development, characterization and testing of nickel
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Development, Characterization and Testing of
Nickel Titanium Based High Temperature Shape
Memory Alloys
By
Saif ur Rahman
Supervised By
Dr. Mushtaq Khan
School of Mechanical and Manufacturing Engineering (SMME)
National University of Sciences and Technology (NUST)
(2016)
ii
Development, Characterization and Testing of
Nickel Titanium Based High Temperature Shape
Memory Alloys
Saif ur Rahman
2011-NUST-DirPhD-ME-45
This work is submitted as a PhD thesis in partial fulfillment of the
requirement for the degree of
PhD in Mechanical Engineering
Supervisor: Dr. Mushtaq Khan
School of Mechanical and Manufacturing Engineering (SMME)
National University of Sciences and Technology (NUST)
Sector H-12 Islamabad, Pakistan
(2016)
iii
Certificate
This is to certify that the work in this thesis has been carried out by Mr. Saif ur Rahman and
completed under my supervision in the School of Mechanical and Manufacturing
Engineering (SMME), National University of Sciences and Technology (NUST), H-12
Islamabad, Pakistan.
Supervisor: ______________
Assistant Prof. Dr. Mushtaq Khan
School of Mechanical and Manufacturing Engineering
Submitted through
Principal/Dean
School of Mechanical and Manufacturing Engineering (SMME)
National University of Sciences and Technology (NUST), Islamabad
iv
Dedication
To my respected and loving mother
Farosha
v
Acknowledgments
In the Name of Allah, the Most Beneficent, the Most Merciful. All the praises and
thanks be to Allah, the lord of the entire Universe.
I would like to extend my gratitude to my supervisor, Dr. Mushtaq Khan for his
guidance and encouragement throughout my research. I would also like to thank him
for pushing me to complete my thesis in time. Thank you Mushtaq Khan, It is great
to have my supervisor.
Dr. Aamir Nusair Khan, my co-supervisor is acknowledged for his contribution,
help and support in alloys development and characterization. Thanks Aamir Nusair
Khan for your guidance, unlimited help and support in the completion of this thesis.
Thanks to my Guidance and Evaluation Committee (GEC) members Dr. Liaqat Ali
and Dr. Syed Hussain Imran Jaffery for their guidance and fruitful discussion. I
would like to extend my gratitude especially to external GEC member Dr. Riaz
Mohammad for selection of the current research area. His continuous guidance and
encouragement enabled me to complete my thesis successfully.
I would like to extend my gratitude to my DG Syed Tahir Hassan Hashmi and
Director Syed Muzaffar Ali for their support and encouragement throughout my
study duration.
Thanks to technical staff Tahir Mehmood Khan, Liaqat Ali, Muhammad Israr
and Muhammad Waseem for their efforts in alloys development and
characterization.
I would also like to thank my PhD colleague Malik Mansoor Muhammad for his
valuable input and assistance.
Of course I cannot forget my parents, my loving parents who have brought me up to
this stage. Without their prayers, love and endless support, I would not have
accomplished this task. Thank you always for your love.
vi
At the end, thanks to my dearest wife Rashida and loving children Shumila Saif,
Somayya Saif, Moneeba Saif, Waqar Ahmad Saif and Fatima Saif. Words are not
enough to express my gratitude for your love and patience.
Saif ur Rahman
July 2016
vii
Abstract
TiNi-based shape memory alloys are well known for their excellent shape memory
and superelastic properties. TiNiPd alloys are considered as the better high
temperature shape memory alloys due to high transformation temperatures, small
hysteresis, reasonable strain recovery and comparable workability. However, by
further increasing the transformation temperatures i.e. by increasing the Pd content,
thermal hysteresis also increases. This has an adverse effect on the actuation
behavior of the alloy. At high temperature the critical stress for slip deformation of
TiNiPd alloys also decreases, which increases the permanent deformation and
reduces the strain recovery in the alloy. In order to prevent increase in thermal
hysteresis, reduce permanent deformation and increase strain recovery in TiNiPd
alloy, Ni has been replaced by 5 at%, 10 at% and 15 at% Cu. Four alloys;
Ti50Ni25Pd25, Ti50Ni20Pd25Cu5, Ti50Ni15Pd25Cu10 and Ti50Ni10Pd25Cu15 (all in atomic
%) have been developed and characterized for their microstructure, phase
transformation temperatures, mechanical and shape memory properties in solution
treated condition. By increasing the Cu content, the transformation temperature of
the TiNiPdCu alloys significantly increased, whereas thermal hysteresis decreased.
Similarly, the microhardness, yield and fracture strength also increased. Shape
memory properties like strain recovery and work output also improved. Therefore,
TiNiPdCu alloys showed improved transformation temperatures, strain recovery and
critical stress for slip deformation through solid solution strengthening mechanism.
The TiNiPdCu alloys were also aged at different aging temperatures i.e. 400°C,
500°C, 600°C and 700°C for 3 hours to investigate their transformation
temperatures, mechanical and shape memory properties and compared with the
solution treated samples. By aging the Ti50Ni25Pd25 and Ti50Ni20Pd25Cu5 alloys, the
transformation temperatures, mechanical and shape memory properties slightly
increased. After aging the Ti50Ni15Pd25Cu10 and Ti50Ni10Pd25Cu15 alloys, the
transformation temperatures and shape memory properties significantly decreased,
however the mechanical properties were improved.
viii
It can be concluded that aging of Ti50Ni25Pd25 and Ti50Ni20Pd25Cu5 alloys is
beneficial to increase their transformation temperatures and shape memory
properties. However it has an adverse effect in terms of transformation temperatures
and strain recovery by aging the Ti50Ni15Pd25Cu10 and Ti50Ni10Pd25Cu15 alloys.
ix
Publications
1. Saif ur Rehman, Mushtaq Khan, Aamer Nusair Khan, Syed Husain Imran
Jaffery, Liaqat Ali, and Aamir Mubashar (2015), Improvement in the Mechanical
Properties of High Temperature Shape Memory Alloy (Ti50Ni25Pd25) by Copper (Cu)
Addition, Advances in Materials Science and Engineering [Article in Press]
2. Saif ur Rehman, Mushtaq Khan, Syed Husain Imran Jaffery, Liaqat Ali,
(2015), Effect of Aging on Phase Transition Behavior of Ti50Ni15Pd25Cu10 High
Temperature Shape Memory Alloys, Advanced Materials Research, Vol. 1101, pp
177-180.
3. Saif ur Rehman, Mushtaq Khan, Aamer Nusair Khan, Liaqat Ali, Syed
Husain Imran Jaffery (2015), Two-step Martensitic Transformation in an Aged
Ti50Ni15Pd25Cu10 High Temperature Shape Memory Alloys, Acta Physica Polonica
A, Vol. 128, pp B-125 – 127.
4. Saif ur Rehman, Mushtaq Khan, A. Nusair Khan, Liaqat Ali, Sabah Zaman,
Muhammad Waseem, Liaqat Ali, Syed Husain Imran Jaffery (2014), Transformation
Behavior and Shape Memory Properties of High Temperature Shape Memory Alloy
Ti50Ni15Pd25Cu10 at Different Aging Temperatures, Materials Science and
Engineering - A, Vol. 619, pp. 171-179.
5. Saif ur Rehman, Mushtaq Khan, A. Nusair Khan, M. Imran Khan, Liaqat Ali,
Syed Husain Imran Jaffery (2014), Effect of Precipitation Hardening and
Thermomechanical Training on Microstructure and Shape Memory Properties of
Ti50Ni15Pd25Cu10 High Temperature Shape Memory Alloys, Journal of Alloys and
Compounds, Vol. 616, pp. 275-283.
x
Conference Presentations
1. Saif ur Rehman, Mushtaq Khan, Liaqat Ali Syed Husain Imran Jaffery, Aamir
Mubashar (2015), Effect of Cu Addition on Microstructure and Transformation
Temperatures of Ti25Ni25Pd25 High Temperature Shape Memory Alloys,
International Conference on Mechanical, Aeronautics and Production Engineering
(ICMAPE) London, 20th April, United Kingdom.
2. Saif ur Rehman, Mushtaq Khan, Syed Husain Imran Jaffery, Liaqat Ali, (2015),
Effect of aging on Phase Transition Behavior of Ti50Ni15Pd25Cu10 High Temperature
Shape Memory Alloys, International Conference on Nano and Materials Science, Jan
24-26, Zhuhai, China.
3. Saif ur Rehman, Mushtaq Khan, Aamer Nusair Khan, Liaqat Ali, Syed Husain
Imran Jaffery, (2014), Two-step Martensitic Transformation in an Aged
Ti50Ni15Pd25Cu10 High Temperature Shape Memory Alloys, International Conference
on Computational and Experimental Science and Engineering (ICCESEN-2014), Oct
25-29, Antalya, Turkey.
xi
Table of Contents
Dedication ------------------------------------------------------------------------------------ iv
Acknowledgments --------------------------------------------------------------------------- v
Abstract ------------------------------------------------------------------------------------- vii
Publications ---------------------------------------------------------------------------------- ix
Conference Presentations ------------------------------------------------------------------ x
Table of Contents --------------------------------------------------------------------------- xi
List of Figures -------------------------------------------------------------------------- xviii
List of Tables ---------------------------------------------------------------------------- xxix
List of Abbreviations ------------------------------------------------------------------- xxx
Chapter – 1 Introduction and Overview .................................................................. 1
1.1 Shape memory alloys ............................................................................................. 1
1.2 Phase transformation temperatures ........................................................................ 3
1.3 Effect on transformation temperatures due to ternary alloying with NiTi ............. 4
1.4 High temperature shape memory alloys ................................................................. 5
1.4.1 TiNiPd based high temperature shape memory alloys ................................... 5
1.5 Problem statement .................................................................................................. 7
1.6 Research aim .......................................................................................................... 7
1.7 Research objectives ................................................................................................ 8
1.8 Research methodology ........................................................................................... 8
1.8.1 Material selection ........................................................................................... 8
1.8.2 Material processing and heat treatment .......................................................... 8
xii
1.8.3 Sample preparation and characterization ....................................................... 9
1.9 Thesis outline ......................................................................................................... 9
Chapter – 2 Background ......................................................................................... 11
2.1 Introduction .......................................................................................................... 11
2.2 Discovery of shape memory alloys ...................................................................... 11
2.3 Mechanism of shape memory effect and superelasticity ..................................... 12
2.4 Crystallography of NiTi-based shape memory alloys .......................................... 14
2.5 High temperature shape memory alloys ............................................................... 15
2.5.1 Nickel (Ni) based high temperature shape memory alloys .......................... 17
2.5.1.1 Nickel Aluminum (NiAl) high temperature shape memory alloys ...... 18
2.5.1.2 Nickel Manganese (NiMn) high temperature shape memory alloys .... 19
2.5.2 Copper (Cu) based high temperature shape memory alloys ......................... 22
2.5.3 Nickel Titanium (NiTi) based high temperature shape memory alloys ....... 25
2.5.3.1 Nickel Titanium Hafnium (NiTiHf) and Nickel Titanium Zirconium
(NiTiZr) high temperature shape memory alloys ............................................. 25
2.5.3.2 Titanium Nickel Platinum (TiNiPt) high temperature shape memory
alloys ................................................................................................................. 29
2.5.3.3 Titanium Nickel Paladium (TiNiPd) high temperature shape memory
alloys ................................................................................................................. 31
Chapter – 3 Equipment Setup and Material Processing ...................................... 41
3.1 Introduction .......................................................................................................... 41
3.2 Development of shape memory alloys ................................................................. 41
3.2.1 Cleaning of constituent elements ................................................................. 43
3.2.2 Melting of constituent elements ................................................................... 43
3.2.3 Homogenization ........................................................................................... 44
3.2.4 Chemical analysis ......................................................................................... 45
xiii
3.3 Sample preparation............................................................................................... 45
3.3.1 Sample dimensions ....................................................................................... 45
3.3.2 Solution treatment and aging ........................................................................ 46
3.4 Materials characterization .................................................................................... 47
3.4.1 Optical microscopy ....................................................................................... 47
3.4.2 Scanning Electron Microscopy (SEM) ......................................................... 48
3.4.3 X-Ray Diffractometry (XRD) ...................................................................... 48
3.4.4 Differential Scanning Calorimetry (DSC) .................................................... 49
3.5 Mechanical Testing .............................................................................................. 50
3.5.1 Microhardness testing ................................................................................... 50
3.5.2 Isothermal tensile testing .............................................................................. 50
3.6 Measurement of shape memory properties .......................................................... 52
3.6.1 Equipment setup ........................................................................................... 52
3.6.2 Constant stress thermal cycling tests ............................................................ 53
3.7 Summary .............................................................................................................. 56
Chapter – 4 Effect of Copper Addition and Aging on Microstructure of TiNiPd
Alloys ......................................................................................................................... 57
4.1 Introduction .......................................................................................................... 57
4.2 Microstructure analysis of solution treated TiNiPdCu alloys with varying Cu
percentage .................................................................................................................. 57
4.2.1 Second phase precipitates ............................................................................. 57
4.2.2 Grain size ...................................................................................................... 60
4.3 Effect of aging temperature on microstructure of TiNiPdCu alloys with varying
Cu percentage ............................................................................................................. 62
4.4 Phase analysis of solution treated TiNiPdCu alloys with varying Copper
percentage .................................................................................................................. 68
4.5 Phase analysis of aged TiNiPdCu alloys with varying Copper percentage ......... 70
xiv
4.5.1 Phase analysis of aged 0Cu alloy ................................................................. 70
4.5.2 Phase analysis of aged 5Cu alloy ................................................................. 71
4.5.3 Phase analysis of aged 10Cu alloy ............................................................... 72
4.5.4 Phase analysis of aged 15Cu alloy ............................................................... 74
4.6 Summary .............................................................................................................. 76
Chapter – 5 Effect of Copper Addition and Aging on Transformation
Temperatures of TiNiPd Alloys .............................................................................. 78
5.1 Introduction .......................................................................................................... 78
5.2 Effect of Cu addition on phase transformation temperatures............................... 78
5.3 Effect of aging on phase transformation temperatures ........................................ 80
5.3.1 Effect of aging on phase transformation temperatures of 0Cu alloy ............ 81
5.3.2 Effect of aging on phase transformation temperatures of 5Cu alloy ............ 82
5.3.3 Effect of aging on phase transformation temperatures of 10Cu alloy .......... 84
5.3.4 Effect of aging on phase transformation temperatures of 15Cu alloy .......... 87
5.4 Effect of thermal cycling on phase transformation temperatures ........................ 89
5.4.1 Effect of thermal cycling on phase transformation temperatures of 0Cu alloy
............................................................................................................................... 89
5.4.2 Effect of thermal cycling on phase transformation temperatures of 5Cu alloy
............................................................................................................................... 92
5.4.3 Effect of thermal cycling on phase transformation temperatures of 10Cu
alloy ....................................................................................................................... 94
5.4.4 Effect of thermal cycling on phase transformation temperatures of 15Cu
alloy ....................................................................................................................... 96
5.5 Summary .............................................................................................................. 98
Chapter – 6 Effect of Copper Addition and Aging on Mechanical Properties of
TiNiPd Alloys .......................................................................................................... 100
xv
6.1 Introduction ........................................................................................................ 100
6.2 Effect of Cu addition on hardness ...................................................................... 100
6.3 Effect of aging on hardness ................................................................................ 101
6.3.1 Effect of aging on hardness of 0Cu alloy ................................................... 101
6.3.2 Effect of aging on hardness of 5Cu alloy ................................................... 102
6.3.3 Effect of aging on hardness of 10Cu alloy ................................................. 103
6.3.4 Effect of aging on hardness of 15Cu alloy ................................................. 105
6.4 Effect of Cu addition on mechanical strength .................................................... 107
6.4.1 Effect of Cu addition on mechanical strength in martensite phase ............ 107
6.4.2 Effect of Cu addition on mechanical strength in austenite phase ............... 109
6.5 Comparison between the mechanical properties of martensite and austenite
phases in solution treated condition ......................................................................... 111
6.6 Effect of aging on mechanical strength .............................................................. 114
6.6.1 Effect of aging on mechanical strength in martensite phase ...................... 114
6.6.2 Effect of aging on mechanical strength in austenite phase ........................ 116
6.7 Comparison between the mechanical properties of martensite and austenite
phases in 600°C-aged condition ............................................................................... 118
6.8 Summary ............................................................................................................ 121
Chapter – 7 Effect of Copper Addition on Shape Memory Properties of TiNiPd
Alloys ....................................................................................................................... 123
7.1 Introduction ........................................................................................................ 123
7.2 Shape memory properties of TiNiPd alloys with varying Cu percentage .......... 123
7.2.1 Shape memory properties of 0Cu alloy ...................................................... 123
7.2.2 Shape memory properties of 5Cu alloy ...................................................... 127
7.2.3 Shape memory properties of 10Cu alloy .................................................... 130
7.2.4 Shape memory properties of 15Cu alloy .................................................... 133
7.3 Effect of Cu addition on transformation temperatures ....................................... 136
xvi
7.4 Effect of Cu addition on transformation strains ................................................. 137
7.5 Effect of Cu addition on recovery ratio and work output .................................. 139
7.6 Summary ............................................................................................................ 141
Chapter – 8 Effect of Aging on Shape Memory Properties of TiNiPd Alloys .. 142
8.1 Introduction ........................................................................................................ 142
8.2 Shape memory properties of 600°C-aged TiNiPd alloys with varying Cu
percentage ................................................................................................................ 142
8.2.1 Shape memory properties of 600°C-aged 0Cu alloy .................................. 142
8.2.2 Shape memory properties of 600°C-aged 5Cu alloy .................................. 145
8.2.3 Shape memory properties of 600°C-aged 10Cu alloy ................................ 148
8.2.4 Shape memory properties of 600°C-aged 15Cu alloy ................................ 151
8.3 Comparison of shape memory properties between solution treated and 600°C-
aged 0Cu alloys ........................................................................................................ 154
8.3.1 Comparison of transformation temperatures .............................................. 154
8.3.2 Comparison of transformation strains ........................................................ 155
8.3.3 Comparison of recovery ratio and work output .......................................... 156
8.4 Comparison of shape memory properties between solution treated and 600°C-
aged 5Cu alloys ........................................................................................................ 157
8.4.1 Comparison of transformation temperatures .............................................. 157
8.4.2 Comparison of transformation strains ........................................................ 158
8.4.3 Comparison of recovery ratio and work output .......................................... 159
8.5 Comparison of shape memory properties between solution treated and 600°C-
aged 10Cu alloys ...................................................................................................... 160
8.5.1 Comparison of transformation temperatures .............................................. 160
8.5.2 Comparison of transformation strains ........................................................ 161
8.5.3 Comparison of recovery ratio and work output .......................................... 162
xvii
8.6 Comparison of shape memory properties between solution treated and 600°C-
aged 15Cu alloys ...................................................................................................... 163
8.6.1 Comparison of transformation temperatures .............................................. 163
8.6.2 Comparison of transformation strains ........................................................ 164
8.6.3 Comparison of recovery ratio and work output .......................................... 165
8.7 Summary ............................................................................................................ 167
Chapter – 9 Summary of Results and Discussion ............................................... 168
9.1 Introduction ........................................................................................................ 168
9.2 Effect of Cu addition on microstructure............................................................. 168
9.3 Effect of Cu addition on transformation temperatures ....................................... 169
9.4 Effect of Cu addition on mechanical and shape memory properties ................. 169
9.5 Effect of aging on microstructure and transformation temperatures ................. 170
9.6 Effect of aging on mechanical and shape memory properties ........................... 171
Chapter – 10 Conclusions and Recommendations for Future Work ............... 173
10.1 Summary of experimentation ........................................................................... 173
10.2 Conclusions ...................................................................................................... 174
10.2.1 Effect of Cu addition ................................................................................ 174
10.2.2 Effect of aging .......................................................................................... 175
10.3 Recommendations for future work .................................................................. 176
References: .............................................................................................................. 177
