the strain can recover only by reverse
transformation on subsequent re-heating. In the temperature regime between A s and A f
the martensites are partly unstable. The Md temperature is the maximum temperature
that martensite phase can be stress induced. Above the M4 temperature slip plasticity is

the only inelastic deformation mechanism in austenitic NiTi.

2.3.3 Superelasticity and Phase Transformations

Superelasticity stems from the stress-induced phase transformations. Although the
parent phase is invariably the B2 austenite, there are three low-temperature phases that
can emerge after phase transformations, depending on chemical composition and
thermal-mechanical history. The superelasticity associated with each low temperature

phase is discussed briefly here.

20

BZ-BI 9’ Transformations Stress induced B2-B19’ phase transformation
is the most important to the superelasticity of NiTi. The mechanism of superelasticity is
the stress-induced martensite phase transformation upon loading and reverse phase
transformation upon unloading. There is direct experimental evidence in Cu-Al-Ni
single crystals to support this mechanism [6 7]. In uniaxial loading of a Cu-Al-Ni single
crystal, a particular variant of martensite was stress-induced with the habit plane that was
most favorable under the applied external stress. The parent single crystal changed into
a single crystal of that martensite variant at the end of the stress plateau. Upon
unloading the reverse transformation occurred with the same habit plane and the strain
recovered completely upon unloading [6 7]. The recoverable strain also agreed with
crystallographic predictions.

The verification of the mechanism of superelasticity in NiTi has some complications.
Solution-treated single crystal NiTi has low critical stress for plastic slip, which is the
case of line B discussed in the section 2.3.2 and Figure 9. When such a specimen is
subjected to external stress at temperature above A f, plastic flow started before the onset
of stress-induced phase transformations and thus excluded superelasticity. However
aged single crystal Ni-rich NiTi had higher critical stress for plastic slip due to
precipitation strengthening. Superelasticity was achieved for aged NiTi single crystals,
which are shown in Figure 10. But the aged single crystals did not give full recovery
strain because untransformed regions were left due to the presence of precipitates. As a

result, the recoverable strain was smaller than theoretical value [68].

21

How superelasticity from B2-B1 9’ martensite phase transformation is affected by

thermomechanical treatments is discussed in 2.3.5.

BZ—R Transformation The B2-R transformation is an independent
martensite phase transformation between B2-R-B19’ transformations. However if the
martensite phase start temperature MS is higher than R phase start temperature R3, the
R-Bl9’ phase transformation is suppressed. Superelasticity is also attainable with
R-phase transformation, although the recoverable strain is about 0.5%. Ling and
Kaplov studied the shape memory effect fi'om R—phase transformation in bending tests
coupled with X-ray diffraction measurements [69, 70]. More systematic work on
superelasticity and shape memory effects were done by Miyazaki et a1. by tensile tests on
both 50 at.% Ti- 47 Ni—3 Fe polycrystals [71] and aged 50.5at.%NiTi single crystals [72].
The stress-strain curves of aged single crystal NiTi specimens at different temperatures
are shown in Figure 11.

R-phase transformation temperatures are different from that of B2-B19’
transformations. It is interesting to note that the recoverable strain increases with
decreasing temperature. The reason is that the rhombohedra angle of R-phases changes
with temperature. Figure 12 shows the critical stress for stress-induced R-phase
transformation increased quite rapidly with increasing temperature, because the
recoverable strains is so small that the slope of da/ dT in the Clausius-Clapeyron

equation (discussed in 2.2.2) becomes very large. The available shape memory or

22

recoverable strain of R-phase transformation is small (less than 1% in polycrystals). But
the small transformation strain may contribute to the small temperature hysteresis of
R-phase transformation ( 1 -2K), which is a very useful property for actuator applications
[4]-

BZ-BI 9 Transformations B2-B19 (cubic to orthorhombic) transformation
exists in NiTiCu SMAs, when the Cu content exceeds more than 8%-10%. The
theoretical recovery strain associated with B2—B1 9 transformation was calculated for 40
at.% Ni-10Cu-50Ti alloy as a function of orientation by Nam et al. [73]. The maximum
recoverable strain of 5.3% is obtained at <110>p orientation, which is much smaller than
the maximum 10.5% strain in B2—B19’ transformation. But there is no supporting data
from single crystal NiTiCu experiments.

Polycrystal NiTiCu SMAs have smaller temperature hysteresis in superelasticity,
compared to binary NiTi (Figure 13). The temperature hysteresis decreases with
increasing Cu content, the minimum being 4K for 20at% Cu. It was suggested that
volume change in B2-B19 phase transformation decreases with increasing Cu content up
to 12.5 at.% Cu [73]. Further addition of Cu would increase the volume change. The
small volume change may be one factor for the small temperature hysteresis but that
alone cannot explain all the experimental data. Another mechanism proposed is the
parent phase strengthening by solution of Cu atoms [74]. The frictional stress for twin

boundary movement is suppressed if the introduction of dislocations is avoided.

23

2.3.4 Mechanical and Microstructural Behaviors in Superelasticity

The loading and unloading of a superelastic NiTi sample involves complicated
changes in microstructure and multiple stages. In stress-strain curves, the loading stage
has a plateau region, after stress has reached a critical level. The plateau region
involves stress-induced martensite phase transformation that is triggered by a resolved
shear stress over the most favorably oriented martensite habit plane. An internally
twinned structure is formed within a single martensite plate. Such a martensite plate can
be called a correspondence variant pair (CVP) because two related martensite single
crystal variants develop. This is the first stage of superelastic loading that is
well-understood.

The second stage is characterized by the strain hardening with slightly increasing
superelastic strain, which was noticed by Miyazaki in multistage loading of superelastic
NiTi alloy ( Figure 14) [68]. It is suggested the detwinning process may explain this
behavior [75 -77]. The detwinning is that one martensite variant of the CVP grows at the
expense of the other and finally saturates the whole martensite plate. The detwinning
process has been observed in TEM of loaded martensitic NiTi alloy (Figure 15 and [78]),
which occurred after the alignment of martensite twins (CVPs) with external stress was
complete. Finally the martensite plate consists of only one variant instead of twins.
Dislocations are generated in the late stage or after the detwinning process has finished
(Figure 15). From then on slip-plasticity is the main deformation mechanism [75, 78],

which is the third stage. In-situ TEM observation of stress-induced martensite

24

transformation and plastic deformation was carried out by J iang et al. [79]. However
they did not intend to study the detwinning in NiTi alloy, although the TEM micrographs
showed some evidence of detwinning (Figure 16).

Tension-compression asymmetry was found present in single crystal and textured
polycrystalline NiTi SMA [76, 77, 80-82]. The reason is that the CVP with the largest
resolved shear stress component is induced preferentially than the rest 23 CVPs, which
can be described by the Schmid law [76]. The tension-compression asymmetry arises
from the different Schmid factors between tensile and compressive directions. Figure
17 shows the tension-compression asymmetry of [211] orientated single crystal NiTi
alloy [76]

In polycrystal NiTi samples without texture, the stress induced martensite
transformation plateau is found to be inhomogeneous, like Luders deformation, that
involves nucleation at some point within the specimen and the subsequent propagation,
according to direct observation with an infrared imaging radiometer by Kyriakides and

Shaw (Figure 18, [83]).

2.3.5 Effects of Thermo-Mechanical Treatments on Superelasticity
Thermomechanical treatments, mainly by slip-plastic deformation and aging
treatments, are important to the superelasticity and shape memory effects.

Thermoelastic phase transformations are reversible but thermomechanical treatments

25

introduce irreversible changes to the microstructure. Their main purpose is to increase
the critical stress for plastic slip by strengthening the parent phase. As discussed in
2.3.2, low critical stress for slip in solutionized single crystal NiTi makes the realization
of superelasticity impossible. Higher critical stress for slip means that superelasticity
can be obtained over a wider temperature range and with larger recoverable strain. The
full potential of NiTi SMA can only be brought out after suitable thermomechanical
treatments.

Cold-working followed by annealing enhances superelasticity. Cold work
introduces dislocations into the specimen; post annealing is a recovery process during
which dislocations relax and rearrange. The critical stress for plastic slip is increased
after this process. As a result, the plateau stress of superelastic loading can also be
increased by loading at higher temperature. The temperature and stress range for
superelasticity is thus increased. Figure 19 shows a series of stress-strain curves from
solution-treated 49.8at.% NiTi wires. There is only partial superelasticity at temperature
above Af. After uniaxial cold drawing and post annealing at 673K for 1h, good
superelasticity was achieved (Figure 20). Figure 21 shows the TEM images of
specimens at different stages of thermomechanical treatments. NiTi alloy was in a
recovered but not recrystallized state after annealing at 673K, and led to the best
superelasticity. Increasing the annealing temperature to 773K resulted in
recrystallization, which again lead to lowered critical stress of slip and decreased

superelasticity.

26

Aging treatment enhances superelasticity by precipitation strengthening in the parent
phase. Figure 22 shows the stress-strain curves of 50.6at.% NiTi wire that were
solutionized and aged at 673K for 1h. Superelasticity was achievable only in a narrow
temperature range. Figure 23 shows that after solution treatment and aging at 773K for
1h, much better superelasticity was obtained. TEM study (Figure 24 ) showed that
precipitates were small but densely distributed in the matrix after aging at 673K;
precipitates were larger and less densely distributed after aging at 773K. Small and
densely distributed precipitates were more effective to strengthen the parent phase, which
results in improved superelasticity.

Combined effects of cold-working and aging enhance superelasticity. Combined
use of work-hardening and age-hardening is more effective to increase the critical stress
for slip, which is presented in Figure 25. Solutionized 50.6at.%NiTi wires were cold
drawn and aged at 673K for 1h. Compared to Figure 22, complete superelasticity was
available at higher stress level. Besides, the combined treatment was useful for

increasing cyclic stability of superelasticity [84].

2.4 The Two-way Shape memory Effect
2.4.1. Introduction
The two-way shape memory effects (TWSME) is the reversible shape change during

martensite and reverse phase transformations, without the stimulus of external stress. In

27

conventional one-way shape memory effect, the martensite transformation does not result
in any macroscopic shape change because the martensite variants are self-accommodated
to minimize elastic and transformation strain.

The two-way shape memory effect is not an intrinsic effect but can only be obtained
by certain thermomechanical treatments, termed as “training”. It was first noted in 1965
by Wang et al [85]. In 1974 Nagasawa et a1. did the first detailed research on this
“reversible shape memory effect” in NiTi and CuZn SMA[86]. After the training
procedures, microstructural and mechanical anisotropies are introduced into the material.
The martensite variant formation follows these structural anisotropies during phase
transformations to minimize internal energy. Macroscopic strain would appear because
martensite variants form that have preferential alignment. Reverse austenite
transformation cancels the shape change. More details about the TWSME mechanism

will be discussed in 2.4.2.

2.4.2. Training Methods and Mechanisms of TWSME

Various training methods have been developed to realize TWSME: (l) constrained
aging, (2) thermomechanical cycling, (3) plastic deformation by slip. There are
variations in configurations of each category. The mechanism associated with each of

these training methods is discussed below.

28

Constrained Aging Training Figure 26 illustrates the process of an
constrained aging training of NiTi springs [8 7]. The spring was aging at 773K under
constraint, which was an over stretched position. After aging, the spring was deformed
in martensite phase to a crushed position. The reverse transformation expanded the
spring while the martensite transformation shortened the spring, completing the two-way
cycle. However the authors did not propose any possible precipitates that might be
induce the TWSME.

Isalgue reported that the aging process under external stress created the oriented
growth of y precipitates in Cu-Zn-Al SMA [88]. Zhang further found that oriented
strain field around these 7 precipitates were distributed along one direction in austenite,
regardless of the grain orientation [89]. But the 7 precipitates are embedded in the
austenite phase coherently. While in the martensite phase, 7 precipitates did not result
in any strain fields. The reversible strain fields around precipitates that were produced
by constrained aging may explain the TWSME.

Fukuda et al. aged a single crystal 51 .at% NiTi under compressive stress along [111]
[90]. Ni4Ti3 precipitates were formed coherently to the matrix along [111]. They
explained the resulting TWSME from an energy approach: the volume fraction of
precipitates was about 9%, based on TEM study of the size, shape and distribution of the
precipitates. The matrix was expanded elastically about 0.2% on average along the [111]
direction. The elastic energy accumulated in the matrix was estimated to be about 0.3

MJ/m3; it was about 1.5% of the latent heat of the B2 --> R transformation. The

29

thermally induced R-phase was composed of a single variant and the spontaneous strain
along the [111] direction was about 0.8% of expansion. The maximum work from the
spontaneous shape change was estimated to be about 0.17 MJ/m3. Chang et al. reported
constrained aging at room temperature [91], where NiTi wires were first solutionized and
aged at elevated temperature and then bended in round shape in room temperature for 24
h. This, however, is more likely to be the deformation training method.

Precipitates formed in constrain aging training can be observed using TEM [89, 90].
It has been shown that the oriented strain fields around precipitates in NiTi can cause
martensite variants to align in preferential directions. The stress fields around the
precipitates can be recovered again after the reverse austenite phase transformation [89].
Molecular dynamics simulation indicated that oriented tension stress fields in the middle
line of the dislocation band are favorable sites for the martensite nucleation and growth of

preferred variants [89].

Thermomechanical Training Thermomechanical training methods have
different variations, depending on the researcher’s interpretation of mechanisms of
TWSME. The most common setup was cycling between temperatures above A f and
below M; when the specimen was subjected to external tension, compression or bending
stress [92-95]. Pons et al. used the following training cycles: cooling down below the
Mftemperature under a constant tensile uniaxial stress of 35 MPa, and then unload and

heat up to A fwith zero stress [96]. This low stress was used to train a Cu-based shape

30

memory alloy. Prader et al. kept the tensile strain constant while thermal cycling
between Mr and M4 temperatures. There is no consensus about the optimal training
method.

Explanations of the development of the TWSME often involve either the presence
after training of a low proportion of locally stabilized martensites in the parent phase, or
the presence after training of an internal stress field in the material. Stalmans [92] et al.
carried out a well-cited research about thermomechanical cycling on Cu-Zn-Al SMA and
discussed the mechanism of TWSME. They disproved these two previous explanations.
Following their study, the most important effect of the dislocation arrays which are
introduced during the training is not to develop internal stresses but rather to give rise to a
microstructural anisotropy inducing a thermodynamic anisotropy. The thermodynamic
anisotropy means that after training, the crystallographically equivalent martensite
variants are no longer thermodynamically equivalent: the trained martensite variants are
thermodynamically preferred to the non-trained variants. The authors argued that this
thermodynamic anisotropy must have come from certain changes in microstructures as
the result of training. The microstructural basis of the thermodynamic anisotropy was
termed microstructural anisotropy. However, the exact nature of the microstructural
anisotropy was not explained clearly. The authors speculated that dislocations arrays
generated during the training should contribute to it. Manach et al. carried out a similar
thermomechanical training procedure but in the pure shear loading conditions [97].

They supported Stalmans conclusion that internal stress field had only a weak influence

31

on the development of TWSME, and dislocation arrays introduced by training created not
only internal stresses but also a thermodynamic anisotropy. They further suggested the
anisotropy is due to the friction energy which is after training no longer equal for all

martensite variants.

One cycle slip-plasticity deformation training The slip-plastic
deformation training was actually the training method used in the original discovery [85]
and initial study of TWSME [86]. The specimens were deformed in the martensite
phase to introduce a certain level of dislocations [98-103]. Slip-plasticity training has
the advantage of requiring fewer training cycles. Figure 27 shows the TWSME
obtained from the uniaxial tensile deformation training. As discussed previously, the
mechanism of TWSME from slip-plasticity training is believed to be that the
microstructural anisotropies (oriented precipitates, dislocations, residual martensites) and
their associated mechanical anisotropies would select martensite variants in specific
orientations that are energetically favorable.

Although dislocations can be observed in transmission electron microscopy (TEM),
they look like densely tangled black lines. Although possible, it is difficult to get
information such as strains and strain gradient from direct TEM observation of
dislocations. The localized stress fields can not be observed in TEM directly. because
these obstacles, the mechanism of TWSME has not been clarified using TEM.

The one-to-one correspondence between TWSME and microstructure “memory

32

effect” was found in Au-Cd alloy, where the specimen exhibited same variant
configuration upon heating and cooling (Figure 28) under optical microscope [104].
Similar results are more difficult to get from NiTi alloy due to the fine size of martensite

variants.

2.4.3 Phase Transformation Temperatures of TWSME

The phase transformation temperatures are affected by the presence of TWSME.
Thermomechanical training of NiTi wires lowered the A S and A f temperatures
significantly but had a less influence on MS and Mftemperatures [105]. Stalmans et al.
showed that external stress opposing the direction of TWSME shifted all transformation
temperatures downwards [106].

If a NiTi bar was trained by slip-plasticity method using uniaxial tensile stress, the
resulting TWSME would be: shrinkage during heating and elongation during the cooling.
Assume there is a tensile residual stress field in the NiTi bar that is responsible for
inducing the TWSME (Figure 29). Applying external compressive stress to NiTi bar
decreases the internal tensile stress, if the level of the external compressive stress is not
higher than that of the internal tensile stress. A schematic diagram is shown Figure 30.
The actual stress inside the specimen is decreased by the external stress. According to
the Clausius-Clapeyron equation, lowered stress also lowers the phase transformation

temperatures.

33

However, the external compressive stress is in the opposite direction of the TWSME.
During cooling the external compressive stress impedes the shrinkage of NiTi bar. If
large enough, the compressive stress can suppress the TWSME completely.

In experiments, both lowering phase transformation temperatures and suppressing of
TWSME were observed when applying external stress in the opposite direction of
TWSME (Figure 31) [106]. In Figure 31, strain change vs. temperature is presented for
the cooling (3) and heating (b) direction, for various external stress levels. The TWSME
with zero external stress is also shown. Note in Figure 31 the compressive stress is in
the opposite direction of TWSME, which suppresses completely the TWSME at
-52.lMPa. The phase transformation temperatures are elevated by the tensile stress and
lowered by the compressive stress in the opposite direction of TWSME.

Stalmans et al. also proposed a thermodynamic explanation for this interesting
phenomenon [106]. The Clausius-Clapeyron equation derived in 2.2.2 can be expressed

as

5:. AH?
AT TAef’

 

(eq. 1.10).

The enthalpy change from austenite phase to martensite phase is positive (heat flows
into NiTi): AH f > O , which is why heat flow jumps in the cooling DSC curves when

phase transformation occurs (Figure 32). The left hand side of the C-C equation 3%:-

M

A
TAs‘;w '

 

must has the same sign as the right hand side If a normal SMA specimen was

loaded by tensile stress and cooled from austenite to martensite phase, the phase

34

M

transformation strain was also tensile. Since AH j" > 0, A8;M > 0 and T Ag“ > 0 ,
A

 

the phase transformation temperatures were shifted upwards (AT > 0) compared to free
standing specimen, because the tensile stress A0 >0.
When heating from martensite to austenite phase, the AH f; < 0 . The recovery of

phase transformation strain is in the opposite direction of tensile strain: A8; < 0. As a

A

AH
result AH :4 < 0, A3; < 0 —) T A; > 0 . There was no change to the tensile stress,
M

 

so the phase transformation temperatures were still increased (AT > 0).

In TWSME, however, the situation is different depending on the direction of external
stress. Consider a specimen that was trained to possess TWSME giving elongation
during austenite-martensite transformation on cooling (As): > 0) and shrinkage during
martensite-austenite transformation on heating ( A8}; < O ). Tensile stress still elevates
phase transformation temperatures as described in the previous case. But compressive

stress lowered the phase transformation temperatures becomes: for cooling from austenite

M

AH
to martensitephase, AHf>0,A8f >0 and TAefM >0. Since Aa'<0
A

 

(compressive), so AT < 0 that the left hand side of (eq. 1.10) was positive as the right
A

M
TAef,

 

hand side; is still positive when heating fi'om martensite to austenite phase,

because AH 3', <0, A8; <0. Since Aa < 0 (compressive), AT < 0. The key point
is that the external stress was in opposite direction against strain change in TWSME as

compared to the case in SME, which caused the negative temperature shift. Figure 31

35

shows the strain-temperature curves in Stalmans’ study that exhibited the lowering phase

transformation temperatures.

2.5 Contact Mechanics and Tribological Study

The subject of contact mechanics may be said to have started in 1882 with the
publication by Heinrich Hertz of his classic paper 0n the contact of elastic solids [107].
His interest in the problem was aroused by the question whether the force of holding
glass lenses could have a significant influence on the pattern of optical interference
fiinges. However, developments in the theory did not appear in the literature until the
early 20th century, stimulated by engineering developments on the railways, in marine
reduction gears, and in the rolling contact bearing industry.

Now contact mechanics are used widely in science and engineering disciplines like
geology, astronomy, defense, materials and mechanical engineering, where impact and
contact is investigated. Most importantly contact mechanics serves as the theoretical

basis to tribological testing methods like indentation, wear and scratch tests.

2.5.1 Elastic Contact Problems
Analytical solutions are given for many elastic contact problems. The elastic

contact problems of spherical and sharp indenters into a half space are especially

36

important because they are directly related to the measurement of hardness. The mean

contact pressure Pm is simply normal load P on the indenter over contact area 7:02 :

P
pm: 2
7m

 

In the elastic spherical indentation, the contact radius a can be calculated:

3PR 1/3
a:
(a) ,

where R is the radius of indenter, and E . is the reduced modulus, a function of the

 

modulus of the indenter and the substrate:

1 1 l
- +
E

 

*
E E indenter substrate

On the contact surface the contact pressure is not uniformly distributed. The maximum

contact pressure p0 is expressed as

3P 6PE*2 ”3
_ 27m2 _ 71'3R2

 

P0

The vertical displacement of the surface within the contact radius, measured with respect
to the original specimen free surface, is given by:

2
— Usubstrate ”p0 (2‘12 _

l
u, (r) = -— 4a

r2), rSa,

substrate

where Umbsmm, is the Poisson’s ratio of substrate [108]. The stress distribution is
depicted in Figure 33 [108]. The maximum shear stress is located at a distance of half
contact radius below surface. Figure 34 presents the photo-elastic fringes pattern of

shear stress fi'om indenters of various geometries. They all show concentric

37

hemispherical contours. For uniform pressure and cylindrical contact, the highest stress
is located at a distance from the contact surface. This distance is much shorter in the

case of point loading. The flat punch shows stress concentration at the contact edges.

2.5.2 Elastic-Plastic Contact Problems and Instrumented Indentation

In metals, slip-plasticity develops when the elastic contact stress exceeds the
material’s yield stress. As the indenter is pressed deeper into the substrate, a plastic
zone appears below the indenter tip, surrounded by an elastic region. This is called the
“contained mode”. The plastic zone grows with indentation depth and finally breaks out
to the fi'ee surface. From then on the displaced material is free to escape by plastic flow
to the sides of the indenter. This is the uncontained mode of indentation. It begins

when mean contact pressure pm reaches about triple the yield stress:

P
p :——:CO’.Id.
m 7:02 ”8

When p," is between 1 and 3, the plastic zone is contained by elastic materials [108].
The mean pressure is also defined as the hardness of the material.

Indentation is a traditional method to measure the hardness of materials.
Instrumented indentation is the advanced form of indentation, where both displacement
and load of the indenter is monitored continuously. It has been developed during the

past decade driven by the need to measure modulus and hardness of thin films efficiently.

38

Nix et al. first proposed that the contact radius could be calculated from
load-displacement curves, assuming that the initial part of the unloading curve is linear
[109]. However when the “sink-in” or “pile-up” situation is present (Figure 35), the
Doemer and Nix method is not accurate enough because the contact radius is larger in the
“pile-up” situation and smaller when the “sink-in” occurs. This method was later
improved by Oliver and Pharr [110], whose method is now used to calculate the contact
radius.

Analytical solutions for most elastic-plastic contact problems are not readily
available, especially with hardening present. But most metals are elastic-plastic solids
with hardening behavior. Both finite element analysis and semi-analytical models are
commonly used in analyzing elastic-plastic contact problems.

A semi analytical: spherical cavity model was proposed by Johnson to explain the
elastic-plastic indentation process [108]. In this spherical cavity model, the contact
surface of the indenter is encased in a hemispherical “core” with a hydrostatic
component fi. Outside the core it is assumed that the material will respond the same
way as in an elasfic half space containing a spherical cavity under hydrostatic pressure F.
The elastic-plastic boundary is somewhere outside the spherical cavity. The hydrostatic

pressure fl below the indenter can be expressed as:

W+4(1-20)

-2 Y
—3{l+ln[ 60w) 1}.

 

*VIEI

39

where Y is the yield stress of indented material. The .3 is the angle of indenter. For
the Berkovich indenter, 8:19. 3°.