xviii
List of Figures
Fig. 1.1 Mechanism of martensitic transformation ----------------------- 2
Fig. 1.2 Dependence of martensite start temperature, Ms on Ni-content
in binary NiTi alloys ---------------------------------------------------
4
Fig. 1.3 Dependence of Ms temperature on the addition of Pd-content by
replacing Ni, in an equi-atomic TiNiPd system --------------------
6
Fig. 2.1 Mechanism of shape memory effect can be observed by
following the path a-b-c-a, while superelasticity can be realized
by following the path c´-c- c´ -----------------------------------------
13
Fig. 2.2 Different transformation paths in TiNi based alloys --------------- 15
Fig. 2.3 Dependence of Ms temperature on aluminum content in NiAl
alloy ----------------------------------------------------------------------
18
Fig. 2.4 Dependence of martensite peak temperature, Mp on either Ni or
Mn contents in NiMnAl alloy ----------------------------------------
20
Fig. 2.5 Composition dependence of the martensite start temperature, Ms
on aluminum content in CuAlNi alloys -----------------------
24
Fig. 2.6 Composition dependence of the martensite peak temperature,
Mp as a function of hafnium content in NiTiHf alloys --------
26
Fig. 2.7 Composition dependence of the martensite peak temperature,
Mp as a function of zirconium content in NiTiZr alloys -------
28
Fig. 2.8 Composition dependence of the transformation temperature as a
function of platinum content in TiNiPt alloys ---------------
30
Fig. 2.9 Change in martensite start temperature, Ms with respect to
palladium content in equi-atomic TiNi50-xPdx alloys ----------
33
Fig. 2.10 Martensite start temperature, Ms as a function of Ti/(Ni,Pd)
ratio ---------------------------------------------------------------------
35
xix
Fig. 2.11 (a) Specific work output for a Ti50.5Ni19.5Pd30 alloy as a function
of applied stress loaded in both tension and compression and (b)
the corresponding transformation strain versus applied stress
38
Fig. 2.12 Work output for a series of TiNiPd and TiNiPt alloys as
function of the transformation temperature range (Mf to Af)
39
Fig. 3.1 Process flow chart presenting the sequence of operations for
materials processing and characterization ----------------------
42
Fig. 3.2 Schematic representation of homogenization and solution
treatment processes -------------------------------------------------
46
Fig. 3.3 Schematic representation the aging process at various aging
temperatures ------------------------------------------------------------
47
Fig. 3.4 Measurement scheme of transformation temperatures from DSC
heating and cooling cycles -------------------------------------------
50
Fig. 3.5 Special gripping arrangement for holding of 0.3 mm thick
samples ------------------------------------------------------------------
51
Fig. 3.6 Internal view of lever arm creep and stress rupture tensile
testing system ----------------------------------------------------------
53
Fig. 3.7 Measurement scheme of transformation temperatures,
recoverable and irrecoverable strains from typical strain-
temperature curve ------------------------------------------------------
55
Fig. 4.1 SEM images showing the second phase precipitates formed
along the grain boundaries in solution treated samples of (a)
0Cu, (b) 5Cu, (c) 10Cu and (d) 15 Cu alloys --------------------
58
Fig. 4.2 EDS spectrums shown for solution treated samples of 0Cu,
5Cu, 10Cu and 15Cu alloys------------------------------------------
59
xx
Fig. 4.3 Optical micrographs (at 200X) of (a) 0Cu (b) 5Cu (c) 10Cu and
(d) 15Cu alloys solution treated at 900°C for 1 hour ----------
60
Fig. 4.4 Optical micrographs (at 500X) of (a) 0Cu (b) 5Cu (c) 10Cu and
(d) 15Cu alloys solution treated at 900°C for 1 hour ----------
61
Fig. 4.5 Effect of increasing Cu-content on grain size -------------------- 61
Fig. 4.6 Back-scattered SEM images presenting the microstructure and
grain boundaries in 0Cu alloys after aging for 3 hours at
temperature of (a) 400°C, (b) 500°C, (c) 600°C and (d) 700°C
63
Fig. 4.7 Back-scattered SEM images presenting the microstructure and
grain boundaries in 5Cu alloys after aging for 3 hours at
temperature of (a) 400°C, (b) 500°C, (c) 600°C and (d) 700°C
64
Fig. 4.8 Back-scattered SEM images presenting the microstructure and
grain boundaries in 10Cu alloys after aging for 3 hours at
temperature of (a) 400°C, (b) 500°C, (c) 600°C and (d) 700°C
65
Fig. 4.9 Back-scattered SEM images presenting the microstructure and
grain boundaries in 15Cu alloys after aging for 3 hours at
temperature of (a) 400°C, (b) 500°C, (c) 600°C and (d) 700°C
66
Fig. 4.10 XRD profiles at room temperature for the samples solution
treated of 0Cu, 5Cu, 10Cu and 15Cu alloys ---------------------
69
Fig. 4.11 XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 0Cu
alloys --------------------------------------------------------------------
70
Fig. 4.12 XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 5Cu
alloys ---------------------------------------------------------------------
72
xxi
Fig. 4.13 XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 10Cu
alloys ---------------------------------------------------------------------
73
Fig. 4.14 XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 15Cu
alloys ---------------------------------------------------------------------
75
Fig. 5.1 DSC heating and cooling cycles of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys -------------------------------------------------
79
Fig. 5.2 Effect of Cu addition on transformation temperatures ------- 80
Fig. 5.3 DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 0Cu alloy -------
81
Fig. 5.4 Effect of aging temperatures on transformation temperatures as
a function of increasing aging temperatures of 0Cu alloy ----
82
Fig. 5.5 DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 5Cu alloy ------
83
Fig. 5.6 Effect of aging temperatures on transformation temperatures as
a function of increasing aging temperatures of 5Cu alloy ----
84
Fig. 5.7 DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 10Cu alloy -----
85
Fig. 5.8 Effect of aging temperatures on transformation temperatures as
a function of increasing aging temperatures of 10Cu alloy ----
86
Fig. 5.9 DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 15Cu alloy -------
87
Fig. 5.10 Effect of aging temperatures on transformation temperatures as
a function of increasing aging temperatures of 15Cu alloy ----
88
xxii
Fig. 5.11 DSC curves representing the transformation temperatures
during five thermal cycles of 0Cu alloy ------------------------
90
Fig. 5.12 Effect of thermal cycles on transformation temperatures of 0Cu
alloy ----------------------------------------------------------------------
91
Fig. 5.13 DSC curves representing the transformation temperatures
during five thermal cycles of 5Cu alloy ---------------------------
93
Fig. 5.14 Effect of thermal cycle on transformation temperatures of 5Cu
alloy ---------------------------------------------------------------------
93
Fig. 5.15 DSC curves representing the transformation temperatures
during five thermal cycles of 10Cu alloy -------------------------
95
Fig. 5.16 Effect of thermal cycle on transformation temperatures of 10Cu
alloy ----------------------------------------------------------------------
95
Fig. 5.17 DSC curves representing the transformation temperatures
during five thermal cycles of 15Cu alloy -------------------------
97
Fig. 5.18 Effect of thermal cycle on transformation temperatures of 15Cu
alloy ---------------------------------------------------------------------
98
Fig. 6.1 Microhardness of solution treated samples of 0Cu, 5Cu, 10Cu
and 15Cu alloys ------------------------------------------------------
101
Fig. 6.2 Microhardness of 0Cu alloy aged at different temperatures 102
Fig. 6.3 Microhardness of 5Cu alloy aged at different temperatures 102
Fig. 6.4 Microhardness of 10Cu alloy aged at different temperatures 103
Fig. 6.5 Microhardness of 15Cu alloy aged at different temperatures 106
Fig. 6.6 Tensile stress-strain curves of solution treated 0Cu, 5Cu, 10Cu
and 15Cu alloys, tested in martensite phase (Mf – 50°C) ---
108
xxiii
Fig. 6.7 Tensile stress-strain curves of solution treated 0Cu, 5Cu, 10Cu
and 15Cu alloys, tested in austenite phase (Af + 50°C) ------
110
Fig. 6.8 Effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite yield stress, tested at (Mf – 50°C) and
austenite yield stress, tested at (Af + 50°C) ----------------------
111
Fig. 6.9 Effect of partial substitution of Ni by Cu in solution treated 0Cu
alloy on martensite fracture stress, tested at (Mf – 50°C) and
austenite fracture stress, tested at (Af + 50°C) ------------------
112
Fig. 6.10 Effect of partial substitution of Ni by Cu in solution treated 0Cu
alloy on martensite fracture strain, tested at (Mf – 50°C) and
austenite fracture strain, tested at (Af + 50°C) ------------------
113
Fig. 6.11 Tensile stress-strain curves of 600°C-aged 0Cu, 5Cu, 10Cu and
15Cu alloys, tested in martensite phase (Mf – 50°C) ----------
115
Fig. 6.12 Tensile stress-strain curves of 600°C-aged 0Cu, 5Cu, 10Cu and
15Cu alloys, tested in austenite phase (Af + 50°C) ------------
117
Fig. 6.13 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite yield stress, tested at (Mf – 50°C) and
austenite yield stress, tested at (Af + 50°C) ----------------------
119
Fig. 6.14 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture stress, tested at (Mf – 50°C) and
austenite fracture stress, tested at (Af + 50°C) -----------------
119
Fig. 6.15 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture strain, tested at (Mf – 50°C) and
austenite yield stress, tested at (Af + 50°C) ----------------------
120
Fig. 7.1 Strain-temperature curves representing the shape memory
properties of solution treated 0Cu alloy at stress levels of 100 –
500 MPa ---------------------------------------------------------------
124
xxiv
Fig. 7.2 Change in transformation temperatures of solution treated 0Cu
alloy at stress levels of 100 – 500 MPa --------------------------
125
Fig. 7.3 Recovered and irrecoverable strains of solution treated 0Cu
alloy at stress levels of 100 – 500 MPa ---------------------------
125
Fig. 7.4 Recovery ratio and work output of solution treated 0Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
126
Fig. 7.5 Strain-temperature curves representing the shape memory
properties of solution treated 5Cu alloy at stress levels of 100 –
500 MPa ---------------------------------------------------------------
127
Fig. 7.6 Change in transformation temperatures of solution treated 5Cu
alloy at stress levels of 100 – 500 MPa ---------------------------
128
Fig. 7.7 Recovered and irrecoverable strains of solution treated 5Cu
alloy at stress levels of 100 – 500 MPa --------------------------
129
Fig. 7.8 Recovery ratio and work output of solution treated 5Cu alloy at
stress levels of 100 – 500 MPa -------------------------------------
129
Fig. 7.9 Strain-temperature curves representing the shape memory
properties of solution treated 10Cu alloy at stress levels of 100
– 500 MPa ---------------------------------------------------------------
130
Fig. 7.10 Change in transformation temperatures of solution treated 10Cu
alloy at stress levels of 100 – 500 MPa --------------------------
131
Fig. 7.11 Recovered and irrecoverable strains of solution treated 10Cu
alloy at stress levels of 100 – 500 MPa --------------------------
131
Fig. 7.12 Recovery ratio and work output of solution treated 10Cu alloy
at stress levels of 100 – 500 MPa -----------------------------------
132
Fig. 7.13 Strain-temperature curves representing the shape memory
properties of solution treated 15Cu alloy at stress levels of 100
– 500 MPa --------------------------------------------------------------
133
xxv
Fig. 7.14 Change in transformation temperatures of solution treated 15Cu
alloy at stress levels of 100 – 500 MPa -------------------------
134
Fig. 7.15 Recovered and irrecoverable strains of solution treated 15Cu
alloy at stress levels of 100 – 500 MPa -------------------------
135
Fig. 7.16 Recovery ratio and work output of solution treated 15Cu alloy
at stress levels of 100 – 500 MPa ----------------------------------
135
Fig. 7.17 Comparison of martensite start temperatures of solution treated
0Cu, 5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500
MPa ----------------------------------------------------------------
136
Fig. 7.18 Comparison of austenite finish temperatures of solution treated
0Cu, 5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500
MPa -----------------------------------------------------------------
137
Fig. 7.19 Comparison of recovered strain of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys under stress level of 100 – 500 MPa
138
Fig. 7.20 Comparison of irrecoverable strain of solution treated 0Cu,
5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500 MPa
138
Fig. 7.21 Comparison of recovery ratio of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys under stress level of 100 – 500 MPa
140
Fig. 7.22 Comparison of work output of solution treated 0Cu, 5Cu, 10Cu
and 15Cu alloys under stress level of 100 – 500 MPa --------
140
Fig. 8.1 Strain-temperature curves representing the shape memory
properties of 600°C-aged 0Cu alloy at stress levels of 100 – 500
MPa -------------------------------------------------------------------
143
Fig. 8.2 Change in transformation temperatures of 600°C-aged 0Cu
alloy at stress levels of 100 – 500 MPa ---------------------------
144
Fig. 8.3 Recovered and irrecoverable strains of 600°C-aged 0Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
144
xxvi
Fig. 8.4 Recovery ratio and work output of 600°C-aged 0Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
145
Fig. 8.5 Strain-temperature curves representing the shape memory
properties of 600°C-aged 5Cu alloy at stress levels of 100 – 500
MPa --------------------------------------------------------------------
146
Fig. 8.6 Change in transformation temperatures of 600°C-aged 5Cu
alloy at stress levels of 100 – 500 MPa --------------------------
146
Fig. 8.7 Recovered and irrecoverable strains of 600°C-aged 5Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
147
Fig. 8.8 Recovery ratio and work output of 600°C-aged 5Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
148
Fig. 8.9 Strain-temperature curves representing the shape memory
properties of 600°C-aged 10Cu alloy at stress levels of 100 –
500 MPa ----------------------------------------------------------------
149
Fig. 8.10 Change in transformation temperatures of 600°C-aged 10Cu
alloy at stress levels of 100 – 500 MPa --------------------------
150
Fig. 8.11 Recovered and irrecoverable strains of 600°C-aged 10Cu alloy
at stress levels of 100 – 500 MPa ----------------------------------
150
Fig. 8.12 Recovery ratio and work output of 600°C-aged 10Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
151
Fig. 8.13 Strain-temperature curves representing the shape memory
properties of 600°C-aged 15Cu alloy at stress levels of 100 –
500 MPa ----------------------------------------------------------------
152
Fig. 8.14 Change in transformation temperatures of 600°C-aged 15Cu
alloy at stress levels of 100 – 500 MPa --------------------------
152
Fig. 8.15 Recovered and irrecoverable strains of 600°C-aged 15Cu alloy
at stress levels of 100 – 500 MPa -----------------------------------
153
xxvii
Fig. 8.16 Recovery ratio and work output of 600°C-aged 15Cu alloy at
stress levels of 100 – 500 MPa --------------------------------------
153
Fig. 8.17 Comparison of transformation temperatures of solution treated
and 600°C-aged 0Cu alloy under stress level of 100 – 500 MPa
155
Fig. 8.18 Comparison of transformation strains of solution treated and
600°C-aged 0Cu alloy under stress level of 100 – 500 MPa
156
Fig. 8.19 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 0Cu alloy under stress level of 100 –
500 MPa ----------------------------------------------------------------
157
Fig. 8.20 Comparison of transformation temperatures of solution treated
and 600°C-aged 5Cu alloy under stress level of 100 – 500 MPa
158
Fig. 8.21 Comparison of transformation strains of solution treated and
600°C-aged 5Cu alloy under stress level of 100 – 500 MPa
159
Fig. 8.22 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 5Cu alloy under stress level of 100 –
500 MPa ----------------------------------------------------------------
160
Fig. 8.23 Comparison of transformation temperatures of solution treated
and 600°C-aged 10Cu alloy under stress level of 100 – 500
MPa ---------------------------------------------------------------------
161
Fig. 8.24 Comparison of transformation strains of solution treated and
600°C-aged 10Cu alloy under stress level of 100 – 500 MPa
162
Fig. 8.25 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 10Cu alloy under stress level of 100 –
500 MPa ----------------------------------------------------------------
163
xxviii
Fig. 8.26 Comparison of transformation temperatures of solution treated
and 600°C-aged 15Cu alloy under stress level of 100 – 500
MPa ---------------------------------------------------------------------
164
Fig. 8.27 Comparison of transformation strains of solution treated and
600°C-aged 15Cu alloy under stress level of 100 – 500 MPa
165
Fig. 8.28 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 15Cu alloy under stress level of 100 –
500 MPa ----------------------------------------------------------------
166
xxix
List of Tables
Table 3.1 Chemical compositions of the four alloys given in weight
percent -----------------------------------------------------------------
44
Table 3.2 Chemical composition of homogenized alloys ------------------- 45
Table 3.3 Dimensions of samples for various characterizations ----------- 46
Table 4.1 Compositional analysis of the overall alloy and second phase
precipitate in solution treated condition for 0Cu, 5Cu, 10Cu
and 15Cu alloys ------------------------------------------------------
59
Table 4.2 Compositional analysis of black and white precipitates formed
in 10Cu and 15Cu alloys after aging for 3 hours at various
aging temperatures ---------------------------------------------------
67
Table 6.1 Yield stress, fracture stress and fracture strain calculated from
stress-strain curves of Fig. 6.6 for solution treated alloys
tested in martensite phase (Mf – 50°C) --------------------------
109
Table 6.2 Yield stress, fracture stress and fracture strain calculated from
stress-strain curves of Fig. 6.7 for solution treated alloys
tested in austenite phase (Af + 50°C) -----------------------------
111
Table 6.3 Yield stress, fracture stress and fracture strain calculated from
stress-strain curves of Fig. 6.11 for 600°C-aged alloys tested
in martensite phase (Mf – 50°C) -----------------------------------
116
Table 6.4 Yield stress, fracture stress and fracture strain calculated from
stress-strain curves of Fig. 6.12 for 600°C-aged alloys tested
in austenite phase (Af + 50°C) -------------------------------------
118
xxx
List of Abbreviations
SMA Shape Memory Alloy
SME Shape Memory Effect
TWSME Two Way Shape Memory Effect
PE Pseudo elasticity
TT Transformation Temperature
ΔT Thermal Hysteresis
εrec Recovered Strain
εirr Irrecoverable Strain
εf Fracture Strain
σy Yield Stress
σf Fracture Stress
σDT Detwin Stress
WQ Water Quenched
Ms Martensite Start Temperature
Mp Martensite Peak Temperature
Mf Martensite Finish Temperature
As Austenite Start Temperature
Ap Austenite Peak Temperature
Af Austenite Finish Temperature
1
Chapter – 1
Introduction and Overview
1.1 Shape memory alloys
Shape Memory Alloys (SMAs) are the materials which are capable to undergo
martensitic transformation and exhibit two unique properties, i.e. Shape Memory
Effect (SME) and Super-Elasticity (SE). In SME, the material is deformed in the
low temperature martensite phase, it recovers its original shape by the mechanism of
reverse transformation upon heating to a specific temperature called the reverse
transformation temperature. In SE, the material recovers a significant amount of
strain above its reverse transformation temperature. SE is an isothermal phenomenon
which is associated with a large nonlinear recoverable strain upon the removal of
externally applied load. For example NiTi wires recover about 8% strain which is
about 40 times larger as compared to conventional metals like steel having only 0.2%
strain [1, 2]. Shape recovery takes place at a particular temperature and thus the
SMA can function both as a sensor and as an actuator. Therefore SMAs are often
called smart or intelligent materials. Due to both Shape Memory and Super-Elasticity
properties of the SMAs, they are being used in various applications such as
automobile industries, high value consumer electronics, pipe couplings, medical
implants and guide wires and antennae for cellular phones. Apart from that, SMAs
are also used extensively in Micro Electro-Mechanical Systems (MEMS).
The properties of SME and SE are the result of a reversible martensitic
transformation which is a diffusionless solid state phase transformation process that
can be activated by temperature change or magnetic field. In such type of
transformation, the atoms of the material move cooperatively or by shear-like
mechanism. The high temperature phase (parent phase) is usually cubic and the
lower temperature phase (martensite) has a lower symmetry. The schematic of
martensitic transformation is shown in Fig. 1.1 [3].
2
Fig. 1.1 Mechanism of martensitic transformation [3]
By cooling the SMAs below some specific temperature, the process of martensitic
transformation exhibits by shear-like mechanism, as shown in the Fig. 1.1. From the
same schematic it can be seen that the structure of martensites present in regions A
and B are the same, however they are oriented in different directions. This difference
of the direction is called variants of the martensite. Due to having a lower symmetry,
considerable number of martensite variants can be formed from the same parent
phase. Similarly by increasing the temperature above the critical value, the available
martensites become unstable and the reverse transformation takes place. In this way
the martensites transform back to the original orientation as was present in the parent
phase and recovers its original shape. Thus the formation of martensite due to
temperature change is called the temperature-induced martensite which provides the
origin for shape memory effect. The phenomenon of SE takes place when the SMA
is deformed above the reverse transformation temperature. At high temperature, the
parent phase is stable in the absence of external load. However by applying load
(induced stress), the parent phase becomes unstable and consequently the martensites
are formed which are oriented in the direction of applied load. Upon removal of
external loading, the martensites become unstable and transform back to the parent
phase in original orientation. Thus the formation of martensite due to induced stress
within the material is called the stress-induced martensite, which is responsible for
the superelastic behavior.
3
1.2 Phase transformation temperatures
From the application point of view, it is important for SMAs to know the
temperatures at which the forward and reverse transformation starts and ends. The
temperatures at which the alloy starts and completes transformation from austenite
phase to martensite phase during forward transformation (cooling from high
temperature to low temperature) are called martensite start (Ms) and martensite finish
(Mf) temperature respectively. Conversely, during reverse transformation (heating
from low temperature to high temperature), the temperatures at which the alloy
begins and ends the transformation from martensite phase to austenite phase are
called austenite start (As) and austenite finish (Af) temperature respectively.
In the NiTi-based shape memory alloys, the phase transformation temperatures are
very sensitive to the Ni content. By increasing the Ni content above 50% (or
decreasing Ti-content below 50%), the transformation temperatures decrease very
rapidly [4]. However, decreasing Ni content below 50% (i.e. Ti content above 50%),
the transformation temperatures remain almost constant as shown in Fig. 1.2. This
shows the importance of the relationship between the transformation temperatures
and the alloy composition in case of NiTi-based shape memory alloys. The Ni-rich
NiTi shape memory alloys normally show superelasticity at room temperature since
the transformation temperatures are below room temperature. These alloys are
mainly used in applications where superelastic effect is required to function i.e.
orthodontic wires, stents, springs and guide wires etc. The Ti-rich NiTi shape
memory alloys have relatively higher transformation temperatures but have less
sensitivity towards the composition. These alloys normally show shape memory
effect at room temperature. The Ti-rich NiTi shape memory alloys are mainly used
as actuator type of applications where the shape memory effect is required to
function. The binary NiTi alloys can only be used at temperatures below 100°C,
since the highest Ms that can be achieved is about 77°C [4].
4
Fig.1.2 Dependence of Ms temperature on Ni-content in binary NiTi alloys [4]
1.3 Effect on transformation temperatures due to ternary alloying
with NiTi
There have been numerous efforts to improve or modify the shape memory and other
properties of NiTi-based shape memory alloys through ternary alloying additions. It
is well known that the ternary alloying additions often affect the transformation
temperatures. Among all the tested elements very few have been reported to increase
the transformation temperatures i.e. Pt, Pd, Au, Zr and Hf [4-7]. All the other
elements i.e. Cr, Mn, Fe, Ag, V, Co, Al, Si etc, generally cause a decrease in
transformation temperatures [5]. Cu addition in NiTi makes the transformation
temperatures less sensitive towards the compositional changes. Addition of Cu
exceeding 5% can change the transformation route into B2 (austenite) – B19
(orthorhombic) – B19´(monoclinic) and further addition causes a slight increase in
the transformation temperatures [6].
5
1.4 High temperature shape memory alloys
The shape memory alloys have already become an important class of materials for
various biomedical applications because of their unique SME and SE. Recently, their
potential has been realized in the areas of aerospace, energy exploration and
automotive industries to be used as solid state actuators [7]. Shape memory alloys
with high transformation temperatures can simplify the design and improve the
efficiency of many mechanical components which are required to operate at
temperatures higher than 100°C in automotive, aerospace, manufacturing and energy
exploration industries. Currently, the practical usage of shape memory alloys is
limited to temperatures below 100°C [8]. This is actually the transformation
temperature limit of the two most commercially successful shape memory alloy
systems; the NiTi binary alloys and Cu based ternary alloys. The low transformation
temperatures of these systems put a limit on their usage and they cannot be used in
those applications where the minimum requirement, in terms of transformation
temperatures, exceeds 100°C. Generally, the shape memory alloys with
transformation temperatures exceeding 100°C are categorized as high temperature
shape memory alloys.
1.4.1 TiNiPd based high temperature shape memory alloys
TiNiPd based high temperature shape memory alloys have been investigated since
their discovery by Eckelmeyer in 1976 [9]. Their potential as a successful high
temperature shape memory alloy system lies in the attractive combination of
essential characteristics which they possess i.e. high transformation temperatures,
adequate workability, reasonable level of strain recovery, small hysteresis and
acceptable mechanical strength. Initially the main research focus was to improve the
transformation temperatures of TiNiPd based high temperature shape memory alloys
but in the recent years the focus has been shifted towards the improvement of their
work output, dimensional and microstructural stability and high temperature cyclic
characteristics. The transformation temperatures of this system can be altered by
replacing Ni with Pd. If the concentration of Ti is held constant at nearly 50%, the
relationship between the transformation temperatures and relative concentration of
Ni and Pd is parabolic [10], as shown in Fig. 1.3. It can be seen that replacing Ni
with Pd up to 10% causes a decrease in the transformation temperatures and then
6
they start to increase upon further replacement of Ni with Pd with an approximate
rate of 15K/at%.
Fig.1.3 Dependence of Ms temperature on addition of Pd-content by replacing Ni, in
an equi-atomic TiNiPd system [10]
The parabolic dependence of the transformation temperatures on composition
originates from the change in the structure of martensite. In Ti50Ni50-xPdx alloys the
B2 (cubic) phase transforms into B19 (orthorhombic) martensite when Pd
concentrations remain higher than 10%, and it transforms into B19ʹ (monoclinic)
martensite or R (rhombohedral) phase when Pd contents are lower than 10%.
Because of the complete miscibility of the TiNi and TiPd systems, it is possible to
develop a continuous range of transformation temperatures from room temperature to
over 500°C by adjusting the Pd contents in TiNiPd alloy. The TiNiPd alloys with Pd
contents higher than 10% are of much importance especially for high temperature
shape memory applications. In these alloys, the B2 (cubic) phase normally
transforms into B19 (orthorhombic) martensite.
7
1.5 Problem statement
For the shape memory alloys to be used as an actuator at high working temperature,
their transformation temperatures must be increased to a reasonable limit, but at the
same time it must have the property to operate with low thermal hysteresis.
Unfortunately, by increasing the transformation temperatures due to increase in Pd
content in TiNiPd, thermal hysteresis also increases due to weakening of
compatibility between the B19 martensite phase and B2 parent phase.
Similar to any other metallic material, the critical stress for slip deformation
decreases with increasing temperature in TiNiPd based high temperature shape
memory alloys. The slip deformation can easily occur simultaneously with the
reorientation of martensite variants by stressing the alloy at high temperatures. Also
the creep deformation becomes a serious problem especially when the martensitic
transformation temperature range overlaps the creep deformation temperature (~
370°C) for these alloys. Undesirable recovery and recrystallization of the
microstructure is another important concern especially in case of
thermomechanically treated TiNiPd alloys. These factors actually put a limit on the
maximum transformation temperature range up to which the transformation
temperatures of TiNiPd based alloys can be raised by changing the Pd contents. The
condition becomes further worse when the material has to work under a certain load
where transformation temperatures are increased according to the Clausius-
Clapeyron relationship.
1.6 Research aim
The aim of this research is to further increase the transformation temperatures of the
TiNiPd based high temperature shape memory alloys while maintaining the thermal
hysteresis at reasonable low level and to improve their mechanical and shape
memory properties.
8
1.7 Research objectives
A number of objectives established for this research work were:
1. To develop TiNiPdCu alloys with varying Cu percentage.
2. To characterize the microstructure, transformation temperatures, mechanical and
shape memory properties of baseline ternary TiNiPd alloy in solution treated
condition.
3. To investigate the effect of Cu addition by various composition in place of Ni in
baseline ternary TiNiPd alloys on microstructure and other properties by comparing
their results.
4. To increase the phase transformation temperatures and decrease the thermal
hysteresis of TiNiPd alloy by substitution of Cu in various composition.
5. To improve the mechanical, shape memory properties and dimensional stability of
TiNiPd alloy.
6. To study the effect of aging at various aging temperatures on microstructure and
other properties of baseline ternary TiNiPd alloy and resultant quaternary TiNiPdCu
alloys.
7. To compare the results obtained from solution treated and aged samples.
1.8 Research methodology
The methodology established for this research work can be divided into three phases
which are explained in detail below:
1.8.1 Material selection
High purity constituent elements are the basic and most essential requirement for
developing the shape memory alloys. Therefore constituent elements; Titanium,
Nickel, Palladium and Copper of highest purity (99.98 – 99.99 weight percent) were
selected for this purpose.
1.8.2 Material processing and heat treatment
Material processing includes the cleaning of constituent elements in Ultrasonic
Cleaner, precise weighing up to one-tenth of mg, melting in Vacuum Arc Melting
Furnace and casting of 20 g buttons of four alloys. Heat treatment cycles consist of
9
homogenization at 950°C for 2 hours, solution treatment at 900°C for 1 hour and
aging at different aging temperatures for 3 hours. Material processing and heat
treatment cycles are explained in detail in Chapter – 3.
1.8.3 Sample preparation and characterization
Samples were prepared by wire Electrical Discharge Machine (EDM) for various
characterizations. The samples of four alloys were characterized for their
microstructure by Optical Microscope (OP) and Scanning Electron Microscope
(SEM). Phase analysis was carried out by X-Ray Diffractometer (XRD). Differential
Scanning Calorimeter (DSC) was used to find out the phase transformation
temperatures. Mechanical properties were investigated by using the Mechanical
Testing System (MTS). Tensile Creep and Rupture Testing Machine (TCRTM) was
used to find out the various shape memory properties. Experimental setup for
complete characterization is explained in detail in Chapter – 3.
1.9 Thesis outline
The current study is focused on the improvement in transformation temperatures,
mechanical and shape memory properties of TiNiPd based high temperature shape
memory alloys through quaternary alloying addition of Cu and precipitation
hardening. The chapter wise distribution of the current study is given as follows:
Chapter – 2 gives an overview of the mechanisms and different aspects of shape
memory alloys, background and literature review.
Chapter – 3 describes materials and methods that include the material processing,
melting, heat treatment cycles, characterization techniques and experimental set up
for investigation of various properties of TiNiPdCu alloys.
Chapter – 4 presents the effect on microstructure of ternary baseline TiNiPd alloy by
addition of Cu and aging at different aging temperatures in TiNiPdCu alloys.
Chapter – 5 explains the change in phase transformation temperatures of TiNiPd
alloy due to quaternary alloying addition of Cu and aging at different aging
temperatures in TiNiPdCu alloys.
10
Chapter – 6 details the effect on mechanical properties of TiNiPd alloy due to
quaternary alloying addition of Cu and aging at different aging temperatures in
TiNiPdCu alloys.
Chapter – 7 focuses on the variation of shape memory properties of TiNiPd alloy by
addition of Cu in place of Ni.
Chapter – 8 discusses the effect of aging at different aging temperatures on shape
memory properties of TiNiPdCu alloys.
Chapter – 9 summarizes the results and discussion of all experimental work carried
out in Chapter – 4 to Chapter – 8.
Chapter – 10 enlists the major conclusions drawn from the current study and suggests
the future work.
It is expected that this research will be beneficial for the improvement of high
temperature shape memory properties of TiNiPd based alloys and will open some
new directions of research in the area of high temperature shape memory alloys.
11
Chapter – 2
Background
2.1 Introduction
Significant progress has been made both in the scientific understanding and
application of shape memory alloys (SMAs) since the discovery of these
multifunctional materials. Owing to the unique behaviors of shape memory effect
and superelasticity, SMAs have become a major materials class of choice in the
biomedical industry and are beginning to spread through other technological areas.
There is a recent revival of interest in SMAs, driven primarily by the aerospace and
automotive industries, for their potential to operate as solid state actuators.
Shape memory effect is a phenomenon whereby a deformed material could recover
its predeformed shape after being heated. When this procedure is performed against
some biasing force, the material is capable of doing work due to change in its shape.
Superelasticity is an isothermal phenomenon where the material is able to recover
high amount of strain triggered by mechanical stress. These two behaviors are the
result of reversible martensitic transformation; a diffusionless solid state phase
transformation mechanism that can be activated by temperature, stress and magnetic
field.
2.2 Discovery of shape memory alloys
The first shape memory alloy which was discovered in 1932 by a Swedish Physicist
Arne Olander was Au-Cd [11]. He observed that the plastically deformed Au-Cd
returned back to its original shape upon heating. The same shape memory effect was
then observed in 1938 by Greninger and Alden B. in a CuZn alloy [12]. In 1958,
Chang and Read demonstrated that this unique property can perform a mechanical
work [13]. They used an Au-Cd shape memory alloy to lift a weight at the Brussels
World’s Fair. The real interest in the shape memory alloys was developed in 1962
after the discovery of NiTi, exhibiting shape memory properties by William Buehler
et al. at the U.S. Naval Ordinance Research Laboratories [14]. The name for this
12
alloy was kept as ‘Nitinol’, derived from NiTi Naval Ordinance Laboratories. This
discovery was proved to be a revolution in the field of shape memory alloys due to
superior qualities as compared to previously discovered shape memory alloys. Since
the NiTi alloy has many complicated structures associated with it, the understanding
of the mechanism for SME was not possible till early 1980s. From the discovery of
SME in CuAlNi [15] alloy that made it possible to relate the SME with the
thermoelastic martensitic transformation.
2.3 Mechanism of shape memory effect and superelasticity
The mechanism of shape memory effect and superelasticity is explained by a
schematic illustration of reversible martensitic transformation, as shown in Fig. 2.1.
The parent phase is stable at high temperatures as shown at 2.1(a) as compared to the
martensite phase as shown at 2.1(b) which is a low temperature phase. When the
temperature of the parent phase is decreased then the martensite phase starts to form
at a temperature called martensite start temperature (Ms) and completely transforms
into martensite when the temperature exceeds the martensitic finish temperature
(Mf). In this case the martensite is called thermally induced martensite. To minimize
the strain energy caused by the difference between the crystal structures of parent
and martensite phases, an invariant habit plane is created between the two
transforming phases. The martensite can have different variants, each with a unique
crystallographic orientation. In Fig. 2.1 [16] only two variants, V1 and V2 are shown
for simplicity. These variants are twin related and the martensite is called twinned
martensite. The transformation strain generated during transformation from parent to
martensite is accommodated by a self-accommodation process in martensite. Self-
accommodation is realized through twinning created between the different variants
[15, 17]. The macroscopic volume in Fig. 2.1(b) remains similar as compared to that
of parent phase shown at 2.1(a) mainly due to the self-accommodation of
transformation strain in twinned martensite. When the twinned martensite is
subjected to an externally applied stress, the different variants of martensite reorient
themselves to accommodate the most favorable variant and try to transform into a
single variant which in this case is V1 as shown at 2.1(c). The martensite at this point
is called deformed or detwinned martensite. The reorientation of different variants
into a single variant actually creates the transformation strain.
13
Fig. 2.1 Mechanism of shape memory effect can be observed by following the path
a-b-c-a, while superelasticity can be realized by following the path c´-c- c´ [16]
Accordingly, shape change occurs and the volume at 2.1(c) becomes greater than that
at 2.1(a) or 2.1(b). In martensite state, the SMA is very easy to deform by exerting
force due to production of twin boundary. If the force is removed at this stage the
deformation in the martensite remains, representing the plastic deformation. Hence,
when the stress is removed, no change is occurred in the shape of deformed or
detwinned martensite because the martensite is stable at this temperature. If the
temperature of detwinned martensite is increased and it exceeds the austenite start
temperature (As) then the martensite starts the reverse transformation process and it
completes when the temperature exceeds the austenite finish temperature (Af) and
martensite completely transforms into the parent phase. As a result, the original
macroscopic shape is restored as shown at 2.1(a). This mechanism following the path
a-b-c-a is called the shape memory effect.
14
Superelasticity is another aspect of shape memory behavior. In Fig. 2.1 the
superelasticity can be realized by following the path c´-c- c´. If the temperature is
kept higher than Af, the martensite phase can be induced by the direct application of
load and is called stress induced martensite. The stress induced martensite contains
the most favorable variant as shown in 2.1(c). By applying the load the
transformation strain is created and the volume of the material is increased. The
stress induced martensite is thermodynamically unstable above Af. As a results when
the applied load is removed the materials reverts back to austenite phase and restores
the original shape. This phenomenon is called superelasticity.
2.4 Crystallography of NiTi-based shape memory alloys
Since the discovery of NiTi shape memory alloys in 1962, extensive investigations
on the important aspects of this alloy i.e. mechanism of shape memory effect,
crystallography, thermo-mechanical behavior, ternary alloying additions, fabrication
methods etc. have been carried out and a significant understanding of this system has
been achieved.
In binary NiTi based shape memory alloys the martensitic transformation occurs due
to the solid to solid phase transformation of B2-parent phase into a B19´ martensite
phase. The parent phase has a cubic-B2 ordered structure. The lattice parameter of
the parent phase is a˳ = 0.3015nm. The B19´ martensite phase on the other hand
possesses a monoclinic crystal structure, with the following lattice parameter: a =
0.2898nm, b = 0.4108 nm, c = 0.4646 nm and β = 97.78° [18, 19]. Other than a direct
transformation route from the cubic-B2 parent phase to monoclinic-B19´ martensite
phase, the parent phase can also adopt a two stage transformation path either by first
transforming into an R-phase and then into monoclinic-B19´ phase or into an
orthorhombic-B19 phase and then into monoclinic-B19´ phase as shown in Fig. 2.2
[6].
The appearance of R-phase transformation is observed mainly due to the effects of
composition (in ternary TiNiFe and TiNiAl alloys), aging (low temperature aging of
Ni rich NiTi alloys) or thermomechanical treatment of near equiatomic NiTi alloys
[20-27].
15
Fig. 2.2 Different transformation paths in NiTi based alloys [6]
R-phase possesses a trigonal crystal structure and exhibits a very narrow
transformation temperature hysteresis (about 1 – 3°C) [28]. The R-phase
transformation is martensitic and thermoelastic in nature and it gives rise to shape
memory effect and superelasticity [29, 30]. The appearance of R-phase
transformation actually suppresses the transformation to B19´ martensite and as a
result the Ms temperature decreases. By adding certain ternary alloying elements in
NiTi alloys i.e. Pd, Pt, Au and Cu, the crystal structure of the martensite phase can be
changed from monoclinic to orthorhombic. In case of orthorhombic martensite; the β
angle is reduced to 90° [6, 31, 32].
2.5 High temperature shape memory alloys
Current practical uses of SMAs are limited to temperatures below 100°C. This is the
transformation temperature limit of the two most commercially successful SMA
systems i.e. the near equiatomic NiTi binary and Cu based ternary alloys. During
thermomechanical processes required to produce stable shape memory or
superelastic behavior, the transformation temperatures are further reduced. Naturally,
such limitation obstructs the utility of SMAs in high temperature applications. On the
other hand, the unique properties of SMAs become more beneficial at high
temperatures, since it is preferable to adopt single piece adaptive and multifunctional
components over more complex multicomponent assemblies due to the higher
probability of wear or damage and the greater weight and volume required by the
latter. These issues have activated several studies on possible SMAs with
transformation temperatures above 100°C. This class of materials is simply referred
to as high temperature shape memory alloys (HTSMAs). Till now, despite intensive
16
research efforts in recent years, HTSMAs have yet to be utilized commercially in
appreciable amounts due to a number of unresolved issues.
In recent years many material scientists have paid close attention to the development
of SMAs which can operate at temperatures above 100°C because of the demand
from automotive, robotic and aerospace industries [33]. One of the simplest
examples of SME is unconstrained recovery in which the SMA wire recovers its
original shape after being deformed or twisted in any shape. The same concept has
been utilized to put forward SMA based antenna arrays. The antennae for a satellite
or a space rover can be rolled into a compact state for storage while the unit is being
transported, but when the unit is in action the antenna can be deployed by passing a
small amount of current through the SMA based antenna to prepare the antenna to
function. Most orbital satellite utilizing this application would also require a high
temperature shape memory alloy with transformation temperatures more than 100°C
to avoid unintended actuation in case of exposure to direct sunlight. Another
application area of SMAs is during constrained recovery in which the SMA
component is prevented from going back to its original shape upon heating, which
usually generates very high stresses [34]. This property can be used for fastening
multiple components, connecting pipes or tubes etc.
Due to material property degradation of currently available SMAs, they cannot be
used for above mentioned application because of their risk of failure due to material
strength [34].The high demand for HTSMAs is from the interest of utilizing them in
solid state actuators. SMA based actuators have been found to have higher energy
densities than pneumatic actuators and DC motors and at par with hydraulics while
having the advantage of weighing significantly less. Along with the advantage of
being lighter in weight, being frictionless and noiseless make SMA based solid state
actuators a formidable candidate for application in weight critical systems such as jet
turbine engines, spaceships and any other aerospace application. Hence, the
development of HTSMA will play a pivotal role in the realization of the use of SMA
in real life applications in aerospace such as clearance control in the compressor and
turbine section of jet engine, variable area and geometry inlets for subsonic jets, self-
damping components in fuel line clamps, down-well flow control valves, electrical
appliances and actuators close to engine parts in automobiles. Currently available
17
NiTi based SMAs can only operate at temperatures lower than 100°C due to their
low transformation temperatures or poor mechanical behavior at high temperatures
[34]. In the mentioned higher temperature range the martensitic transformation
occurs at a temperature range where the diffusion controlled processes like
decomposition, recrystallization, recovery, etc. could take place and might
deteriorate the shape memory behavior.
Some of the most common approaches suggested to meet the industry requirements
of HTSMAs include the following [35, 36].
Development of new metallic alloys with high transformation temperatures.
Addition of ternary or quaternary element(s).
Optimization of various thermo-mechanical treatments and process of SMAs.
To distinguish HTSMAs from SMAs, many definitions have been put forward. One
of the most widely accepted definition is given by Firstov et al. [33] where he
describes HTSMAs as alloys of which the reverse transformation temperature As
starts above 117°C in stress free conditions following any thermomechanical
treatment. Some of the most extensively researched HTSMAs are: Ni-based [37-42];
Co-based [43-46]; Zr-based [47-53]; Ti-based [54-57]; Pt-based TiPt, TiPtIr [58, 59];
Cu-based [60-64] TaRu, NbRu [65, 66] and NiTi-based TiNiPd, TiNiPt, TiNiAu,
TiNiZr, TiNiHf [67-71]. Some of the above mentioned HTSMAs are studied in detail
as under.
2.5.1 Nickel (Ni) based high temperature shape memory alloys
Ni based HTSMAs have a long history, being the first SMAs to be extensively
studied with transformation temperatures above 100°C [72, 73] . The two main
systems are NiAl and NiMn, and both have been thoroughly investigated. Relatively
low cost and a fairly thorough knowledge of related structural materials have
provided a firm basis for the study of Ni based HTSMAs.
18
2.5.1.1 Nickel Aluminum (NiAl) high temperature shape memory alloys
In 1971, Enami and Nenno first identified both shape memory and superelastic
behavior in Ni63.2Al36.8 (at%) [37]. Au and Wayman showed in 1972 that ΔT was
only 10°C and the transformation was thermoelastic. As shown in Fig. 2.3, the Ms of
NiAl increases with decreasing Al concentration by ~130°C/at%Al and reaches
100°C at Ni64Al36 [74] and 370°C in Ni66Al34 [72]. Ternary alloying in small
amounts up to 5 – 10 at% with V, Cr, Mn, Fe, Zr, Mo, Ta or W decreases
transformation temperatures moderately, alloying with Nb, Ti or Si decrease them
significantly [75] while Co, Cu or Ag additions have no effect on or slightly increase
transformation temperatures [73]. Maximum εrec in binary solution treated Ni63.1Al36.9
was reported to be 3 – 4% by Au and Wayman. When aluminum is removed to
increase transformation temperatures above room temperature, the shape memory
properties become very poor in binary NiAl alloys. No more than 0.5% εrec was
possible in solution treated Ni63.8Al36.2 prepared by powder metallurgy [40]. Even in
this small strain range, recovery rate is commonly less than 50%.
Fig. 2.3 Dependence of Ms temperature on aluminum content in NiAl alloy [73]
19
The lack of satisfactory shape memory behavior in NiAl HTSMAs has several roots;
poor tensile ductility, phase decomposition processes, and texture. While Jee et al.
[40] reported 9.5% strain to failure in solution treated Ni63.8Al36.2 in compression.
Kim et al. [72] was only able to achieve 0.5% strain to fracture in bending for
solution treated Ni63.5Al36.5, which failed by transgranular fracture in the elastic
range. Two reasons exist for this poor ductility; easy intergranular fracture due to
large grain sizes and inherently poor fracture toughness, and transgranular fracture,
due to the insufficient number of slip systems in the β´ martensite to accommodate
polycrystalline deformation at room temperature [76]. Ternary additions of Fe and
Co are highly effective in increasing ductility [77] mainly as a consequence of the
precipitation of ductile γ fcc phase. Although boron addition suppresses intergranular
failure, it does not improve ductility by itself because the alloy will fail instead by
transgranular fracture. A combination of boron and Fe or Co addition is most
effective in increasing ductility since both intergranular and transgranular failure
modes are suppressed. Grain refinement processes, such as rapid solidification, are
effective in further improving ductility in NiAl alloys containing Fe or Co.