The representative strain proposed by Tabor is also used to characterize spherical
indentation into elastic-plastic solids [I 1 1, 112]. The representative strain is a depth-

sensitive parameter that is defined as:
a
8, = 0.2 — ,
R

where a is the contact radius and R the indenter radius. When the representative strain
is equal to the true strain, the hardness H is about 2.8 times of the true stress (Figure 36,
[1 I 1, 1 12]):

E- = 2.8, when 8, = 8
Y

Tabor arrived at this conclusion after comparing the experimental data of spherical
indentation and uniaxial deformation of copper and steel. The indentation and uniaxial
deformation data set collapsed into a single line (Figure 36) when applying the above two
equations. The correlation between spherical indentation experiments and uniaxial
’ tensile data, although an empirical law, indicates that using the concept of representative
strain, spherical indentation experiments may possibly be used to measure the true
stress-true strain relationship.

Finite element analysis is an effective tool for investigating elastic-plastic contact
problems [113]. Figure 37 shows the finite element modeling of the expansion of

plastic zone under spherical indentation [114]. Sinclair et al. carried out a detailed finite

40

element analysis of the spherical indentation on elastic-plastic solids [115]. They
confirmed Tabor's simple experimental correlations for work-hardening materials using
the representative strain, and found generally good agreement between theory and other
physical evidence. However they also found there was high strain gradient near the

contact edge that was not predicted before.

2.5.3 Indentation-induced Shape Memory Effect

Shape memory effect under contact loading conditions has received very little
attention previously. This is partly due to the complexity of the problem and the
absence of suitable characterization methods. Ni et al. reported the shape memory and
superelastic effects in NiTi alloy during spherical and sharp indentation [116, 11 7] using
a well-controlled instrumented indentation method. A recovery ratio r] was defined to

quantify the indent recovery:

77 = hrnax - hresidual
hmax ’

where hmx is the maximum indentation depth, and hm

Mm, is the indent depth after

recovery. Spherical indents in NiTi alloy recovered almost completely (7] =100%) due

to the shape memory or superelastic effect, if the indentation representative strain was
less than 5% [116, 117]. Spherical indents in martensitic NiTi recovered after being

heated above Af (Figure 38) while in austenitic NiTi the indents recovered during

41

unloading (Figure 39 (a)). The recovery ratio was only about 45% during sharp
indentation (Berkovich), because slip-plasticity impeded the martensite-austenite
transformation and affected SME adversely. Deeper spherical indentation also resulted
in decreased recovery ratio.

Gall et al. did Vickers indentation on NiTi single crystals and observed the existence
of dislocations in the vicinity of indents in TEM Gigure 40). The same study also
suggested that the precipitates size and crystal orientation also affect the recovery ratio
[118, 119]. Indentation on single crystal CuAlNi SMA also showed
indentation-induced SME [120, 121].

The above research shows robust shape memory and superelastic effects under
contact loading conditions in NiTi and Cu—based SMA. This new phenomenon of shape
memory alloys could be used to engineer novel applications in surface engineering,
information storage and optics. It also prompts new questions about the mechanism of

the shape memory effect.

2.5.4 Wear and Scratch Damage

Wear is the removal of material from a solid surface as a result of mechanical action.
Wear can be divided into four categories: adhesive wear, abrasive wear, corrosive wear
and surface fiacture wear [122, 123]. In addition, there are several compound types of

wear: erosion, fretting, and cavitation erosion [122, 123].

42

Adhesive wear arises as a result of a process by which isolated spots on two sliding
surfaces adhere together momentarily. When shear occurs, the shear takes place at
some place other than the original interface. As a result, an adhering particle is formed,
which may come loose at a later stage. Adhesive wear is considered the purest and most
important form of wear because it is the only one which is always present. Unlike other
types of wear, adhesive wear can not be completely eliminated.

A quantitative expression of adhesive wear was derived by Archard [124-126] :

.4“?
3H’

V
where V is the wear volume formed after sliding through a distance x; P is the normal
load and H the hardness of the tested material. The non-dimensional constant It
represents the probability that a junction between two surfaces would lead to the
formation of a wear particle. It would be ideal but actually difficult to relate the wear
coefficient k to basic material properties (modulus and hardness) and surface properties
(surface roughness and lubrication conditions).

Abrasive wear is produced by a hard protuberance plowing a groove in a softer
surface so that the material in the groove is worn away. A similar expression to

Archard’s equation is also obtained:

_ tanepx
3H ’

 

where tan 0 is the tangent of the average effective roughness angle for the hard surface,

43

or for hard abrasive grains at an interface [127]. Values of fa—h—é under various
circumstances have been tabulated [127]. This equation can explain some wear
situations, for example, the reduction in wear rate when the abrasive hardness and the
tested material hardness are nearly equal; the reduction in wear rate when smaller
abrasive particles are used [127].

Corrosive wear is a form of wear where corrosion of a surface first occurs, then
sliding occurs and removes the corrosion products, allowing further corrosion.
Quantitative information on this form of wear is not readily available. Because the wear
rate is governed by so many factors: the rate of corrosion, the rate of removal of
corrosion products, the geometry of the sliding surfaces, and whether the whole of the
corrosion products are removed all at once or in part.

Surface fracture wear is common in rolling elements like ball and roller bearings,
wheels and rails, and gears. In this case, as the stresses come and go during continued
rolling service, cracks grow in or below the surface of the element, and eventually a
spalled particle is produced. This phenomenon is closely related to fatigue failure.

Wear tests are used to measure wear resistance. ASTM standards exist regarding
various test materials and experimental configurations [128]. The pin-on-disk or
ball-on-disk wear test is used most widely [129]. More details will be discussed in
Chapter 3.

The scratch test is a semi-quantitative method to evaluate a coating or thin films

strength and interfacial adhesion. Small-radius spherical indenters are used most widely

44

as the scratchers. The radius is between 100 and 200 um. Different scratch methods,
like progressively increasing load, constant load, and repeated scratch are available on
various testing platforms.

Coating failure occurs when the scratch load is increased to a certain value termed
the critical load. The failure can happen at the interface between coating and substrate,
or within the coating itself, or a combination of both. The critical load is the most
important numerical indication of interfacial adhesion from a scratch test. Although
scratch tests are used widely and are easily understood intuitively, and there are ongoing
efforts to establish a standard [13 0]. It remains difficult to interpret scratch test results
quantitatively and consistently. Description of the sudden failure of the coating is
beyond the established realm of continuum elasticity and plasticity.

Acoustic emission (AE) signals result either fiom the sudden release of elastic
energy, or from surface interactions such as friction and adhesion. Sudden energy
release can occur during unstable crack growth, high-speed phase transformations, or
plastic instabilities. A transducer to measure the acoustic emission is attached to the
scratch slider and its output can be correlated to events which occur along the scratch
track. At constant load, as the scratch is created a baseline AE response is established
due to plastic deformation and friction/adhesion. Individual AE responses fiom fiacture
are superimposed upon this signal. The magnitude of the acoustic emission signal
depends on the size of the crack produced since the energy in the AE signal scales with

the energy released in the process [131]. For coatings thicker than a micron, the area of

45

adhesive failure is large compared to the area of through-thickness cracks. Therefore, a
large jump in the acoustic signal is thus a reasonable indication of adhesive failure.

Besides establishing the critical load, microscopic observation of the failed region
also provides important information on the failure mechanism. Bull et al. proposed
several failure modes from the appearance of the failed hard coating, such as through-
thickness cracking, coating detachment, chipping, buckling, wedge spallation [132, 133].
The geometry and distribution of coating cracks also depends on the mechanical

properties of coating and substrate.

2.5.5 Tribological Study on Bulk NiTi Shape Memory Alloy

A number of researchers have investigated the erosion wear behavior of NiTi alloys
in different wear modes and compared NiTi alloys to conventional engineering materials
like steels, Ni-based and Co-based tribo-alloys [134-I44]. It is found that NiTi alloy has
a better erosion wear resistance than conventional wear-resistant materials, especially
during cavitation erosion processes [14], I 43]. In addition, NiTi alloy is resistant to
corrosive wear [145].

The wear resistance of conventional tribo-materials strongly depends on their
mechanical properties such as hardness, modulus, hardening [146, 147]. Each of the
mechanical properties contributes in greater or lesser degree to the wear resistance in

different wear modes. It has been suggested that the high wear resistance of NiTi alloy

46

can be attributable to superelasticity [142]. Richman et al. demonstrated the superior
cavitation resistance of NiTi alloy compared to a number of standard and wear-resistant
materials including cast martensitic stainless steels and Co-based alloys (Stellite) [141].
Shida and Sugimoto observed remarkable erosion resistance of NiTi alloys during a
water-jet erosion test [142]. Liang et al. observed that NiTi alloy showed high
resistance to wear under impact during a sand blasting test. They compared NiTi alloy
to a wear resistant alloy Ni-hard—4 (CrggNisngpFe) and showed that NiTi has a lower
erosion rate than the Ni-hard-4 alloy. They noticed that there was a strong correlation
between the wear resistance and the superelastic strain.

Bulk superelastic NiTi also shows excellent adhesive and abrasive wear resistance.
Li et al. carried out pin-on-disk and block-on-ring wear tests on superelastic NiTi and
found the adhesive wear resistance was superior to steels and other alloys [148-150].
But few researchers compared the wear behavior of superelastic materials at different
temperatures. Qian et al. studied the microscopic wear resistance of SE ( superelastic )
NiTi using a Berkovich indenter under uN load at various temperatures, which is an
abrasive wear condition [144]. The wear resistance of NiTi is significantly better in the
temperature range around A f. This temperature dependence of wear resistance was
shown in Figure 41. It is clear that at temperatures 50K above A f=289K the wear

performance deteriorated.

47

2.5.6 Hard Coatings and Interlayers

Hard coatings are widely used to reduce wear in tribo-contacts. Material loss by
abrasive wear in tribo-contacts is minimized by using a coating with hardness greater
than that of the counterface. Material transfer to the counterface by adhesive
interactions and the accompanying material loss may also be reduced. Lubricating the
contact, even if it is only boundary lubrication, serves the same purpose.

Chromium nitride (CrN) has a FCC NaCl lattice structure, and exhibits high hardness
(~20 GPa), low electrical resistivity (640 ufl-cm) and high melting point (1553K)
[151-158]. The outstanding properties of CrN have generated practical applications in
the protective coating, tribological, and superconductive disciplines. Reactive
magnetron sputtering combined with a negative substrate bias is the most common
method of making chromium nitride fihns. Argon gas is normally used as sputter gas
and N2 as the reactive gas. The application of nitride hard coatings on cutting and
forming tools has been successfully established over the past decade.

Hard coatings like CrN normally contain a high degree of intrinsic compressive
stress owing to lattice distortion and CTE mismatch between coating and substrate.
Residual compressive stress in the GPa range is not unusual. The residual stress
increases with coating thickness, and thick CrN coatings tend to suffer from premature
interfacial failure. It is difficult to produce single-layer nitride coating thicker than
6-7urn without encountering adhesion problems, which limits the expansion of CrN

applicability on into other areas.

48

Diamond-like carbon (DLC) is a metastable form of amorphous carbon containing a
significant fiaction of sp3 bonds. It has a high hardness (12-17 GPa), chemical inertrress,
optical transparency, and it is a wide band gap semiconductor [159]. DLC films have
widespread applications as protective coatings in areas including optical windows,
magnetic storage disks, car parts, biomedical coatings, and micro-electro-mechanical
devices. The most common industrial process for deposition of DLC is direct-current or
radio-frequency sputtering of a graphite electrode in an Ar plasma. Other methods like
ion beam assisted sputter deposition, pulsed laser deposition and cathodic arc deposition
are also used. DLC has a hardness around 12-17 GPa, slightly lower than that of CrN.
The high compressive stress in DLC also limits the maximum thickness of the coating.
DLC coatings are notable for their low fiiction coefficient compared with other hard
coatings [160-163]. The friction coefficient of DLC on steel can be below 0.1 without
lubrication, close to the fiiction coefficient of steel on steel with lubrication [164, 165].

Interlayers have long been used to boost interfacial adhesion, moderate intrinsic
stress and increase performance. For example, the problem of high residual stress in
electroplated nickel was overcome by introducing a Cu interlayer [166-168]. NiP
interlayers were used to increase corrosion resistance and tribological performance of
TiN and TiAlN coatings because the porosity of these PVD coatings was sealed off by
the NiP interlayer [169-1 72]. Interestingly, the TiAlN/NiP structure on steel drill bits
did not meet expectations because NiP can not withstand high impact forces during

service, which shows the importance of the mechanical properties of interlayers [169].

49

There is more discussion about interlayers between hard coatings and soft substrates in

section 2.6.

2.6 Tribological Study of Hard Coatings on Soft Substrates
2.6.1 Experimental Research

PVD hard coatings on soft lightweight alloy substrates are intended for potential
application in machine components in automobile, aerospace and other industries. Here,
the tribological system is closed, so that wear debris including coating fiagrnents
circulates within the system and can cause severe consequences. The control of wear
resistance and coating properties are more stringent for coated machine components. It
represents a wider application area and more complicated problem than traditional hard
coating:hard substrate system for tool materials.

The economic and engineering significance of hard coating:soft substrate tribological
systems is being recognized, leading to a rapidly expanding research literature on this
topic since 2000 [1 73-196]. A rich selection of coatings have been applied to aluminum
substrates by various methods: nitride coatings by PVD [181, 185, 191, 194, 196] and are
ion plating [1 74, I 76, 183-185, I 90], ferrous coating by thermal spray [182], ceramic
coating by micro-arc oxidation [185], plasma spray [192] and Mo-WC coating by laser
surface cladding [189]. The majority of research efforts are devoted to coating

processes and tribological characterization.

50

Experimental studies offer valuable information and solid support for the hard
coating: soft substrate tribological systems. Zhou et al. used are ion plating to deposit
TiN and CrN onto 6061 and 7075 aluminum alloy substrates [183, 184]. They
compared the tribological performances of hard coatings on these two alloy substrates
and suggested the substrates’ mechanical properties had an influence on the tribological
properties. They pointed out that the plastic deformation of substrate during wear tests
can cause hard coating fragmentation. Miki et al. depositied TiN on aluminum alloy
substrates by are ion plating (AIP) method and an ion beam mixing (IBM) method [190].
They focused on the measurement of residual stress in hard coatings using X-ray
diffraction. Different residual stress was obtained by altering substrate bias voltage
during the deposition process with residual stresses of 2.3 GPa developed at lower N2
pressures. However they found that wear resistance of nitride coatings was not sensitive
to the compressive stress. Kuruppu compared tribological performance of CrN and
Cr2N coatings on A2 tool steel, 52100 stainless steel, and 2024 aluminum substrates,
with and without Cr interlayers [194]. The coated alurrrinurn samples showed slightly
lower scratch resistance than hard substrates. Multilayered coatings on aluminum
substrates showed better scratch resistance than a single coating structure. The substrate
deformation was apparent in some wear tests, but the hard coatings did not wear or
delarninate significantly. Murakami and Zhou et al. also investigated a similar problem.
But they found single layer CrN on 2024 aluminum substrates has higher scratch

resistance than multilayer structures but suffered dramatic coating delamination during

51

wear tests [1 74, 183, 184].

Xia et al. studied diamond-like carbon coated 2024 aluminum alloy samples with
four different intermediate layers [180, 186]. Results showed that a thicker interlayer
improves tribological performance. Specifically, stable fiiction coefficients for these
four samples were nearly same and about 0.1, but the number of sliding cycles giving a
stable fiiction coefficient increased with increasing interlayer thickness. The wear rate
of the samples also decreased with increasing interlayer thickness. The same group also
investigated gradient layers of DLC/T iN/T i on nitrogen-implanted 2024 aluminum
substrates. The results indicate that the gradient layer shows a gradual change in
hardness, compact surface appearance and good tribological properties. Xu et al. used a
laser cladding method to coat 2024 aluminum alloys substrates with composite
Ti-Al-Fe-B coating [1 73]. The coated samples exhibited improved wear resistance
compared to bare aluminum samples. Cain et al. investigated TiN coated 6061
aluminum alloy samples for electrical contact resistance improvement [195]. These
were intended for electromagnetic shields with better corrosion and wear protection, and
lower contact resistance at mated contact joints, which is important for contacts which are
periodically operated in aggressive environments, such as the doors and access panels on
mobile communications facilities used by the military in the field. It was concluded that
TiN:Al contacts have significantly lower electrical contact resistance than comparable
chromate conversion coatings or bare aluminum alloy contacts, when cycled through

several thousand contact closures.

52

Some general points can be summarized from these studies on hard coating:soft
substrates systems.

(1) Hard coatings on aluminum substrates generally exhibit satisfactory interfacial
adhesion and scratch resistance, although in some cases poor interfacial adhesion
was regarded as the reason for coating delamination during wear tests.

(2) Hard coatings provide wear and scratch resistance to soft substrates as expected.
Thick coatings offer better tribological properties than thin coatings.

(3) Substrate deformation occurred during wear and scratch tests, which may contribute
to coating failure.

(4) An interlayer can provide better tribological properties, but it depends closely on the

process, interlayer properties and thickness.

2.6.2 Theoretical Research

Soft and compliant substrates pose a fundamentally different contact situation than
hard and stiff substrates. A hard and stiff substrate does require a hard coating for
abrasion resistance. The purpose of the coating is to guard against diffusive material
transport under sliding contact conditions. The soft and compliant substrates have a
lower load bearing capacity, which can partly be compensated in design by limiting
contact stresses. However it is necessary to compare the differences in contact stress

between coated hard substrates and soft substrates, to understand the benefits and

53

limitations of soft substrates. So far theoretical research on this problem lags in
comparison with numerous experimental studies.

Souza et al. used frrrite element analysis to model spherical indentation on TiN and
SiC coated 6061 aluminum substrates [1 79, 188]. They focused on circular crack
patterns in hard coatings on soft substrates generated during spherical indentation, and
compared modeling with experiments [188]. The number and location of circular
cracks was found to be related to several factors: overall contact geometry, the friction
coefficient between coating and indenter, the modulus mismatch between coating and
substrate, the level of residual stress and coating fiacture toughness. Another study
attempted to relate pile-up from spherical indentation in TiNzAl with plastic strain [1 79],
but without much success. The relation between pile-up and plastic strain is an
important topic in instrumented indentation. These could have been more informative
studies if the authors had shifted their attention to comparison of the contact stress
situations of hard coating on soft substrates and hard substrates.

Rarnalingam et al. used the displacement formulation method to compare the elastic
contact stresses at the interface between coatings and substrates [197, 198]. It is found
that stiff (high modulus) substrates and compliant (low modulus) substrates contribute to
a quite different stress condition. Different loading conditions are also considered:
normal point load (Figure 42), shear point load (Figure 43), normal strip load and shear
strip load. Point loading represents the concentrated contact loading where stress and

strain is localized, while strip load is the distributed contact loading over a length scale an

54

order of magnitude larger than the coating thickness. Results are summarized in the

following table, which are compared to the case with stiff (high modulus) substrates.

 

O Displacements are distributed in compliant substrates over
larger lateral dimensions than in stiff substrates. It is good
for compliant substrates because the strain gradient is
smaller.

0 The maximum normal bearing stress 0' y at

Normal line load coating:substrate interface is lower in compliant substrates;
(Figure 42 and the maximum interfacial shear stress Txy is also lower in
Figure 43) compliant substrates. It is good for compliant substrates

because it is beneficial for coating adhesion.

O The interfacial tensile stress 0', is higher in compliant
substrates, but can be lowered by increasing coating
thickness. It is good for compliant substrates because

increasing coating thickness could have a beneficial effect.

 

 

O Displacements are distributed in compliant substrates over

Shear Line Load
larger lateral dimensions than in stiff substrates. It is good

for compliant substrates.

 

 

 

55

 

O Interfacial shear stress 7,9. is lower in compliant
substrates; interfacial tensile stress 0'Jr is higher in
compliant substrates than in stiff substrates. It is good for
compliant substrates.

0 Stiff substrates have a higher positive 0' y , which tends to
lift the coating fiom substrates, while 0'), is moderated in
compliant substrates. It is good for compliant substrates

because it is beneficial for coating adhesion.

 

O 0' 3 becomes compressive instead of tensile, regardless of

Normal Strip Load substrates modulus. 0', is more compressive in compliant
(contact lengthle
times coating substrates than in stiff substrates. It is not good for
thickness)

compliant substrates.

 

 

0 Higher tensile stress 0' x at coating:substrate interface in
compliant substrates. It is not good for compliant substrates

Shear Strip Load

0 Lower lifting stress 0', in compliant substrates. It is good

for coating adhesion on compliant substrates.

 

 

 

Bilayer composite structures, which involve an interlayer between a stiff coating and

the substrate, have also been considered [197]. It was found that an interlayer reduced

the maximmn tensile stress 0'I up to 75%. Decreases in interfacial tensile stress are

56

accompanied by increases in shear and bearing stresses at the interface, but the rate of
increase is low. For shear traction, similar but smaller gains in reducing tensile stresses
are obtained with stiff interlayers. The lift-off (positive) 0' y is enhanced by
introduction of an interlayer. The authors suggested even a thin interlayer can be
helpful in lowering the maximum film tensile stress in heavily loaded conditions.
Tensile and shear stresses are affected more strongly than the bearing or lifting stresses.
The Ramalingarn study showed that there were both advantages and disadvantages of
compliant substrates as support structures compared with stiff substrates. But since only
elastic contact was considered, the plastic deformation on both soft and compliant
substrates encountered during experimental research still posed a major problem.
Introduction of an interlayer has the positive effect of reducing interfacial stress.

Experimental research has also pointed to this direction.

2.6.3 Preliminary Study of CrN:NiTi:Al Tribological System

Bulk SE NiTi has shown excellent wear resistance [134-144]. The wear and scratch
track in martensitic NiTi also showed recovery after austenite phase transformation [2].
This corresponds to the indentation-induced shape memory and superelastic effect
discussed in 2.5.3. However the high cost of NiTi alloy production and processing
would inhibit its tribological application. What is an economic and versatile way to

utilize the tribological advantages of NiTi shape memory alloy? The NiTi interlayer

57

approach may well be the right answer. Ni et al. studied the tribological performances
of CrN:NiTizAl tribological systems, where CrN is nitride hard coating, NiTi is the
interlayer, and 6061 aluminum alloy is the substrate [1-3]. It was found that a SE NiTi
interlayer greatly improves the wear resistance compared to SMA NiTi or Cr interlayers
(Figure 44). In scratch tests, the CrN coating shows no delamination with SE NiTi
interlayer, while severe damage in CrN coating without NiTi interlayer was obvious
Figure 45.

These studies provided no systematic information on how the evolution of
martensite-austenite phase transformations affects tribological performance. This
knowledge can only be obtained from wear and scratch tests carried out over the
complete transformation temperature range. Another key problem is how
coating:interlayer thickness ratio affects tribological performance. The residual stress of
CrN coating and NiTi interlayer interact with each other, and with contact stresses during
tribological loading. Residual stresses also tend to increase with coating and interlayer
thickness. Detailed knowledge of this problem can be obtained from tribological tests
over a matrix of coating:interlayer thickness ratios, and over a temperature range that

covers the phase transformation temperatures.

58

2.7 Constitutive Modeling of Shape Memory Effect
2.7.1 One dimensional Constitutive Model of SME

Computational modeling of the shape memory effect has been developed with
increasing applications of shape memory alloys. The macroscopic mechanical behavior
of shape memory alloys is usually modeled following either a phenomenological or a
micromechanical approach. One-dimensional phenomenological models are suitable for
engineering practice, because they make use of measurable quantities as parameters and
are often computationally efficient to implement. Uniaxial phenomenological models
have been proposed by many researchers: Tanaka and Nagaki [199], Liang and Rogers
[200], Ivshin and Pence [201-203], and Brinson and Huang [204], Boyd and Lagoudas
[205, 206]. The main factor distinguishing these one—dimensional models is the
particular kinetics governing phase transformations. Because the transformation strain
is large compared to the elastic strain, the phenomenological transformational kinetics
play a significant role.

The micromecharrics approach tries in various ways to follow very closely the
crystallography, using thermodynamics to describe the phase transformation [207-211].
These models often consider martensite variants as transforming inclusions and use
micromecharrics to calculate the interaction energy due to the phase transformation.
Stresses and strains are obtained as volume averages in which many possible martensite
variants are included. The micromechanical models are much more complicated than

phenomenological models and are computationally demanding. But based more closely

59

on physical metallurgy, these models may yield a more accurate constitutive law.
One-dimensional phenomenological models start with how to model the mixture of

martensite and austenite phases. Both parallel (Figure 46) and series (Figure 47)

approaches can be adopted [204]. In the parallel model, the relationship between stress

and strain is expressed as:

0':(1—§)O'A +‘f'0M
GM = M(8-§6)$0A =EA(5_§5)
0:(§°EM +(1—6).EAX8_;6)

In the series model, it is:

8:581” +(l-§)£A

 

£M=5+i,€A=£-
EM A
_ («e—:6)
_5_ ti
(EM+ EA)

where 8M , 0' M and 8A , a", are the strain and stress contribution from the martensite
phase and austenite phases, respectively. The 0' and e are the total stress and strain.
The 2,‘ is the martensite volume fiaction, and 6 is the phase transformation strain. E M and
EA are the modulus of martensite and austenite phases.