Similar to Cu based HTSMAs, phase instabilities and decomposition cause
significant problems in NiAl SMAs. Two types of microstructural decomposition
processes are relevant. In alloys with high Al or high Fe/Co content, γ phase appears
after quenching or aging. Although the γ phase improves ductility as mentioned
previously, it does not undergo martensitic transformation and is therefore
detrimental to reversibility of the shape memory behavior [78]. The appearance of
the γ phase can be avoided by appropriate composition selection, and training can
overcome some of the deleterious effects of the γ phase in dual phase alloys.
2.5.1.2 Nickel Manganese (NiMn) high temperature shape memory alloys
The second well known NiMn alloy system is similar to the NiAl system in many
aspects. This system is based on equiatomic binary NiMn with Ms of 674°C [38].
Transformation is observed only for equiatomic NiMn and Mn rich compositions,
proceeding from B2 austenite to L10 tetragonal martensite (θ phase) [79]. Increasing
Mn content in Mn-rich compositions, decreases transformation temperatures slightly
[80]. Although the thermal hysteresis in this system is very large (75 – 150°C), the
20
martensitic transformation was observed to be thermoelastic in nature based on
microstructural studies of the resulting martensite structure. In addition, binary NiMn
alloys are extremely brittle, particularly in Mn-rich compositions such as Ni45Mn55,
where cracks were found in a specimen right after quenching.
Fig. 2.4 Dependence of martensite peak temperature, Mp on either Ni or Mn contents
in NiMnAl alloy [38]
Ternary alloying imparts profound changes to both phase transformation and
mechanical properties of binary NiMn alloys, and most are based on the NiMnAl
system, where a wide range of transformation temperatures can be obtained. It can be
noted from Fig. 2.4 that increasing the Ni content, Mp of NiMnAl alloy increases.
Although Adachi and Wayman [79] concluded that Ti, Al, Cu or B additions did not
improve ductility, results from later studies contradicted this finding. Instead, it has
been reported that elongation to failure in compression at room temperature
increased from less than 2% strain in Ni50.2Mn49.8 to above 10% for a
Ni40.3Mn50.1Ti9.6 (at%) alloy and over 15% strain in Ni39.7Mn50.3Al10 [81]. Ductility in
bending was worse than in compression in all cases, but up to 10% strain in bending
21
was still possible without cracking in Ni39.7Mn50.3Al10. The improvements due to Ti
addition were argued to be caused by a decrease in grain size, and improvements
from Al addition to be a result of the formation of the ductile γ phase. Replacement
of Ni by Al, Ti, Cu or Fe reduces transformation temperatures more than replacement
of Ni by Mn in identical amounts [82]. These ternary additions up to 10 at% also
lowered ΔT to 25°C. Addition of titanium or aluminum also led to changes in the
martensite structure. Kainuma et al. [38] indicated that in the NiMnAl system with
more than 30 at%Mn, an additional monoclinic 7R martensite can be found. Jee et al.
[40] reported that 10 at%Ti addition caused a complex 4R orthorhombic martensite
to form in the Ni–Mn–Ti system.
Shape memory properties of Ni–Mn based alloys are far superior to those of the NiAl
system. Specimens deformed in compression to 3% total strain yielded εrec of 2.8%,
and 5% εrec was obtained from 8% total applied strain in Ni45.1Mn44.9Ti10. In bending
thermal cycling experiments under constant stress of 290 MPa, 3.5% εrec (88%
recovery rate) of ~4% applied strain was also found in Ni42.5Mn50Ti7.5. Shape
memory reversibility is inferior in NiMnAl alloys compared to NiMnTi because of
non-transforming and soft γ phase precipitates. Maximum of 0.6% εrec with complete
recovery is possible in bending thermal cycling experiments of Ni39.7Mn50.3Al10, and
εrec of 1.5% is possible in the same alloy with 75% recovery rate [39]. Yang and
Mikkola [83] investigated the shape memory recovery rate of several NiMn based
ternary and quaternary alloys in the solution treated conditions under compression at
room temperature. While none of these alloys are capable of full recovery, both
NiMnTi and NiMnAl alloys showed reasonable recovery rates of around 70% up to
very high applied strain levels.
Similar to NiAl alloys, diffusion based decomposition processes also occur in NiMn
alloys when the Ni content is less than 50 at% for binary NiMn, ternary NiMnFe and
NiMnAl, and quaternary NiMnAlFe alloys [41]. Decomposition of γ phase during
aging in Ni60Mn16Al19Fe5 causes large shifts in the transformation temperatures.
Transformation temperatures in this alloy are also argued to be affected by
quenching rate, which is most likely related to the formation of γ phase during
quenching [84].
22
2.5.2 Copper (Cu) based high temperature shape memory alloys
Cu-based SMAs are the second commercialized SMAs after NiTi-based SMAs.
Although in polycrystalline form, the mechanical and shape memory properties of
Cu-based SMAs are inferior as compared to NiTi-based SMAs, however due to the
lower cost of Cu-based SMAs and ease in development process, their application is
comparable with NiTi-based SMAs in many fields. Among the two main Cu-based
binary alloy systems, (CuZn and CuAl), CuAl is considered to be more suitable
candidate for high temperature shape memory applications. As this alloy system
exhibits higher transformation temperatures and better microstructural stability as
compared to CuZn.
In the binary CuAl system, when the content of Al is reached closer to the eutectoid
concentration, the parent phase becomes unstable and decomposes to equilibrium α
and β phases. Hence for stabilization of parent phase, Ni is added. However by
ternary alloying of Ni, transformation temperature lowers by ~ 20°C/wt-% [63]. For
single crystal CuAlNi alloys, the shape memory and superelastic properties are
remarkable. In single crystal Cu81.8Al14Ni4.2 (wt-%) at deformation temperature of
205°C, perfect superelastic recovery for the applied strain of 17% is reported, due to
a distinctive stress induced martensitic transformation in two steps [60, 64]. For the
other single crystal CuAlNi alloys, complete superelastic and shape memory
recovery of about 8% is a normal behavior for different compositions as stated above
[61, 85]. However for the polycrystalline CuAlNi alloys, mechanical and shape
memory properties are significantly degraded due to transgranular fracture produced
by large grains of about 1 mm. The other problems for these alloys are the elastic
anisotropy and brittle γ phase precipitation at the grain boundaries, lowering the
ductility of material [86]. Moreover exposure at high temperature for long time
creates instability in the transformation behavior [71]. Moreover transformation is
strongly dependent on Al content as shown in Fig. 2.5. By adding Aluminum content
in place of Cu, the Ms temperature quickly drops below room temperature.
Several efforts have been made to increase the ductility by micro alloying of fourth
elements. In this regard titanium [87, 88], boron [89], vanadium [90] and zirconium
[7] have been added in small amounts, reduced the grain size significantly and
23
resultantly the improved ductility was observed. Due to reduction in diffusion of the
constituent elements, titanium is significantly effective for reducing the grain size as
compared to other elements [88]. By increasing the nickel content, precipitation of
brittle γ phase can be lowered, however significant increase can causes reduction in
ductility as well. Use of manganese in place of nickel is more effective in improving
the ductility and in lowering the brittle γ phase precipitation [91].
In solution treated condition, the shape recovery of CuAlNi alloys is poor when the
stress is applied at room temperature, however it improves by increasing the
deformation temperature [92]. Thus it was concluded that the deformation by stress
induced martensite just above Af resulted in less plastic deformation as compared to
the deformation by martensitic orientation below As [93].
Superelasticity for some CuAlNi alloys has been reported when these are deformed
above Af. It was found by Morris [91] that complete superelastic recovery can be
obtained for 2% applied strain in a solution treated Cu80.96Al12Ni4Mn3B0.04 at 150°C
and shape recovery of 4.5% can be obtained for Cu79.96Al12Ni4Mn4B0.04 at 150°C.
To improve the ductility and phase stability in CuAl binary alloys, niobium [62, 94-
96] and silver [97-99] have been added in place of nickel recently. By addition of
niobium, ductility was improved to 12.5% tensile strain at room temperature by
reducing the grain size in Cu84.44Al13Nb2.56 alloy with Ms of 305°C [95]. However, by
adding silver, the ductility was reduced to less than 1% in. Cu83.6Al10.6Ag58 alloys
[99]. Addition of either Nb or Ag reduces the decomposition process, and also
increases transformation temperatures [97, 99] Ms temperature as high as 313°C can
be reached in Cu86.23Al13.5Nb0.27 [95] , and 376°C in Cu87.5Al9.8Ag2.7 [97].
Although the increase in transformation temperatures may seem appealing at first
glance, it is important to note that the temperature for the activation of
decomposition processes is only slightly increased. For higher transformation
temperature CuAlAg alloys, the martensitic transformation vanishes after the alloys
have been heated to above their Af temperature (over 400°C), due to phase
decomposition in this temperature range [99]. Cyclic stability also suffered and as a
result the transformation temperatures shifted by 25°C after five DSC cycles in
24
CuAlNb alloys [62]. It is very difficult to suppress decomposition processes at such
high temperatures and the useful lifetime and stability of these alloys are very
limited. The shape memory behavior of these alloys are also unknown; while it was
mentioned that the Cu86.23Al13.5Nb0.27 alloy is capable of 98% recovery rate after
bending at room temperature, the applied strain level was not mentioned [95].
Fig. 2.5 Composition dependence of the martensite start temperature, Ms on
aluminum content in CuAlNi alloys [61]
Further alloying the CuAlNb system with nickel, chromium and cobalt can lower
transformation temperatures to levels where decomposition occurs more slowly, as
well as increase the intrinsic stability of the austenite and martensite [94]. While the
ductility of these quaternary alloys are worse than ternary CuAlNb alloys, tensile
elongation to failure in all compositions is still around 6–8% at room temperature
and remains superior to identically processed CuAlNi alloys [96]. Transformation
temperatures remain stable in nickel containing quaternary compositions with Af of
179°C after 30 thermal cycles, and decomposition did not affect the behavior until
after 1000 min aging at 300°C [94]. Quaternary additions have little effect on shape
25
memory behavior, with the exception of chromium, which decreases the maximum
recoverable strain. Shape memory behavior has not been studied in the CuAlAg
alloys. Other potential Cu Based SMAs include CuAlCo and CuAlZr with Mp around
250°C and thermal hysteresis of nearly 100°C and show poor cyclic and thermal
stability [7].
2.5.3 Nickel Titanium (NiTi) based high temperature shape memory alloys
NiTi alloys can be alloyed with various elements like Pd, Pt, Au, Zr, Hf, etc. to form
various ternary and quaternary alloys. The transformation temperatures (TTs) are
very strongly dependent on the composition in NiTi alloys [100]. In Ni-rich
intermetallic compounds, even a slight increase of the Ni content away from the
stoichiometry leads to a sharp decrease in the TTs while the Ti rich side of the
intermetallic compound is much less sensitive to the mentioned compositional
variation primarily due to precipitation of Ti2Ni particles which counteract the
compositional variation of the matrix. Addition of ternary element within the range
of less than 10% usually decreases the Ms or usually has very little effect.
Substitution of Fe or Co in place of Ni; or Mn, Cr, V when substituted for Ti
decreases the TTs very significantly. Addition of more than 10% Pd, Pt, Au, Hf and
Zr to NiTi increases the transformation temperature [101].
2.5.3.1 Nickel Titanium Hafnium (NiTiHf) and Nickel Titanium Zirconium (NiTiZr)
high temperature shape memory alloys
The NiTiHf and NiTiZr systems are possible alternatives to the much more costly
precious metal containing alloys like NiTiPd, NiTiPt and NiTiAu. In this case, the
ternary alloying elements (Hf and Zr) are substituted at the expense of titanium. A
prime motivation behind studying NiTi (Hf,Zr) alloys is the relatively low raw
material cost of Hf and Zr, at least compared to Pt, Pd and Au, and their greater
influence on transformation temperatures, which allows them to be used in smaller
concentrations. Since the mechanical and shape memory properties of NiTiHf and
NiTiZr alloys are very similar, they will be discussed concurrently in this section.
Additions of hafnium above 3 at% as shown in Fig. 2.6, increase the transformation
temperatures of the binary NiTi system [102]. This increase is only ~5°C/at% up to
26
5–10 at% hafnium, then elevates to over 20°C/at% thereafter. In NiTiZr alloys, any
increase in transformation temperatures occurs above 10 at% zirconium as shown in
Fig. 2.7. At this level, further zirconium addition increases Ms at a rate of ~18°C/at%
[53]. Therefore, hafnium elevates transformation temperatures more effectively than
does zirconium at identical concentrations [93]. In the solution treated state, the
martensitic peak temperature Mp is 190°C in Ni49Ti36Hf15 and 160°C for
Ni49.5Ti35.5Zr15 [103]. Transformation temperatures are not notably affected by a
change in nickel content in either system as long as the alloys were (Ti,Hf/Zr) rich,
but dropped steeply when nickel content is increased beyond the equiatomic (50 at%)
composition [104] consistent with the behavior of Ni–Ti alloys in general.
Fig. 2.6 Composition dependence of the martensite peak temperature, Mp as a
function of hafnium content in NiTiHf alloys [102]
27
Most research efforts have focused on (Ti,Hf/Zr) rich alloys with under 20 at%Hf/Zr.
In these alloys, a single stage transformation of cubic B2 parent phase to monoclinic
B19´ martensite takes place, [105] whereas in higher Hf/Zr content compositions
approximately above 20–25 at%, orthorhombic B19 martensite appears. Ti2Ni, [106]
or Ti4Ni2Ox (x˂1) type precipitates are found in (Ti,Hf) rich NiTiHf alloys depending
on interstitial oxygen level. The situation is more complicated in NiTiZr alloys. In
addition to Ti2Ni type precipitates, a phase with relatively low melting point named
λ1 appears in alloys in excess of 10 at%Zr in (Ti,Zr) rich compositions when cooled
slowly from solution [107]. NiTiZr appears to be physically less stable than NiTiHf
alloys with identical Ti and Ni content, and the volume fraction of the precipitates in
the NiTiZr alloys is generally higher. In both systems, precipitate volume fraction
increases as alloy composition moves deeper into the (Ti, Hf/Zr) rich region of the
ternary phase diagram. Both NiTiHf and NiTiZr are somewhat brittle in tension at
room temperature [108]. For example, Ni49Ti36Hf15 failed in tension at 7% strain
[109] at room temperature; fortunately, ductility is improved at higher temperatures,
and in the same alloy deformed at 260°C in full austenite, fracture did not occur until
30% tensile strain. Ductility worsens sharply with increasing Hf and Zr
concentrations [110]. Addition of 0.1 at% boron did not improve the ductility of
Ni49Ti36Hf15. Shape memory behavior of NiTiZr alloys is in general inferior to those
of NiTiHf alloys. Full recovery of up to 1.8% applied strain in bending at room
temperature was possible for Ni50Ti35Zr15, and up to 1.6% applied strain in bending
could be fully recovered in solution treated Ni50Ti30Zr20 alloys. A full recovery of
1.8% applied strain in Ni50Ti35Zr15 and an 85% recovery rate for 2.8% applied strain
in solution treated Ni49.5Ti35.5Zr15 were also observed after room temperature
deformation in compression.
Meng et al. [106] reported that aging at 700°C severely lowered the transformation
temperatures of Ni49Ti36Hf15. After 20 hours aging, Ms and As dropped by
approximately 70°C and 40°C respectively; these changes were accompanied by the
precipitation of the (Ti,Hf)2Ni phase. An increased σy and ultimate tensile strength in
martensite of the aged alloy provide further evidence of precipitate formation.
During the course of aging, these precipitates grew in size and volume fraction, and
the authors concluded that 20 h were needed for precipitates to reach the peak aged
condition and impart the best improvement in mechanical and shape memory
28
properties. Further aging coarsened the precipitates, and the strengthening effect was
lost.
Fig. 2.7 Composition dependence of the martensite peak temperature, Mp as a
function of zirconium content in NiTiZr alloys [102]
Aging in Ti rich compositions of NiTiHf alloys have shown a clear improvement in
recovery rate due to precipitation hardening when the alloy is near or at the peak
aged condition. Although Meng et al. [106] noted that shape memory recovery of
the material after aging was inferior to that found by Wang et al. [108], the initial
applied strain was larger in the former study, and improvements of recoverability of
the aged specimen over an unaged one was clearly demonstrated. More recently,
Meng et al. [36, 111] successfully achieved high transformation temperatures in Ni
rich NiTiHf SMAs through aging.
One other concern with the NiTiHf/Zr HTSMAs is that they generally have wide ΔT,
an unfavorable characteristic for actuator applications. Similar to binary NiTi alloys,
the martensite structure of the composition range of NiTiHf/Zr currently being
studied as HTSMAs is monoclinic B19´. Applying a similar strategy that proved
29
effective for the reduction of ΔT in binary NiTi [6], Meng et al. [112] alloyed NiTiHf
HTSMAs with copper and sought to change the martensite structure from B19´ to
B19, but found copper additions of 1 to 5 at% to Ni49Ti36Hf15 did not change the
martensite structure, and actually increased ΔT instead. In addition, the 5 at% copper
addition had no effect on transformation temperatures while 3% copper caused a
slight decrease. Although stress free thermal cycling in the alloy with 5 at% copper
produced double transformation peaks in DSC during heating, which was reported to
be R-phase transformation from XRD and DSC analyses, there was no obvious
improvement on the cyclic thermal stability in terms of the shift in transformation
temperatures [112].
2.5.3.2 Titanium Nickel Platinum (TiNiPt) high temperature shape memory alloys
Transformation temperatures in the TiNiPt system can be increased by replacing
nickel with platinum, but only after a threshold value of approximately 10–15% Pt is
reached [113] as shown in Fig. 2.8. At ≤10% Pt, the martensite structure is B19´
(monoclinic), and transformation temperatures are relatively insensitive to Pt content
or decreases slightly with a minimum observed at 5 – 10 at%Pt. However, at higher
levels of Pt, at least 16 at% or greater, the martensite formed is B19 (orthorhombic)
[114] and transformation temperatures increase linearly with Pt concentration until
transformation temperature near 1000°C is reached for the binary TiPt alloy. The
transformation temperatures for (Ni,Pt) rich compositions do not decline sharply with
deviations from stoichiometry and a (Ni,Pt)3Ti2 precipitate phase readily occurs in
alloys slowly cooled from elevated temperatures.
Lindquist and Wayman [115] recorded a rather small ΔT for their TiNiPt alloys,
except for the highest Pt content studied at 30%. This particular alloy exhibited ΔT
of over 80°C, which is consistent with the later data from Rios et al [113] who
reported a rather wide ΔT (40 – 80°C) for alloys containing 30% or more Pt. On the
other hand, ΔT is generally less than 20°C for alloys with 25% or less Pt, which is
desirable in any applications requiring active control of the transformation or where
cycle time and frequency is critical. Lindquist and Wayman [115] attempted to
measure the recoverable shape memory strains in TiNiPt alloys but were limited by
the extremely low room temperature ductility of their alloys. Hosoda et al. [69] have
30
also measured the room temperature tensile properties of Ti50Ni50-xPtx alloys from 0
to 50 at% Pt and found that ductility generally decreased as the amount of Pt
increased, but that σy followed a much more complicated dependence on
composition. Nickel actually appears to harden TiPt alloys considerably, while Pt
additions to NiTi soften the monoclinic martensite phase, at least initially. Compared
to the TiNiPd alloys, the stress free shape memory response of TiNiPt alloys has not
received a similar level of attention, with the exception of the attempts by Lindquist
and Wayman [115] in tension and Meisner and Sivokha [114] for a Ti50Ni34Pt16 alloy
in torsion. Instead, more effort has been focused toward the constant stress thermal
cycling behavior of TiNiPt alloys including the properties of Ti50Ni30Pt20 bar, fine
rod and wire. Alloys with transformation temperatures up to ~300°C show
reasonable work output with good reversibility for near stoichiometric compositions
and Pt contents up to 20%, even in the as processed, i.e. hot extruded, condition with
no prior training.
Fig. 2.8 Composition dependence of the transformation temperature as a function of
platinum content in TiNiPt alloys [113]
As with the other precious metal additions to NiTi, the high initial material costs will
have to be offset by the increased benefits in using HTSMAs at the system level,
31
probably limiting bulk material to high performance and demanding applications,
such as those existing in the aerospace industry. While the data in most cases are still
cursory, the TiNiPt alloys with 20% Pt or less show promise for such demanding
applications, since they have high transformation temperatures, good work output
and dimensional stability below ~350°C, low ΔT, good thermal stability and
excellent oxidation resistance below ~500°C. Certain optimization studies are
required to be performed on the TiNiPt system to further improve the properties and
consider for commercial consideration.
2.5.3.3 Titanium Nickel Paladium (TiNiPd) high temperature shape memory alloys
Interest in the TiNiPd system as potential HTSMAs was derived from three sets of
studies: the comprehensive study of phase transformations in binary B2 titanium
alloys [116], the discovery of high transformation temperatures in the TiPd binary
systems by Donkersloot [117] and the discovery of ternary alloying effects on the
transformation temperatures of binary NiTi SMAs by Eckelmeyer [9]. Based on the
results from these studies, palladium was added to the TiNi system in order to
increase transformation temperatures, and nickel to the TiPd system to improve
shape memory behavior. TiNiPd HTSMAs have received the most rigorous attention
over the years. Initial focus was centered on improving their high temperature shape
memory behavior, but more recently, the focus has shifted towards improving their
work output, as well as dimensional and microstructural stability. In this system,
transformation temperatures can be altered by replacing nickel with palladium. If the
concentration of titanium is held constant at nearly 50%, the relationship between the
transformation temperatures and relative concentration of nickel and palladium is
parabolic, as shown in Fig. 2.9. A minimum in transformation temperatures occurs at
approximately 10% Pd, although the exact composition of this minimum is still
subjected to debate. In compositions with palladium concentrations greater than the
palladium concentration at this minimum, replacing nickel with palladium increases
transformation temperatures by approximately 15°C/at%. [68, 118, 119]. On the
other hand, if the concentration of palladium is lower than the composition at the
minimum, replacing nickel with palladium actually lowers the transformation
temperatures by 4°C/at% [115]. This parabolic dependence of the transformation
temperatures on composition stems from the change in the structure of martensite.