As pointed out by Brinson and Huang [204], the key difference between various
constitutive models of shape memory materials is the kinetic equations governing the
phase transformations. The kinetic equation specifies how the martensite volume
fraction changes with stress, temperature, and loading history. For instance, in the

Boyd-Lagoudas model [205, 206], the kinetic equation takes the form of exponential

60

frmction:

ézl-exp[aM(Ms —T)+bM0']

M ln('01) bM __q:
(M, —M,)’ CM '

where a

The three-dimensional model developed in the present work is based on the
one-dimensional Ivshin-Pence model [201, 202], whose kinetic equation is more

complex.

2.7.2 One-dimensional Ivshin-Pence Model

The Ivshin-Pence model is a one—dimensional phenomenological model of shape
memory effects. The transformation kinetics were carefully constructed to capture
various aspects of the shape memory effect.

The shape memory alloy is treated as homogenous mixture of martensite and
austenite phases. The total strain a is expressed as:

e=(1—a)eM +a8A,

where 8M and 8 A are the strain contributions from martensite phase and austenite
phase. The variable a is the austenite volume fraction, used in the Ivshin-Pence

model instead of the martensite volume fiaction f. The variable 8,, is purely elastic

but 8,, has both an elastic component (El) and transformational component (6):
M

6]

What distinguishes the Ivshin-Pence model from other 1D constitutive models is its

transformation kinetics that govern the hysteresis exhibited by the austenite phase
fraction a .

The rate form for the austenite volume fiaction is:

 

 

d_a_ a(tK) _(_6_fl_ 6T +_ 613 do do: —.<_ 0
d—t amax(,6(T(tK),o(tK))) 6?: 6T 6t an dt dt

1‘1: 1—a(tK) amm ____(_6_,B _6__T+_6_fl_ do da>0
dt l-amin(fl(T(tK ),o(tK))) ddfl 6T 6t 6o dt dt

where ,6 characterizes neutrality curves that is a function of stress and temperature but do
not involve hysteresis of phase transformations. The austenite volume fiaction a is an
explicit function of ,B. The variables amax and “min are envelope functions describing
the outermost hysteresis behavior for the martensite phase transformation and the reverse
austenite phase transformation, respectively. The variable t,{ is the previous time
when the transformation directions switch.

,8(T,o) is a function of temperature and stress which characterize the driving force

of phase transformations:

,6(T,o)= T+ 1 {(E“ -E,.)q2 ~50}.

(53 —s;’,) 215MB,

 

where S 3 and S j), a re the specific entropies of the austenite and martensite phases at

reference temperature, E M and E A are the Young’s modulus of the respective

martensite and austenite phase, and 6 is the martensite phase transformation strain.

62

The envelope functions can take various forms. The following form is used by the

original authors and also in this study:

a 3,, =O.5+0.5tanh(k,,6+k2)
a:mm =0.5+O.5tanh(k3,6+k4)'

The austenite volume fraction (1 can be integrated using amax and am)":

 

 

d__a_ a(tK) amam _(6__,6_ 6T +_6__,6 do da —.<_
dt am(,6(T(tK),0(tK))) dd: 6T 6t +60 dt ’dt

da_ 1—a(tK)___(6_fl_amin6T +_ 6,6 do)d _> 0
dt l—amin(,B(T(tK),o(tK))) dds ar 6t 6o dt dt

The k parameters are determined by the phase transformation temperatures:

amax(fl(M 30)) = amin(fl(As’O)) = W:
am(,6(MS,0)) = amin(,8(Af,0)) =1—t/l ’

The parameter W is an arbitrary small number because w = 0 is a numerical

singularity. In Ivshin-Pence model, {/1 = 0.02 is used [201, 202].

2.7.3 Three Dimensional Constitutive Models of SME

A number of three-dimensional constitutive models have been proposed based on
existing one-dirnensional models. For phenomenological models, the expansion fi'om
one to three dimensions is based on generalized plasticity theory, assuming the
martensite-austenite phase transformation also holds for the assumptions of classical
plasticity theory. These assumptions are: 1) yield stress is insensitive to hydrostatic

stress; 2) inelastic deformation due to the shape memory effect, like plastic deformation,

63

is volume conserving; 3) the plastic strain increment in each direction is proportional to
the stress deviator component in that direction.

Current three-dimensional constitutive models of SME do not have enough
experimental support from multidirectional loading [212-21 7]. Only one-dimensional
models are well calibrated with uniaxial experimental data that is more readily available.
The unique behavior of indentation-induced SME may thus be a suitable proving ground
for three dimensional models. Modeling of indentation-induced SME also offers better
understanding of complex SME under contact loading conditions.

Lubliner and Auricchio et al. proposed a framework of three dimensional
constitutive model of shape memory effect based on generalized plasticity [21 7].
Generalized plasticity is an intemal-variable model of rate-independent inelasticity that
includes conventional or classical plasticity as a special case. It was developed in order
to account for the behavior of elastic-plastic solids in which, following initial plastic
loading and elastic unloading, the reloading is not necessarily elastic up to the state at
which unloading began; such solids include graphite, some stainless steels and some
minerals. The model is based on some fundamental axioms and on results fi'om
elementary set theory and topology, which makes it also applicable to other mechanical
behavior like shape memory effect.

In the case of shape memory effect, the transformation strain rate tensor 8;- was

determined by:

. . J—

e;=,i—3—S)

2J7,”

where Si] was the deviator stress, J; the second invariant of stress, and A is the effective
transformation (inelastic) strain, whose rate form was determined by the kinetics of

martensite-austenite phase transformations.

2.8 Summary

This chapter has briefly reviewed recent work in various research areas involved in
this dissertation. The literature on shape memory alloys, shape memory and superelastic
effect, contact mechanics, and constitutive modeling of shape memory materials is very
extensive. It is difficult to cover these topics completely in this chapter. More
attention is given to the superelastic effect, the Clausius-Clapeyron equation, and the
Ivshin-Pence constitutive model due to their importance in this work.

Although there is also abundant literature on two-way shape memory effect, most of
them are focused on training methods instead of reaching deep into the mechanism of the
two-way shape memory effects. The few that focused on the mechanism and were well
recognized were reviewed. There presently exists no literature about two-way shape
memory effects from training methods involving indentation.

For hard coatings on soft and compliant substrates, more and more research papers

have been published in recent years, reflecting a growing interest in these tribological

65

systems. The majority are devoted to coating deposition and tribological tests. Very
few concern mechanics and modeling. Both experimental and theoretical work on this

tribological system has been reviewed.

66

Chapter 3

Methods and Materials

3.1 Tribological Study

3.1.1 Substrate Materials

Both M2 tool steel and 6061 T6 aluminum alloy. were used as substrate materials.
The M2 tool steel was supplied by Crucible Service Center (Syracuse, NY.) The 6061
T6 aluminum was supplied by United Aluminum Corp. (North Haven, C.T.)

Two mm thick M2 steel plates were cut into 25 mmx 25 mm square pieces by
electrical discharge machining. The steel substrates were polished to 50 nm surface
roughness with 0.25 pm grit size diamond paste. The surface roughness was measured
by optical surface profilometer. The samples were cleaned ultrasonically in acetone and
then in methanol for 10 minutes before being introduced into the deposition chamber.
3mm thick 6061 T6 aluminum alloy plates were cut into 30mm><30mm square pieces and
then mechanically polished to 50-70nm surface roughness using 0.05 pm grit size
alumina suspension. The polished aluminum substrates were then cleaned with acetone

and methanol.

 

‘ The T6 temper in the as-supplied material was later changed by processing.

67

3.1.2 NiTi Interlayer and Hard Coating Deposition

NiTi thin films were deposited by DC magnetron sputter deposition. The sputter
cathode compositions were Ni 46.5 Ti 53,5, Ni 43 Ti 52, and Ni 33,5 Ti 535 Cu 8 (atomic
percentage) for various experiments. The thin films compositions was Ni 433, Ti 51.7, Ni
51,3 Ti 43,2 and Ti 51.7 Ni 4;; Cu 6,] respectively, measured using a wet chemical analysis
(the materials were dissolved with nitric and hydrofluoric acid and measurements were
made by inductively coupled plasma optical emission microscopy by Wangyang Ni at
GM R&D Center [218]). The detailed deposition parameters are listed in Table 3
through Table 6.

For samples intended for later diamond-like carbon (DLC) deposition, a Cr layer of
0.3 pm thick was deposited onto NiTi interlayers in the same chamber, without breaking
vacuum. The purpose was to insulate the NiTi interlayer fi'om contamination.

The thicknesses of deposited films were measured by the spherical abrasion method
(CALOTEST, CSM Instruments SA)T, and Wyko NT surface profilometer at film edges
produced by masking. The surface roughness of the deposited thin films was measured
to be 35i10 nm with a Wyko NT surface profilometer.

The NiTi:Al samples were then transferred to a TEER UDP 550 unbalanced
magnetron sputter system for deposition of the hard coatings. Surfaces were sputter-

cleaned for 600 seconds before deposition of CrN, and 3.6 ks before deposition of DLC.

 

I A rotating steel sphere wears through the coating, which gives a rim and then into the substrate which gives a crate.
Coating thickness is determined from the diameters of the rim and the crate, measured optically. More details may be
found at http://www.csm-instruments.com

68

400 V was applied to the samples during sputter cleaning. CrN coatings were deposited
on both M2 steel and 6061 T6 aluminum substrates, with and without NiTi interlayers.
DLC coatings were only deposited on the aluminum substrates with and without NiTi
interlayers. A Cr layer of 0.1 pm thickness was deposited before the deposition of CrN
and DLC to facilitate bonding, which was indispensable for CrN and DLC coatings
deposition.

CrN coatings were deposited by sputtering two chromium targets of 99.99% purity in
a nitrogen- containing argon environment. Argon of 99.999% purity was used as the
sputtering gas. The base pressure of the system was 8><10‘4 Pa and the working partial
Ar pressure was 0.13 Pa. The partial pressure of nitrogen was 0.13 Pa. The deposition
power was 600W with 4 A at 150 V. The DLC coating was deposited by sputtering
from one chromium and two graphite targets. The base pressure was 1.33 X 10'4 Pa and
working Ar pressure was 0.13 Pa. The deposition power was 525 W with 3.5 A at 150
V.

NiTi interlayer and CrN layer thicknesses were measured at masked thin film edges
by a surface profilometer. CrN and DLC coatings were also deposited onto reference
aluminum and steel substrates without NiTi interlayers. Reference samples went
through one same thermal history so that the hardness of the substrates was the same as
the samples with a NiTi interlayer. The hardness of the as-received 6061 T6 aluminum
alloy was 94 HRE (62 HRB), which dropped to 21 HRE after deposition of NiTi

interlayers and CrN coatings.

69

3.1.3 Special Treatments of DLC:NiTi Interfaces

Special treatments were given to create a strong DLC:NiTi interface, because the
CrN:NiTi interface was found to be weaker than desired during scratch tests. A 0.3 pm
Cr layer was deposited onto NiTi interlayers after the deposition and annealing process,
without breaking vacuum. The Cr layers shields the NiTi interlayers from any potential
contamination. After the samples were set into the TEER deposition system, the Cr
layers were subjected to 3.6 ks sputter cleaning. It had been calibrated that the 3.6 ks
sputter cleaning would remove about 0.1 um Cr layer. As a result, about 0.2 urn thick
Cr layers still remain on the NiTi interlayers. Before the DLC deposition, another 0.1
um thick Cr layer was deposited onto the sputter cleaned Cr surfaces. Thus,.the DLC
coatings and NiTi interlayers connected at the Cr-Cr interface, while the CrN coating and

NiTi interlayers had a Cr-NiTi interface.

3.1.4 Characterization of NiTi Interlayer Properties

The phase transformation temperatures of NiTi interlayers on aluminum and steel
substrates were estimated using temperature scanning X-ray diffraction (XRD). For
room temperature martensitic NiTi interlayers, the measurement was carried out on a
Scintag 2000 diffractometer (Cu Ka radiation) using a thermoelectric module for heating.
The temperature of the NiTi interlayer was measured by a K-type thermocouple that was

adhered to the interlayer surface. An XRD scan was taken after each designated

70

temperature was stabilized. Room temperature austenitic NiTi interlayers needed a
special cooling stage to reach sub-ambient temperature. The thermoelectric cooler was
not capable of reaching steady state below 283K. Therefore additional measurements
were carried out on a Rigalcu AF C6S (Cu Ko. radiation) diffractometer with a low
temperature sample stage which was cooled by liquid nitrogen. Since the thin films
were attached to the substrates during the temperature controlled XRD measurements, the
measured transformation temperatures included the effects of residual stresses in the thin
films. Results of temperature scanning XRD shown in Figure 48 shows that room
temperature martensitic NiTi on steel substrates has A f350i5K and Ms (the martensite
start temperature) 3 10:t5 K. Figure 49 shows that room temperature austenitic NiTi has
Af310Kzt5K and Ms 270i5K. The results of NiTizAl samples are shown in Figure 50,
Figure 51 and Figure 52.

The phase transformation stress of NiTi thin fihns on aluminum substrates was
measured by the wafer curvature scanning method, which is based on the Stoney equation
[58, 60, 219]. The surface reflectivity of DLC and CrN hard coatings are not good
enough for laser scanning stress measurement, but good enough for optical profilometer.
The optical profilometer can also measure curvature over longer length, which is more
than 20mm. Longer length provides more accurate curvature measurement. However
the laser scanning stress measurement is capable of providing real time stress

measurement vs. changing temperature, from which residual stress evolution can be

71

obtained. The optical profilometer method can only take “snap-shot” of the curvature at

certain stable temperature. This first-order approximation is valid for to < ts / 100 .

The Stoney equation calculates the stress oc in the thin film:

 

2
oc = E"‘ K
6t ’

C

where EC is the biaxial modulus of the substrate, tc and ts are the thicknesses of thin film
and substrate, respectively. K is the curvature of thin fihn:substrate bilayer, which is
measured by the laser scanning method.

The aluminum substrates that supported NiTi thin films were plastically deformed
during thin film deposition and annealing by the intrinsic residual stress in the NiTi thin
fihn (section 4.1.1). As a result the total curvature measured was not caused entirely by
the residual stress in the NiTi thin film. Part of the curvature came from plastic
deformation of the aluminum substrates. However the Stoney equation is still valid for
measuring the relative residual stress change in the NiTi thin film. Let Kplasfic be the
curvature from plastic deformation; K1 and K2 are total curvatures measured at two

different temperatures; a; and of are the residual stresses in the NiTi thin films at

 

 

two different temperatures:
2
1 _ Ests _
ac — 61c (Kl Kplastic)
2
2 _ Ests
0'6 _ 6! (K2 —Kplastic)

C

. . 2 2 l
The change in resrdual stress A0] = O'c - O'c between two temperatures can be

72

calculated by Stoney equation:
2
1 Ests

6t

C

 

2_ 2 _
A01 —O-C— C—

(Kz‘Kl).

A resistance-heating furnace was used to heat the samples, which were placed on the
surface of the heater type thermocouple was inserted in a tight-fitting hole on the side of
the furnace. Because the thermocouple did not contact the samples, actual temperature
could have been lower than the measured temperatures.

The CrN hardness and elastic modulus were 20i2GPa and 200i20GPa, measured
from nine indentation tests at different locations on the coating; the DLC hardness and
modulus were 10:1:1GPa and 160i20GPa; both were measured by nanoindentation on
Nano Indenter XP (MTS Systems Corp.).

Temperature scanning indentations were also done on NiTi interlayers on steel
substrates to verify indentational superelasticity at elevated temperatures. The
indentation experiments were carried out on a Micro Materials (Wrexham, U.K.) Nano
Test 600 system with a Berkovich diamond indenter. The samples were immersed in
liquid nitrogen for 10 minutes before experiments to obtain a fillly martensite phase in the
interlayers. The indenter and substrate holder had resistance-heating ability. The
temperature at the indenter and sample surface was measured by two thermocouples and
controlled by independent temperature controllers. After the designated temperature
was reached the experimental set-up was left overnight to reach thermal equilibrium and

minimize thermal drift. The highest temperature used was 473K, far below the typical

73

aging temperature of 773 K for NiTi. As a result, the microstructure and mechanical
properties of thin fihns are not expected to change by the heating step. The maximum
load was set at 25 N so that the indentation depth would not exceed 10% of the film
thickness in order to minimize the substrate influence on the indentation response of the
films. The indentation speed was 5 nm/s, which was small enough to minimize any

dynamic effects. Load-depth curves are shown in Figure 53.

3.1.5 Tribological Testing Methods

Scratch and wear tests were used to evaluate tribological performance. The scratch
tests were performed on a Micro-Scratch Tester (CSM Instruments SA), which can also
record acoustic emissions from the scratch. A 107 um tip radius conical diamond
indenter was used as the scratcher. (Conical indenters actually have a small-radius
spherical tip.) If the indent or scratch dimension was small compared to the radius of
the spherical tip, the conical indenter functions as a spherical indenter. A thermoelectric
cooler heated the samples from below and the sample’s surface temperature was
measured by attaching a thermocouple.

The wear tests were performed on an ISC-200 Tribometer (Implant Sciences Corp.)
equipped with heating sample holder. The wear counterpart was a 200 um tip radius

diamond spherical indenter under 0.49 N and 0.98 N loads, and a 3.175mm radius WC

74

ball under 0.98N, 1N and 5N loads. The geometry of the wear tracks and scratch tracks
were characterized by Wyko NT surface profilometer and an Hitachi SC-2500C SEM.
The measurement of wear rate was made according to ASTM standard G 99

pin-on-disk wear test [220]:

Volume of the Wear Track
Total moving distance of the Ball

Wear Rate =

 

Volume of the Wear Track=(Cross-sectional area) X (Circumference)

Total moving distance of the Ball = (Number of Total circles) >< (Circumference)

3.2 Training and Characterization of The Two-way Shape Memory Effect

Both bulk NiTi and thin film NiTi SMAs were used in the two-way shape memory
study. The bulk NiTi metal was purchased from Special Metals Corp. (New Hartford,
New York). The nominal composition was N15031 Ti49,69. The phase transformation
temperatures were measured by differential scanning calorimetry using TA Instruments
modulated 2920 DSC. The DSC scanning curves are presented in Figure 32. The
phase transformation temperatures were: A s =350K, A f=404K; Ms =344K, and Mf
=287K.

The as-received material was electrical-discharge machined to small pieces with
dimensions of 2.45cmX2.45cm>< 1cm. Surface roughness was reduced to 0.05pm in
three steps of mechanical polishing using 3, 1, and 0.5 um grit size diamond paste.

The NiTi thin films were sputter-deposited on to a (100) silicon wafer at 673K in 45

75

minutes, followed by annealing at 773K for 3.6ks. The film thickness was measured to
be 4.5 pm. The chemical composition was Ni 43.3 Ti 5L7 by wet chemical analysis.
The phase transformation temperatures, A s =343K, A f=390K; MS =329K, and Mf=282K,
were estimated by the temperature scanning wafer curvature method
A simple “ball-and-clamp” method (Figure 54) was used to induce cyclic

thermomechanical training to the NiTi sample. The NiTi sample was first cooled in
liquid nitrogen for 5 min to ensure fully martensitic structure. A 3.175 mm diameter
tungsten carbide ball was then clamped into the sample using a steel c-clamp with a fixed
number of rotations of the clamping screw. The indentation depth was 170nm measured
by an optical surface profilometer. The representative strain was about 9%. The
whole fixture was then placed in a resistance-heating oven for 2 minutes to reach
423i10K. Afterwards, it was quenched into ice water for another 2 minutes, which
concluded one training cycle. A total of 30 training cycles were performed. Since the
number of rotations of the c-clamp screw is kept the same during the training procedure,
the thermomechanical training is considered under a constant strain condition.

The profile of indents was measured using Wyko NT optical surface profilometer.
A thermoelectric cooler (Marlow Industries Inc., Dallas, Texas) was placed below the
sample for heating and cooling. A thermocouple was taped onto the side of the sample
to measure one temperature. The profiles of the indents were taken after the sample was
heated to 400i5K and measured again after the sample was cooled to 300i3K. During

the process the position of the sample was kept the same.

76

Another training method, using a single-cycle deformation approach, was carried out
by the following procedures. A 3 X3 matrix of spherical indents was created on the bulk
sample surface using a 1.59 mm diameter steel ball indenter under a 100kg load. Also a
scratch was made with a 107m tip radius spherical indenter under 15N load, and
nanoindents were created on the NiTi thin films by a 10pm tip radius spherical diamond
indenter using a Nano XP indenter (MTS). Before indentation, the samples were cooled
in liquid nitrogen for 5min to ensure a fully martensitic structure. The samples were
then heated to 423K on a hot plate for 10min to let the spherical indents recover, i.e.,
transform from the martensite to austenite. They were cooled again in liquid nitrogen
for 5 min. The profiles of the indents and scratches were measured using a Wyko NT
optical surface profilometer (V eeco Instruments, Inc.). A thermoelectric device
(Marlow Industries, Inc.) was placed below the sample for heating and cooling. A
thermocouple was taped to the back side of the sample to measure temperature. The
profiles of the indents were taken after the sample was heated to 400i2K and taken again
after it was cooled to 300i2K.

The residual indents and scratches on the bulk NiTi were removed by mechanical
polishing: first using 240, 400 and 600 grit size silicon carbide grinding paper, and then
with 6 um and 1 um grit diamond paste. The indents on the surface of NiTi thin film
were polished using 0.05pm alumina suspensions. The profiles of surface protrusions
were then measured using the Wyko surface profilometer together with the

thermoelectric device for heating and cooling.

77

In the layered planarization experiment, the depth of successively polished layers
was measured by comparing of the thicknesses of the polished NiTi sample and a

reference sample, using optical surface profiler.

3.3 Finite Element Modeling of Spherical Indentation on Superelastic NiTi
Spherical indentation-induced superelasticity in NiTi was simulated in FEA and
compared to the results of experiments, including load-displacement curves and recovery
ratios. A three dimensional constitutive model is developed based on the
one-dimensional Ivshin-Pence constitutive model [202, 203] and generalized plasticity
theory [21 7], to describe the shape memory effect under contact loading conditions.
Details can be found in section 2.7. The von-Mises slip-plasticity is also included in
this model to capture the mechanical response of NiTi SMA in large deformation.
Implementation of the model in a commercial finite element program (ABAQUS) is done

by writing a user material subroutine.

3.3.1 Formulation of Three-dimensional Constitutive Model

78

First the 3D constitutive model is presented in terms of total strain rate Eij . This

total strain rate is decomposed into elastic part 85 and a transformation strain (inelastic)

T
part8ij I

. .T .5
56=EU+3U. (1)

Both the martensite phase (with volume fiaction f) and austenite phase (with volume

fraction a) contribute to the elastic response, with 5 + a = 1. The elastic strain rate is:

85: = Ell-7:1 old
Eli-'21 =(1-a)Ei},n-a + “Elly-n. (2)
where E 1:41, ij k I , 2,1,)“ are the respective stiffness tensor for martensite and austenite
phase, respectively.
For superelastic behavior, the transformational strain is dominant. The

simplification that E Maj/Cl = E Asijkl = Ey-k] can be made so that:

E _ -l
81-]. — Elmo/r1, (3)

1
The E5” is represented in terms of Young’s modulus E and Poisson’s ratio 0.
If loading a tensile specimen of superelastic NiTi to 6% total strain using 400MPa
tensile stress, it is assumed that 80% of the structure is stress induced martensitic phase.

Without the simplification, the elastic strain is:

79

400Wa 400Wa

35:20% +80%——=1.1%,
80GPa 3OGPa
With the simplification that E Mj-jk, = E Mk, = Egg, , the elastic strain is
5’3 = W :- 0. 5% .
SOGPa

There is 1.1%-O.5% = 0.6% difference in the elastic strain. This difference is 10% of the
6% total strain. So the simplification causes about 10% difference in the model’s
behavior.

Lubliner and Auricchio proposed a 3D frame work of the shape memory effect based
on the generalized plasticity theory [21 7]. The transformation strain rate is determined

by:

.T . J3
1 (2 172 J), (4)

where Si]- is the deviator stress, J 2 the second invariant of stress, and A the effective
transformation (inelastic) strain, whose rate form is determined by the austenite volume
fiaction rate:
i=—56. (3
Here a is again the austenite volume fraction, whereas 6 is a material constant
that gives the maximum transformation strain that can be induced by the shape memory
effect. This maximum transformation strain 6 follows fi-orn crystallographic

considerations and is typically about 5%. The austenite volume fraction a = a(x,t) is a

material dependant function of current von Mises stress and the overall loading history.

80

The a value is calculated using the Ivshin and Pence constitutive model. Combining

the elastic part and transformation part, the complete strain rate was expressed as:

. .fi _1.

ei'=—6a —S~ +E.. O'kl
J (2J2 l1) ,1” (6)

This constitutive relation between stress and strain is strictly in accordance with
classical von Mises plasticity theory. The assumptions and limitations of the classical
plasticity theory, like insensitivity to hydrostatic pressure and tension-compression
asymmetry, were also implied in the model. Experimental results showed some shape
memory materials had tension-compression asymmetry [77], and the phase
transformation behavior shifted with elevating hydrostatic pressure [221, 222].

Inclusion of these factors into the current model will be the topic of future studies.

3.3.2 Formulation of Slip Plasticity

Slip plasticity dominates the deformation of SE NiTi if stress-induced martensite
phase transformation is exhausted as the dominant deformation mechanism. Indentation
usually involved strains larger than 7 %. Plasticity was thus necessary to be included in
the constitutive model for indentation modeling. A simple but effective algorithm is
used in this constitutive model to shift the deformation mechanism from superelasticity to
von Mises plasticity where ever the austenite volume fiaction has decreased to 5 % (95 %

transformed from austenite to martensite). At this threshold superelasticity is replaced

81

by von Mises plasticity as the mechanism of plastic deformation. The strain rate is still
decomposed in the same way as equation (1) with the elastic part expressed as equation
(3) and inelastic part as (4). Equation (5) is replaced by an alternative expression so as
to model von Mises plasticity with isotropic hardening. The complete form of equation

(5) is expresses as:

r 1 J5 .

— S..o.--, a_<_0.05, von Mises lastici
hZJJ—Z ,1 J ( P W)

9 (7)
l _ 6a, (a > 0.05, superelasticity)

 

pe
ll

 

where h is the plastic hardening slope.
After the superelastic constitutive model is switched to the slip-plasticity model, a is
taken as fixed. The plastic strain continues to change according to equation (7) as long

as the plastic loading is in effect while .7 2 > 0. The von Mises stress and plastic strain is

increasing as plastic deformation proceeded.

Unloading occurs whenJ 2 < 0 . Initial unloading from the plastic deformation
stage is elastic until the von Mises stress returns to the value associated with the previous
transition from superelasticity to slip plasticity. Further unloading is governed by the
constitutive model of superelasticity where the austenite volume fraction increases
(a > 0 ).

The constitutive model is implemented in ABAQUS finite element program by
customizing the user material subroutine. It is first tested under simple loading

conditions of uniaxial tension, biaxial compression and shear. The stress-strain curves

82

from this test are presented in Figure 55. The constitutive model can simulate the
superelastic behavior in multi-direction loading directions. Plastic deformation in
uniaxial loading results in the residual strain, which came from residual martensites.
The materials properties used are representative of typical SE NiTi, but not for specific

samples, which were summarized in Table 8.

3.3.3 Finite Element Program Implementation

The finite element implementation is accomplished using the User Material Interface
provided by ABAQUS. The strain increment is passed to the user material subroutine
and the corresponding stress increment calculated from the constitutive relations using
the return mapping algorithm. The state variables, including austenite volume fiaction

and effective plastic strain, are also updated. The strain increment can be expressed as:

J5

day. = —6da(—S,-j) + Egfildou,

2w:

Multiplying both sides by elastic modulus gives the trial stress increment as:

73

l
2

The austenite volume fiaction increment is expressed as:

do
da- 6a eff do,j

6oefi. do,j

 

(9)

83

where 0' eff is the von Mises stress (effective stress). 6a / 50817 is determined by

the one dimensional Ivshin-Pence model. The rate form of austenite volume fiaction in

Ivshin-Pence model is expressed as:

 

 

 

 

 

 

 

 

 

d3: a(tK) do:max (6)6 6T + 6,3 610.5) da < 0
dt am(,B(T(tK ),o,fl(t,, ))) dfl 6T 6t 6o dt ’ dt ’
_c_l_a_ l—a(tK) mm(6,8 6T +66 610.17 fl) do >0
dt l-am (,6(T(tK),0',fl(tK))) (its 6T 6t 6o dt ’ dt ‘
In the modeling of superelasticity at a constant temperature, the above equation becomes:
6a = a(tK) docmax 6,8 da 50
6a ={ l—a(tK) }damm an da >0
50.1; l-am (fl(T(tr),a(tr))) dB 60.1; dt _

For isothermal phase transformation, the switching constant t), are the times of stress

reversal. In the 3D generalization considered here the switching constants become

times of von Mises stress reversal (equivalent to J 2 reversal). Insert equation (9) into

(8), the elastic trial stress is expressed as:

t _ 5680' :17” CE

After Cijkl is calculated, the real stress increment dog. can be obtained by

correction of the elastic trial stress increment d0;-

The indentation model was comprised of 15000 axisymmetric 4-node bilinear

elements. Part of the mesh around contact region is shown in Figure 56. The indenter

84

was modeled as rigid body sphere using the corresponding feature in ABAQUS to
increase computation efficiency. The fiiction coefficient between the indenter and the
substrate is assumed to be zero. A reference elastic-plastic material was simulated
under the same indentation conditions for comparison. Its elastic modulus and
Poisson’s ratio were the same as that of SE NiTi; its von Mises plasticity duplicated the

loading stage of SE NiTi. The reference material properties are summarized in Table 8.

3.3.4 Validation and Discussion

The absolute and normalized load-displacement curves of indentations are shown in
Figure 57 for three different indentation depths. The recovery ratio r] is defined as a
measure of the indent depth recovery:

h max — h residual
’7 — h ,

max

 

where hm is the maximum indentation depth and hmidual the residual indent depth after

unloading. The recovery ratio 1] is smaller at deeper indentation depths for both SE

NiTi and reference material. Experimental results has shown complete spherical indents

recovery (:7 = 100%) on SE NiTi [11 7], which were reproduce computationally here at
shallow indentation depth, by the modeled indentation-induced superelasticity. The
reference elastic-plastic materials’ recovery ratios are much smaller than that of the SE

NiTi at each indentation depths.

85

Tabor introduced the representative strain for spherical indentations:
a
8, = 0.2 — ,
R

where a is the contact radius and R the indenter radius [112], which is also reviewed in
2.5.2. Ni et al. found that the spherical indentation-induced SME has 100% recovery
when a, s 5% [116]. This is comparable to results of uniaxial loading of SE NiTi:
when uniaxial strain was smaller than 5%-7%, the recovery of deformation strain upon
unloading can be complete. The recovery ratio vs. representative strain from FEA
modeling is shown in Figure 58. However, the complete recovery of r] =100% is
obtained only at the lower representative strain of 1.5%. The cause of this discrepancy
may be due to the constitutive modeling. Another likely cause is the stress-strain
concentration at the contact edge, discussed below.

Figure 59 compares two computational load-displacement curves with experimental
results from Ni et a1 [11 7]. The load-displacement curves are from two modeled NiTi
materials with same properties but different phase transformation hysteresis A, where A =
(A f +A s)/2-(MS+Mfi/2. The modeled SE NiTi with A=30K is used only in Figure 59, to
show the effect of phase transformation hysteresis on load-displacement curves. It is
obvious that larger A also increases the hysteresis of indentation load-displacement
curves. The two computational curves match the experimental results quite closely.
The A=15K curve is more close to the experimental results in the loading stage while

A=30K curve is more close in the unloading stage. In contrast, the experimental SE

86

NiTi’s A is more than 70 K [11 7]. Better material calibration is needed to obtain a
closer match between experiments and the finite element model.

The distribution and size of the indentation-induced phase transformation region is
difficult to determine from experimental measurements, but is easily available fi'om finite
element modeling. Figure 60 depicts the indentation-induced phase transformation
regions at peak load and at the end of unloading. The indentation depth is shallow so
that the recovery ratio is 100 %. The contours of the transformed area consist of
concentric hemi-spherical rings with maximum value at 0.6 a (contact radius) and size of
1.8 a. This corresponds to the contours of von Mises stress by the one dimensional
constitutive relation. The shallow indentation does not exhaust the stress induced
martensite phase transformation, where maximum austenite volume fiaction is less than
60 %. Reverse transformation upon unloading lead to a perfect indent recovery and
there are very few residual martensites (Figure 58 (b)). In comparison, the reference
material has a recovery ratio of 56 %.

Figure 61 (a) shows that a deeper indentation resulted in a larger transformed area of
2.2a. The maximum transformed area is still located at 0.60 but inclined towards the
contact edge, because there is stress-strain concentration at contact edge at deeper
indentation depth. Although this concentration is not predicted by contact mechanics,
the situation is also present in the modeling using reference materials and in results from
other researchers [1 15]. The plastic strain contour in Figure 61(b) shows the largest

plastic strain initiated at the contact edge. The plastic deformation impedes the reverse

87

phase transformation during the unloading stage and the indent recovery is not complete

()7 = 71%). Contours of residual martensite expressed as austenite volume fraction is

shown in Figure 61(c), which has an intriguing irregular pattern.I More residual

martensite segregates around the contact edge and near surface area and fade into
austenite phase around 1.5a. The size of the indentation-induced transformed region is
also plotted in Figure 58. The transformed region is approximately two times the
contact radius, which increases slightly with indentation depth.

In summary, spherical indentation on SE NiTi is investigated using finite element
method, which incorporates a three dimensional constitutive model of the shape memory
effect with plasticity:

l) The computational load-displacements curves from spherical indentation modeling
capture indentation-induced superelasticity, which has exceptionally large indent
recovery ratios compared to elastic-plastic materials.

2) Martensite phase transformation is driven by contact stresses, whose hemispherical
contours have an outer limit of approximately two times the contact radius for indents
with perfect recovery, and increase in size at deeper indentation depths.

3) For shallow indentation depths, the stress-induced martensite phase transformation is
the deformation mechanism and no plasticity is involved. The indents recover

perfectly after unloading, leaving very few residual martensites. For deep

 

‘ The contours did not change with different mesh densities. Thus the pattern is not likely to come from numerical
errors in the FEA.

88

4)

indentations, plasticity develops first around contact edges. Plasticity impedes the
reverse austenite to martensite phase transformation and results in partial indent
recovery, where much more residual martensites are left after unloading.

The computational load-displacement curves match experimental results closely.
There is disparity between the phase transformation hysteresis of computational and
experimental SE NiTi. The perfect indent recovery predicted by the finite element
model occurs at the lower representative strains than experimental values. Closer
agreement between experiments and finite element models would need more accurate

input parameters from better material calibration.

89

Chapter 4

Results and Discussion

4.1 Tribological System with NiTi Interlayers

The tribological performance of hard-coatings, such as chromium nitride and
diamond like carbon, depends intimately on mechanical properties of the support
structure beneath coating, which was review in section 2.6. In this work, the hypothesis
of using superelastic NiTi interlayers to moderate contact stress and strain, and shield the
soft substrates from plastic deformation, was tested by experiments.

To investigate how the mechanical properties of NiTi interlayers affect the
tribological performance of the whole coating system, it is best that the tribological
performances can be compared with martensite and austenite (superelastic) NiTi
interlayers, within one sample. The temperature scanning methods complemented
previous wear test results at room temperature by showing that in one single sample the
tribological performance could change dramatically with the phase transformations of
NiTi interlayers.

The various CrN:NiTi:Al specimen forms prepared in this study are designated by the
following notation:

ACrN: B NiTiAFzAl ,

where the superscript A is coating thickness (um) and B is interlayer thickness (um).

The subscript “AF” is the austenite finish temperature A fof NiTi interlayer. For
instance, a sample with a 3.9 um thick CrN coating and a 15 um NiTi interlayer with A f
353 K on an aluminum substrate would be

3'9CrN:liNiTi353:Al.
A sample with a 2.5 pm thick DLC coating and a 21 um NiTi interlayer with A f368 K on
an aluminum substrate would be

Z'SDLC:21NiTi363:Al.

4.1.1 Residual and Phase Transformational Stress in NiTi Interlayers

The measurement of phase transformational stress is to prove that NiTi interlayers on
the aluminum substrates, even covered with hard coatings, can change the mechanical
behavior of the whole composite structure significantly, due to martensitic phase
transformations. It has been shown in temperature scanning X-ray diffraction (Figure
50 to Figure 52) that the NiTi thin fihns on aluminum substrates undergo thermoelastic
martensite phase transformation. Figure 62 to Figure 66 show that NiTi interlayers,
attached to the aluminum substrates and further covered by the hard coatings, could have
the stress-strain response similar to those from flee-standing thin fihns.

Figure 62 shows the stress evolution in NiTi thin films on aluminum substrates
continuously over three heating-cooling cycles, while Figure 63 shows the details of the

first heating-cooling cycle. Figure 64 shows the slope of the residual stress change vs.

91

temperature due to CTE mismatch. Figure 65 shows the residual stress change due to
phase transformation.

The residual stress in NiTi interlayers comes fi'om two origins: the intrinsic stress
from PVD process, and the extrinsic stress from coefficient of thermal expansion (CTE)
mismatch, which was discussed in section 2.3. 1. The aluminum substrates’ CTE is
larger than that of NiTi thin films. Thermal stress change is tensile during heating and
compressive during cooling. The CTE mismatch between NiTi thin films and silicon
substrates gives -1 .0 MPa/K thermal stress rate (derivation can be found in section 2.3.1).

The residual stress can be relaxed partly or completely by martensite transformation,
and restored by the reverse phase transformation. The transformational stress relaxation
rate is governed by the thermodynamics of martensite phase transformation, which is
described by the Clausius-Clapeyron equation in section 2.2.2. The transformational
stress relaxation and restoration cause a significant but reversible wafer curvature
change/displacement within a relatively small temperature hysteresis of about 80K.

The most significant problem encountered in the stress measurement is the plastic
deformation of the 6061 T6 aluminum substrates occurs during deposition or heat
treatment process where temperatures reached around 800K. This temperature is more
than 2/3 of the melting point of aluminum (933K). The intrinsic compressive stress in
NiTi thin fihns could plastically warp the aluminum substrate at this temperature to give
a large plastic curvature. Were this curvature elastic, it would require 2-3GPa residual

stress in the NiTi thin films, which is not possible for the NiTi thin films. It is difficult

92

to distinguish the plastic curvature and elastic curvature in the total curvature.

The residual stress shown fi'om Figure 62 to Figure 65 is compressive (negative).
The actual temperature could be lower than the measured temperatures because the
thermocouple was inserted in a tight-fitting hole on the side of the furnace, not contacting
the samples.

There are some fine details about the phase transformations revealed in the stress
evolution. First there are nonlinear stress relaxations before the phase transformations
in both heating and cooling directions. The stress relaxation flattens to the theoretical
1.01MPa/K curvature to a smaller or even negative values before actual phase
transformation starts.

The R phase transformation might also be present in the second cycle (red curves) in
the cooling direction (Figure 62). There was a stress relaxation about 20MPa before the
martensitic phase transformation that can be caused by R-phase transformation. But the
complicated precursor behaviors are not the focal point of the residual stress
measurements here.

Relaxation of the residual stress is also obvious in Figure 62 between three
consecutive heating-cooling curves, because the stress loops do not overlap each other.
The stress relaxation might also come from the creep of the aluminum substrate. But
generally the phase transformation behavior of NiTi thin films on aluminum substrate is
quite strong and stable.

Figure 66 shows the curvature of 2'SDLC:'2NiTi3(,3:Al sample at different

93

temperatures, measured by optical profilometer scanning, from which the phase
transformation stress in the NiTi interlayer can be measured. M; and A; temperatures
were reached because further cooling and heating about 10 K did not change the
curvature significantly.

To support tribological experimental data presented from 4.1.2 to 4.1.7, it is
necessary to consider the residual stress in the composite structure of hard
coating:interlayerzsubstrate. Because of the residual stress in DLC coating, the stress in
the NiTi interlayer, and their coupling, plus the plastic curvature in aluminum substrate,
it is very difficult to obtain full details of the real residual stress. One important piece
of information that can be obtained in Figure 66 is the residual stress change due to the
phase transformation of the NiTi interlayer, with the covering DLC coating. The phase
transformation relaxed 450MPa compressive stress in the NiTi interlayer when cooling
from austenitic phase to martensitic phase.

This is presented as proof that the NiTi interlayer has robust phase transformation

that can affect the mechanical response of the tribological system.

4.1.2 Wear Resistance of CrN Coatings on 6061 Al Substrates with NiTi
Interlayers
The wear resistance was measured in terms of wear loss at different temperatures for

samples with and without NiTi interlayers. The definition of wear loss is introduced in

94

2.5.4. Figure 67 shows the temperature scanning wear test results on
3'9CerlsNiTi353:Al and 3""CerAl samples using the 200m tip radius spherical indenter
under 0.49N and 0.98N loads. 3 '9CrN215N iTi353:Al had two orders of magnitude
improvement in wear resistance at temperatures above A ;, compared to 3""CerAl. For
3“S’CerlsNiTi353zAl, the superelastic (austenitic) NiTi interlayer gave two orders of
magnitude improvement in wear resistance as compared to the martensitic interlayer.
The improvement in wear resistance decreased as the temperature was decreased towards
room temperature. At room temperature with a martensitic interlayer, the wear
resistance of 3“S'CerisNiTi353zAl was slightly better (40% less for 0.49N load ) than that
3'9CrN2Al. The wear resistance of samples without a NiTi interlayer changed very little
over the temperature range.

The CerAl samples’ wear results had no temperature dependence while CrN:NiTizAl
samples had a strong temperature dependence. Since the only difference between
CrN:NiTi:Al and CerAl samples is the NiTi interlayer, the temperature dependence can
only be explained by the temperature dependant mechanical properties of the NiTi
interlayer. The improvement in wear resistance was thus apparently due to the
superelasticity of the NiTi interlayer at temperatures above A ;. The superelasticity of
the NiTi interlayer accommodated the contact strain, which was recovered after the wear
counterpart was moved. Much less plastic deformation was left with each passing of the
wear counterpart. The cyclic wear test amplified this effect. The wear loss of

3'S’CrN:'5NiTi:Al (Figure 67) with SE interlayer is 40 times smaller than the same sample

95

with a martensite interlayer.

SEM images of the wear tracks where are shown in Figure 68 (0.98N load) and
Figure 69 (0.49N load). The wear tracks from 0.49N shown in Figure 69 were similar
to that in Figure 68, only the extent of deformation was much less. The better wear
resistance is obvious in the less deformed wear tracks that were made above A ;. 0.98N
load gave especially deep wear tracks in 3'9CerAl in Figure 68, with part of the CrN
coatings delaminated. The wear track in 3'9CerAl also showed large spherical craters
lying across the wear tracks.

It is natural that improvement in wear behavior with superelastic interlayers under
0.49N load is more significant than that under 0.98N. Under 0.49N load, the
improvement in wear loss with SE NiTi interlayers compared to 3‘9CerAl samples was
higher than the case under 0.98N load. This is because lower load produced less
deformation, and the percentage of the deformation that could be recovered by the
superelastic response of the NiTi interlayer was larger, which means the superelasticity of
NiTi interlayer was more prominent.

During the wear test, the motion of spherical indenter pushed the soft substrate
material forward and piled the material ahead of the indenter by plastic deformation.
But with the increasing piling-up, resistance to the motion of the indenter also increased,
which eventually forced the indenter to “j ump” across the piled material to a flat area.
However wear tests on samples with superelastic NiTi interlayers did not show this

behavior because the superelastic interlayer significantly reduced the plastic deformation

96

of the aluminum substrate.

Wear tracks with SE NiTi interlayers in Figure 68 had a smooth and shallow
geometry, showing strong evidence of superior wear resistance. Wear tracks produced
at lower temperature with martensitic NiTi interlayer had a resemblance to the wear
tracks in CrN:Al, also showing slight delamination of CrN and substrate deformation. It
is quite remarkable that the wear resistance changed so dramatically in the same sample
when temperature was changed less than 100K.

The temperature scanning wear tests were repeated with a NiTiCu interlayer. The
ternary NiTiCu interlayer has a narrower hysteresis than NiTi interlayer, which can be
seen in the X-ray difiraction data (Figure 51). Tribological behavior can be potentially
by influenced by the hysteresis of superelastic interlayers.

Figure 70 shows the SEM images of the wear tracks at different temperatures. The
wear track produced with the presence of superelastic NiTiCu interlayer was narrower
and shallower than wear track produced with a martensitic interlayer. Figure 71 shows
the wear data on l‘llCergNiTiCu3432Al samples. At room temperature, martensitic
NiTiCu interlayer gave around 2 times improvement in wear resistance compared to the
sample without interlayer, while 3'QCerlsNiTi:Al had only about 2 times improvement to
compared to 3'9CrN:Al at room temperature. This may be because the
l'7CrN:9NiTiCu343:Al had larger ratio of interlayer to hard coating thickness (5.3) than
3'9Cer'5NiTi353zAl (3.8).

At temperature above A ;, the wear resistance of l‘7CrN:9NiTiCu343:Al is about 10

97

times higher than that the sample without interlayer. But the improvement in wear
resistance with a superelastic interlayer was only 5 times compared to a martensitic
interlayer. This number was smaller than the 80 times improvement in
3‘S’CerlsNiTi353zAl samples when the NiTi interlayer transformed from martensite to
austenite. This can only be explained by the difference in the mechanical properties
between NiTi and NiTiCu interlayers. The superelasticity of the NiTiCu interlayer may
be worse than that of the NiTi interlayer, which means the NiTiCu interlayer had smaller
recoverable strain. The smaller recoverable strain limited the improvement in wear
resistance. It would be best if there are tensile test results of both NiTiCu and NiTi thin
films so that the differences in mechanical properties could be compared.

The SE NiTi interlayer thickness affects the wear resistance of CrN coating
significantly. Figure 72 shows at room temperature how CrN and NiTi interlayer
thickness affect wear resistance. The wear test was done using 6.35mm diameter WC
under 5N load. For aluminum substrates, thicker CrN and NiTi interlayers improved the
wear resistance. It is quite easy to understand that thicker CrN coating has better wear
resistance, because thicker CrN coating bears more contact stress due to its high hardness,
providing the increased residual stress associated with the thicker coating does not
compromise interfacial adhesion.

CrN coating’s intrinsic compressive stress scales with thickness [223-225]. Thick
CrN coatings on steel substrates usually suffer performance degradation for the large

intrinsic stress weakens the interface between CrN coating and substrate. This situation

98

is proved in the tribological tests of CrN:NiTi:M2, which is discussed in the next
paragmph.

The martensitic NiTi interlayer also increases the wear resistance, compared to
samples without interlayers. There is 40 times improvement between
l'7CrN:'5NiTi353:Al and l'7CerAl at room temperature. Where did this improvement
come from? In scratch test results discussed in 4.1.6, it is obvious that the interfacial
adhesion between CrN and NiTi at room temperature is worse than adhesion between
CrN and Al. The improvement in wear resistance can not come fiom the adhesion of
CrN coatings. It is more likely that the improvement in wear resistance comes from the
better mechanical properties of martensitic NiTi interlayers compared to the soft
aluminum substrates. The fatigue resistance of martensitic NiTi interlayer may be
superior to those of the aluminum substrates. The martensitic NiTi interlayer may be
able to accommodate large strain by aligning martensite variants and twins without
introducing significant defects. The yield stress of NiTi interlayer (~200 MPa) [51, 52]

is close to that of the fully annealed aluminum substrate (HRE 62, ~180MPa).

4.1.3 Contact Stress Induced Phase Transformation in NiTi Interlayer
The experimental results discussed in the previous sections clearly show that contact
stress induced phase transformation in NiTi interlayer has a strong influence on

tribological performance. Here two topics will be discussed about the contact stress

99

induced phase transformation in NiTi interlayer that will provide some interesting
information. The first is about contact loading conditions that trigger the phase
transformation. The second is about finite element modeling of spherical indentation on
CrN:NiTizAl.

Similar to uniaxial loading of SE NiTi, whether there is superelastic effect in NiTi
interlayer is also dependant on the contact load conditions. If the contact stress is too
small, the NiTi interlayer behaves like normal elastic material, and there is no difference
in wear resistance between martensite and austenitic NiTi interlayer. Figure 73 shows
the temperature scanning wear tests on l'7CrN:15NiTi353:Al and ”CerAl samples using a
3.175mm radius WC ball under 0.98N. There is no improvement of wear resistance
when the NiTi interlayer is superelastic at temperature above A ;. The reason is that
under same 0.98N load, 3.175mm radius WC ball gave much less contact stress than the
200m radius diamond indenter used in section 4.1.2. The contact stress was not large
enough to initiate the superelastic effect in NiTi interlayer. In other words, the contact
strain can be accommodated by the elastic deformation of CrN coating and NiTi
interlayer and superelasticity is irrelevant.

In Figure 73, the martensitic NiTi interlayer improves the wear resistance of the CrN
coating by nearly 100 times. This improvement is higher than the 40 times
improvement in l'7Cerl5NiTi353zAl and l'7CrN: Al discussed in section 4.1.2 with Figure
72 because the load used here was decreased to 0.98N instead of 5N, but the wear

counterpart remained a 6.35mm WC ball. It has been discussed that thicker martensitic

100

NiTi interlayers resulted in better wear resistance. So under lower wear load, the
thicknesses of NiTi interlayers were virtually increased because the contact stress
contours are geometrically similar (lower load gave smaller stress contours).

Figure 74 plots the elastic contact stress in elastic half space under 3.175mm radius
sphere and 200m radius sphere as a function of load using the equation introduced in

2.5.1:

3P [6PE *2 )“3

 

p0 _ 27m2 _ n3 R2

1 1 1 ’
_*. = +
E E indenter E substrate

where P is normal load, a is contact radius, R is the indenter radius, E. is composite
modulus. The 3.175mm radius sphere under 0.49N has 558MPa elastic contact stress.
This number is further attenuated by the CrN coating, and its intrinsic compressive
residual stress, which in the end is not strong enough to trigger the superelastic response
of NiTi interlayer. The elastic deformation of the CrN:NiTizAl composite structure
would be enough to support the contact stress around 550 MPa.