32
On the higher palladium concentration side of the minimum, B2 austenite transforms
to B19 orthorhombic martensite, and at the lower palladium concentration side, it
transforms to B19´ monoclinic martensite or R phase. The composition at which the
transformation temperatures are at a minimum corresponds to the point of the
structure transition. Because of the complete mutual miscibility of the TiNi and TiPd
systems, the relationships between transformation temperatures and composition
hold over all ranges of palladium concentration. This enables the access to a
continuous range of transformation temperatures from room temperature to over
500°C by adjusting the amount of palladium in the alloy. Otsuka et al [120] reported
poor shape recovery of binary Ti50Pd50 HTSMA, concluding that this was primarily
caused by the low σy of both austenite and martensite. They proposed three possible
solutions to improve its shape memory behavior by strengthening the material
through:
Solid solution hardening via ternary alloying
Thermomechanical processing
Precipitation hardening or some combination of these approaches.
As the following discussion demonstrates, these solutions have met with varying
degrees of success. While the main effect of nickel addition to Ti–Pd alloys is the
reduction in transformation temperatures, it also has an indirect effect of improving
shape memory behavior by lowering the temperature range at which the alloy would
be used. At a lower temperature, σy is at a higher level. Khachin [116] observed full
recovery of 4% applied strain in high temperature torsion experiments on
Ti50Ni13Pd37. However, Lindquist and Wayman [115] studied the same alloy at room
temperature under tension, and were only able to obtain 40% recovery of 6% applied
strain. The reason for this discrepancy was not resolved, but it is likely due to the
differences in the way the materials were processed and tested, which can have a
remarkable effect on recovery rate in TiNiPd alloys. Although the σy of TiPd alloys
can be raised indirectly through nickel addition, it is still too low for perfect shape
recovery.
33
Fig. 2.9 Change in martensite start temperature; Ms with respect to palladium content
in equi-atomic TiNi50-xPdx alloys [68]
Several researchers addressed the issue of low σy in TiNiPd HTSMAs by further
solid solution strengthening. Suzuki et al. [68] and Yang et al. [121] added small
amounts of boron to Ti50Pd30Ni20 and Ti50.7Ni22.3Pd27 alloys, respectively, but neither
caused a notable reduction in εirr with boron additions of up to 0.2%. Micrometer
sized Ti2B particles were found along the grain boundaries, but they were too large
and non-uniformly distributed to possibly function as particles for precipitate
hardening. Instead, boron acted as a grain refiner by reducing the grain growth rate in
these alloys. In identically solution treated or hot rolled specimens, 0.12 at% boron
reduced the grain size from ~40 μm down to 10 μm in the Ti50.7Ni22.3Pd27 HTSMA.
Presumably for the same reason, 0.2% boron addition to Ti50Ni30Pd20 doubled the
room temperature tensile elongation to failure from 8 to 16% strain and increased the
ultimate tensile strength from 460 to 800 MPa for a sample deformed in martensite at
a temperature of 170°C. It is not clear why the grain refinement effect of boron did
not improve the shape memory behavior, since σy increases with grain size
refinement. In other studies on solid solution hardening, 5% gold or platinum
34
replacing palladium in Ti50.5Ni19.5Pd30 increased σy and mildly enhanced cyclic
stability, but had little effect on total recoverable strain. From these results, it appears
that increasing σy alone will not necessarily always result in improvements in shape
memory behavior. Another recent study similarly observed little impact on the shape
memory recovery rate after alloying Ti50.6Ni19.4Pd30 with 1 wt% cerium [122], no
reason for this observation was given.
Although boron and cerium do not appear to improve the shape memory behavior, a
recent study by Atli et al [123] showed that 0.5 at% scandium addition to
Ti50.5Ni24.5Pd25, replacing titanium, was more effective in this aspect. Although the
scandium lowered transformation temperatures by about 6 – 10°C, it was able to
reduce εirr from constant stress thermal cycling experiments under tension by half at
stress levels above 200 MPa without adversely affecting εrec. This improvement was
believed to be caused by the solution hardening effect of scandium, but the effect of
scandium on the σy of martensite and austenite was not explicitly shown.
Additionally, scandium addition also reduced ΔT and improved cyclic stability. After
10 thermal cycles at 200 MPa, the cumulative εirr was reduced by ~20% in the
scandium containing specimen.
Investigating TixPd30Ni70-x (x = 48.5 to 51.0%) alloys encompassing both sides of the
equiatomic composition, Shimizu et al. [124] realized (Fig. 2.10) that with
decreasing titanium content, transformation temperatures decrease only slightly in Ti
rich compositions, but drops off dramatically on the Ni/Pd rich side, such that Ms
declines to room temperature in a Ti48.5Pd30Ni21.5 alloy. This is similar to the
composition dependence of transformation temperatures in binary NiTi SMAs, and is
rationalized by the solubility of excess titanium or nickel atoms near the equiatomic
composition.
35
Fig. 2.10 Martensite start temperature; Ms as a function of Ti/(Ni,Pd) ratio [125]
Solubility for extra titanium atoms in near equiatomic NiTi is almost negligible, and
although the solubility for extra nickel atoms is also small, it is possible to
accommodate excess Ni concentrations under 1 at% in solution. Therefore, in
titanium rich compositions, the extra titanium atoms tend to immediately form
second phases, and do not affect the composition of the matrix and thus the
transformation temperatures of the alloy. However, a small concentration of extra
nickel atoms can dissolve in the matrix, changing its composition and transformation
temperatures. Utilizing precipitates in Ti rich compositions after proper heat
treatments, Shimizu et al. [124] were able to demonstrate that Ti50.6Pd30Ni19.4
outperformed the equiatomic Ti50Pd30Ni20 alloys with a recovery rate of 90% versus
78% respectively, on samples deformed to 6% applied strain at 200°C under tension.
The improvements were attributed to the homogeneous distribution of fine
Ti2(Ni,Pd) type precipitates formed during the annealing process.
Shirakawa et al. [126] and Nagasako et al. [127] studied the phase transformations
in the (Ni,Pd) rich TiNiPd alloys, and compensated for the decrease in
36
transformation temperature from the nickel rich matrix by increasing the palladium
content. While the two way shape memory behavior of a Ti48Pd31Ni21 alloy annealed
at 400°C was qualitatively described, no quantitative data on either one way or two
way shape memory behavior were available. Furthermore, it should also be possible
to maintain reasonably high transformation temperatures in nickel rich compositions
without increasing Pd content by applying appropriate heat treatments. The reported
precipitates are nickel rich in nickel rich compositions, and were identified to be
Ti2(Ni,Pd)3 and Ti3(Ni,Pd)4 types, so as these precipitates grow, nickel content in the
matrix should be reduced, thereby increasing transformation temperatures.
Martensite reorientation, and thus the shape memory behavior was absent in 21–29%
cold rolled Ti50Pd30Ni20 alloy during subsequent deformation at 173°C when post-
rolling annealing temperatures were below As. Moreover, shape memory behavior
deteriorated if the post-rolling annealing temperatures were very high. Golberg et al.
[128] concluded that annealing at a temperature below As does not allow the
preferentially oriented martensite inherited from the rolling process to reset to a self-
accommodated martensitic morphology. Conversely, annealing far above Af initiated
microstructural recovery, destroying the work hardening effects of cold rolling. An
ideal post-rolling annealing temperature was proposed to be above As but below the
recrystallization temperature. A follow-up study by Xu et al. [119] examined the
recovery and recrystallization processes over a full range of cold rolled and annealed
Ti50(Ni,Pd)50 alloys with palladium contents ranging from 0 to 50 at%. They
observed that recovery started at 450°C in Ti50Pd30Ni20 alloy and recrystallization
began at 550°C in Ti50Pd40Ni10 alloy. Physically, these diffusion driven mechanisms
reduce εrec by increasing grain size and reducing dislocation density from cold
working. If recrystallization takes place in martensite, reverse transformation
temperatures (As and Af) increase due to the loss of internal twins in the martensite
[129]. This occurs because during phase transformation, elastic energy is stored in
the internal twins of martensite as elastic strain. Recrystallization replaces these
martensites with the ‘new’ strain free martensite, and the stored elastic energy is lost.
Since stored energy is a driving force for the reverse transformation, its loss results in
the increase in As and Af.
In particular, recovery and recrystallization temperatures in the range of 450 – 600°C
prevent effective thermomechanical treatment of alloys with high Pd contents and
37
high transformation temperatures since recovery would occur within the operational
temperature range of the alloy. Recovery should also cause thermomechanical
training and TWSME in these HTSMAs to be almost impossible. This factor has
severely limited the development of higher temperature TiNiPd HTSMAs with Pd
contents beyond about 35 at%. Similarly, thermal cycling can also affect the stability
and properties of SMAs either by introducing a buildup of internal stresses or
through their relaxation. For example, thermal cycling of Ti50.6Pd30Ni19.4 under 200
MPa increased transformation temperatures by up to 16°C [130] and stabilized after
about 30 cycles.
In contrast, thermal cycling under no stress reduced transformation temperatures by 5
– 10°C, with alloys of higher Pd content exhibiting a greater shift in transformation
temperatures during cycling [131]. It was suggested that dislocations introduced
during cycling suppressed the martensitic transformation during stress free thermal
cycling, but the oriented internal stress fields generated from isobaric thermal cycling
assisted the external applied stress, increasing transformation temperatures [130].
TiNiPd alloys containing less than ~35% Pd have acceptable work characteristics but
tend to suffer from rather high εirr during thermal cycling under load. This
irreversibility worsens with increasing applied stress and increasing Pd concentration
at constant stress. Therefore, a significant focus of the recent research on TiNiPd
alloys has been in pursuit of mechanisms that reduce εirr. In particular, Bigelow et al
[132] have found that alloying with Pt or Au reduced εirr considerably in the absence
of thermomechanical training, but no significant reduction of εirr from these alloying
elements occurred in trained specimens. Not only are these approaches useful in
reducing εirr in TiNiPd alloys, they also have no detrimental effect on recoverable
strain and thus relative work output for these materials.
The amount of εirr for a given alloy and stress level increases as the maximum
temperature to which the sample is thermally cycled increases, with all other
considerations being constant, thus placing another criterion on the upper
temperature limit for these materials. However, certain alloying additions such as 5%
Pt increases the maximum temperature capability of TiNiPd alloys by ~30°C, thus
providing some measure of protection against overheating.
38
Fig. 2.11 (a) Specific work output for a Ti50.5Ni19.5Pd30 alloy as a function of applied
stress loaded in both tension and compression and (b) the corresponding
transformation strain versus applied stress [125]
Similar to conventional NiTi SMAs, the work output for TiNiPd alloys peaks at
some optimum stress level [125] as shown in Fig. 2.11a. This peak in work output is
a result of competing factors. As the stress level increases, more martensite variants
become favorably aligned and detwinned, forming more single variant morphology
and accommodating larger strains. At the same time, the increasing stress level
approaches σy of the alloy and causes irrecoverable deformation to occur instead of
martensitic transformation, and negatively affects the recoverability of martensite.
Consequently, these factors cause εirr to reach a maximum and then rapidly decrease
with increasing stress as shown in Fig. 2.11b for the Ti50.5Ni19.5Pd30 alloy. Therefore,
even though the applied stress continues to increase, εirr diminishes and its product
with applies stress, which is work output, reaches a maximum at a particular stress
level as well.
The maximum work output for various TiNiPd alloys as a function of the
transformation temperatures range, such that the full range over which the
transformation occurs (Mf to Af) for each alloy is clearly indicated [133], is shown in
Fig. 2.12. Large and consistent work output of between 8 and 11 J/cm3 is achieved
for alloys with transformation temperatures between 100 and 300°C. However, for
39
alloys with transformation temperatures above this range, the work output drops off
almost catastrophically so that alloys with transformation temperatures in the
neighborhood of 500°C are capable of essentially zero work output, placing another
limit on the viability of HTSMAs beyond just the need for a high transformation
temperature. This is a natural consequence of the reduction of σy at higher
temperatures and the onset of creep deformation [134]. The experiments by Bigelow
et al [135] are very useful in revealing the critical factors that would lead to the
development of superior alloys for actuator type applications. These are: a low σDT; a
high σy to any types of slip or other processes that would lead to irreversible
deformation of the martensite phase, and a high σy of austenite. Therefore, the key to
develop dimensionally stable HTSMAs with good work output is to prevent the
plastic deformation processes while not greatly affecting reorientation/detwinning
stress.
Fig. 2.12 Work output for a series of TiNiPd and TiNiPt alloys as function of the
transformation temperature range (Mf to Af) [133]
The primary methods for accomplishing this are the same as those originally
suggested by Otsuka et al. [136] for improving the stress free shape memory
40
behavior of TiPd alloys, namely: solid solution hardening, the use of
thermomechanical processing, and precipitate strengthening.
The only study on the direct effect of creep on the shape memory behavior was
carried out on Ti50Ni10Pd40 alloy with Ms of 388°C and Af of 411°C [134].
Significant amount of εirr was accumulated during tensile thermal cycles between 300
and 520°C under 200 MPa. When the heating and cooling rate was decreased from
20 to 2°C/min, εirr after one heating and cooling cycle almost doubled; an indication
of high creep activity. When the stress level was reduced to 100 MPa for heating
cooling cycles conducted within the same temperature range, εirr levels were much
lower, and did not appear to be a function of heating cooling rate. This finding was
further supported by constant temperature creep experiments under 100 and 200 MPa
at 520°C. After 2.5 hours, total creep strain of 1.2% was recorded for the specimen
under 100 MPa, while the specimen under 200 MPa showed creep strain of 12.4%.
Clearly, the combination of 200 MPa and maximum temperature of 520°C was
sufficient to activate unacceptable levels of creep activity. Although the effect of
creep under 100 MPa is comparatively much smaller, the net effect remains very
significant. Because operating temperatures in SMA components are dictated by the
transformation temperature, the Ti50Ni10Pd40 HTSMA cannot be used above 100
MPa in an actuator type application.
In summary, extensive research and development of TiNiPd alloys have produced
material with reasonable shape memory behavior, low thermal hysteresis and good
work output. On the other hand, reversibility and corresponding dimensional stability
still need to be further addressed and are the primary focus of several continuing
research efforts. However, the current system is an extremely viable alloy for
HTSMA applications in the 150 – 300°C ranges.
41
Chapter – 3
Equipment Setup and Material
Processing
3.1 Introduction
This chapter provides detailed description about the constituent elements, its
processing, equipment setup and different experimental procedures followed during
the course of the present research. A total of four alloys (SMAs) were developed by
using non-consumable electric arc furnace in the presence of argon. The prepared
alloys were then homogenized and water quenched. The homogenized button was
sliced into 0.4 mm thick strips by wire electrical discharge machine (wire-EDM).
The 0.4 mm thick strips were cold rolled by 25% and reduced their thickness to 0.3
mm. Samples for differential scanning calorimetry (DSC), X-ray diffraction (XRD),
microstructural studies, mechanical properties and shape memory measurement were
cut using wire EDM from each alloy. The detail of each process is presented in the
subsequent sections.
3.2 Development of shape memory alloys
The alloys were developed by melting of pure constituent elements. High purity
constituent elements; 99.98 wt.% Ti, 99.98 wt.% Ni, 99.99 wt.% Pd and 99.99 wt.%
Cu were used for the preparation of these alloys. The mass of constituent elements
was taken exactly as calculated from the conversion formula, i.e. atomic percent to
weight percent up to four decimal points of a gram i.e. one-tenth of mg. The
sequence of operation and material processing is shown in Fig. 3.1 and explained in
the following sections.
42
Fig. 3.1 Process flow chart of the sequence of operations for materials processing and
characterization
Mechanical Testing (Hardness + Stress-Strain Diagrams)
Thermo-mechanical Cycling Experiments (Stress of 100 – 500 MPa)
X-Ray Diffractometry (XRD)
Differential Scanning Calorimetry (DSC)
Preparation of Samples for Scanning Electron Microscopy (Polishing)
Preparation of Samples for Optical Microscopy (Polishing + Etching)
Solution Treatment of Samples at 900°C for 1 h and Water Quenched
Preparation of Samples for Different Characterization (Wire EDM)
Homogenization of Cast Buttons at 950°C for 2 h and Water Quenched
Melting of Constituent Elements in Vacuum Arc Melting Furnace
Precise Weighing of Constituent Elements (one-tenth of mg)
Cleaning of Constituent Elements by Ultrasonic Cleaner (in CH3OH)
Arrangement of Pure elements Ti, Ni, Pd, Cu (99.98 wt.% or more)
43
3.2.1 Cleaning of constituent elements
The weighed elements were cleaned by an ultrasonic cleaner in methanol (CH3OH)
for 5 minutes. After cleaning, the materials were taken to electric arc furnace for
melting.
3.2.2 Melting of constituent elements
Titanium is very reactive at high temperature and form stable chemical compounds
with oxygen or nitrogen [137, 138]. To avoid the formation of oxides or nitrides and
to get the desired composition, the melting was done by evacuating the furnace from
air (oxygen) and then partially filling it with 99.99% pure argon gas.
The alloys were prepared in the form of buttons (20 g each) in non-consumable DC
electric arc melting furnace in the presence of argon using tungsten electrode and
water cooled copper crucible. Usually the DC electric arc melting furnace is used for
melting of highly oxidizing materials like Ti. The water cooled copper crucible was
cleaned by grinding paper and acetone to remove the contamination layers from the
surface. The constituent elements were placed separately in the two hemi-spherical
cavities and pure titanium in the third cavity as getter. The function of the getter is to
absorb oxygen content and humidity from the furnace chamber before melting of the
elements. The furnace chamber was vacuumed at 1x10-2
mbar and then flushed with
99.99% pure argon gas. This process was repeated twice to remove oxygen from the
furnace chamber. For the third time the chamber was vacuumed at 1x10-4
mbar and
then small amount of argon gas was supplied to the chamber. The existence of argon
gas in the chamber is necessary to provide medium for production of electric arc
between tungsten electrode and copper crucible. Moreover it also provides shielding
medium to prevent oxidation.
Before melting the constituent elements, first the getter was preheated and melted to
absorb oxygen from the chamber. Then both the required alloys were preheated by
supplying current (60 100 A) and melted by increasing the current to 250 A, while
the voltage of 10 V was kept constant. The required alloys in form of circular buttons
with elliptical cross-section were melted six times and flipped over after each
melting to get maximum alloying homogeneity. After melting, the buttons were
44
weighed again and the mass loss was found to be less than 0.1%, therefore, the actual
composition was considered the same as before melting process i.e. nominal
composition. The same process was repeated for all four alloys. The required mass
of constituent elements for the desired four alloys is shown in Table 3.1. Hereafter
the four alloys are called 0Cu, 5Cu, 10Cu and 15Cu according to their Cu contents in
atomic percent.
Table 3.1 Chemical compositions of the four alloys given in weight percent
Alloy
Composition
(at%)
Constituent Elements (g)
Total Mass (g) Ti Ni Pd Cu
Ti50Ni25Pd25 7.3415 3.4998 8.1587 0.0000 20.0000
Ti50Ni20Pd25Cu5 7.31428 3.58646 8.12852 0.97074 20.0000
Ti50Ni15Pd25Cu10 7.2873 2.6799 8.0985 1.9343 20.0000
Ti50Ni10Pd25Cu15 7.2605 1.7800 8.0687 2.8908 20.0000
3.2.3 Homogenization
To make uniform chemical composition throughout the volume of developed alloys,
homogenization process is necessary to be carried out. Each button of the developed
alloys was kept in quartz tube, evacuated by rotary and diffusion pumps till vacuum
of 1x10-3
mbar, filled the quartz tube with pure argon gas and then sealed. Tube
furnace was heated to a temperature of 950°C and then charged the quartz tubes into
the furnace. The charging of alloy buttons at high temperature was done to avoid
precipitation process which can occur during the slow heating process. After keeping
the alloy buttons at 950°C for 2 hours, they were fast quenched in ice water to avoid
the formation of precipitates during slow cooling. To protect from oxidation, the
alloy buttons were quenched without breaking the quartz tubes. The homogenization
process is shown in Fig. 3.2.
45
3.2.4 Chemical analysis
After homogenization, the chemical composition of the four alloys was determined
by Energy Dispersive Spectroscopy (EDS). The chemical composition of impurities
(O, N, C and S) ingressed during handling and melting of constituent elements was
found by LECO Oxygen/Nitrogen Determinator and LECO Carbon/Sulphur
Determinator. The chemical compositions of all four alloys are summarized in Table
3.2.
Table 3.2 Chemical composition of homogenized alloys
Element
Composition (atomic percent)
Analysis Method 0Cu 5Cu 10Cu 15Cu
Ti 50.08 50.02 50.08 50.11
EDS Ni 25.03 19.87 15.21 9.84
Pd 23.83 25.13 23.72 23.85
Cu 0.00 3.92 9.93 15.14
O 0.022635 0.022526 0.022715 0.022526 O/N Determinator
N 0.043252 0.044535 0.043686 0.044535
C 0.005623 0.005269 0.005356 0.005269 C/S Determinator
S 0.000031 0.000027 0.000024 0.000027
3.3 Sample preparation
3.3.1 Sample dimensions
The homogenized buttons were sliced into 0.4 mm thick strips by wire Electrical
Discharge Machine (EDM). The 0.4 mm thick strips were cold rolled by 25% and
reduced their thickness to 0.3 mm. Samples for Differential Scanning Calorimetry
(DSC), X-Ray Diffractometry (XRD), microstructural studies, mechanical testing
and shape memory measurement were cut using wire EDM. Wire EDM was used for
dimensional accuracy, avoid material loss and reduce machining / cutting stresses.
Dimensions of the samples for various characterizations are shown in Table 3.3.
46
Table 3.3 Dimensions of samples for various characterizations
Characterization Length (mm) Width (mm) Thickness (mm)
DSC 2 2 0.3
XRD 10 5 0.3
Optical Microscope 10 5 0.3
SEM 10 5 0.3
Shape Memory Effect 25 3 0.3
Mechanical Properties 25 3 0.3
3.3.2 Solution treatment and aging
For solution treatment, all the prepared samples were sealed in different quartz tubes
after being evacuated and then filled with pure argon gas. Tube furnace was heated to
a temperature of 900°C and then charged with the quartz tubes. All the samples were
soaked at 900°C for 1 hour in the tube furnace and then fast quenched in cold water
without crushing the quartz tubes. The solution treatment process is shown in Fig.
3.2.
3.3.3 Aging
Tem
per
atu
re (
°C)
Time (h)
2 h at 950°C 1 h at 900°C
Ho
mogen
izti
on
So
luti
on
Tre
atm
ent
Fig. 3.2 Schematic representing homogenization and solution treatment processes
47
After solution treatment, some of the samples were aged at various aging
temperatures for investigation its properties in aged condition and comparing with
that of solution treated condition. The samples from 0Cu, 5Cu, 10Cu and 15Cu
alloys were again sealed in quartz tubes adopting the same procedure as mentioned
above. The sealed tubes were kept in the tube furnace after reaching the required
aging temperatures. All the four alloys were aged at various temperatures ranging
from 400°C to 700°C for time duration of 3 hours as shown in Fig. 3.3. After
finishing the soaking time, the samples were quenched in cold water without
breaking the quartz tubes to avoid direct contact of air and water.
Fig. 3.3 Schematic representing the aging process at various aging temperatures
3.4 Materials characterization
3.4.1 Optical microscopy
Upright Metallurgical Microscope of SINOWON manufacturer China model UMS –
300 was used for metallography. The microscope is capable to magnify the image to
1000X. The samples were first mechanically rough ground by 200, 400 grit papers
and then fine ground by 800 grit emery paper and finally fine polished by 0.5 micron
alumina slurry. After polishing, the samples were etched using reagent in the given
volume ratio 10%HF, 15%HNO3, 75%H2O to reveal the microstructure and grain
boundaries. During this process the samples were dipped in the solution for 5
Agin
g T
emp
eratu
re (
°C)
Aging Time (h)
3 h
600°C
500°C
400°C
700°C
48
minutes while stirring them regularly inside the solution. Moreover, after time
duration of each minute, the samples were washed properly by flowing water and
then dried by hot or compressed air. The samples were regularly checked in the
optical microscope to observe etching after every minute until the grain boundaries
and microstructure were become clearly visible.
3.4.2 Scanning Electron Microscopy (SEM)
For SEM, the samples were first mechanically ground using emery paper of grit
200,400 and 800. Fine polishing was carried out by using alumina slurry of 0.5
micron size. The prepared samples were then loaded in Field Emission Scanning
Electron Microscope (FE-SEM) of TESCAN made in Czech Republic, model
MIRA3 XM. It is operated at 200 V to 30 KV. It is a high resolution SEM and can
generate images of 1.2 nm while its magnification reaches to 1,000,000X at 30 KV.
Images were obtained at different magnification from various locations of the
samples. Energy Dispersive Spectroscope (EDS) attached with FE-SEM was used to
determine quantitatively the chemical composition of constituent elements in each
phase. The chemical compositions of the bulk matrix and other phases present in the
form of precipitates were analyzed by the combine study of the images of SEM and
data from EDS analysis.
3.4.3 X-Ray Diffractometry (XRD)
In the present study XRD patterns were recorded using a computer controlled X-Ray
Diffractometer model JDX-99C JEOL Japan which was operated at 40 KV and 20
mA. The radiation used was Cu Kα (λ=1.5406 Å) at room temperature with incident
angle 2Ɵ from 20° to 80°. The resulted patterns were then compared with the
standard Powder Diffraction File (PDF) of the International Centre for Diffraction
Data (ICDD) database. By comparing the d-spacing of the corresponding diffraction
angle and intensity of diffracted peak, the atomic structure of the different phases
was determined.
49
3.4.4 Differential Scanning Calorimetry (DSC)
DSC is a common technique used for determination of the thermal properties of
specimens with high degree of precision. It is a differential method which calculates
heat flow directly from the samples during heating, cooling or holding isothermally
by measuring the amount of energy required to maintain the specified temperature.
During heating or cooling, if the phase transition occurs in the sample, the heat
released or absorbed is abruptly changed and their corresponding peaks are observed.
In the present study transformation temperatures under stress free condition were
determined by using NETZCH Differential Scanning Calorimeter Germany model
214 Polyma. Using Liquid Nitrogen, it automatically controls temperature, ranging
from – 170°C to + 600°C with heating/cooling rate of 0.001 K/min to 500 K/min.
Each sample was cycled twice through full transformation range from 100°C to
250°C to ensure the reproducibility of thermoelastic transformations. It was observed
experimentally that the transformation temperatures stabilize after two complete
cycles in the solution treated samples and then remain almost constant in subsequent
cycles. The heating and cooling rate for DSC cycles was kept as 5°C per min.
Fig. 3.4 shows a typical DSC plot representing the phase transformation
temperatures. Two peaks are shown, marking the forward and reverse
transformations, where the forward transformation is exothermic and the reverse
transformation is endothermic. The temperatures at start, highest and end points of
the peak formed during forward transformation are known as martensite start,(Ms),
martensite peak(Mp) and martensite finish (Mf) temperatures respectively. Similarly
the temperatures at start, lowest and end points of peak formed during reverse
transformation are known as austenite start (As), austenite peak (Ap) and austenite
finish (Af) temperatures respectively. Phase transformation temperatures (Ms, Mf, As
and Af) of all samples are measured from the intersection of the base line and the
linear portions of exothermic or endothermic peaks as shown in Fig. 3.4. The
difference between austenite finish; Af and martensite start; Ms is called thermal
hysteresis.
50
3.5 Mechanical Testing
3.5.1 Microhardness testing
Microhardness tester of Wilson Wolpert (Model 401) was used to find out the
hardness of the samples. Pyramid indenter of 136° tip was used with 100 g load
applied for 10 seconds dwell time. The hardness values of each sample were
determined from the average of 5 measurements taken at different locations on each
sample.
3.5.2 Isothermal tensile testing
Isothermal tensile tests were performed to measure the mechanical behavior of
solution treated and aged samples taken from the four alloys. The testing was
performed on Instron Tensile Testing Machine Model 1195. In house designed and
manufactured holding grips were used to grip the samples as shown in Fig. 3.5. Load
cell of 100 KN was used and strain measurement in tension was carried out by the
movement of crosshead directly. Samples were heated by induction process and a K-
type thermocouple was attached directly to the central portion of the samples by
Fig. 3.4 Measurement scheme of transformation temperatures from DSC heating
and cooling cycles
50 100 150 200 250
He
at F
low
(w
/g)
exo
up
Temperature (°C)
Reverse Transformation
Forward Transformation
Heating
Cooling
As A
f
Mf M
s
Ap
Mp
51
wrapping the copper wire for temperature measurement. After gripping the sample
and attaching thermocouple, each sample was heated at required isothermal test
temperature while controlling load at 0 N. The samples were allowed to stay at
isothermal test temperature for 5 minutes to stabilize the thermal fluctuations. The
samples were strained at a strain rate of 1 x 10–4
mm/sec. For all alloys, only two
isothermal tensile testing were performed; one at temperature 50°C below Mf and
one at temperature 50°C above Af. All the samples were loaded continuously until
breakdown of samples, cooled down to room temperature and then unloaded.
Fig. 3.5 Special gripping arrangement for holding of 0.3 mm thick samples
Tightening Nut
Holding Jaw
Sample for mechanical and
shape memory testing Complete holding arrangement
52
3.6 Measurement of shape memory properties
3.6.1 Equipment setup
Shape memory properties were measured on Creep and Rupture Tensile Testing
System of 2330 series equipped with Applied Testing System (ATS) frame controller
and Windows Computing Creep System (WinCCS) software as shown in Figure 3.6.
Samples of 0.3 mm thickness were gripped in the sample holder as shown in Fig. 3.5.
Loading was done step by step by increasing from zero to the required maximum
load so that strain rate may be maintained as slow as possible. For tensile strain
measurement, extensometer shown in Fig. 3.6 was used. The extensometer is
consisted of dual rod-in-tube located at 180 degree apart, to which a pair of LVDT is
attached as shown in Fig. 3.6. This extensometer can give reading up to + 50% and -
20% strain in the tension and compression respectively with accuracy of thousandth
of mm. In the present experiments, the strain measurement was carried out by
loading the sample in tension.
During thermomechanical testing, the samples were heated using split tube furnace
equipped with resistive heating nichrome element which can raise the temperature of
the furnace up to 1010°C. Split tube furnace can raise the temperature of the samples
with variety of heating rate having power supply of 3300 W. In the present research,
heating rate of 10°C/min was maintained in all experimentation. The samples were
air cooled with approximate cooling rate of 102°C/min.
For temperature sensing, K-type thermocouple was used and controlled by barber
colman controller. Thermocouple was directly attached to the middle of each sample
by thin copper wire to precisely measure the sample’s temperature.
53
3.6.2 Constant stress thermal cycling tests
These tests were performed by loading the sample in tension, while keeping the
strain rate at lowest level. In these experiments the sample was loaded initially to a
stress of 50 MPa in martensite state and then heated to Af + 50°C (where Af is the
austenite finish temperature under no load). At Af + 50°C, the sample was loaded to
a stress of 100 MPa in an austenite state isothermally followed by cooling to Ms –
50°C (where Ms is the martensite start temperature under no load), while holding the
stress at constant level. The initial half heating cycle at 50 MPa was not shown in the
Fig. 3.7, because this initial behavior was not the representative of thermal cycles at
higher stresses and could not be used for evolution of shape memory characteristics.
Fig. 3.6 Internal view of lever arm creep and rupture tensile testing system
54
During cooling of sample from Af + 50°C to Ms – 50°C, initial linear response
indicate the complete austenitic phase without any deformation. However by
decreasing the temperature further, the sample was started to expand at specific
temperature and continued to remain in the expanding condition till approaching of
another specific temperature. The former and latter specific temperatures are called
martensite start and finish temperatures respectively. Between the phase
transformation temperatures, the alloy is consisted of austenite and martensite phases
and completely transformed to martensite phase at martensite finish temperature.
Upon further cooling, the strain-temperature curve was again become linear showing
no further deformation. By heating the sample from Ms – 50°C to Af + 50°C, the
initial response of the sample was found to be linear and overlap the linear portion of
cooling curve. This behavior indicates the absence of viscoplastic deformation in the
complete martensite state.
During reverse transformation process, when the temperature was increased further,
at specific temperature the sample was started to decrease in length. This behavior
was remained continuously till approaching of another specific temperature, where
the strain-temperature curve was again become linear. The temperatures at which the
shrinkage of the sample was started and then ended are called the austenite start and
finish temperatures respectively. Between the transformation temperatures, the alloy
is again consisted of martensite and austenite phases and completely transformed to
austenite phase at austenite finish temperature. When the sample was heated further,
the strain-temperature curve was again become linear. This process was repeated at
stress levels of 200, 300, 400 and 500 MPa.
Depending on the sample area, loads (in N) were calculated to produce equivalent
stress of 100, 200, 300, 400 and 500 MPa. Transformation temperatures i.e.
Martensite start (Ms), Martensite finish (Mf), Austenite start (As) and Austenite
finish (Af) temperatures under biased load were calculated from the intersection of
the linear portion of the transformation region and linear portion of martensite and
austenite region as shown in Fig. 3.7. Recovered strain (εrec) or transformation strain
was measured by subtracting the strain at the end of heating cycle from the strain at
the beginning of heating cycle. Similarly irrecovered strain (εirr) was measured by
subtracting the strain at the end of the heating cycle from the strain at the beginning
55
of the cooling cycle. The measurement scheme of transformation strains is shown in
Fig. 3.7.
The work output (J/cm3) and the recovery ratio exhibited by the alloy can be
calculated by using equations 3.1 and 3.2 respectively.
Work output = (εrec x σ) ------------- 3.1 [139]
Recovery Ratio = (εrec /( εrec + εirr)) ------------- 3.2 [140]
Where σ is the applied stress at which recoverable strain has been achieved.
Fig. 3.7 Measurement scheme of transformation temperatures, recoverable
and irrecoverable strains from typical strain-temperature curve
0 100 200 300 400
Stra
in (
%)
Temperature (°C)
Ɛirr
Ɛrec
As
Af
Mf
Ms
Heating
Cooling
56
3.7 Summary
In this chapter, material processing starting from the raw materials to the final stage
of alloys melting, heat treatment, samples preparation, characterization, mechanical
testing and constant stress thermal testing were presented. Titanium, Nickel,
Palladium and Copper of highest purity were taken and weighed precisely to get the
required alloys in accurate composition. The constituent elements were cleaned in
methanol using the ultrasonic cleaner. The cleaned constituent elements were melted
in vacuum arc melting furnace for six times to improve alloying homogeneity. Four
alloys of required composition; 0Cu, 5Cu, 10Cu and 15Cu alloys each of 20 g button
were homogenized in sealed quartz tubes separately at 950°C for 2 hours.
After homogenization, each button was sliced into 0.4 mm thick strips using wire
EDM. The same strips were then cold rolled and its thickness was reduced to 0.3
mm. The strips of each alloy were then cut into the required sample dimensions. All
the samples of four alloys were solution treated in sealed quartz tubes at 900°C for 1
hour and cold water quenched without crushing the quartz tubes. After solution
treatment, some samples of each alloy were aged at 400°C, 500°C, 600°C and 700°C
for 3 hours in sealed quarts tubes and then water quenched.
Microstructural characterization like grain size and its structure was carried out by
optical microscope in solution treated condition. By using the SEM, the size, shape
and color of second phase precipitates were found whereas chemical composition
was found by EDS in both conditions; solution treated and aged. Phase analysis of all
four alloys was carried out by XRD for both solution treated and aged conditions.
Similarly stress free transformation temperatures in both conditions of solution
treated and aged were found by DSC. Mechanical properties like microhardness and
stress strain curves were obtained for all four alloys using microhardness tester and
tensile testing machine respectively. Shape memory properties of all four alloys were
found in both conditions of solution treated and aged by using the creep and rupture
testing machine.
LVDT Pushing Supporting
57
Chapter – 4
Effect of Copper Addition and Aging
on Microstructure of TiNiPd Alloys
4.1 Introduction
This chapter presents the effect on the microstructural properties due to varying
percentage (0, 5, 10 and 15 atomic%) of copper (Cu) addition with TiNiPd alloys in
solution treated condition. It also deals with the microstructure analysis of TiNiPdCu
alloys with varying Cu percentage, aged at different aging temperatures (400°C,
500°C, 600°C and 700°C for 3 hours) which in turn changes the final microstructure
of the alloys. Microstructural analysis of these alloys was carried out by using optical
microscopy, Scanning Electron Microscopy (SEM), X-Ray Diffractometry (XRD)
and Energy Dispersive Spectroscopy (EDS).
4.2 Microstructure analysis of solution treated TiNiPdCu alloys with
varying Cu percentage
4.2.1 Second phase precipitates
Fig. 4.1 shows the SEM images of the 0Cu, 5Cu, 10Cu and 15Cu alloys in the
solution treated condition (detail of solution treatment is given in Section 3.3.2).
From Fig. 4.1, it is observed that second phase precipitates of black colors with low
density randomly distributed along the grain boundaries are present in all four alloys.
The average size of these precipitates found in the four alloys was measured to be 1.0
– 2.8 μm. These precipitates had almost the same shape which seems to be of circular
or elliptical type.
This observation shows that addition of Cu had no effect on the precipitate size.
However the density of the precipitates decreased as the Cu content in the alloys
increased. For the 0Cu alloy, the volume fraction of the precipitates was estimated to
be ~2.6%, decreased to ~1.5% for 15Cu alloy. To confirm the chemical composition
of the overall and second phase precipitates, the EDS analysis was carried out. The
58
EDS spectrums are shown in Fig. 4.2, whereas the chemical composition is shown in
Table 4.1.
Fig. 4.1 SEM images showing the second phase precipitates formed along the grain
boundaries in solution treated samples of (a) 0Cu, (b) 5Cu, (c) 10Cu and (d) 15Cu
alloys
10 μm
a
10 μm
b
10 μm
c
10 μm
d
59
The compositional data shown in Table 4.1 suggested that second phase precipitates
formed at the grain boundaries are Ti2Ni(Pd) in 0Cu alloy and Ti2Ni(Pd,Cu) in the
remaining three alloys. Such type of precipitates were commonly observed in Ti-rich
NiTi-based alloys [139]. It was believed that the Ti2Ni(Pd) second phase precipitates
were formed during solidification, because these precipitates had lower melting point
than the matrix [141]. Same type of precipitates were also reported to be formed in
ternary TiNiPd alloys having Pd contents more than 10% [139].
Table 4.1 Compositional analysis of the overall alloy and second phase precipitate in
solution treated condition of 0Cu, 5Cu, 10Cu and 15Cu alloys.
Alloy Name Analysis Region Ti (%) Ni (%) Pd (%) Cu (%)
0Cu alloy Overall alloy 49.2 23.5 27.3 0
Ti2Ni 57.5 27.4 15.1 0
5Cu alloy Overall alloy 49.4 19.1 26.8 4.7
Ti2Ni 56.8 22.3 17.3 3.6
10Cu alloy Overall alloy 49.9 14.2 24.7 11.2
Ti2Ni 66.8 16.4 12.2 4.6
15Cu alloy Overall alloy 49.0 9.8 27.4 13.8
Ti2Ni 59.3 15.4 15.2 10.1
0Cu 5Cu
Fig. 4.2 EDS spectrums shown for solution treated samples of 0Cu, 5Cu, 10Cu
and 15Cu alloys
10Cu 15Cu
60
From these observations it was concluded that by addition of Cu, the basic
microstructure of quaternary alloys did not change. The alloying element Cu
remained in the matrix of solid solution and did not generate any other second phase
precipitate, as was reported in case of Scandium (Sc) addition to TiNiPd [142, 143]
in form of Sc2O3. Similarly second phase precipitates of TiB2 is also reported by
Suzuki et al. [68] to be formed when 0.2% Boron (B) was added in Ti50Pd30Ni20 in
place of Ni.
4.2.2 Grain size
Fig. 4.3 and 4.4 show the optical micrographs of the 0Cu, 5Cu, 10Cu and 15Cu
alloys at magnification of 200X and 500X respectively in the solution treated
condition.
Fig. 4.3 Optical micrographs (at 200X) of (a) 0Cu (b) 5Cu (c) 10Cu and (d) 15Cu
alloys solution treated at 900°C for 1 hour
a b
d c
61
a b
d c
Fig. 4.4 Optical micrographs (at 500X) of (a) 0Cu (b) 5Cu (c) 10Cu and (d) 15Cu
alloys solution treated at 900°C for 1 hour
0
10
20
30
40
Gra
in S
ize
(μm
)
Cu (at%)
0 5 10 15
Fig. 4.5 Effect of increasing Cu-content on grain size
62
From these figures it can be observed that all the alloys are consisted of single phase
having typical twinned martensite structure with clearly visible grain boundaries.
Grain size was estimated by using an intercept method, described as follows. Many
straight lines all the same length were drawn through micrographs that show the
grain structure. The grains intersected by each line were counted, the line length was
then divided by an average of the number of grains intersected, taken over all the line
segments. The average grain diameter is found by dividing this result by the linear
magnification of the micrographs. The average grain size of the 0Cu, 5Cu, 10Cu and
15Cu alloys was found to be 13 μm, 16 μm, 18 μm and 20 μm respectively as shown
in Fig. 4.5. It can be concluded from the above observation that addition of Cu in
place of Ni resulted in an increase of grain size. The average grain size, 13 μm of
0Cu alloy increased to 20 μm in 15Cu alloy showing about 35% increase in grain
size.
The increase in grain size due to replacement of Ni by Cu can be attributed to the
decrease in Ni content. When the Ni content decreased in the alloy, the formation of
the second phase precipitate (Ti2Ni) also decreased. Hence, due to decrease in
formation of second phase precipitates, grain growth increased when Cu content was
increased from 0% to 15%. In case of 15Cu alloy, due to low density and
consequently less pinning effect of second phase precipitates, the grain growth was
greater as compared to grain growth in 0Cu alloy. Therefore, due to reduction in
pinning effect, the grain size of the alloys increased as the Cu content increased.
4.3 Effect of aging temperature on microstructure of TiNiPdCu
alloys with varying Cu percentage
To investigate the effect of different aging temperature on the four alloys, the
samples were analyzed by SEM with backscattered electron detector. Fig. 4.6 (a–d)
represents the back-scattered SEM images of 0Cu alloy aged under the conditions
mentioned in Section 3.3.2. These images show that only the second phase
precipitates, which were formed during solidification process, were present.
However, there was no evidence of precipitation on the grain boundaries as well as in
the grain interior due to the aging effect.
63
Fig. 4.7 (a–d) shows the back-scattered SEM images of 5Cu alloy aged under the
same conditions. It can be observed from these images that the second phase
precipitates were available along the grain boundaries. From Fig. 4.7 a, it can be seen
that due to aging at 400°C, the precipitation process was not started. Increasing the
aging temperature to 500°C, very fine precipitates of bright color were seen along the
grain boundaries. By further increasing the aging temperature to 600°C, the density
of the precipitates also increased. The size of the precipitates was further increased
when the aging temperature was increased to 700°C. These precipitates were found
along the grain boundaries as well as within the grain interior.