But on the other hand, the 200m radius sphere gave a very high elastic contact stress
(>2000MPa), which was impossible to support elastically. It has to be mediated by the
deformation in the coating:interlayerzsubstrate system. The superelastic NiTi interlayer
can accommodate the deformation superelastically while a martensitic NiTi layer and the
aluminum substrate cannot, which is why the wear resistance improvement is present

only under higher contact stress/strain loading conditions.

101

4.1.4 Wear Resistance of DLC Coatings on Soft Substrates with NiTi Interlayers

Temperature scanning wear tests were performed on the DLC:NiTizAl and DLCzAl
samples. The DLC (diamond-like-carbon) coating has different mechanical properties
and residual stress compared to the CrN coatings. How these differences change the
tribological behaviors with the presence of NiTi interlayers would provide valuable
information on this complex problem, along with more convincing evidence that SE NiTi
interlayer could enhance tribological performance.

First the DLC:NiTizAl samples also showed improvement in wear resistance with the
superelastic NiTi interlayers at elevated temperature. Figure 75 presents the wear
resistance of2'5DLCz9NiTi363:Al, 2-5I)Lc:‘5NiTi363:A1, 2‘5DLC:2‘NiTi363:Al and
2'SDLCzAl samples at different temperatures. The NiTi interlayers have 9pm, 15m,

21 um thicknesses. The SEM images of the wear tracks are compared in Figure 76.

The 21 um SE NiTi interlayer improved the wear resistance 340 times compared to
2'SDLC:AI samples at temperature above A ;. The 15pm SE NiTi interlayer improved the
wear resistance about 60 times compared to 2'SDLCzAl. This improvement is similar to
the 3'9CerlsNiTi353zAl with 15 um SE NiTi interlayer, which had a 80 times
improvement compared to 3'9CerAl. 2'SDLCz9NiTi363:Al sample only improved wear
resistance about 4 times compared to 2"(’DLC:A1 sample.

The thicker NiTi interlayers greatly increased wear resistance. The
Z'SDLC29NITI3682A1 sample with a thin 9pm SE NiTi interlayer did not improve the wear

resistance significantly, while 21 um SE NiTi interlayer produced a much better

102

improvement. Thicker NiTi interlayer (21 um) even improved the wear resistance of
DLC coating at room temperature in martensite phase quite significantly (10 times),
compared to the DLC:Al sample. The improvement in wear resistance as a function of
interlayer thickness is not linear, but logarithmic.

The reason of why thicker interlayers are better lies in the fact that contact stress and
strain decays with depth. The basics of contact mechanics are reviewed in section 2.5.1
and 2.5.2. The contact stress and strain fields are hemispherical contours, the intensities
of which decay with depth. For a given coating thickness, a thicker interlayer means
larger portion of the contact strain is supported by the coating and interlayer, while the
substrate is subjected to smaller stress and strain. Smaller contact stress and strain in
aluminum substrates gave smaller plastic deformation.

It is true that thicker NiTi interlayer costs more in deposition time and target material
but the increase in cost is linear, which justifies the benefits of thicker NiTi interlayer.

The DLC coating is known to have higher compressive stress than CrN coating but
lower modulus and hardness. But the low friction coefficient of DLC coating makes it
favorable in many applications. However no significant change in fiiction coefficient
was observed with superelastic NiTi interlayer and martensitic NiTi interlayer in this

work.

103

4.1.5 Wear Resistance of CrN Coatings on Hard Substrates with NiTi Interlayers

The wear results from CrN:NiTi:M2 system are quite different fiom that fi'om
CrN:NiTizAl system, which vividly illustrates the importance of the mechanical
properties of supporting substrates.

Figure 77 shows that the addition of a NiTi interlayer affects adversely on the wear
resistance of CrN hard coatings. The increase in wear loss was biggest for 1.7 pm thin
CrN coating and thick 12 um NiTi interlayer, showing that large interlayer/coating
thickness ratio is not suitable for hard substrates. On the other hand, small
interlayer/coating thickness ratio for 6 pm NiTi interlayer and 3.9 um CrN coating only
caused moderate increase in wear loss compared to 3.9 um CrN coating without an
interlayer. The best result is seen in a thick coating/interlayer combination: 3.9 um CrN
coating and 12 um NiTi interlayer, which is slightly better than the case without the NiTi
interlayer.

Figure 78 presents the same data in Figure 77 but as a function of NiTi interlayer
thickness. Although a thicker NiTi interlayer improves the wear resistance of samples
with 2.8 um and 3.9 pm CrN coatings, the wear resistance is still not improved compared
with cases without the NiTi interlayers. For samples without interlayers, the wear
resistance was highest for samples with the intermediate 2.8um CrN coating.

The wear results on CrN:NiTi:M2 systems indicate that the NiTi interlayer may not
be suitable for improving the wear resistance of hard coating/hard substrate combinations.

The modulus and yield stress of NiTi interlayers are substantially lower than the CrN and

104

M2 tool steel. As a result, strong interfacial acflresion must be realized for a relatively
weak interlayer between hard coating and hard substrate to function as an adhesive layer.
However, scratch test results in section 4.1.6 shows the evidence of a weak interface
between CrN and martensitic NiTi. Although the substrate is aluminum in that case, the
CrN and NiTi interface is still the same here on steel substrates. The weak interface
may very well be the reason for the worse wear resistance of CrN:NiTi:M2 compared to
CerM2.

The interfacial adhesion problem can be solved, as shown in the DLC:NiTizAl
system discussed in section 4.1.7. It may yet be possible to improve the wear resistance
of CrN:NiTi:M2 by promoting the interfacial adhesion between CrN and NiTi in a future

study.

4.1.6 Interfacial Adhesion between CrN and NiTi

Figure 79 shows temperature scanning scratch tracks on a 3"”Cer9NiTi353zAl sample.
There is a significant temperature dependence of the scratch resistance. At temperature
above A ;there was minimum coating delamination and failure but there is very
significant coating delamination and failure at temperature around MS in
3'9Cer9NiTi353zAl. Figure 80 shows the scratch tracks of 3'9CerAl samples that were
made at the same temperature. There is no change in scratch resistance with changing

temperature. This comparison proves that the temperature dependence of the scratch

105

resistance of 3'S’CergNiTi353zAl sample only comes fi‘om the NiTi interlayer, which has
temperature-dependant mechanical properties. It is direct evidence that the
stress-induced superelasticity of NiTi interlayer can improve scratch resistance and lead
to better tribological performance.

The scratch resistance can be quantified by a “critical load”, which is the normal
scratch load where interfacial failure occurs. The critical load is determined after
scrutiny of various experimental data: the appearance of scratch tracks, changes in
friction coefficient, and acoustic emission signals. In scratch tests on 3'9Cer9NiTi353zAl
samples, the acoustic emission signals, shown in Figure 81 and Figure 82, are good
indication of the critical coating failure because of the sudden jump in acoustic emission
when dramatic coating delamination happened. The critical coating failure load
determined from acoustic emission can also be verified by inspecting SEM images of the
scratch tracks (Figure 83 and Figure 84). The critical load determined fi'om acoustic
emission and SEM images agreed each other well. It is shown in Figure 85 that
3'9CergNiTi353zAl has improved scratch resistance above A ;. The critical load jumped
from 1.8N at room temperatures around MS to 4.2N at temperatures above A ;. For
3'9CerAl, the scratch resistance remains relatively stable across the temperature range.

However it is noticeable in Figure 79 that 3'9CrN:9NiTi353:Al sample suffered more
severe coating delamination than 3'9CerAl sample at temperature below A ;. The EDS
obtained from delarninated region showed that the failure occurred at the CrN:NiTi

interface (Figure 83) and CrN:Al interface (Figure 84), which suggested that the

106

CrN:NiTi interface was weaker than CerAl interface. Considering CrN and NiTi
deposition took place at two different vacuum systems, the exposure of NiTi surface to
air before deposition of CrN may contribute to the poorer interfacial adhesion. However
the problem of insufficient adhesion is improved in subsequent DLC:NiTizAl

experiments.

4.1.7 Interfacial Adhesion between DLC:NiTi

The interfacial adhesion between DLC and NiTi was improved by two methods:
interface chemistry and a sputter cleaning process. After deposition of NiTi and heat
treatment, a lum pure Cr layer was deposited onto the NiTi surface without breaking
vacuum. The NiTizAl samples were then transferred to the TEER deposition system for
DLC deposition. Aggressive sputter cleaning (3600 ks) was employed to remove about
0.5 um of the Cr. Another thin Cr layer (0.1um) was then deposited onto the sputtered
Cr surface before DLC deposition. The potentially weak interface between DLC and
NiTi was replaced by cleaned and thin Cr surface. The reference samples without NiTi
interlayer also went through this preparation procedures.

Scratch tests were used again to measure the interfacial adhesion of
Z'SDLCIISN1T13631AI and 2'SDLCzAl samples. The phase transformation temperatures of
the NiTi interlayer were determined by temperature scanning X-ray diffraction (Figure

52). Figure 86 shows the scratch track on the DLCzAl sample produced by 107nm tip

107

radius spherical indenter and 7N linear progressive load, at room temperature. Figure
87 shows the scratch tracks on “DLC: ”NiTiggngl produced under the same conditions.
There is very severe DLC delamination and failure in the 2'SDLC:Al sample (Figure 86),
while no significant coating fracture occurs in 2‘SDLC:”NiTi3(58:Al (Figure 87) except at
the edges of the scratch track. It is evident the interfacial adhesion between DLC and
NiTi is much better than that between DLC and Al, even at room temperature with a
martensitic NiTi interlayer. The delamination and dramatic failure of DLC coatings on
aluminum substrates without NiTi interlayer at 7N scratch load made it unnecessary to
conduct the scratch test at higher load.

Figure 88 shows the scratch track on 2'SDLC:‘5NiTi363:Al produced by the very high
30N scratch load. This load is more than 4 times higher than the case in Figure 86 and
Figure 87. Under the high scratch load, the DLC coating failure is inevitable, which is
manifest as fine cracks inside the scratch track. However the strong interface adhesion
kept the fractured DLC coating attached to the NiTi interlayer. Otherwise the fractured
coating would have the fragmented appearance in Figure 86 and Figure 79. The
delamination interface was determined by EDS spectrum shown in Figure 89.

How does interfacial adhesion affect the appearance of the scratch tracks? Consider
the scratch of 3'9CrN29NiTi353:Al in Figure 79. The destructive nature of scratch test
inevitably leads to some damage to the coating, CrN or DLC. A strong interface keeps
the coating adhered to the interlayer or substrate, so that the fractures, which

accommodate the contact strains, are uniformly spaced over the scratch track. This is

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the case in the scratch tracks at temperature above A; in Figure 79, also in
2'5DLC2'5NiTi3éngl (Figure 87 and Figure 88). If the interface is weak, the coating
would be fiactured and then delaminated, damaged by the moving scratch head, resulting
in the highly damaged look in scratch tracks at temperatures around Ms as seen in
3‘S'CergNiTi353zAl (Figure 79), and also in the 2‘SDLC:A1 sample (Figure 86). Whether
there are dislodged coating fragments are good indication of whether there is good
interfacial adhesion.

The temperature scanning scratch test in the 2""DLC:'5NiTi;;(53:Al sample did not show
significant temperature dependence, which is different than the temperature scanning test
on 3‘9Cer9NiTi353:Al sample. The reason is that the interface between DLC and NiTi is
intrinsically so strong that the influence of SE NiTi interlayer on the interfacial adhesion
is greatly reduced. Due to the good adhesion, scratch at temperature around Ms did not
cause delamination of DLC coating; the enhanced adhesion with SE NiTi interlayer at
temperatures above A; still gives similar scratch tracks because fiacture in the DLC is
something SE NiTi interlayer can not prevent. So under the improved adhesion between
DLC and NiTi, there is no significant temperature dependence in the scratch tracks in

25DLC:"NiTi363:Al sample.

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4.1.8 Finite Element Modeling of Indentation on CrN:NiTi:Al and CrN:Al

Single-cycle spherical indentation on 3'S’CerlsNiTizAl and 3'9CrN:Al samples is
simulated in FEA. Although the indentation loading is simpler than the cyclic sliding
loading in the actual ball-on-disk wear tests, spherical indentation captures the main
features of the contact loading conditions. Figure 90 shows the load—displacement curve
from the indentation simulation. The maximum indentation depth in 39(3er lsNiTizAl
and 3'9CrN:Al was 2.1um and 2.5um, respectively, because the SE NiTi interlayer had
higher yield stress than the fully annealed aluminum substrate. After unloading, the
3'9CrN:'5NiTi:Al had only 1 pm residual depth while the 3'9CerAl had 1.9pm residual
depth. The shallower residual depth in 3“S’Cerfi’NiTizAl probably results fiom a
combination of strain-recovery in the SE NiTi interlayer and a decrease in the amount of
plastic deformation occurring in the aluminum substrate. The result is generally
consistent with the experimental result showing that the 3"’lCerlsNiTizAl provided better
wear resistance.

Figure 91 and Figure 92 show the strain on and eyy calculated by FEA of
3"’CerlsNiTizAl and 3'9CrN:Al, in the loaded and unloaded state. The contours for the
loaded state of 3'9CerlsNiTizAl showed that the highest strain was at the NiTizAl
interface, but without much strain mismatch across the interface; the strains in
CrN:NiTi:Al were smaller compared to strains in 3'9CerAl. After unloading, the strain

contours in CerA1 were similar to that at the loaded state.

110

Figure 93 shows the shear strain (exy) contours in 3'9CerlsNiTi:Al and 3‘9CrN:Al.
A high shear strain of about 6.5% appears at the position under the edge of contact, which
indicates that the material was being pushed outside of the contact zone. The hard CrN
coating was not affected by this shearing flow due to its high hardness. As a result, the
high shear strain gradient was located beneath the hard coating interface. In 3'S’CerAl,
the majority of the shear strain was plastic deformation; in 3'9CrN:'5NiTi:Al, the NiTi
interlayer accommodated most of the shearing strain superelastically. There were also a
shear strain gradient across NiTizAl interface, which was caused by different material
properties across the interface. Overall, the shearing strains in aluminum substrate were
less in 3'9CrN:‘5NiTi:Al both in the loaded and unloaded state, as compared to ”CerAl.

Figure 94 depicts the martensite volume fraction from stress induced phase
transformation in loaded and unloaded states. There was a highly transformed area
around the contact edge due to the stress concentration there. The lower center near the
interface between NiTi and Al also showed more stress-induced martensite, which
accommodated the strains from the vertical displacement of the spherical indenter.

Figure 95, Figure 96 and Figure 97 show the stress components in 3'9CerlsNiTi:Al
and ”CNN. The highest stress always was located within the CrN coating, although
in the terms of strain, the CrN coating always had the smallest amount of deformation.
This shows that the hard coating on softer substrate can shield the substrate from high
contact stress, but not from deformation. On the other hand, the differences in stress

components in 3'9CrN:15NiTi:Al and 3""CerAl are not as significant as that in strains.

111

This means that the addition of NiTi interlayer can shield the softer substrate from large
plastic deformation. With the hard coating supporting contact stress and the SE NiTi
accommodating contact strain, the finite element analysis proves that the CrN:NiTi is a

useful combination for the protection of softer substrates.

4.2 Indentation-induced Two-way Shape Memory Effect

The Two-Way Shape Memory Effect (TWSME) has been achieved in NiTi and other
shape memory alloys under uniaxial and bending loading, which has been reviewed in
section 2.4. In this work TWSME under contact loading conditions is demonstrated for
the first time by two different training methods.

Textured surfaces can have different tribological properties, for instance, lower
fiiction coefficient, compared to smooth surfaces [226-230]. Surfaces with reversible
texture that is controlled by temperature can then have variable friction coefficient and
other controllable tribological and optical properties. Such surfaces can be realized
using indentation-induced TWSME.

The experimental demonstrations of indentation-induced TWSME are presented in
sections 4.2.1 and 4.2.2. Before the discussion of the mechanisms of
indentation-induced TWSME in 4.2.5, two experiments are introduced and discussed in

4.2.3 and 4.2.4, which reveal some interesting aspects of this complicated problem.

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4.2.1 Indentation-induced TWSME from Thermomechanical Training

Thermomechanical training of shape memory alloy is accomplished by cycling
between temperatures above A; and below M; when a specimen is subjected to external
stress or constraint, which has been reviewed in section 2.4.2. In this study a complex
contact loading condition was used instead of simple axial loading conditions. However
the experimental setup was simple. The details are described in section 3.2. Briefly, a
3.175mm diameter tungsten carbide sphere was pressed into a martensitic NiTi alloy
sample to a certain depth and then the sample was cycled between temperatures below
M S and above A ;.

After the treatment, spherical indents exhibit two-way depth recovery with a shallow
indent depth observed at high temperature and deep indent depth at low temperature.
The reversible depth change is about 45% of the total indent depth and is stable over
many cycles. Figure 98 (a) shows the 3D indent profiles taken at temperatures above A;

and below Ms, respectively. Figure 98 (b) shows the cross section profiles of the heated

and cooled indent. DA was the depth after heating the sample to 400:1:5K. Since this
temperature was near A ;, the NiTi was in the austenite phase. DM was the depth after
cooling the sample to ambient temperature of 30013K. Since the ambient temperature
was close to M;, the sample was in the martensitic structure.

After the first heating step, the indent depth decreased from the original indent depth
of about 170 pm to 80 um (Figure 99), an approximately 53% indent depth recovery.

This is reminiscent of the one-way indentation depth recovery reported by Ni et al., who

113

showed that the shape memory effect exists under complex loading conditions produced
by spherical and sharp indentations [116, 231]. After cooling to the ambient
temperature of 300K, the indent depth increased from D; of about 80 um to DM of about
140 um, an approximately 75% increase in the depth of indent. The subsequent
heating-cooling cycles produced nearly constant indent depth ratio, (DM-DA)/DM, of about
45% (Figure 100).

Figure 99 and Figure 100 show that the two-way indent depth change was quite
stable over the five heating-cooling cycles tested. Both DA and DM slightly increased
over the thermal cycles, with a small decrease in the two-way depth change DA-DM over
the five thermal cycles. It is so far unclear the fatigue behavior of the two-way depth
change (DA-DH) over more thermal cycles. The depth changed most significantly
between the first and second thermal cycles. DA increased 2.5um; DM increased 6.3 um
and DA-DM increased 3.7um. But their values stabilized over the next 4 cycles with DA
increasing 0.75um/cycle, DM 1.4 uni/cycle and DA-DM 0.65um/cycle.

Figure 101 shows the ratio of (DM-DA)/DM, which compares the depth change to the
depth of the indent at martensite phase. The smallest ratio of depth change, around 0.5,
is at the center of the indent. Although the ratio increases to 1 further away from the
indent, both DM and DA become small and dominant by the surface roughness, causing the
fluctuation in the ratio. Ideally the (DMDJ/DM should approach zero for surfaces far
away from indents, because the surface does not move between austenite phase and

martensite phase. But (DH-DA) is not zero because: 1) small surface roughness change

114

in NiTi between martensitic and austenitic phase, 2) measurement error between two
measurements.

The “ball-and-clamp” training used in this study was used to induce the low
temperature shape into the sample by constraining the high temperature recovery.
Under this “constant-strain” training procedure, the stress was not constant during the
heating and cooling cycles. When heating temperatures above A ;, the stress under the
sphere increased as the recovery was constrained. When cooling below M;the stress
was decreased. Since the displacement of the clamped ball was constant, the only effect
that thermal cycling had created was the alternation between high and low contact stress.
The resulting “sink-in” of the area around the indent indicated that additional dislocations
were generated during the training, which most likely occurred during the heating stage
when contact stress increased.

Undoubtedly plastic deformation was introduced in the martensite phase during the
initial ball clamping or indentation process. The initial contact radius was about 0.8 mm.

The representative strain 8}; :

a, = 0.2-5’- : 0.2x 0'8
R 1.5875

=0.10

 

The 10% representative strain was high enough to introduce slip-plasticity in the
NiTi sample. These dislocations would have a spherically symmetric distribution due to

the symmetry of the spherical indentation. When first heated above A ;, the intensified

contact stress due to constrained recovery must have generated additional dislocations

115

into the austenite phase. The process should be equivalent to the process of an
unconstrained recovery followed by indentation to the same depth in the austenite phase.
These dislocations first introduced in the austenitic phase are different in mechanics
compared to the previous set of dislocations from indentation in martensitic NiTi. It is
also very possible that these dislocations would move when the sample was cooled back
to martensitic phase, because internal stress and strain field changed. When the sample
was heated back to austenitic phase, new dislocations might be introduced because the
previous dislocations in the austenitic phase have moved due to phase transformations.
The training cycles might stabilize these two sets of dislocations so that no more
additional dislocations are introduced during heat-cooling cycles. However it is also
possible that the stabilization is not achieved during training so that new dislocations are
introduced during each training cycle, which might contribute to the “sinking-in” effect
described below.

It is noticeable in Figure 98 that the area around the indent had a “sinking-in” effect
in the martensite phase but was leveled in the austenite phase. The radius of the
sinking-in area was about 1.8 mm, which was about 2.6 times of 700 um indent radius.
This sinking-in area and its reversible change over heating-cooling cycles indicate that
the training procedure has affected not only the microstructure beneath the sphere but
also a large portion of material around the indent. Since the surrounding area is not
directly under the indenter, either propagation of dislocations or residual stress field

causes the “sink-in” effect. Additional research is necessary to determine the

116

mechanisms of indentation-induced TWSME.

4.2.2 One—cycle Deformation Training and Reversible Surface Topography

The One-cycle deformation training method is an efficient and effective alternative
training method to the thermomechanical training because it does not involve cyclic
heating and cooling. Section 2.4.2 reviews previous research on the one-cycle
deformation training method. The experimental procedures were described in section
3.2 of Chapter 3. First, spherical indents or scratches were made on the surface of a
NiTi alloy in its martensite phase. Second, the indented or scratched surface was
polished mechanically to restore a flat surface in a step referred to as “planarization”.
Two-way reversible circular and line protrusions were thus produced by altering the
temperature to drive martensite-to-austenite phase transformations. Figure 102
schematically illustrates the preparation method and the proposed mechanisms
responsible for the two-way reversible surface protrusions [57].

Spherical indentation was performed on a martensitic NiTi sample so that after being
heated above A ;, the indent recovered partially. The experimental details are described
in 3.2 of Chapter 3. Figure 103 showed that residual indent had two-way reversible
depth change with heating and cooling: indents were shallower in the austenitic condition
and deeper in the martensitic condition. The average reversible depth change was
13.6:1:0.3 um over five thermal cycles tested, which is about 24% of the total indent depth.

But the reversible depth change is not constant: it increased after first two thermal cycles.

117

This “self—training” effect is discussed in section 4.2.4.

Reversible surface topography can be realized fi'om the indentation-induced TWSME.
Figure 104 shows a 3X3 matrix of circular protrusions. Figure 105 shows the sectional
profile of the reversible surface protrusion and its height over first five thermal cycles.
Figure 106 shows the actual shape of the circular surface protrusion, with same length
scale in the horizontal and vertical directions. Although in Figure 106 it seems the
circular surface protrusion is not very significant because the aspect (height/ width) ratio
is low, it is a strongly visible phenomenon.

Surface morphology can be characterized by several different kinds of surface
fluctuations: form error, waviness, and roughness (Figure 107). Surface roughness is
generally referred to the shortest-wavelength irregularities of a surface, and waviness is
the intermediate-wavelength deviation, and form is the longest-wavelength deviation.
Reversible surface protrusions can be potentially created to change the roughness,
waviness or form error, depending on the loading conditions of the indentations. The
surface protrusion shown in Figure 106 creates a surface form error, which is a
long-wavelength deviation from a flat surface. Figure 108 shows that imposing surface
protrusions close to each other and Figure 111 shows that nanoscale protrusions can be
created. These indicate that close-packed micro- and nanoscale protrusions can be made
on the NiTi surface to change surface waviness and roughness.

Figure 108 shows a reversible GM logo on the surface of a 2.54 cm diameter NiTi

SMA sample. The images were taken from a optical microsc0pe, but also visible to

118

naked eye. The letters were made by joining circular surface protrusions closely
adjacent to each other. The indenter and indentation load are the same as those in
making circular protrusions shown in Figure 104, except that the indents are close to each
other with around 0.8 mm spacing.

Comparing Figure 106 with Figure 108, it seems the reversible surface protrusions
are more significant optically than in physical dimensions. The reason is that the
reversible surface protrusions are big change in terms of surface form. The polished flat
surface of NiTi SMA had a roughness about 60 nm. The reversible surface protrusions
have a height above 10 um, which is very big compared to the surface roughness.

It is also shown in Figure 108 that making overlapping indents does not significantly
diminish the reversible surface protrusions. Although the stress and strain field of each
indent should be affected by the adjacent indents, it seems there is no detrimental effect.
Currently it is unclear whether this would change fatigue behavior.

Figure 109 shows a line protrusion after an indented and planarized bulk NiTi sample
was heated to the austenite phase. These protruding structures disappeared when the
sample was cooled down to the martensite phase. The height of the circular and line
protrusions was 13.6:t0.1um and 0.83:0.05um over the first five heating-cooling cycles,
respectively.

The geometry and size of reversible surface topography (RST) can be readily

controlled by indenter shape and applied load. Furthermore, intricate patterns of RST

119

can be efficiently laid out by arranging positions of indentations or length and direction
of scratches accordingly.

The reversible surface topography was also realized in sputter-deposited NiTi thin
film on silicon substrates. This means the reversible surface topography could be
readily integrated into the micro-electro-mechanical systems (MEMS). Figure 110
shows the same effect achieved in a 4.5 pm thick sputtered NiTi thin film deposited on a
(100)-silicon substrate. The height of the circular protrusion was 33.2:h7 .2 run over the
first five temperature cycles. However the protrusion deteriorated over the first five
thermal cycles more significantly than RST on bulk’NiTi samples (Figure 112), which is
discussed in section 4.2.3.

Figure 111 shows the reversible line protrusions on NiTi thin films, which were
planarized from a scratch that was produced by manually sliding a razor blade on the
NiTi thin film. The line protrusions had three parallel lines. This may had come fi'om
some imperfections in the razor blade. There was also a non-reversible “x” shaped line
protrusions nearby, which were from scratches in mechanical polishing.

Indeed, not every surface deformation in NiTi SMA could give reversible surface
protrusions. Spherical indentation and sliding were shown to be effective loading
conditions. But it is possible that more “blunt” indenters, like flat end punch, or more
“sharp” indenters, like Berkovich indenter, could not successfully induced TWSME,
which is for future research to clarify.

The RST in NiTi thin fihn on silicon substrate shown in Figure 110, deteriorated over

120

the first five thermal cycles (Figure 112). It is quite different from the RST in bulk NiTi.
This can be explained by the effect of the NiTi thin film thickness. The indent diameter
is about 7 pm. Therefore 7 pm NiTi thin fihn was needed to fully contain the
deformation zone that supported the TWSME (see section 4.2.3 for detail). But the
NiTi thin fihn thickness used in the experiment had a thickness of only 4.5 pm, less than
the contact diameter size that needs to support reversible surface protrusion. The
compromised deformation region underneath indent caused the deterioration of reversible

surface protrusion in Figure 112.

4.2.3 Layered Planarization

Insight into the mechanism underlying cyclic TWSME fiom indentation of
NiTi-based shape memory alloys can be obtained by experiments in which successively
deeper surface-removal steps are performed, until the reversible displacement effect is
found to vanish. Results of such an experiment are shown in Figure 113. An indent
was again made in martensitic NiTi, with a 1.59 mm diameter ball, and the specimen
polished just sufficiently to remove any trace of the indent, while maintaining the
temperature below Mf. After heating past A fthe height of the thermally-reversible
surface protrusion was determined, and subsequently measured again as additional
surface layers were successively polished away in each case, while the specimen was

martensitic. The protrusion amplitude was found to monotonically decrease fi'om initial

121

maximum of just over 11 pm until it vanished after ~05 mm (extrapolated) of the surface
had been removed. Protrusion formation thus apparently involves transformational
behavior in a subsurface zone nearly 40 times the protrusion height, indicating that the
effect can be produced with an average of ~2.3% transformational strain in the subsurface
volume, well within the known capability of TWSME in NiTi alloys [4].

Furthermore, when the contact mechanics of indentation are considered, the depth at
which the cyclic protrusion amplitude vanished is consistent with existing theories on the
mechanisms of TWSME in NiTi alloys. Figure 114 adapted from Sarnuels and
Mulhearn [232] shows theoretical predictions and experimental results that give the
distribution of plastic strain beneath a spherical indenter in Cu-30%Zn brass. Contours
are shown for plastic strains of 7%, 1% and 0%, the latter contour corresponding to the
elastic-plastic boundary. Of particular significance is the location of the 7% contour,
occurring at a depth approximately equal to the radius of the contact perimeter of the
indent. In the present experiment, the initial indent contact radius was 450 +/- 25 pm.
Therefore, the surface removal experiment, in which the effect vanished after removal of
~0.47 mm (roughly equal to the initial indent contact radius), indicates that cyclic
protrusion formation depends on TWSME occurring in material that experienced more
than 6 to 7% plastic strain (in the martensite phase) during indentation. This
corresponds closely to the maximum strain (6-8%) that can be tolerated by thermal
martensite before the onset of plastic flow via dislocation production. It is therefore

concluded that the initial indentation event created, within a near-surface zone whose

122

depth corresponded to the indenter radius, sufficient numbers of dislocations in the
martensite that could later cause a shape strain to occur during the austenite-to-martensite
transformation. At greater depth below the indent, where smaller plastic strains
occurred, little dislocation production is expected since strains in this region can be
accommodated entirely by martensite variant-boundary motion (deformation twinning).
This deeper material effectively self-accommodated the transformational shears on
cooling, and was therefore unable to contribute significantly to TWSME in the present

context.

4.2.4 Small Hysteresis and “Self Training” of Reversible Surface Protrusion
The relaxation and recovery of surface protrusion has a much smaller temperature
hysteresis compared to the phase transformation hysteresis of annealed NiTi alloy sample.
Figure 115 showed the surface protrusion height vs. temperature, for planarization of a
spherical indent that was produced under 60kg load using 1.587mm diameter ball. The
protrusion started to rise before the nominal A S temperature. Before the temperature
reached (A s +A ;)/2, surface protrusion height has already reached 90% of maximum
height, with a rate 0.26um/K. During the cooling stage, the surface protrusion
decreased slowly before MS but started a rapid descent well above MS. Again 90% of

the protrusion height was diminished before (MS+M;)/2 with a rate 0.31 rim/K.

123

As can be seen in Figure 115, the martensite-austenite phase transformation
temperatures shifted to lower temperatures in the heating direction; the
austenite-martensite transformation temperatures shifted to higher temperatures in the
cooling direction. The hysteresis of the reversible surface protrusion was 30:1:3K
compared to the phase transformation hysteresis of 60i5K, from DSC measurement of
free standing sample (Figure 32).

How could, in TWSME, the martensite—austenite transformation temperatures shifted
downwards to lower temperatures in the heating direction? Stalmans et al. noticed the
lowering of phase transformation temperatures in TWSME when the sample was
subjected to external stress that was “opposing” the direction of TWSME, and provided a
plausible explanation from the Clausius-Clapeyron equation [106] ( reviewed in detail in
section 2.4.3) .

The main reason is that the directions of the transformation strain in TWSME are set
by the training process; the external stress posed afterwards can either be “along” or
“opposing” the direction of TWSME. The former elevates phase transformation
temperatures and the later suppresses phase transformation temperatures. For instance,
if a specimen was trained to possess such TWSME: shrinkage occurs during
austenite-martensite transformation in the cooling direction (Ate;w < 0) and elongation
occurs during martensite-austenite transformation in the heating direction (A813 > 0 ).
The tensile stress is in the opposite direction of TWSME because it suppresses the strain

change in the cooling direction.

124

The Clausius-Clapeyron equation can be used to explain this by simply considering

the sign of enthalpy, strain and stress. The C-C equation is:

Ao _ AHfi’
AT Maj"

 

Tensile stress lowers the phase transformation temperatures as follows: on in cooling

M

from austenite to martensite phase, A1124 > 0, A87 < 0 and :54” < 0. The left hand
8A

 

side of the Clausius-Clapeyron equation :—; must has the same sign as the right hand

M
A

 

side, Since Ao > 0 (tensile), AT must be negative to give :—; < 0 so that

M.
8A

the left hand side of C-C equation was negative. Here the tensile stress lowers the phase

transformation temperature in the cooling direction.

A

In the heating from martensrtlc to austemtlc phase, TA ”f, 18 stlll negatlve because
8M

 

AH A”; <0, A8,”; >0. Both enthalpy and strain change flip sign. Note that in one-way
SME, the strain change does not change sign. Since Ao > 0 (tensile), AT is still
negative.

Therefore, small hysteresis can be obtained in indentation-induced TWSME only if
the residual stress meets the following requirements:

It must be in the opposite direction of TWSME during heating so that phase
transformation temperatures are lowered;

It must not be in the opposite direction of TWSME during cooling so that phase

transformation temperatures are elevated.

125

If the same TWSME specimen in the previous paragraph were subjected to
compressive stress during cooling (A324 < 0 ), the austenite-martensite phase
transformation temperatures would be shifted upwards in the cooling direction; if the
stress was changed to tensile during heating (As A: > 0 ), the martensite-austenite
transformation temperatures were shifted downwards in the heating direction. Thus
small hysteresis can be achieved by moving the transformation temperatures towards
each other.

In the above situation, the direction of the stress can actually assist the deformation:
tensile stress with positive strain change and compressive with negative strain change.
This could stabilize the TWSME.

A similar situation may be what actually happened in the small hysteresis behavior of
reversible surface topography, which is shown in a schematic in Figure 116. It is known
that residual stress exists in elastic-plastic indentation [108]. The distribution and extent
of the residual stress in indentation-induced TWSME is very complicated, considering the
coupling of martensite phase transformation and planarization process. However the
residual stress may very possibly act in the direction that elevates the phase transformation
temperatures during heating and suppresses the transformation temperatures during
cooling, creating the small hysteresis.

Other experimental evidence comes from the “self training” effect of the
indentation-induced TWSME. As can be seen in Figure 103 (reversible depth change

from one-cycle deformation training) and Figure 105 (reversible circular surface

126

protrusions), initial thermal cycles after training process did not lead to immediate
deterioration of the TWSME, but instead acted to strengthen the TWSME slightly. It
was in sharp contrast to the fatigue behavior of uniaxial TWSME, where the two-way
effect deteriorated significantly during the first several thermal cycles [101]. It has been
pointed out in the previous paragraphs that for TWSME to have small hysteresis, the
direction of external stress must be in directions to assist the deformation. The residual
stress in indentation-induced TWSME may very possibly act this way, strengthening the

TWSME, curbing the deterioration, and narrowing the temperature hysteresis.

4.2.5 Mechanism of Indentation-induced TWSME

The successive layered planarization experiment revealed that a zone underneath
indent with the size of the indent diameter is mainly responsible for the reversible surface
topography. Theoretical calculation based on contact mechanics and experimental
measurements by Samuel and Mulhearn [232] also pointed out this zone is a highly
deformed region with at least 7% plastic strain.

In one-cycle deformation training of TWSME, it is believed that the dislocation
structure and the associated stress field create rrricrostructural, thermodynamic and stress
anisotropy within the martensite phase, as discussed in section 2.4.2. The
thermodynamic anisotropy means that after training, the crystallographically equivalent

martensite variants are no longer thermodynamically equivalent: the trained martensite

127

variants are thermodynamically preferred to the non-trained variants. The
thermodynamic anisotropy must have come fi'om certain changes in microstructure as the
result of training. The microstructural basis of the thermodynamic anisotropy was
termed microstructural anisotropy. As a result, martensite variants in certain
orientations are thermodynamically favored over others when cooled from the austenite
phase. The preferred nucleation and growth of these directional martensite variants
enable the SMA to “remember” its low temperature shape.

As is shown section 2.4.2, slip-plasticity is effective to introduce TWSME under
uniaxial loading. In indentation-induced TWSME, dislocations should play a similar
important role to induce the reversible effect. Comparing to simple loading conditions
such as tension and compression, indentation and scratching are effective to introduce
dislocations locally and symmetrically. Several authors have suggested that
indentations are accommodated by circular loops of “geometrically necessary
dislocations” under spherical indenters [233, 234] (Figure 117). It is assumed that the
indentation is accommodated by circular loops of geometrically necessary dislocations
with Burgers vectors normal to the plane of the surface. As the indenter is forced into
the surface of a single crystal, geometrically necessary dislocations are required to
account for the permanent shape change at the surface. These geometrically necessary
dislocations lead to spherically symmetric microstructures and residual stresses beneath
the indents, which may induce TWSME.

On the other hand, the small hysteresis of the reversible surface topography and

128

“self-training” effect indicated the residual stress played a significant role in altering
phase transformation temperatures and stabilizing the two-way effect. However, the
distribution and extent of the residual stress in indentation-induced TWSME are very
complicated, especially considering the coupling with the martensite phase
transformation, which undoubtedly changes the residual stress significantly.

The residual stress and strain fields in an as-indented martensitic NiTi should be close
to those in an elastic-plastic material. But how the stress and strain fields change after
the first recovery is unclear. How those change in the following thermal cycles would
be a useful finite element analysis to shed some light on indentational TWSME.

Indentational shape memory effect has not been modeled in this work due to the
difficulty determining appropriate boundary conditions. Therefore a finite element
analysis with general approximation was made here, which is an indentation in an
austenitic NiTi. The reason is that the stress-strain field in austenitic NiTi, which
includes the effects of stress-induced phase transformation, may be close to those after a
thermally driven phase transformation in an austenitic NiTi.

Figure 118 presents a finite element analysis of the ow residual stress field (normal to
the surface) in austenitic NiTi after spherical indentation. There exits a highly
compressive stress “core” extending fiom one quarter of the indent diameter to one
indent diameter beneath the residual indent. The compressive stress also agrees with
Figure 116, where there should be a compressive stress at temperature above A fto pull

the surface protrusion down. When cooling down from the austenite phase to the

129

martensite phase, this compressive stress in the vertical direction may be able to help pull
material back from surface, which would create a deeper indent. How this residual
stress evolves in phase transformations by heating and cooling would help provide more

information.

130

5.1

Chapter 5

Conclusions and Future Work

Conclusions: Hard Coating Systems with NiTi Interlayer

Based on temperature-controlled scratch and dry sliding wear experiments on CrN

and DLC coated 6061 T6 aluminum, tested both with and without NiTi interlayers, the

following conclusions may be drawn:

1)

2)

3)

Magnetron sputter deposition is capable of producing excellent adhesion of NiTi
on aluminum substrates, without elaborate surface preparation or other
adhesion-promoters.

The NiTi thin films on aluminum substrates display robust martensitic
transformational behavior, as shown in temperature-scanning X-ray diffraction
and wafer-curvature measurements. The wafer-curvature measured on
DLC:NiTizAl also shows strong phase transformational stress relaxation and
recovery in the NiTi interlayer.

Superelastic NiTi interlayers on very soft aluminum substantially improves
adhesion of nitride hard coatings, as indicated by an increase in the critical load
for delamination as measured in temperature-controlled scratch tests. The

marked increase in performance at test temperatures near A fstrongly suggests

that stress-induced A->M transformation is responsible for the observed

131

4)

5)

6)

7)

improvement.

The adhesion of hard coatings to martensitic NiTi interlayers can be improved
by deposition of a protective Cr layer and subsequent aggressive sputter cleaning.
This procedure could also be replaced by depositing NiTi interlayers and hard
coatings within the same vacuum chamber.

NiTi interlayers dramatically improve the wear behavior of CrN and DLC films
on soft aluminum and can increase wear resistance by a factor of 100-3 00. The
best performance is noted for test temperatures near or above A f. Martensitic
NiTi interlayers can improve the wear resistance by a factor of 40—100.

Thicker NiTi interlayers are more effective. Performance monotonically
improves as the interlayer thickness is increases to 21 um. Thicker CrN
coating also increases wear resistance monotonically for thicknesses up to 4 pm.
It is unclear now whether there is an upper limit to hard coating and interlayer
thickness on aluminum substrates.

The NiTi interlayers between CrN and hard M2 steel substrates do not improve
wear resistance significantly, compared to the samples without NiTi interlayers.
It is suspected that the weak CrN:NiTi interface is responsible for the worse
wear resistance of CrN:NiTi:M2. However the interfacial adhesion can be
improved, which might lead to better wear performance of CrN:NiTi:M2 in
future studies. Placing an interlayer that is softer and more compliant than the

substrate can lead to larger strains in the hard coating, and larger interfacial

132

‘8)

mismatches.

Finite element analysis proves that the use of superelastic NiTi interlayers can
shield the soft aluminum substrate fi'om contact strains, while the CrN hard
coating can support contact stress. The combination of CrN hard coating and

superelastic NiTi interlayer successfillly protects the aluminum substrates.

5.2 Conclusions: Indentation-Induced TWSME

1)

2)

The two-way shape memory (TWSME) effect can be realized at the micro- and
nano-scale using spherical indentation. Both thermomechanical cycling and
slip-plastic deformation were proven to be effective training methods to induce
the two-way shape memory effect. Thermomechanical cycling results in larger
reversible indent depth change and a “sink-in” effect. The slip-plastic
deformation is an efficient one-cycle training method and there is no “sink-in”
effect. TWSME does not decay significantly over initial thermal cycles.
Reversible surface protrusions (RSP) are produced fi'om planarizing spherical
indents with TWSME. The height of the RSP is the same as the two-way
indent depth change, which means the planarization process does not affect the
TWSME. Layered planarization experiment shows that the RSP arises from
transformational activity in a zone with the size of the contact radius beneath the

indent, which is about the boundary of 7% plastic strain.

133

3)

The residual stress underneath the indents may possibly strengthen the TWSME.
The RSP has a smaller temperature hysteresis than that of the bulk NiTi SMA.
Second is the “self-training” effect, which means the RSP does not decay but
strengthens slightly during initial thermal cycles. From the above two
phenomena and thermodynamic analysis, it is concluded that the residual stress

from contact loading may be able to promote the TWSME and RSP.

5.3 Suggestions to Future Work

1)

2)

3)

Mechanisms of indentation-induced shape memory, superelasticity and two-way
shape memory effect. New experimental tools like synchrotron X-ray
diffraction and in-situ TEM indentation can be used to analyze the
microstructures beneath indents in NiTi alloys, along with traditional
experimental techniques like cross-sectioning samples with indents.
Tribological applications of indentation-induced two-way shape memory effect.
Surfaces with reversible surface protrusions could have different fiiction
coefficients compared to flat surface. The difference in friction coefficients is
also dependent on testing conditions like loading configurations and lubrications.
Development of finite element modeling of shape memory effect is the next step
after current modeling of superelasticity. FEA of shape memory effect can

provide stress-strain distributions beneath indents in NiTi before and after phase

134

4)

5)

transformations, which can be used to explain indentation-induced shape memory
and two-way shape memory effect.

Tribological tests on CrN:NiTizAl and DLC:NiTi:Al samples with re-heat treated
aluminum substrates. The temper of 6061 T6 aluminum substrates were
destroyed during NiTi interlayer deposition, which led to low hardness (HRE 62).
It is possible to re-heat treat the aluminum alloy to restore the hardness to HRB
60, after the deposition of NiTi interlayers and hard coatings.

Finite element modeling using cohesive elements and user material subroutine
for superelasticity. The cohesive elements can be used to model interfacial
delamination between coatings, interlayers and substrates. It is possible to show
in FEA that superelastic NiTi interlayers discourage interfacial delamination
under contact loading, which supports experimental results. Parametric studies
can also be carried out to find the optimal mechanical properties of hard coating,

NiTi interlayer and substrate to maximize tribological performance.

135

Appendices

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Table 1. Comparison of crystallographic data of NiTi martensite a, Otsuka et a1 [20];

b, Hehmann and Sandrock [235]; c, Michal and Sinclair [21]; d, Kudoh et a1 [22].
088 (a) HS (b) MS (c) KTSO (d)

a (run) 0.2889 0.2883 0.2885 0.2898

b (nm) 0.4120 0.41 17 0.4120 0.4108

c (nm 0.4622 0.4623 0.4622 0.4646

[1 (°) 96.8 96.8 96.8 97.78

v (nm’) 0.05463 0.05449 0.5455 0.05479

Atoms per unit cell 4 4 4 4

SPace 8mm ”’0 P21/m P21/m P21/m

Table 2. Lattice structure and parameters of precipitates in NiTi alloys

Name Ti3N14 Ti2Ni3 TiNi; TiN i2

Space group R3 I4/mmm D024 Fd3m

(rhombohedral) (tetragonal) (hexagonal) (cubic)

 

Lattice a = 0.67nm a = 0.3095nm a = 0,501,222), a = 1.132nm
parameter a = 113.8° c =1.3585nm c = 0.8307nm
c/ a = 1.6284

 

 

 

 

 

 

 

Reference [236, 23 7] [238] [239] [240]

 

136

Table 3.

steel substrates, before CrN deposition

Deposition parameters for NiTi interlayers on 6061 T6 substrates and M2

 

 

 

 

 

 

 

 

 

Target Deposition Substrate- Deposition
Ar Pressure
Composition Power Target Distance Temperature
(Pa)
(at.%) (W) (mm) (K)
53.5 Ti 46.5Ni 200 0.35 63.5 483
Film
Deposition Time Base Pressure NiTi Film Heat
Composition
(hour) (Pa) Thickness (um) treatment
(at.%)
800K for
51. 7 Ti 48.3 Ni 0.45,].5,2.5 5.2><10'5 4.5, 9, 15
7200 ks

 

 

 

 

 

Table 4. Deposition parameters for NiTiCu interlayers on 6061 T6 substrates,
before CrN deposition
Target Deposition Substrate- Deposition
Ar Pressure
Composition Power Target Distance Temperature
(Pa)
(at'/o) (W) (mm) (K)
53.5 Ti 8 Cu
200 0.35 63.5 743
38.5 Ni
Film
Deposition Time Base Pressure NiTi Film Heat
Composition
(hour) (Pa) Thickness (pm) treatment
(at.%)
51.7 Ti 6.] Cu 793K for
1.5 5.3 x 10'5 9
42.2 Ni 7200 ks

 

 

 

 

 

 

137

 

 

Table 5.

Deposition parameters for NiTi interlayers on 6061 T6 substrates, before

 

 

 

 

DLC deposition
Target Deposition Ar Pressure Substrate- Deposition
Composition Power (Pa) Target Distance Temperature
(at.%) (W) (Hill!) (K)
53.5 Ti 46.5
200 0.53 63.5 523
Ni
Film
Deposition Time Base Pressure NiTi Film Heat
Composition
(hour) (Pa) Thickness (um) treatment
(at.%)
773K for
51.7 Ti 48.3 Ni 25,35 2.0x10“ 9,15,21
9000 ks

 

 

 

 

 

 

 

 

 

 

Table 6. Deposition parameters for NiTi interlayers on M2 tool steel substrates,
before CrN deposition
Target Deposition Substrate- Deposition
Ar Pressure
Composition Power Target Distance Temperature
(P8)
(at.%) (W) (mm) (K)
53.5 Ti 46.5 Ni
200 0.33 63.5 723
51.2 Ti 48.8 Ni
Film
Deposition Time Base Pressure NiTi Film Heat
Composition
(hour) (Pa) Thickness (pm) treatment
(at.%)
51.7 Ti 48.3 Ni 793K for
1,2 4.3x 10'5 6,12
48.2 Ti 51.8 Ni 7200ks

 

 

 

 

 

 

138

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

Table 7. Phase transformation temperatures of thin film and bulk NiTi used in
TWSME study
. , Martensite
Austenite start Austenite finish Martensite start
finish
(As) (A1) (Ms)
(M1)
Bulk NiTi 350K 404K 344K 287K
Thin Film
343K 390K 329K 282K
NiTi
Table 8. Material parameters of NiTi and reference materials used in finite element
modeling. (The “Ref” stands for reference material)
Austenite Austenite Martensite
start , start Martensite finish
finish
Phase
transformation (111;)
318 K 338 K 323 K 303 K
k1 *2 k3 *4
k3 parameters
0.1946 -7.784 0.1946 -10.622
. . Elastic Plastic . . Yield
Transformation Testing Modulus hardening Porsson’s stress
Strain Temperature , , K slope It ratio
(NlTl and Ref)
(Ni’l‘i and Ref) (Ref)
5% 90K 80 GPa 10 GPa 0.3 533.5
MPa

 

 

 

 

 

 

I39

 

 

 

 

 

 

 

 

 

 

 

2000
8’ .
Q 00 _ 1455
3 15 1380
e u
or no
a. 1304
E 1118 Ni
1- TiNi
looo -
. Ni3'fi2
“‘8“ . Nuria‘l' }Metastable
500 . . . . . I I . . . . .
T 20 “ZN' 40 , 6o "'3“ 80 Ni
Atom % Nl
Figure 1. Phase diagram of a NiTi alloy by Massalski[241], to which the phase

equilibrium between the B2 and Ti3Ni4 phases are added [242].
T" .. i ‘ (C)

   

Figure 2. Transmission electron micrographs of Tl3Nl4 precipitates in 51.3at%NiTi
thin films which were solution-treated at 973 K for 3.6 ks followed by age-treatment at
773 K for (a) 3.6 ks; (b) 36 ks and (c) 360 ks, at 673 K for (d) 3.6 ks; (e) 36 ks and (f)
360 ks, and at 573 K for (g) 3.6 ks; (h) 36 ks and (i) 360 ks (after reference [51]).

140

Parent Phase Martensite

 

I ’ e— martenslte matrix

e—twln

Invariant plane (habit plane)

 

 

 

Figure 3. Schematic illustration of parent and martensite with invariant plane as a
habit plane.

141

 

 

(010)82

 

 

 

(100182

Figure 4. Stereographic projection of 24 habit plane variants from WLR theory, for
<01—1> Type II twinning.