Fig. 4.6 Back-scattered SEM images presenting the microstructure and grain
boundaries in 0Cu alloys after aging for 3 hours at temperature of (a) 400°C, (b)
500°C, (c) 600°C and (d) 700°C
10 μm
a
10 μm
b
10 μm
c
10 μm
d
64
The backscattered SEM images shown in Fig. 4.8 (a–d) demonstrate the precipitation
behavior of 10Cu alloys aged under the same condition. Fig. 4.8a presents the
microstructure of the sample aged at 400°C. It can be seen that the precipitation
process has been started and very fine precipitates were formed on the grain
boundaries. It can also be noted that the precipitation was not observed in the grain
interior, confirming that the grain boundaries are working as nucleation sites. Fig.
4.8b shows the microstructure of the alloy aged at 500°C, indicating the population
of two types of precipitates randomly distributed along both sides of grain
boundaries. The size of these precipitates was larger as compared to the precipitates
formed at 400°C. As the aging temperature was increased to 600°C, the density of
10 μm
a
10 μm
b
10 μm
c
10 μm
d
Fig. 4.7 Back-scattered SEM images presenting the microstructure and grain
boundaries in 5Cu alloys after aging for 3 hours at temperature of (a) 400°C, (b)
500°C, (c) 600°C and (d) 700°C
65
the precipitates had further increased as shown in Fig. 4.8c. At 600°C, both types of
precipitates were observed to be formed heterogeneously on the grain boundaries as
well as in the grain interior. The shape of the precipitates observed in Fig. 4.8b and c
has thread-like structures with variable length.
Fig. 4.8d represents the microstructure of the alloy aged at 700°C where the
formation of low density and a relatively large size precipitate can be seen. The
precipitation behavior of 700°C-aged alloy was observed to be completely different
from the behavior observed in case of aging at 500°C and 600°C. Formation of nano-
scaled precipitates also reported by Khan et al. [140] to be formed during annealing
Fig. 4.8 Back-scattered SEM images presenting the microstructure and grain
boundaries in 10Cu alloys after aging for 3 hours at temperature of (a) 400°C, (b)
500°C, (c) 600°C and (d) 700°C
10 μm
a b
10 μm
d
10 μm
c
10 μm
66
at 550°C, 600°C and 650°C for 1 hour after cold deformation by 40% in the same
alloy. However the precipitates found in annealing process were of nano-scale sizes
(100 – 300nm) with different morphologies (elliptical and rod-like shapes) as
compared to precipitates formed in the present study. Moreover, the precipitation
behavior during annealing was observed to be completely different from that of
aging process. Cold deformation followed by annealing resulted in the formation of
heterogeneous nucleation sites due to deformation-induced defects [6], whereas in
case of age hardening, the high energy grain boundaries were responsible as
nucleation sites for the initiation of precipitation process.
a
10 μm 10 μm
b
c
10 μm
d
10 μm
Fig. 4.9 Back-scattered SEM images presenting the microstructure and grain
boundaries in 15Cu alloys after aging for 3 hours at temperature of (a) 400°C, (b)
500°C, (c) 600°C and (d) 700°C
67
In Fig. 4.9 (a–d), backscattered SEM images of aged 15Cu alloys are shown. Fig.
5.9a shows the microstructure of the sample aged at 400°C. It can be seen that very
fine precipitates were formed on both sides of grain boundaries. However the size of
the precipitates was greater as compared to 400°C-aged sample of 10Cu alloy as
shown in Fig. 4.8a. Fig. 4.9b shows the microstructure of the alloy aged at 500°C,
representing two types of precipitates, however the size of these precipitates was
larger as compared to the precipitates formed at 400°C for the same alloy. As the
aging temperature was increased to 600°C, the density of the precipitates had further
increased as shown in Fig. 4.9c. At 600°C, both types of precipitates were observed
to be formed heterogeneously on the grain boundaries as well as in the grain interior.
The shape of the precipitates observed in Fig. 4.9b and c has thread-like structures
with variable length. Fig. 4.9d represents the microstructure of the alloy aged at
700°C where the formation of low density and a relatively large size precipitates can
be seen. The precipitation behavior of 700°C-aged alloy was observed to be
completely different from the behavior observed in case of aging at 500°C and
600°C.
Table 4.2 Compositional analysis of black and white precipitates formed in 10Cu and
15Cu alloys after aging for 3 hours at various aging temperatures
It has been reported [144] that in backscattered SEM images, a precipitate with
brighter contrast represents the existence of heavier elements i.e. elements with
higher atomic number, and a precipitate with darker contrast shows the existence of
lighter elements i.e. elements with lower atomic number. It can be seen from the
periodic table that Ti has lower atomic number than Ni, Pd and Cu and thus it is
lighter as compared to other constituent elements. Therefore the presence of darker
and brighter contrast represented the existence of Ti-rich and Ti-lean precipitates,
respectively [144]. To confirm the exact composition of the two newly born
precipitates as a result of aging at various temperatures, EDS analysis was carried out
Alloy Name Analysis Region Ti (%) Ni (%) Pd (%) Cu (%)
10Cu alloy Black precipitates 60.6 9.3 23.8 6.3
White precipitates 41.1 11.4 26.7 20.8
15Cu alloy Black precipitates 57.3 10.6 24.5 7.6
White precipitates 42.5 10.7 27.3 19.5
68
and their data was summarized in Table 4.2. It can be confirmed from back-scattered
SEM images shown in Fig. 4.8, 4.9 and from EDS analysis that the darker and
brighter contrast particles are Ti2Pd (Ti-rich) and TiPdCu (Ti-lean) precipitates. It
can also be observed that in both types of precipitates, the Ni concentration is lower
than the target composition. So it is expected that by formation of these precipitates,
the Ni content was increased and Ti content was decreased in the matrix.
Comparing the precipitation behavior of 0Cu, 5Cu, 10Cu and 15Cu alloys, it can be
easily observed that stoichiometric (Ti : Ni+Pd = 50% : 50%) TiNiPd alloys were not
responding to any precipitation process within the aging temperature ranges from
400°C to 700°C. The precipitation behavior in 5Cu alloys is not much different from
the 0Cu alloy. However, very fine precipitates of brighter contrast were observed at
aging temperatures of 500°C, 600°C and 700°C as shown in Fig. 4.7 (b – d). By
adding 10% and 15% Cu in place of Ni, the precipitation behavior of stoichiometric
TiNiPd alloys remarkably changed. In both alloys, precipitation started along the
grain boundaries at lower aging temperature. As the temperature increased, the
density and size of precipitates also increased till aging temperature of 600°C. At
700°C aging temperature, the size of the precipitates further increased, however their
density decreased. It was also observed that the size of the precipitates formed in
15Cu alloys was bigger as compared to the precipitates of 10Cu alloys at
corresponding temperatures.
4.4 Phase analysis of solution treated TiNiPdCu alloys with varying
Copper percentage
Phase analysis was carried out using X-Ray Diffractometer with Cu Kα (λ = 1.5406
°A) radiation. The X-Ray Diffraction (XRD) profiles were measured in the 2Ɵ range
of 20 to 80 degree at room temperature in martensite phase. It is observed during
XRD analysis that the maximum required peaks are detected between the diffraction
angle of 30 and 60 degree, therefore XRD profiles are shown in the same range.
Fig. 4.10 shows the XRD profiles of solution treated TiNiPdCu alloys with varying
Cu percentage. All peaks of XRD profiles indicated only the B19 orthorhombic
martensite phase. From this observation, it can be confirmed that the alloys were
69
consisted of single B19 martensite phase at room temperature. Presence of second
phase particles of Ti2Ni(Pd) in 0Cu alloy and Ti2Ni(Pd, Cu) in the remaining three
alloys were not detected as these precipitates were seen in the SEM images shown in
Fig. 4.1.
The reason of unavailability of corresponding peaks of Ti2Ni(Pd) in 0Cu alloy and
Ti2Ni(Pd, Cu) in the remaining three alloys could be their low volume fraction i.e.
~2.6% to ~1.5% as discussed in Section 4.2.1. By adding Cu in TiNiPd alloys,
consistent shift of B19 martensite peaks was observed towards low angle of 2Ɵ. It
can be seen in Fig. 4.10 that the strongest B19 (111) martensite peak emerged at ~
43 degree given for 0Cu alloy, shifted towards the lower angle for 5Cu, 10Cu and
15Cu alloys. The decrease in diffraction angle by increasing the Cu addition may be
20 30 40 50 60
Inte
nsi
ty (
CP
S)
Angle (2Ɵ)
0Cu alloy
5Cu alloy
15Cu alloy
B1
9 (
11
1)
B1
9 (
00
2)
B1
9 (
02
0)
10Cu alloy
Fig. 4.10 XRD profiles at room temperature for the samples solution treated
0Cu, 5Cu, 10Cu and 15Cu alloys
70
expected by the increase in the lattice constants. The increase in lattice constants is
attributed to the addition of relatively large atomic radius of Cu (0.128 nm) as
compared to the atomic radius of Ni (0.125 nm).
4.5 Phase analysis of aged TiNiPdCu alloys with varying Copper
percentage
4.5.1 Phase analysis of aged 0Cu alloy
Fig. 4.11 shows the XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 0Cu alloys. It is seen that
four peaks (002), (020), (111) and (022) represent the B19 martensite phase while
one peak (011) shows the B2 austenite phase. Hence it is confirmed that at room
20 30 40 50 60
Inte
nsi
ty (
CP
S)
Angle (2Ɵ)
Solution Treated
Aged at 400°C
Aged at 500°C
Aged at 700°C
Aged at 600°C
B1
9 (
11
1)
B1
9 (
00
2)
B1
9 (
02
0)
B2
(0
11
)
B1
9 (
02
2)
Fig. 4.11 XRD profiles at room temperature for the samples solution treated and
aged at 400°C, 500°C, 600°C and 700°C of 0Cu alloys
71
temperature, the predominant phase is the B19 martensite phase where the retained
austenite B2 phase in low volume fraction is also present in the solution treated
sample. When the alloy was aged at 400°C, all the four martensite peaks are
available at the same intensities as were present in the solution treated sample.
However, the intensity of the peak corresponding to the B2 austenite phase was
increased. By further increasing the aging temperature to 500°C, 600°C and 700°C,
the XRD patterns show the same phases and no change was observed. These
observations confirm that aging at 400°C, 500°C, 600°C and 700°C for 3 hours,
stoichiometric TiNiPd alloys do not respond to any phase change.
4.5.2 Phase analysis of aged 5Cu alloy
Fig. 4.12 shows the XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 5Cu alloys. Here it can be
observed that all the four peaks (002), (020), (111) and (022) represent the B19
martensite phase and no phase other than martensite is detectable. Therefore, it is
confirmed that at room temperature, the material is transformed completely to the
B19 martensite phase. By aging the alloy at 400°C, all the four martensite peaks are
available at the same intensities as were present before aging. By further increasing
the aging temperature to 500°C, 600°C and 700°C, the XRD patterns show the same
phases and no change was observed. From these observations it can be confirmed
that aging at 400°C, 500°C, 600°C and 700°C for 3 hours, TiNiPdCu alloys with 5%
Cu do not respond to any phase change. It was observed from Fig. 4.7 (b – d) that by
aging the 5Cu alloys at 500°C, 600°C and 700°C, very fine precipitates are formed,
however the XRD profiles do not show the existence of those precipitates. The
reason for this observation is the low volume fraction of precipitates which cannot be
detected by XRD analysis.
72
4.5.3 Phase analysis of aged 10Cu alloy
Fig. 4.13 depicts the XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 10Cu alloys. The XRD
pattern for solution treated sample shows that all the five peaks (101), (002), (020),
(111) and (022) represent the B19 martensite phase. By aging the alloy at 400°C, the
XRD profile was slightly changed. At this aging temperature, the height of the
martensite peaks, except (111), was slightly decreased.
20 30 40 50 60
Inte
nsi
ty (
CP
S)
Angle (2Ɵ)
Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
B1
9 (
02
2)
B1
9 (
00
2)
B1
9 (
02
0)
B1
9 (
11
1)
Fig. 4.12 XRD profiles at room temperature for the samples solution treated and
aged at 400°C, 500°C, 600°C and 700°C of 5Cu alloys
73
On the other hand, emerging of new peaks, one at the left side (39.77o) and another at
right side (43.67o) of strongest B19 (111) peak were observed. Reduction in volume
fraction of martensite phase and emerging of new phase had confirmed that
precipitation has been started at aging temperature of 400°C. As the aging
temperature was increased to 500°C and 600°C, the XRD profiles were changed
significantly. At these aging temperatures, (002) and (022) reflexes of B19
martensite phase were completely disappeared due to remarkable decrease in volume
fraction of martensite phase and formation of two types of second phase precipitates
in the matrix were detected. It can be observed from the XRD profiles, that peaks
corresponding to new formed second phase precipitates are showing maximum
height at 600°C which is due to the presence of dense population of second phase
precipitates with higher volume fraction. From compositional analysis given in Table
20 30 40 50 60
Inte
nsi
ty (
CP
S)
Angle (2Ɵ)
Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C B
19
(0
02
)
B1
9 (
02
2)
B1
9 (
10
1)
B1
9 (
11
1)
B1
9 (
02
0)
Ti 2
Pd
TiP
dC
u
Fig. 4.13 XRD profiles at room temperature for the samples solution treated and
aged at 400°C, 500°C, 600°C and 700°C of 10Cu alloys
74
4.2, XRD profile and reference [145], it was confirmed that the peaks located on the
left and right side of (111) reflex are Ti2Pd and TiPdCu, respectively. The
observations resulted from the XRD analysis regarding the precipitation behavior
was found to be completely in accordance with the microstructures given in Fig. 4.8
(a – c). When the aging temperature was further increased to 700°C, the XRD profile
was again significantly changed and the peaks corresponding to Ti2Pd and TiPdCu
were observed to be disappeared. Although it was observed from Fig. 4.8d that both
types of precipitates with low density of relatively coarser size were present in the
alloy aged at 700°C, however it could not be detected in the XRD profile, probably
due to low volume fraction. The XRD profiles for the alloys aged at 700°C were
found to be the same as was observed for solution treated sample, where all peaks
were representing the B19 orthorhombic martensite phase.
4.5.4 Phase analysis of aged 15Cu alloy
Fig. 4.14 presents the XRD profiles at room temperature for the samples solution
treated and aged at 400°C, 500°C, 600°C and 700°C of 15Cu alloys. The XRD
pattern for solution treated sample shows that all the four peaks (002), (020), (111)
and (022) represent the B19 martensite phase. These observations suggest that the
material is completely in martensite phase having the B19 orthorhombic structure at
room temperature.
By aging the alloy at 400°C, the XRD profile was significantly changed. At this
aging temperature, one martensite peak with reflex (022) was completely
disappeared. On the other hand, two new peaks, one at the left side (39.77o) and
another at right side (43.67o) of strongest B19 (111) peak were immerged. The
absence of one martensite peak and emergence of two new peaks had suggested that
volume fraction of martensite phase was decreased and precipitation was started at
aging temperature of 400°C. When the aging temperature was increased to 500°C,
remarkable change in the XRD profile was observed.
75
Fig. 4.14 XRD profiles at room temperature for the samples solution treated and
aged at 400°C, 500°C, 600°C and 700°C of 15Cu alloys
20 30 40 50 60
Inte
nsi
ty (
CP
S)
Angle (2Ɵ)
Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
B1
9 (
11
1)
B1
9 (
02
2)
B1
9 (
00
2)
Ti 2
Pd
TiP
dC
Uu
B1
9 (
02
0)
B1
9 (
10
1)
All the peaks corresponding to martensite phase were totally absent, however the
intensities of the new emerged peaks were observed to be increased. The XRD
pattern obtained at aging temperature of 600°C shows similar profile as was obtained
for 500°C-aged sample with slight increase in intensities. It can be observed from the
XRD profiles, that peaks corresponding to new formed second phase precipitates are
showing maximum height at 600°C which is due to the presence of dense population
of second phase precipitates with higher volume fraction. From compositional
analysis given in Table 4.2 and XRD profile, it can be confirmed that the peaks
located on the left and right side of (111) reflex are Ti2Pd and TiPdCu respectively.
The observation obtained from the XRD analysis regarding the precipitation
behavior was found to be completely in accordance with the microstructures given in
Fig. 4.9 (a – c). When the aging temperature was further increased to 700°C, the
76
XRD profile was again significantly changed. Here the four peaks corresponding to
martensite phase were again observed, however at the same time the two peaks
corresponding to Ti2Pd and TiPdCu were also present. The XRD analysis confirms
the presence of both types of precipitates as was observed from Fig. 4.9d. Here it is
also confirmed that the volume fraction of both precipitates is significantly high and
can be detected through XRD patterns.
By comparing the XRD analysis for 10Cu and 15Cu alloys, it can be concluded that
the size and density of precipitates produced in 15Cu alloys were greater as
compared to those of 10Cu alloys at the corresponding aging temperatures. This
conclusion fully supports by the backscattered SEM images shown in Fig. 4.8 (a – d)
for 10Cu alloy and Fig. 4.9 (a – d) for 15Cu alloy. Moreover the XRD profile
obtained for 700°C-aged sample of 15Cu alloy is significantly different from the
700°C-aged sample of 10Cu alloy. In 10Cu alloy, the XRD pattern shows that all the
peaks available are B19 martensite whereas the XRD pattern of obtained for 15Cu
alloy shows the martensite peaks as well as peaks for Ti2Pd and TiPdCu.
4.6 Summary
In this chapter, the effect on the microstructural properties due to varying percentage
of Cu addition with TiNiPd alloys in solution treated condition was presented.
Moreover, the effect on the microstructure due to aging at different aging
temperatures of TiNiPdCu with varying percentage of Cu was also discussed. It was
observed that the density of second phase precipitates decreased and grain size
increased when Cu was increased in TiNiPd alloys. The grain size of 15Cu alloy was
observed to be increased by 35% as compared to 0Cu alloy. By aging the 0Cu alloys,
no change in its microstructure was noticed. In 5Cu alloys, very fine precipitates
along the grain boundaries were formed at higher aging temperatures. Both 10Cu and
15Cu alloys responded significantly to precipitation process at all aging
temperatures. At low aging temperatures, precipitates of greater size were formed
along the grain boundaries in both alloys. When the aging temperature increased, the
size of the precipitates also increased. At the same time, formation of precipitates
increased not only on the grain boundaries but also spread into the grain interior.
77
However, the size of the precipitates formed in 10Cu alloys was relatively smaller
than those of 15Cu alloys.
During phase analysis it was observed that in solution treated condition, only B19
martensite phase was present in all four alloys. However the peaks representing the
martensite phase shifted consistently towards lower 2Ɵ angle by increasing the Cu in
TiNiPd alloys. By aging 0Cu and 5Cu alloys at different aging temperatures, no
phase change was observed and the phase was remained as martensite phase.
Remarkable change was noticed when 10Cu and 15Cu alloys were aged at different
aging temperatures. At low and intermediate aging temperatures, two types of second
phase precipitates (Ti2Pd, TiPdCu) were also detected, resultantly the volume of
martensite phase was reduced. However, at higher aging temperature, the formation
of those precipitates was again disappeared and only martensite phase was present.
78
Chapter – 5
Effect of Copper Addition and Aging
on Transformation Temperatures of
TiNiPd Alloys
5.1 Introduction
This chapter presents the effect of varying percentage of copper (0Cu, 5Cu, 10Cu
and 15Cu) addition on phase transformation temperatures in TiNiPd alloys.
Moreover, it also includes the changes in transformation temperatures of TiNiPdCu
alloys occurred due to aging at different aging temperatures. These transformation
temperatures were measured by using Differential Scanning Calorimetry (DSC).
5.2 Effect of Cu addition on phase transformation temperatures
Fig. 5.1 shows the DSC heating and cooling cycles of 0Cu, 5Cu, 10Cu and 15Cu
alloys solution treated at 900°C for 1 hour. The phase transformation temperatures
measured from DSC cycles of Fig. 5.1 were re-plotted in Fig.5.2 to compare the
change in transformation temperatures for replacing of Ni with 5%, 10% and 15%
Cu with respect to the baseline 0Cu alloy. All alloys exhibited very similar behavior
with single-stage martensite transformation and demonstrated proper developed
peaks with well calculated transformation temperatures. However the transformation
heat (ΔHc) released (area under cooling curve) during forward transformation cycle
and transformation heat (ΔHh) absorbed (area under heating curve) during reverse
transformation cycle increased as the Cu contents increased in the alloy.
The martensite start temperature; Ms of 0Cu alloy increased by 12.5°C from 142.5°C
to 155°C, while the austenite finish temperature; Af increased by 11°C from 167°C to
178°C upon addition of 5% Cu in place of Ni. Moreover, it can also be noted that,
the thermal hysteresis decreased by 1.5°C from 24.5°C to 23°C. From the above
experimental results it has been observed that substitution of Ni by 5% Cu increased
the transformation temperatures significantly while thermal hysteresis decreased
slightly.
79
Similarly the transformation temperatures of 10Cu alloy further increased and it
became 27.5°C higher than the baseline 0Cu alloy. The Ms temperature of 0Cu alloy
increased by 27.5°C from 142.5°C to 170°C, while the Af temperature increased by
25°C from 167°C to 192°C. Moreover, it was also observed, that thermal hysteresis
decreased by 2.5°C from 24.5°C to 22°C. From these experimental results it was
concluded that substitution of Ni by 10% Cu, increased the transformation
temperatures significantly, while thermal hysteresis decreased slightly.
By replacing Ni with 15% Cu in 0Cu alloy, the transformation temperatures of 15Cu
alloy further increased. The Ms temperature of 0Cu alloy increased by 60.5°C from
142.5°C to 203°C, while the Af temperature increased by 52°C from 167°C to
219°C. It can also be noted that thermal hysteresis decreased by 8.5°C from 24.5°C
to 16°C. From the above experimental results it has been concluded that substitution
of Ni by 15% Cu, the transformation temperatures increased and thermal hysteresis
decreased significantly.