142

 

 

  

 

 

 

m _1 L l l 1 1 l l
350 , r
A 310-4 r-
5
g 250~ -
(B
5.; aoo~ -
Q
E.
1' 150 d o Kormllov
E” + Hanlon
100 - v Melton '-
x Miyazaki
50d 0 Nishida _
In Wasilewski
a Smith
0 r r 1 W r 1 r r
0.48 0.49 0.50 0.51 0.52
Mole-Fraction NI
Figure 5. Ms temperature as a function of Ni content for binary NiTi alloys.

Different data symbols represent data from different authors. The solid line is given by
thermodynamic calculations (after reference [34]).

143

1.2 GPa

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

600
1\
q 1 1.5 W .
~ ~ 1
., I
m w
85 .
E .,
g o ‘L — ““““““ 1.—
-200
400
-600
’2 x 00mm
We:
soo P1>Pz>P3 e
0 473 573 873
Temperature, K
Figure 6. Schematic stress-temperature diagram for isochronal annealing of

sputtered NiTi films on (100) silicon. Dark curves P1, P2 and P3 show the general
intrinsic stress—deposition temperature characteristics for different pressures of the
working gas. Trajectories a, b, c and d are for amorphous films deposited at different
pressures and at temperatures below T a), the minimum in-situ crystallization temperature.
Isochronal crystallization occurs at Tc giving a transient tensile stress spike at x. Curve d
additionally shows the effect of an isothermal hold inserted in the isochronal sequence.
Trajectories e, f and g are for crystalline films deposited above T 04. Cooling from any
point on a trajectory generally proceeds along a path with slope 1.5 MPa/K, to points
such as d‘, a‘ and g‘, allowing a wide range of stress levels to be obtained at T ,1. (after
reference [58])

144

153.0 K 164.1 K

 

 

 

 

   
 

 

 

O
'82.?“ 285K
l93.3l( ~203.0K '
25LOK
? 24I.OK
. w -
g 223.7K —o r
a
w
o l
= Q
a, A A - # A A __-
3 O O 0 O
'5 2765K
5 273.2K
I" 263.4K
(9) 283.7 K
M9190 K
Af=221 K
o 2 4. 2 4o é‘To 2 4‘
Strain (“/o)
Figure 7. Stress—strain curves as a function of temperature of 50.6 at.% NiTi alloy,

which was solution-treated at 1273 K for 1 h. The critical stress for inducing
martensites and that for reverse transformation are indicated by an arrow (—)) and by a
double arrow (—>->) in 0), respectively. (after reference [68])

145

Tensrle Stress (MP0)
5
8

 

 

zoo 04k °
F 0‘0 O/ .
so /
\o O
\ / ’I /
O ’1 /.
o_‘_€} L j 1; v A L
n :50 “s 200 A, 250 300

Temperature (K)

Figure 8. Critical stress as a function of temperature for inducing martensites (open
circles) and for reverse transformation (solid circles), which were plotted from the data in

Figure 7. (after reference [68])

146

   
  

 

 

 

 

I drape memory effect " "
' ".o‘ 0“.“ ~“.:.'o" ':. .3... ..
a . 0.: :.o. o O...- ... .
a .. .- . :-
+§;.I~... '0‘- ., ;".t. ...v.._ $
, . ‘. "P .- .x. . h: 9
. 0 , ., 37:1... IQ crtlcal stress for sllp (B)
-.. . lit-.1 o .:.‘._.. 3.. .. _ '. V ~5‘~~‘-
',.v .' tf";'; ' \i ‘5‘.
' ‘-" .. :'- ' ' '.° ' 3f
Mst AsAf

Temperature —>

Figure 9. Schematic representation for the appearance of the shape memory effect
and superelasticity in temperature—stress space. (after reference [4])

500' Ti-50.3%Ni single crystal
1273 K IQ- 673 I0

400 - Ms=244 K ,.
Af =2e4 K ', ’

Tensile Stress (MPa)
8 or
9 8

100

 

 

 

0 R 4 J Vlfi A j
150 200 250 300 350
Deformation Temperature (K)

Figure 10. Critical stress for inducing martensites and for reverse transformation of
a 50.3at.%NiTi single crystal, which was solution-treated at 1273 K for 1 h followed by

aging at 673 K for l h. (afier reference [243].)

147

80 203 K 220K 243 K
40 J
+ M

O ‘- 0....ng 1 4— 1’

l- 262 K 233 K 289 K 295 K

 

80

    

40

 

0 ,

Stress (MPa)

200 b . Ms=233 K
160 .298 K 302 K 307 K Af= 285 K

Tm=315 K
120

80
40

 

 

o 040 0.4 00.403
Strain (%)

 

0

Figure 11. Stress—strain curves as a function of temperature of a Ti—50.5Ni single
crystal, which was aged at 673 K for l h after solution-treatment. The curves are related
with R—phase transformation or deformation of R-phase. T’m=Rs. (after reference

[72])

148

120- (a) :\

IOO

V

Tensale Stress (MPa)
1)
o 8
_ "" Dfl=szb:-—r>—

 

 

 

20 '
( w—A—a—
0 4 . . - r . . m
)- “-.o~“~
08» (b) ox‘k‘No
. \0
A ‘\
2‘3 04) ‘o
.9 x,
.5.’ ‘ ‘9
U) Q

 

0 . . . . . L . a m
220 240 260 280 300 320

Temperature (K)
Figure 12. (a) Critical stresses for R-phase deformation and superelasticity, and (b)

the recoverable strain of R-phase superelasticity, which is a function of temperature.
(after reference [72])

149

 

4— a = 40 MPG

 

 

 

 

 

 

 

 

 

 

5 at?‘ Cu 12.5 otx Cu
2..
0r- L - -
A 7.5 ates Cu
3 4“ N " 15 am Cu
.5 r T:
a; 2- N -+ «12 K ‘"
o
_ t
4— '0 ‘3’" C“ - 20 am Cu
Mf ’5
2" § _. .11 K "
O3 ' ) -"‘ “‘K
0C 1 Ms , _. Ms ,1
230 270 310 350 250 290 330 370
Temperature (K)
Figure 13. Elongation vs. temperature curves under 40 MPa load for Ti—Ni—Cu

alloys with various compositions indicated.

150

(after reference [73]).

<—Stoge I->(<—Stoge H~—->l<—~—— Stage ll—ll—-->

 
   
 

 

 

 

 

 

IOOO ® a 5!
E
a 800
U
I
2%
___°;. Deformed at243 k
E Ms=190 K
.2 Af=221 K
E 4" \~————-/' '\._‘I —fi]:¥
0 l 20 50
(a) Strain (%)
.Stago 14+— Stage II —-t-—— StageIII ---+—-
[5 A
(or 8s
93 (e) 8,
‘2 0L (Al SP
'8
t—
p
m o
5- ‘
O
O
. .
A O
A A . . A :
o s to Is 20
(b) Total Strain (“16)
Figure 14. Stress—strain curves of 50.6 at.%NiTi alloy, which was deformed at 243

K (>Af). The alloy was solution-treated at 1273 K for l h. Dashed lines represent the
recovered strain upon heating to 373 K. The symbol (x) represents the fracture point. (b)
Plots of various strains in (a) as a function of total strain. (after reference [68])

151

 

6(1) Deformation of a polycrystalline martensitic NiTi SMA under tension

    

  
   

‘JP' ‘v‘ {D "L‘..~' .
".sjw he»
i .. ,. {a .

I...
\
. .

        

 

 

 

cu
CL
:5 400»
vi
to
2
:75
.E’ 200-
a
0
.E
U)
ch

o .

0 3 o 9 12 15
Engineering Strain, %
(8) Twin Bot-fiery um My

 

mmmummrwmdnmmm
don

 

 

comma
(0W ‘ " "“ “"‘,
(ammm

(b) mmmuw‘ ' “fomddomfllon

Figure 15. Deformation stage of martensitic NiTi alloy sample (after reference
[73])-

152

 

Figure 16. TEM images of in-situ tensile loading of superelastic NiTi. (a)
martensite plates 1, 2, 3, 4 formed at 118 um displacement. (b) High density of looped
dislocations (marked by ABCD) formed at 150 um displacement. (c) Magnified photo
of the area between martensite plates 1, 2 and 3, in (b). (d) Diffraction pattern
corresponding to the area indicated by single-headed arrow in (c). (After reference [79].)
The martensite plate 2 shows some detwinning at the lower part; but the plate 3 does not
seems to detwin, even though dislocations have been generated.

153

 

 

 

 

 

 

 

 

 

I . 1 l I
Tensron
12m ._ [211] Orientation __
Single Crystal Ti-50.8at%Ni .
1000 Aged 1.5 hours a 823 K 3% strain
A '— .—
n de/dt=1x10'3
E 800 l— _
O
‘ 600 — -
I; cycle I _ 2 10 16
.00 _ 7 _ -
200 _
o l l
o 0 0
Strain, s (°/o)
1400 _
ompressron
1200 _ [21110rientatlon
Single Crystal Ti-50.8at%Ni
Aged 1.5 hours a 823 K
1000 — 4
‘3 dc / dt = 6x10
E 800 _
c cycle 1 2 10 XX)
g: 600 - —
m 400 —
200 l .-
l
o l.‘
0 0 0
Strain, e (%)
Figure 17. (above) Tensile cyclic stress—strain response of single crystal NiTi

oriented in the [1 12] direction. The test temperature is above the A f transformation
temperature. (below) Compressive cyclic stress—strain response of single crystal NiTi
oriented in the [1 12] direction. The test temperature is above the A f transformation
temperature. (after reference [76])

154

 

a. * ,.,. ,- ‘
”w
“’W

 

Infill“...

Figure 18. Observation of Luders deformation by photography for Ti-50.5Ni alloy.
Black regions represent the parent phase, and lighter colored regions martensite. The

strain increases with increasing numbers from (1) to (13). (after reference [83]).

 

. ..... H421

 

Tensile street (MPa)
3
.
x

g ‘ C.--_‘_‘_I‘ ‘0 l.-_—:
H
332 It 1149.8 at‘llr Ni
1273K 10-673 IQ
M9294 K
Af=318 K
.0 2 4 5 I 101214

Skin 0‘)

Figure 19. Effect of thermomechanical treatment on the stress—strain curves as a
function of deformation temperature for 49.8at.%NiTi alloy. Solution-treatment at 1273
K for 1 h followed by annealing at 673 K for 1 h. (after reference [244])

155

v

 

 

Tenslle Stress (MPa)
«- 3

‘C)

    

Ti=4SJ 1')“- Hi
673 K l0
Ms=251K
M=313K

     
     

8.246.246.

 

00246024

Strain (%)
Figure 20. Effect of thermomechanical treatment on the stress—strain curves as a

fimction of deformation temperature for 49.8at.%NiTi alloy. Annealed at 673 K for l h
immediately after cold-working. (afier reference [244])

156

 

(a) as-rolled

O I. 3 J‘, m_
Figure 21. TEM images of 49.8at.% NiTi alloy, which were annealed for 1 h at

various temperatures indicated after cold-working; (a) as cold-worked, (b) at 673 K, (c) at
773 K and (d) at 873 K. (after reference [245])

157

  

(0)209K (b12llK (c) 221K (6)229K
(8)24OK

Omani

h-.-’ \aoncvi \----’ 6.--.l 6----d

 

 

 
     
   
       
     

 

 

 

 

   

O
001296K
“is"
O.
a [Z
(I)
8
L. \..J
O
6
g soot (ml 324K
I!)
S tkl304K tl)3l4l<
!- , Ti-50.60t%Ni
4oor I273K to -*
773K (hr 10
Ms=ZSOK
200 Af-292K
o 02‘4L680244681l0‘
S tr 0 in {%)
Figure 22. Stress—strain curves as a function of testing temperature of 50.6at.%NiTi

alloy, which was aged at 773 K for 1 h after solution-treatment at 1273 K for l h.
(after reference [245])

158

800~ 0" 0,, Heat treatment
A A l273K 10* 773K [0

O O s - 673K 10
o I = ~573K IO

Ti-50.6 at°/o Ni

     

0')
O
O

Tensile Stress (MPa)
.3
3

 

 

am I-
O
G :30 200 '
Temperature (K)
Figure 23. Effect of aging temperature on the critical stresses for inducing

martensites and for reverse transformation in a 50.6at.% NiTi alloy, which was aged at
various temperatures for 1 h afier solution treatment. (after reference [244])

159

6‘;
(a) 573K m, ,5, ..
"lf$1:\ 1}“: n. kilns" .
5'?! V “1‘5“ png\ I”? _
. .2.; 3} it”; a. ‘ét’:

- . v . .- .
A;.. §\“‘.5}§‘€:': o.‘ .u‘

 

R:
“I? — .

(b) 773K" “‘8 "‘
\ \K x“; \Q Q"\
‘- \\ ‘ ‘&\‘ r\\
s y‘

\ \V\\\ \ $15!. is ' \‘

\\ \ ‘80 K. ;

fk‘V‘h‘ \ ‘ \ \\ \\

‘\ ~‘I

“(it s: \ ‘3.
Figure 24. TEM images and corresponding diffraction patterns of 50. 6at. % NiTi

alloy specimens, which was aged at various temperatures afier solution-treatment at 1273
K for l h; (a and a’) at 673 K, (b and b’) at 773 K. (after reference [245])

 

 

 

 

Tenslle Stress (Ina)

 

 

 

 

Strain fl)

Figure 25. Stress—strain curves as a function of deformation temperature of
50.6at.%NiTi alloy, which was aged at 673 K immediately after cold-working. (afier
reference [245])

161

Anealed at fill) C for 1h, air cooled
WW Dimaeter - 4.5 mm, Pitch .0.1 mm

1 Deformed and constrained atthb shape
and annealed at different temperatures

[\/\/\/\/\/\/|

1 Diameter - 4.5 mm, Pitch - 12 mm

a. Deformed at martensite state
Diameter -4.5mm, Pitch-0 mm

Heated up to temperature higher than M

NWWVl

l Cooled to temperature lower than If
Figure 26. The procedure of processing TWSME TiNi SMAs spring (afier reference
[87])

 

 

 

 

  

 

 

 

 

 

 

1

o 6.2 'llr defamation l 346 K

i
a? ‘ A
.5 ‘2 ' g
s s;
a -3 -

-4 k tramway J J

L memory .

-5 _ —

-6 b L 1 L .1 L L 1 _l L L t __

260 280 300 320 340 360 380

Temperature (K)

Figure 27. The two-way shape memory effect obtained after a single deformation by
6.2% of equiatomic NiTi alloy (afier reference [102]).

I62

a T: 22 C b T=60 C<1 min

 

.

c T: 22 C d T=60 0.30 min

   

Figure 28. Optical microscopy verification of the predicted nricrostructure memory
for an aged AU51Cd 49 martensite and its disappearance after aging in the parent phase.
(a) Domain pattern of the aged martensite (aged at room temperature for six months).
The sample is heated up to (b) parent phase (60 °C), and immediately cooled down to (c)
martensite. Then the sample is heated up again to (d) parent phase (60 °C) and aged for
30 minutes, followed by cooling down to (e) martensite. (after reference [104]).

_-

63

 

Austenite

with tensile
residual stress field

 

l i

Cool

 

 

 

l l

 

 

Heat

 

Martensite

 

 

Figure 29. TWSME due to internal tensile residual stress: elongation during cooling

and shrinkage during heating.

fill.
i i

 

 

 

 

M
NH
(ill
Austenite ’1‘ T

l l

 

 

 

(if)

Cool

Heat

Cool

Heat

 

 

 

 

 

 

Martensite

Martensite

Figure 30. External Stress and Internal Compressive Stress Field. The tensile
external stress increases the internal tensile stress, elevating phase transformation
temperatures but compressive external stress decreases the internal tensile stress.

164

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eogoomm 358“ a mo common cougomaab 05 .3385 wagon 3v can mam—coo A3 wet—6 noise o§eH&EoH-qum ._m onE
3: 22360... 3: 839.030...
. mm . mmm . awn . mum . mmm . mmm . men man
i . ~31 r
l. . . --.----... new .8 . ................ w ........ ll. .. - Lfin .2.
. . .ll itiitttitiitt 4 in“)... iiiiiii
e ix... 174%.? (ii. . 41./v1 “ a
_. . rmd
/ t l E 2 n Eu m
. . n . m
. T3 w 33 a
.0 m.
t (E. u E e
.06 m. 6.0 m.
. m. u
r a
Em. .2 M _E :3 ./.
0" ll‘
ll PHD . E .
lllill E E .
i la .2 VB 3
l. E . pg.
n E _ .
In I film. .2 QB .2
H l E T. E a

 

 

 

 

 

 

 

 

3. .3

165

 

 

    
     
  

   
    
  

 
   

 

 

°‘2 50.51 °C
°" * 13.96°C
g 1 62.15°C
_g 0.0 . 25.41 J/g
'§ 103.09%:
I 25.20J/g
77.05°C
-°-1 * 31.56°C
112.64°C
-0.2 . , T r . , . . . . j . . , . .
-50 0 50 100 150 200
Temperature (°C)
Figure 32. DSC curves of bulk NiTi samples used in TWSME study

 

   
  

 

 

 

 

1.5
oz/p
oalp
' arlp
0.5 -—
-- - -’ - —- -:I’
--“~‘ F a “—4 ”"’
o,/p ‘ ~' . l’ or/pm
Uniform pressure p Hertz pressure
— I .0

—0.5 o —0.5 —1.0 ~15

 

 

Figure 33. Stress distributions at the surface and along the axis of symmetry caused
by (lefi) uniform pressure and (right) Hertz pressure acting on a contacting radius a (afier
K.L. Johnson, Contact Mechanics, Cambridge University Press, 1985, pp.94, [108]).

I67

 

(c) (d)

Figure 34. Two-dimensional photo—elastic fi'inge patterns (contours of shear stress).
(a) point load; (b) uniform pressure; (c)rigid flat punch; (d) contact of cylinders. (after
reference [108], pp. 103).

I68

 

 

 

Figure 35. The “pile-up” (above) where contact radius is larger than linear case and,
“sink-in” (below) where contact radius is smaller.

I69

 

 

 

 

 

200
b
281'
l-
r
5‘
E .
E
g 100 -
E P
D- e
, 2-8?
' B
i.
, / equivalent defer-attain %
i '9 1.5

y I T I

01 0-2 0-3 0.4 0.5 0-6 0-7 0.3 0-9
dID

Figure 36. Representative strain vs. hardness in spherical indentation, and its
correlation with true stress-true strain in uniaxial tension. Metal A is steel and metal B is

copper. (after reference [1 1 I, 1 12])

170

 

Figure 37. Finite element modeling of indentation of an elastic-plastic half space by
a rigid sphere: development of the plastic zone. Broken line: contour of J2. Py is the
normal load needed to first initiate plastic flow. (after reference [114])

I7]

Before Heating Afier Heating
9“'-‘ r' _‘ A

518 N

I‘_‘ "
if. _O‘ I.
,2 t." .‘ 4.1M

      
    

 

Figure 38. Recovery of spherical indents in NiTi alloy by the shape memory effect.
(afier reference [231] )

 

  

 

 

 

 

 
   
 

 

 

 

12
1 .
E 0.8 ‘
I: 0.6
\ Superelastic
Ll.
0.4 ‘ "m
0.2
o
o 02 o 4 0.6 0.3 1 1 2
h/h max
(a)
1.2
1 .
x 0.3 4
E
Superelastic
& 06 m
LL
0.‘
02 . Cu
0 4 , . -
0 0.2 0.4 0.8 0.8 1 1.2
Mb max (b)

Figure 39. Load-displacement curves of (a) Berkovich and (b) spherical indentations
in SE NiTi. (afier reference [11 7] )

172

 

Figure 40. TEM images from the indent-affected zone in 50.9at%NiTi single crystals.
Images (a) to (c) were taken adjacent to the same indent at difiermt magnifications.
(after reference [1 18]).

I73

 

Figure 41. Images of wear marks on NiTi afier 200 cycles of scanning scratch under
different normal loads and temperatures. A fiS 16 °C. (afier reference [144] )

I74

 

-0.80

 

 

 

-0.40

Stress c’y at the fllm-matrlxlnterface

 

 

   

4.00 4.00

 

 

-0.25

 

 

 

 

r
-0.13

 

 

 

 

0.00

0.13

 

Stress Txy at the film-matrix Interface

 

 

0.25 ---i...i...i-.-
4.00 -2.00 0.00 2.00 4.00

Distance from normal line load applied at x=0.0

 

Figure 42. Calculated stresses due to an isolated unit normal line load applied at the
origin as a function of substrate modulus E2. Variation of bearing stress 0,, and shear
stress 1,, at the film-substrate interface are shown for a 1000 nm film with a reference
elastic modulus E1=1. (afier reference [198])

I75

 

 

 

 

P
8

 

 

Stress “x atthe fllm-matrlxlnterface

 

 

 

 

 

 

I

-

8 'vvv'vviv
.

—o— £2=29
+ Ef=50
1"” . . l . L g . . l . . 4.
-2.00 0.00 2.00 4,00
Distance from normal line load applied at x=0.0
0.20 Substrate Modulus E2=0.2

 

 

 

 

 

 

 

 

 

Stress “x at the film-matrix interface

 

 

050 J r . i - . . i - . . i , - .
4.00 -2.00 0.00 2.00 4.00
Distance from normal line load applied at x=0.0

 

 

 

Figure 43. Calculated tensile stress a, at the film-substrate interface for a 1000 nm
film with a reference elastic modulus E1=l for a unit normal line load at the origin.
Substrate modulus E2 is varied in the upper figure. The effect of varying the film
thickness on the flexure stress is shown in the lower figure for a substrate of modulus
E2=0.2. (afier reference [198])

176

‘|.OE+07 -

 

 

l
CrN1-Al ’
A 1.0E+06 i
z_ I
g (I .0E+05 CTN-32'“ 5
\ l
"E 1 .0E+O4
5
c - .
8 1 .0E+03 m 31 A'
2
g 1 .OE+02
3
'l .OE+O‘I
’l .OE+00 ,, " fl *
Figure 44. Wear test results of samples withl um CrN coating and with: (from lefi to

right) no interlayer, with 4 um SE NiTi interlayer (S1), 4 um Cr interlayer and 4pm SMA
NiTi interlayer (S2). (afier reference [3] )

l77

(c) with

-—‘

 

Figure 45. SEM Images of scratch track end with (3) 1pm CrN coating, no interlayer,
(b) Sum CrN coating no interlayer, (c) 1pm CrN coating 4pm SE NiTi interlayer, (d) 1pm
CrN coating and 4pm SMA NiTi interlayer and (e) 1pm CrN coating and 4pm Cr
interlayer. (after reference [3])

I78

  
 
 

 

transformed M f

gig/141 ////// i

 

 

 

 

 

 

 

Figure 46. Vogit (parallel) model for stress induced martensite in austenite phase.
(after reference [204] )

     

I

I /r
/ tmnsfonnedM .
j////////// '
VKKKKKKKKKKK
Ill IIIIIIIIII/i
\\\\ \\\\\\\\\\\
III III Ill/l1
\\\\\\ s\\\\\

Ill/ll III/Ii
\\\\\\ s\\\\\

 

I;////////////.
\\\\\\\\\\~.\\\\
Il/./_/_/./_/./£/_/,/!/’/_/..

‘5
\
\
\
\
\
N.
\
\

 

A

 

 

Figure 47. Series model for stress induced martensite in austenite phase. (after
reference [204] )

I79

A
0)
v

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

U)
.fi 300K
:—l
o
o
U 35 OK
0 ii
C‘.
'3 32 BK
r0
(0
:1:
3 0 OK
38 40 42 44 46
2theta
Figure 48. X ray diffi'action patterns of NiTi thin films (sample #1) at various
temperatures
(b)
433aFe
2 1 8 K
0 L A
6‘ 112R 300R
.5 2 63K
:—i
o
0
U
3 0 3K
or
.5 286K
4.) A
8 2 2 3K
:13
38 40 42 44 46
2theta
Figure 49. X ray diffraction patterns of NiTi thin films (sample #2) at various

temperatures sample

180

NlTizAl A,
W

 

0)
g
300R (.8)
328K
1 10A
349K
.8
AI ‘8
020M 1-1 1M 298K I

 

 

f

3’8'4'0'4'2'4'4'4'6'4'8
Ztheta

Figure 50. Temperature scanning X-ray diffraction of NiTi thin film on 6061T6
substrates, from which phase transformation temperatures were estimated and listed in
TABLE. CrN coating was later deposited onto the NiTi thin fihns.

NiTiCuzAl
318K
323K W
338K WWW

Al 110A Al

Cooling

 

348K

331K MM
002M 012M
298K
' ' ' 4'0

4'0 42 44 45
2 theta

Temperature scanning X-ray diffraction of NiTiCu thin fihn on 6061T6
CrN coatings were later deposited on the NiTi thin films

Heating

 

Figure 51.
substrates.

181

15 um NiTizAl

ii 300K
338K A a.) W
110A
378K I l A
ll

111M

020 M ‘

T l

38 40 42 44 46 48

2 9
Temperature scanning X—ray diffraction of NiTi thin film on 6061 T6
The NiTi thin films were interlayers of DLC coatings

Cooling

 

Heating

 

Figure 52.
aluminum substrates.

 

 

0.9 - — 300K
— 353K
—— 433K
——473K
x 0.6 ‘
(U
E
E
D.
0.3 ~
0.0 r r r l l
0.0 0.2 0.4 0.6 0.8 1.0

h/hmax

Figure 53. The normalized load-depth curves fi'om Berkovich indentation of the
NiTi thin films at different temperatures.

182

  
 

3.175mm WC Ball

NiTi alloy

 

Figure 54. Diagram of thermomechanical training method to create two-way indent
depth change.

Plastic

6004 x4
. , Uniaxial l

BlaXIal '

480 - i
360- ...........

240 4

von Mises Stress (MPa)

120~

 

 

0 1' . . . T
0.000 0.015 0.030 0.045 0.060

Strain
Figure 55. Stress-strain curves of simple loading conditions from the 3D constitutive
model for superelastic effect with plasticity. For the biaxial loading curve, the strain

rate is set the same in two loading directions. The dashed curve represents a high
uniaxial load which causes plastic hardening and permanent plastic strain.

183

 

Figure 56. Finite element mesh of indentation model. The refined region is the
contact region. Total node number is around 15000.

(a)

   

6000
1 — SE Nrri
----- Reference
2
3
E
o 3000 q
_I

15004

 

 

 

0.00 0.03 0.06 0.09
Depth (um)

— SE NITI
----- Reference

max)

04*

0.2 -

Normalized Load (F/F

 

 

‘l I I A I I I
0 0.2 0.4 0.6 0.8 1.0

Normalized Depth (dldm)

 

Figure 57. Absolute (a) and normalized (b) load-displacement curves of SE NiTi and
reference materials from finite element modeling. Indentation-induced superelasticity
enables large or 100% indent depth recovery.

185

 

 

 

 

 

f; 804 ‘

a. | \.

“9" 60

§ 1 Reference

2‘ - 4 m

g 40« 2

O \I —-——-> g

8 A ‘ DC

........ e ._ 2

0: 20- A , , ‘6
Transformed reglon suze g
(in contact radius) 5

0 . . . 0 o
0 1 2 3 4

Representative Strain (%)

Figure 58. Recovery ratios vs. representative strains at different indentation depths,
and the size of indentation-induced phase transformation region at different indentation
depths, measured in contact radius.

 

 

A 1-00‘ FE Modeling
u-é - Experiment
0; 0.75]

'U

(U

3

.0 0.50- A=30K
(D

.5

E

E 0.25

O

z

 

 

0.00 0.25 0.50 0.75 1.00
Normalized Depth (d/dmax)

Figure 59. Comparison of experimental and calculated load—displacement curves.

The experimental results are adapted from Ni et alf‘ Modeling SE NiTi has a phase

transformation hysteresis A (A= (A f- As)/2-(M, - Mf)/2 ) 15K and 30K for the two
computed curves.

186

 

 

 

 

 

 

z/a

 

 

Figure 60. Contours of the austenite phase volume fraction at indentation depth of
0.018um (a), and complete indent recovery after unloading (b), where few residual
martensites are present after unloading.

187

 

 

 

 

 

 

 

 

Figure 61. Contours of the austenite phase volume fraction (a) and plastic strains (b)
at indentation depth of 0.10pm. (c) Contours of the austenite volume fraction afier
unloading. For the partially recovered indent (recovery ratio 71%), there were much
more residual martensites than in FIG. 5(b).

I88

Compressive (Negative)

—- Cycle 1
—— Cycle 2
—— Cycle 3

 
 
  

 

 

100 200 300
Temperature (C)

01

Figure 62. Residual stress evolution of NiTi thin fihn on Al substrates with
temperature, for three heating-cooling cycles.

 

 

0 100 200 300
Temperature (C)

Figure 63. First cycle residual stress evolution of NiTi thin film on A] substrates
with temperature.

189

 

 

 

 

 

l .“p
50 MP3 "I .I "" 'w

l 1 I

70 80 90
Temperature (C)

 

 

Figure 64. The stress/temperature rate of NiTi thin film on Al substrates during
martensite phase transformation.

 

 

 

 

/
" 11 l‘
30 MPa ,“i'
1"”
a “|lfl
l 1’ do/dT=1.6 MPa/K
. I"
l 11
210 260 200

Temperature (C)

Figure 65. The stress/temperature rate of NiTi thin film on Al substrates from CTE
mismatch between thin film and substrates, without the influence of phase
transformation,

190

Z'SDLC:'2NiTim:Al

 

 
     
 
  

 

 

 

 

12- n
E 9‘ Austenite 375K
1 I
.g 64
I
31 Martensite,303K,
R: 15.94m
0 I l I l I
O 5 10 15 20
Width (mm)
Figure 66. Curvature change of Z'SDLC112N1T13632A1 sample between room

temperature around M; and elevated temperature above A f. The curvature change mainly
resulted from residual stress change in NiTi interlayer.

191

100
3"CrN:"NiTi353:Al

_/

    

Deformation (umz)

—I— CrN:NlTizAl, 0.98N

 

0-1 —o— CrN:NlTizAl, 0.49N o
-- l:l-- CerAl, 0.98N
~ O-- CerAI, 0.49N
0.01 - I I ' ' '
300 320 340 360 330
Temperature (K)

Figure 67. Temperature scanning weer deformation of CrN:NiTi:A1 and CrN:Al
samples, under 0.49 and 0.98 N load, using a 200 um tip radius spherical indenter.
Superelastic NiTi improves the wear resistance by limiting plastic deformation to the sofi
aluminum substrates.

192

  
    

 

298K 333K 353K 378K
Figure 68. SEM images of temperature wear tracks of 3‘9CerAl samples (top row)

and 3'S’CerlsNiTimzAl sample (bottom row), under 0.98N load. The wear data was
presented in Figure 67

298K 378K
Figure 69. SEM images of temperature scanning wear tracks of 3“’CerA1 samples

(first row) and 3‘9CrN:'5NiTi353:Al sample (second row), under 0.49N. The wear data
was presented in Figure 67.

193

     

._ - _‘. _»~:-"‘ _. . , 720011111.
298 K 338 K 358 K

Figure 70. SEM images of temperature scanning wear tracks of l"'CrN:$’NiTiCu;i43:Al
samples. The wear data was presented in Figure 71.

1o_o, 200 um tip radius conical indenter, 0.49N Load

 

 

[j ........................... D .............. El """""
NA ‘ 7cmuu
E
3
.5 '
‘5 1.0 q \
E I l
23 ‘-’cm:°NITICu,,,:AI\.
O
a
a)
3
0.1 . 'Y'T‘I‘s “iii“ .
300 320 340 360
Temperature (K)
Figure 71. Temperature scanning wear deformation of l‘7Cer9NiTiCU343zAl and

1'7CrN2Al samples, under 0.49 and 0.98 N load.

194

8.0x10“), ”cum:

3\
j\

 
 
  

0’1

1

‘-’cm:’Niri,,3:AI

H

1 “cl-nun

.U‘ .‘l
0 U1
X X
_L _x
O O
or

4

Wear Loss (nm2)

2.5x105- “CrN:NiTimml

 

0.0 l l I i l
0 4 8 12 16

NiTi Interlayer Thickness (um)

 

Figure 72. The effect of interlayer and CrN coating thickness on wear resistance at
room temperature. The wear test was done using 6.3Smm diameter WC under 5N load.
Thicker coating and interlayer gave better wear resistance.

195

6.35mm WC Ball, 0.98N Load

 

 

 

 

1051 'g T' '
.9 g ”CrN:AI
E I
5 .
3 10“:
o :
_l . .
a .
g 3 l o/.* . "/
1° 1 ‘-7cm:“Nirim:Al
300 330 360 330
Test Temperature(K)
Figure 73. Wear loss vs. temperature using a different loading condition (6.35mm

ball, 0.98N load). Wear resistance was not improved at temperature above

200 pm radius

4000 ~

  

0.98N, 3498MPa

 
 
 
 

3000 -

0.49N, 2776MPa
2000 -

3.175 mm radius

0.98N, 554MPa

0.0 0.6 112 1.8
Load (N)

1000~

Elastic Contact Stress (MPa)

 

 

 

Figure 74. Elastic contact stress in elastic halfspace from 3.175mm radius and
200um radius sphere under different loads.

196

 

 

1 0 1
1 2-5I)Lc:Al
”E 1 _ "Killi; 1:; 1:" """"""" '
:1. 5 O \\\\\
c i I
Q 4 3531-09:'FT‘i;c.-_.:.-5.l
E . .15
E 0 1 _ \ :\'DLC. NiTi333. Al
O o
“5 \\o
D \e
g 0.01 -.
g ' 2-50L0:2‘NITI3,,;AI
1E-3 . . . . ' .
300 320 340 360 A f 380
Temperature (K)
Figure 75. Temperature scanning wear tests of 2‘SDLC:NiTi363:Al samples, using

200um tip radius diamond indenter under 0.49N load. Thick NiTi interlayer and SE
NiTi interlayer at temperature above A I showed better wear.

197

      

2-5DLc:Al 2'5DLC:AI
298K, 50 turns 378K, 50 turns

        

2"DLC:“NiTi3s:Al 2"DLC:‘5NiTim:AI
298K, 100 turns 378K, 100 turns

200 um

   

2"DLC:2‘NiTim:AI 2"DLC:21NiTim:AI
298K, 200 turns 378K, 2000 turns

Figure 76. SEM images of temperature scanning wear tracks of 2‘SDLCzNiTi353zAl
and 2'SDLCzAl samples. The wear data was presented in Figure 88.

198

1.2x106-

 

 

o
9.0x1051 CrN:6NiTi31o:MZ
NA ‘
E. ""
g 6.0x105 ~
3 o
a CrN:12NiTi31o:M2 A
é’ 3.0x105 -
D
CerMZ 0
0.0 . ”'1' """""" D"; . .
0 1 2 3 4 5
CrN Thickness (um)
Figure 77. The wear resistance of CrN:NiTi:M2 samples at room temperature as a

function of CrN coating thickness.

1 .2x106 -

1.7 .
~D--- CrN-M2 ”CrN:NiTiszZ
WA 2'BCrN:hfl2

9.0x1054 --~o--~"‘"CrN:M2

 
   

"3»

2.8 . - - ,
CTN.NIT|31O.M2

6.0x105‘
A
.0

5 3'gCerNiTisz2
3.0x10 -
$.91. .
0.0 '

I I I I I I

o 2 4 6 8101T2

NiTi Interlayer Thickness (um)

Figure 78. The wear resistance of CrN:NiTi:M2 samples at room temperature as a
function of interlayer thickness.

 

Wear Loss (nmz)

 

353K

Figure 79. SEM images of temperature scanning scratch tracks on 3'9CrN:9NiTi353:Al.
Images of the whole scratch tracks (1 mm in length )are placed at the upper right corner.

      

, - 50 run

306 K 331K 353K 375 K 398 K

     

Figure 80. SEM images of temperature scanning scratches tracks on 3'9CerAl.
Images of the whole scratch tracks (1 mm in length )are placed at the upper right corner.

200

298K
333K JV 1 1111'

91% L 1”

403K ”1

1' i 3 4 :7
Normal Load (N)

Figure 81. Acoustic emission of scratch tests on 3‘S’Cer9NiTi3s3zAl. It showed
strong temperature dependence that matched phase transformation temperatures of NiTi

interlayers
ex 101
w
W

1' i 3 4 6
Normal Load (N)

Acoustic Emissions

 

 

Figure 82. Acoustic emission of scratch tests on 3'9CerAl. There is no significant
temperature dependence.

201

  
  

 
 

 

Figure 83. EDS spectrum of the delarninated area in scratch tracks on
3'9C1'N39N1T13531A1, showing that delamination occurred at the interface between CrN and
NiTi interlayer.

Friends-13 cps Canoe: cerium ._

Figure 84. EDS spectrum of the delarninated area in scratch tracks on 3‘9CerAl,
showing the delamination occurred at the interface between CrN and the Cr interlayer.
The Cr interlayer (<um) was deposited before CrN as a precursor layer. It is an
indispensable process of making CrN.

202

 

 

A\
4 4 \A—A
A 3'9CrN:°NiTi353:AI
3, 3 _
.8 IR '\
3 2 _ '\n
a A ”CrN:AI
.2
'5 1 .
3520 360 400 440 430
Temperatue (K)
Figure 85. Critical load of CrN coating in scratch tests as a function of temperature.

The superelasticity of NiTi interlayer improved the scratch resistance of CrN coating.

 

Figure 86. SEM images of scratch track on 2'SDLCzAl sample, showing severe DLC
coating failure and delamination. 7N progressive load was used.

203

v

A .. ‘

f; ‘
—

50 run

 

Figure 87. SEM images of scratch track on 2"SDLCzlsNiTimzAl sample, showing
much less coating failure compared to the previous figure without NiTi interlayer. 7N
progressive load.

204

'17.
'1'

31'
1"

fr???

\ep ..,

Figure 88. SEM images of scratch track on 2'SDLCzlsNiTimzAl sample, under high
30N progressive load. Although there was coating failure, no dramatic delamination
occurred, proving the strong interface between DLC coating and NiTi interlayer.

 

205

Friends-.13 ape

       

 

’ “ .1 l I Q r 13:; W 06- 4 3w,

Figure 89. EDS of delarninated region in scratch tracks on 2'SDLC:'5NiTi;i63:Al
sample, shows the delamination occurred at DLC and (Cr)NiTi interface. The extra Cr
layer on NiTi is represented by the high Cr peak here, compared to the weak Cr peak in
Figure 83.

   

/i~ 7111

 

 

1 .00 -

0.75 -
2 3.9 .15 - -.
; CI'N. NITI.AI 3'9CrN:Al
8 0.50 -
_l

0.25 -

0.00 . . . ,

0 1 2
Depth (pm)
Figure 90. Computational load-displacement curves fiom finite element modeling

of spherical indentation on 3‘9CerlsNiTizA1 and 3'S’CrN:A1, using 200 pm tip radius
spherical diamond indenter.

206

 

Al
Substrate

 

4
I _ .L

Unloaded

: I . 1
Al
Substrate

Figure 91. Horizontal strain component 8,, in 3‘S’CerlsNiTizAl and 3'S’CerAl from
FEA

 

 

 

 

207

CIN

.;.-2 ,._.. ..
- I
~ A . want: H

Al
Substrate

 

 

 

     
 

Loaded Unloaded

Al
Substrate

   

Figure 92.

Vertical strain component a” in 3 '9CrN215NiTi2Al and 3'9CerAl from
FEA

208

 

Substrate

       

Unloaded

Al
Substrate

Figure 93. Shear strain component 6,, in 3'9CrN:'5\NiTi:Al and 3'9CerAl from FEA

 

 

 

A! -
Loaded

     

 

 

 

-r~nar‘iCSItc

Loaded Unloaded

 

 

Figure 94. Martensite volume fiaction in NiTi interlayer from FEA. The blank
region represent the coating and substrate because there is no martensitic phase.

209

 

Al
Substrate

 

 

 

AI
Substrate

   

Figure 95. Horizontal stress component on in 3'S’Cer'sNiTi:Al and 3'S’CerAl from
FEA

210

    

Al
Substrate

 

 

 

Figure 96. Vertical stress component a” in 3‘9CerlsNiTizAl and 3‘9CerAl fiom
FEA

211

 

Al
Substrate

 

 

Al
Substrate

 

 

Figure 97. Shear stress component 6,, in CrN:NiTi:A1 and CerAl from FEA

212

Heated
(Austenite)

Cooled
(Martensite)

 

 

(b) Sectional Profile

 

 

 

 

 

 

0 _ Heated (Austenite)

g 404 / DA
.2
‘5.
0.)
E -80~ \ —
5 ‘1
3 DM

_120 ] Cooled (Martensite)

0:0 0:7 1:4 2:1 2:8

Horizontal Direction (mm)

Figure 98. (a) 3D profile of a heated and cooled indent, (b) Sectional profile of the
heated and cooled indent.

213

Indent Depth (um)

Figure 99.

Indent Depth Change (um)

Figure 100.

 

 

 

-30.
-60 fi D A ,Heated (Austenite)
. ____.. + ——.———-O
-90 - g
0) O
-120 4 if: 8
a: __———I————I-—""
-150 .1 ‘7: '/DM, Cooled (Martensite)
-180 4 'lndent depth before heating

01-1
09.1

0 1 2 3 4
Heating-Cooling Cycles
Two-way indent depth change over several thermal cycles.

 

 

 

 

 

 

 

0‘ - 1.0
20- - 0.8
404 (DM-DAflDM, Indent Depth Change Ratio' 0'6

‘ A A. A. A
4 0.4
604
D A-DM, Indent Depth Change
v/' v V V . 0.2
804
I ' ' i I 0.0
0 1 2 3 4 5 5

Heating-Cooling Cycles

Indent depth change ratio and absolute indent depth change.

214

Indent Depth Change Ratio

(DM-DA)/DM, Indent Depth Change Ratio

Indent Depth Change Ratio
0
9°

 

0.4~

 

 

0:0 0:7 1:4 2:1 2:8
Horizontal Direction (mm)

Figure 101. Indent depth change ratio (DM-DA)/DM vs. horizontal direction, referred
to center. The center of indent had smallest ratio of depth change. Further away fi‘om
the indent, both DM and DA become small and dominant by the surface roughness,
causing the fluctuation in the ratio. Ideally the (Did-DAVDM should approach zero for
surfaces far away from indents because (DM—DA) should equal zero. But DM—DA is not
zero because: 1) small surface roughness change in NiTi between martensitic and
austenitic phase, 2) measurement error between two measurements.

215

(a) “Indenter (d) Polishing

   

 

High Temp ()e Low Temp
F}. _1'1
+111 , m

 

 

 

 

 

(c) @3115? LOW Temp f() in High Temp

[:3 Undeformed Martensite

E] Deformed Martensite

D Austenite

_L Dislocation
Figure 102. Schematic diagrams of experimental procedures and mechanisms for
achieving RSP. (a), indentation is accommodated by the deformed martensite and
slip-plasticity. (b), When heated to the austenite phase, the indent is partially recovered.
(c), when cooled down to martensite phase, the residual indent became deeper. The
indent depth change is reversible with heating and cooling temperatures (Figure 2). The
dislocation structure and its associated stress field can facilitate the grth of oriented
martensite variants during austenite-to-martensite transformation, leading to two-way
indent depth change. (d), The residual indents are removed by polishing. (e), The
oriented martensite region and dislocations can still exist afier the indent is removed. (f),
A reversible surface protrusion is formed upon heating.

216

   
 

 

 

 

 

 

 

 

E -20 a :3:
.2 7 — C4
2 _ H5
.5 — cs
‘5. 4
0)
D _40 _

‘ Cycle 5

“ Cycle 1 l

-60 T r I I I I I I I I I I
O 200 400 600
Width, Microns

Figure 103. Reversible indent depth change. A spherical indent made with a steel
ball pressed relatively deeply into a martensitic sample of NiTi not only recovers a
substantial percentage of the indent depth when heated above the Af temperature, but also
subsequently displays cyclic displacement behavior as the temperature is cycled from
below Mfto above A f.

217

Austenite
(heated)

5.0mm

Martensite
(cooled)

 

 

Figure 104. Thermally-reversible surface protrusions. A 3x3 matrix of
thermally-reversible protrusions produced in a N150,3Ti49_7 alloy by first creating deep
spherical indentations and then planarizing the specimen while it was still in the
martensitic condition. The protrusions appear and disappear as the specimen is heated
fiom below Mf to above the A f temperature.

218

Peak Height over 5 Thermal Cycles

 

 

 

 

 

16 1 ]
I I/. . I
121
E ‘— Heated
3 8 - I
E 1 (
U)
.5 1|
I 4 -‘ ’w
0‘ 1 Cooled
0:0 0:3 1:6 2:4
Width (mm)

Figure 105. Sectional profiles of reversible circular surface protrusions and peak
height over five thermal cycles.

100 um ]

 

 

 

0.5 mm

Figure 106. Actual shape of the circular surface protrusion

219

V 1,1,... . _ ‘ Trace = Form + Primary

  

‘s. 1: : 1' .

<‘ i ' ‘ 1 ..

[M H 1 ‘1 IV 1 ‘ .
l ‘ ’1. l I. u . l1-
1 11.1 III“: "MW 1111'” 1]

Primary= Form Error 4 Texture

1; \r '1 1 11‘
1 ""1" 1’ ’1 ‘ 111‘”.

l , 1:1 I I I. 1 i .
, I J ("1 ' “‘ ,1
1 1‘ 1 r
* H r [1,5]} 1 Form Error In. | ‘ ' ‘
1‘1: .. 1w l .1: ,I-x‘
Texture Waviness + Roughness
1‘ ’
‘ l I . ' I
I. ‘1 .1" l! ‘ 1 J. in ‘ ‘0 ..
III I' I! “ll I 111‘?!’ ‘ Ill “'13“ ll'ifl .11 .11" "‘11: 1 .11 ‘11)I 111‘ l“
’y i " T ’ I ‘
Waviness

  
 

1111111111 1111101101101 1111:

Figure 107. Schematic of different kind of surface fluctuations. Image fiom
http://ru \x'u.prcdcv.00m”smu/introlltm. More detail in reference [246].

 

Martensite Austenite Martensite
(Cooled) (Heated) a (Cooled)

Figure 108. Reversible GM logo made from adjoining reversible surface protrusions.
Imaged in normal bright field optical microscopy.

220

Austenite
(heated)

600 pm

Martensite
(cooled)

Figure 109.

Austenite
(heated

Martensite
(cooled)

 

Figure l 10. 3D profiles of nanoscale circular surface protrusions on NiTi thin
film/ Silicon substrate.

Austenite

 

 

   
 

 

 

 

(heated)
» 84.6
g - 74.5
_ A 64.5
‘ J 126.9 pm ‘
62.6 Hm I/ ’4, : 545
- 44.5
Martensite :
(cooled) _ 34.5
k - 24.5
- 14.5
i 0.0
Figure 111. 3D profiles of reversible nano scale line protrusion on NiTi thin
film/Silicon substrate
481 Peak Height over 5 thermal cycles
l
e '\
5 24‘
E Cooled
E 12‘ — Heated
>
0.1
-12 , r f u
0 10 20 30 40
Horizontal (pm)
Figure 112. Sectional profiles of the nano circular surface protrusions and peak

height over five thermal cycles.

222

_a
N

 

 

 

4\ ’’’’ [ ho
p1. _ _. #15111:sz 1111 ..

--- - —‘

 

 

 

Protrusion amplitude, hp ,microns

 

 

 

O f I ' I T l
O 200 400
Polished layer depth, dpl, microns

V I

Figure 113. Decay of protrusion amplitude with surface-removal. When successive
surface layers of thickness dpl are removed afier initial planarization to a depth ho, the
amplitude of protrusion appearance and disappearance, hp, diminishes and approaches
zero when material has been removed to a depth roughly equal to the radius of the initial
indent contact.

223

 

 

 

 

 

 

 

 

—l

 

l 1/U *k‘9
9% '

(elastic-plastic O
boulndary) \
l ‘K
__ db/D = 0.51

I I

l l

 

Depth, y/D

N
/

 

 

 

 

 

 

 

 

2 l O
Diameter,x/D

Figure 114. Plastic strain distribution under an indent. Theoretical and

experimental results , adapted from Samuel and Mulhearn [232] showing the expected
magnitude of plastic strain under an indent in a ductile metal.

224

 

ATp = 62

 

AM‘ = “6 AMs = +34

_3 l

Cooling JEN

O / :

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

In
.t’
C
D
.d
h
<
_d
4.:
Q.
a) I
CD I
c ' _
.9 : ATn — 28 I
H I I
(U I |
H I
C I
a) i
“o l K
E
~ l
O /O/O/
U \“1 Heating
m
G _
\ I
I
I
< I
AAS = -24 AAf = .4
0 50 100 l 50
Temperature, °C
Figure 115. Height of reversible circular protrusions vs. temperature. The phase

transformation temperature of NiTi sample is also marked. The hysteresis of the
reversible protrusion is much narrower.

225

Austenite

 

llll

Martensite 7

llll

Residual stress

Figure 116. Schematic of Residual Stress in Reversible Surface Protrusion. The
residual stress is in the opposite direction of TWSME during heating from martensite to
austenite phase, and not in the opposite direction during cooling, narrowing hysteresis.
The residual stress also stabilizes the TWSME.

 

 

 

226

Geometrically Necessary Dislocations

 

 

Figure 1 17. Geometrically necessary dislocations. It is assumed that the indentation
is accommodated by circular loops of geometrically necessary dislocations with Burgers
vectors normal to the plane of the surface. As the indenter is forced into the surface of a
single crystal, geometrically necessary dislocations are required to account for the
permanent shape change at the surface.

227

 

Figure 118. Residual stress in y direction (ayy) afier spherical indentation in

austenitic NiTi alloy, by finite element analysis.
direction.

228

The negative sign is compressive

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"Surface Texture, Surface Roughness, Waviness and Lay:ASME B4 6. I -2002 "
(ASME, 2003)

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