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exo
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Temperature (°C)
0.5
0Cu alloy
10Cu alloy
5Cu alloy
15Cu alloy
Cooling
Heating
Fig. 5.1 DSC heating and cooling cycles of solution treated 0Cu, 5Cu, 10Cu and
15Cu alloys
80
The change in transformation temperatures and transformation heats are actually due
to the change in Ni/Pd content, because Cu content does not affect the transformation
temperatures according to Mercier et al. [146]. Suburi, T. [147] reported that
according to the TiNiTiPd pseudo-binary phase diagram, the increase in Pd-content
and decrease in Ni-content, increases the transformation temperatures. Thus the
decrease in Ni/Pd ratio (as Ni content decreased and Pd content remained constant)
resulted an increase in transformation temperatures and transformation heats
absorbed and released during forward and reverse martensitic transformation,
respectively. Moreover, according to Clausius-Clapeyron equation, the increase in
transformation temperature causes an increases in the transformation heats [3].
5.3 Effect of aging on phase transformation temperatures
To investigate the effect of aging on phase transformation temperatures, the
transformation temperatures of four alloys aged at different aging temperatures were
measured. The detail discussion is given as under.
100
135
170
205
240
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Cu (at%)
Ms
Mf
As
Af
0 5 10 15
Fig. 5.2 Effect of Cu addition on transformation temperatures
81
5.3.1 Effect of aging on phase transformation temperatures of 0Cu alloy
Fig. 5.3 represents the DSC heating and cooling cycles of the samples solution
treated, aged at 400°C, 500°C, 600°C and 700°C of 0Cu alloy. Fig. 5.4 shows the
variation of transformation temperatures with respect to the increasing aging
temperatures. The DSC heating and cooling cycles for the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C were showing properly developed peaks.
However the sharpness of the peaks increased as the aging temperature increased
from 400°C to 700°C. Thus the transformation heats released and absorbed during
forward and reverse transformation cycles increased as the aging temperatures
increased. The Ms of the 400°C-aged alloy increased by 8°C from 142°C to 150°C,
whereas Af increased by 9°C from 168°C to 177°C, when the aging temperature
increased from 400°C to 700°C. Moreover, it was also observed that thermal
hysteresis of the 0Cu alloy slightly increased when the aging temperatures increased.
Thermal hysteresis of the 400°C-aged alloy increased by 1°C from 26°C to 27°C,
50 100 150 200 250
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Temperature (°C)
Cooling
Heating Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
0.5
Fig. 5.3 DSC heating and cooling cycles of the samples solution treated, aged at
400°C, 500°C, 600°C and 700°C of 0Cu alloy
82
when the aging temperature increased to 700°C. The increase in transformation
temperatures of the 0Cu alloy due to increase in aging temperatures can only be
attributed to the grain growth [148], as there was no precipitation detected in aging of
0Cu alloy as shown in Fig. 4.6.
5.3.2 Effect of aging on phase transformation temperatures of 5Cu alloy
Fig. 5.5 shows the DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 5Cu alloy. The transformation
temperatures evaluated from DSC cycles of Fig. 5.5 were re-plotted in Fig. 5.6 which
shows the variation of transformation temperatures with respect to the increasing
aging temperatures. It can be observed from Fig. 5.6 that the transformation
temperatures increased when the aging temperatures increased from 400°C to 600°C.
However, further increase in aging temperature i.e. at 700°C, the transformation
temperatures decreased. The Ms of the 400°C-aged alloy increased by 11°C from
163°C to 174°C, when the aging temperature increased to 600°C. Similarly Af
increased by 8°C from 186°C to 194°C. At aging temperature of 700°C, the Ms
decreased by 7°C from 174°C to 167°C while the Af decreased by 2°C from 194°C
100
135
170
205
240
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Aging Temperature (°C)
Ms
Mf
As
Af
ST 400 500 600 700
Fig. 5.4 Effect of aging temperatures on transformation temperatures of 0Cu alloy
83
to 192°C. Moreover, it was also observed that thermal hysteresis of the 5Cu alloy
slightly increased when the aging temperatures increased. Thermal hysteresis of the
400°C-aged alloy increased by 2°C from 23°C to 25°C, when the aging temperature
increased to 700°C.
The increase in transformation temperatures of the 5Cu alloy due to increase in
aging temperatures till 600°C is the result of grain growth [148], because there was
no significant precipitation observed in aging of 5Cu alloy as shown in Fig. 4.7 (a –
c). However, by further increasing the aging temperature to 700°C, the precipitation
process started and low volume fraction of brighter precipitates along the grain
boundaries and grain interiors detected as shown in Fig. 4.7 (d). It has been discussed
in chapter 4 that the brighter precipitates were Ni-lean in nature. Due to the
formation of these precipitates, the Ni-content in the matrix increased. By increasing
the Ni-content, the transformation temperatures decreased [149].
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/g)
exo
up
Temperature (°C)
Cooling
Heating Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
0.5
Fig. 5.5 DSC heating and cooling cycles of the samples solution treated, aged at
400°C, 500°C, 600°C and 700°C of 5Cu alloy
84
5.3.3 Effect of aging on phase transformation temperatures of 10Cu alloy
Fig. 5.7 represents the DSC heating and cooling cycles of the samples solution
treated, aged at 400°C, 500°C, 600°C and 700°C of 10Cu alloy. Fig. 5.8 shows the
variation of transformation temperatures with respect to the increasing aging
temperatures. The DSC heating and cooling cycles in Fig. 5.7 for the samples
solution treated and aged at 400°C, were showing proper developed peaks with well
calculated transformation temperatures. However, the transformation peak for
solution treated sample was sharper as compared to the transformation peak of
400°C-aged sample. When the aging temperature increased to 500°C and 600°C, the
transformation peaks broadened and their height become lower and thus difficult to
calculate martensite finish temperature; Mf and austenite start temperature; As.
Therefore for comparison purpose, only Ms and Af were calculated and drawn in Fig.
5.8.
100
135
170
205
240
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Aging Temperature (°C)
Ms
Mf
As
Af
ST 400 500 600 700
Fig. 5.6 Effect of aging temperatures on transformation temperatures of 5Cu alloy
85
By increasing the aging temperatures from 400°C to 600°C, the transformation
temperatures decreased remarkably. The Ms of the 400°C-aged sample dropped by
16°C from 170°C to 154°C when the aging temperature increased to 500°C.
Similarly the Af of the 400°C-aged sample dropped by 16°C from 192°C to 176°C
when the aging temperature increased to 500°C. By further increasing the aging
temperature to 600°C, the Ms decreased by 16°C from 154°C to 138°C and the Af
decreased by 10°C from 176°C to 166°C. At the same time thermal hysteresis
increased by 6°C from 22°C to 28°C. When the aging temperature further increased
to 700°C, again the transformation peaks of the DSC cycles resulted in well-
developed profile with proper height. Moreover, by increasing the aging temperature,
the transformation temperatures also increased remarkably as shown in Fig. 5.8. The
Ms of the 600°C-aged sample increased by 45°C from 138°C to 183°C when the
aging temperature increased to 700°C. Similarly the Af increased by 36.5°C from
166°C to 202.5°C.
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/g)
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Temperature (°C)
Cooling
Heating
Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
0.5
Fig. 5.7 DSC heating and cooling cycles of the samples solution treated, aged at
400°C, 500°C, 600°C and 700°C of 10Cu alloy
86
The decrease in transformation temperatures of the 10Cu alloy due to increase in
aging temperatures to 500°C and 600°C can be attributed to the formation of Ni-lean
precipitates. In Fig. 4.8, it was shown that, at aging temperatures of 500°C and
600°C, two types of precipitates (Ti2Pd and TiPdCu) were formed in the alloy. By
the compositional analysis of precipitates, it was confirmed that both types of
precipitates were Ni-lean. Due to formation of these precipitates, the Ni
concentration in the matrix increased and thus stoichiometric TiNiPdCu alloy
transformed into off- stoichiometric composition; Ti : (NiPdCu) ≠ 50 : 50. According
to Fuentes et al. [150], in the Ni-rich NiTi shape memory alloy, a change of 0.1% Ni,
varies the transformation temperature by 10°C. The remarkable increase in
transformation temperatures at aging temperature of 700°C can be attributed to the
decrease in precipitation process and grain growth [148]. Due to lower density of Ni-
lean precipitates, the Ti concentration in the matrix was again increased and resulted
in higher transformation temperatures.
100
135
170
205
240
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Aging Temperature (°C)
Ms
Mf
As
Af
ST 400 500 600 700
Fig. 5.8 Effect of aging temperatures on transformation temperatures of 10Cu alloy
87
5.3.4 Effect of aging on phase transformation temperatures of 15Cu alloy
Fig. 5.9 shows the DSC heating and cooling cycles of the samples solution treated,
aged at 400°C, 500°C, 600°C and 700°C of 15Cu alloy. The transformation
temperatures calculated from DSC cycles of Fig. 5.9 was plotted in Fig. 5.10 which
shows the change in transformation temperatures due to change in aging
temperatures.
The DSC heating and cooling cycles in Fig. 5.9 for the samples solution treated and
aged at 400°C, were showing proper developed peaks with well calculated
transformation temperatures. However the transformation peak for solution treated
sample is sharper as compared to the transformation peak of 400°C-aged sample.
When the aging temperature increased to 500°C and 600°C, the transformation peaks
broadened and their height become lower and thus difficult to calculate Mf and As.
By increasing the aging temperatures from 400°C to 500°C and 600°C, the
transformation temperatures decreased significantly. The Ms of the 400°C-aged
sample decreased by 16°C from 184°C to 168°C when the aging temperature
increased to 500°C. Similarly the Af of the 400°C-aged sample dropped by 15°C
50 100 150 200 250
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Temperature (°C)
Cooling
Heating Solution Treated
Aged at 400°C
Aged at 700°C
Aged at 600°C
Aged at 500°C
0.5
Fig. 5.9 DSC heating and cooling cycles of the samples solution treated, aged at
400°C, 500°C, 600°C and 700°C of 15Cu alloy
88
from 206°C to 191°C. By further increasing the aging temperature to 600°C, the Ms
decreased by 3°C from 168°C to 165°C and the Af decreased by 2°C from 191°C to
189°C. At the same time thermal hysteresis increased by 2°C from 22°C to 24°C.
When the aging temperature further increased to 700°C, again the transformation
peaks of the DSC cycles resulted in well-developed profile. Moreover, by increasing
the aging temperature, the transformation temperatures also increased significantly as
shown in Fig. 5.10. The Ms of the 600°C-aged sample increased by 18°C from 165°C
to 183°C when the aging temperature increased to 700°C, whereas the Af increased
by 15°C from 189°C to 204°C.
The decrease in transformation temperatures of the 15Cu alloy due to increase in
aging temperatures to 500°C and 600°C can be attributed to the formation of Ni-lean
precipitates, as discussed in section 5.3. The significant increase in transformation
temperatures at aging temperature of 700°C can be attributed to the decrease in
precipitation process and grain growth [148].
100
135
170
205
240
Tran
sfo
rmat
ion
Te
mp
erat
ure
(°C
)
Aging Temperature (°C)
Ms
Mf
As
Af
ST 400 500 600 700
Fig. 5.10 Effect of aging temperatures on transformation temperatures of 15Cu alloy
89
5.4 Effect of thermal cycling on phase transformation temperatures
To investigate the effect of thermal cycling on phase transformation temperatures, all
the four; 0Cu, 5Cu, 10Cu and 15Cu alloys were thermally cycled for complete five
times. The detailed discussion is given as under.
5.4.1 Effect of thermal cycling on phase transformation temperatures of 0Cu
alloy
Fig. 5.11 represents the DSC heating and cooling curves for the evolution of
transformation temperatures during five thermal cycles for solution treated sample of
0Cu alloy. The calculated transformation temperatures were plotted in Fig. 5.12 to
compare the variation in transformation temperatures due to increasing thermal
cycles. Here all the DSC curves showed sharp peaks with well estimated phase
transformation temperatures. It can be observed from Fig. 5.12 that the
transformation temperatures dropped quickly in the second cycle and then decreased
slightly in the next cycles and continue to be decreased till fifth cycle. The Ms;
142.5°C in the first cycle dropped quickly to 138.5°C in the second cycle, resulting a
net decrease of 4°C. Similarly the Af; 167°C in the first cycle dropped quickly to
159°C in the second cycle, resulting a net decrease of 8°C. Thermal hysteresis in the
second cycle decreased by 4°C from 24.5°C to 20.5°C. The decrease in
transformation temperatures and thermal hysteresis in third thermal cycle reduced as
compared to the second thermal cycle.
90
The Ms decreased by 2°C from 138.5°C to 136.5°C and Af decreased by 3°C from
159°C to 156°C in the third cycle. Similarly thermal hysteresis in the third cycle
decreased by only 1°C from 20.5°C to 19.5°C. In the fourth thermal cycle, the
decrease in transformation temperatures and thermal hysteresis further lowered as
compared to the third thermal cycle. The Ms decreased by 1°C from 136.5°C to
135.5°C and Af decreased by 2°C from 156°C to 154°C in the fourth cycle.
Similarly thermal hysteresis in the fourth cycle decreased by 1°C from 19.5°C to
18.5°C. The decrease in transformation temperatures and thermal hysteresis in fifth
thermal cycle further lowered as compared to the fourth thermal cycle. The Ms
decreased by 0.5°C from 135.5°C to 135°C and Af decreased by 1°C from 154°C to
153°C in the fifth cycle. Similarly thermal hysteresis in the fifth cycle decreased by
0.5°C from 18.5°C to 18°C. The net decrease in Ms and Af was observed to be 7.5°C
and 14°C respectively during five thermal cycles. Similarly the net decrease in
thermal hysteresis was observed to be 6.5°C.
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Temperature (oC)
5th Cycle
1st Cycle
5th Cycle
Cooling
Heating
0.3
Fig. 5.11 DSC curves representing the transformation temperatures during five
thermal cycles of 0Cu alloy
91
It was noticed that, by increasing the number of thermal cycles, the transformation
temperatures of solution treated sample of 0Cu alloy decreased quickly in the second
cycle and then slow down till last thermal cycle. The faster decrease in
transformation temperatures in solution treated sample is due to the generation of
dislocations and other defects during repeated motion of the austenite-martensite
interface. This decrease in transformation temperature is unavoidable for quenched
and annealed TiNi-based alloys, however, this can be prevented by aging and
thermo-mechanical treatment [147]. It was also observed that the first heating cycle
of solution treated sample resulted in higher transformation temperatures and then
decreased faster in the second cycle. The largest drop in austenite transformation
temperatures after first heating cycle is attributed to the increase in the formation of
dislocations [151]. The dislocation densities have been reported to be increased
remarkably during the first thermal cycle and then its rate of formation decreases as
the number of cycles increases [152]. Therefore due to formation of increased
dislocation densities in the first cycle, the austenite transformation temperatures
dropped significantly in the second cycle. This effect takes place in the alloys having
low yield strength as observed in the solution treated sample. The decrease in
thermal hysteresis after the first thermal cycle is, due the stabilization effect which
decreased the As quickly and Ms slowly and resultantly their difference (thermal
hysteresis) become small.
110
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170
200
230
Tran
sfo
rmat
ion
Tem
per
atu
res
(oC
)
Cycle Number
Ms
Mf
As
Af
1 2 3 4 5
Fig. 5.12 Effect of thermal cycles on transformation temperatures of 0Cu alloy
92
5.4.2 Effect of thermal cycling on phase transformation temperatures of 5Cu
alloy
Fig. 5.13 represents the DSC heating and cooling curves for the evolution of
transformation temperatures during five thermal cycles for solution treated sample of
5Cu alloy. The calculated transformation temperatures were plotted in Fig. 5.14 to
compare the change in transformation temperatures due to increasing thermal cycles.
Here all the DSC curves show sharp peaks with well estimated phase transformation
temperatures. It can be observed from Fig. 5.14 that the transformation temperatures
dropped quickly in the second cycle and then decreased slightly in the next cycles
and continue to be decreased till fifth cycle.
The Ms; 155°C in the first cycle dropped quickly to 153°C in the second cycle,
resulted a net decrease of 2°C. Similarly the Af; 178°C in the first cycle dropped
quickly to 173°C in the second cycle, resulted a net decrease of 5°C. Thermal
hysteresis in the second cycle decreased by 3°C from 23°C to 20°C. The decrease in
Af and thermal hysteresis in third thermal cycle reduced as compared to the second
thermal cycle. The Ms decreased by 2°C from 153°C to 151°C and Af decreased by
3°C from 173°C to 170°C in the third cycle. Similarly thermal hysteresis in the third
cycle decreased by only 1°C from 20°C to 19°C. In the fourth thermal cycle, the
decrease in transformation temperatures and thermal hysteresis further lowered as
compared to the third thermal cycle. The Ms decreased by 0.5°C from 151°C to
150.5°C and Af decreased by 1°C from 170°C to 169°C in the fourth cycle.
Similarly thermal hysteresis in the fourth cycle decreased by 0.5°C from 19°C to
18.5°C. The decrease in transformation temperatures and thermal hysteresis in fifth
thermal cycle further lowered as compared to the fourth thermal cycle.
93
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t Fl
ow
(W
/g)
exo
up
Temperature (oC)
0.3
5th Cycle
5th Cycle
1st Cycle
Cooling
Heating
Fig. 5.13 DSC curves representing the transformation temperatures during five
thermal cycles of 5Cu alloy
110
140
170
200
230
Tran
sfo
rmat
ion
Tem
per
atu
res
(oC
)
Cycle Number
Ms
Mf
As
Af
1 2 3 4 5
Fig. 5.14 Effect of thermal cycle on transformation temperatures of 5Cu alloy
94
The Ms decreased by 0.5°C from 150.5°C to 150°C and Af decreased by 1°C from
169°C to 168°C in the fifth cycle. Similarly thermal hysteresis in the fifth cycle
decreased by 0.5°C from 18.5°C to 18°C. During five thermal cycles, the Ms and Af
dropped by 5°C and 10°C respectively, whereas thermal hysteresis decreased by 5°C.
The reasons for decreasing the transformation temperatures and thermal hysteresis by
increasing thermal cycle in 5Cu alloy are given in section 5.4.1.
5.4.3 Effect of thermal cycling on phase transformation temperatures of 10Cu
alloy
Fig. 5.15 shows the DSC heating and cooling curves for the evolution of
transformation temperatures during five thermal cycles for solution treated sample of
10Cu alloy. The calculated transformation temperatures were drawn in Fig. 5.16. It
can be observed that all the DSC curves had sharp peaks with well estimated phase
transformation temperatures. From Fig. 5.16, it can be viewed that the transformation
temperatures dropped quickly in the second cycle and then decreased slightly in the
next cycles and become almost stable in the fifth cycle.
The Ms; 170°C in the first cycle dropped quickly to 167.2°C in the second cycle,
resulted a net decrease of 2.8°C. Similarly the Af; 192.2°C in the first cycle dropped
quickly to 186.7°C in the second cycle, resulted a net decrease of 5.5°C. Thermal
hysteresis in the second cycle decreased by 2.7°C from 22.2°C to 19.5°C. The
decrease in transformation temperature and thermal hysteresis in third thermal cycle
reduced as compared to the second thermal cycle. The Ms decreased by 0.9°C from
167.2°C to 166.3°C and Af decreased by 1.5°C from 186.7°C to 185.2°C in the third
cycle. Similarly thermal hysteresis in the third cycle decreased by 0.6°C from
19.5°C to 18.9°C. In the fourth thermal cycle, the decrease in transformation
temperatures and thermal hysteresis further lowered as compared to the third thermal
cycle. The Ms remained stable at 166.3°C while the Af decreased by 0.9°C from
185.2°C to 184.3°C in the fourth cycle. Similarly thermal hysteresis in the fourth
cycle decreased by 0.9°C from 18.9°C to 18°C.
95
The decrease in Af and thermal hysteresis in fifth thermal cycle further lowered as
compared to the fourth thermal cycle. The Ms remained stable at 166.3°C while the
100 150 200 250
Hea
t Fl
ow
(W
/g)
exo
up
Temperature (oC)
0.3
5th Cycle
5th Cycle
1st Cycle
Cooling
Heating
Fig. 5.15 DSC curves representing the transformation temperatures during five
thermal cycles of 10Cu alloy
110
140
170
200
230
Tran
sfo
rmat
ion
Tem
per
atu
res
(oC
)
Cycle Number
Ms
Mf
As
Af
1 2 3 4 5
Fig. 5.16 Effect of thermal cycle on transformation temperatures of 10Cu alloy
96
Af decreased by 0.7°C from 184.3°C to 183.6°C in the fifth cycle. Similarly thermal
hysteresis in the fifth cycle decreased by 0.7°C from 18°C to 17.3°C. After
completing thermal cycles for five times, the Ms and Af lowered by 3.7°C and 8.7°C
respectively. At the same time, thermal hysteresis decreased by 4.9°C.
It was observed that the martensite start temperatures of solution treated sample of
10Cu alloy decreased quickly in the second cycle and then reduced in third cycle and
remained stable till last thermal cycle. The faster decrease in transformation
temperatures in solution treated sample was due to the generation of dislocations and
other defects during repeated motion of the austenite-martensite interface. This
decrease in transformation temperature is unavoidable for quenched and annealed
TiNi-based alloys, however, this can be prevented by aging and thermo-mechanical
treatment [147]. Rehman et al. [153] reported that by aging the same alloy at 600°C
for 3 hours, the decrease in Ms and Af temperatures is only 2°C, while thermal
hysteresis remained stable at 28°C after five thermal cycles. It was also observed that
the first heating cycle of solution treated sample resulted in higher transformation
temperatures and then decreased faster in the second cycle. The largest drop in
austenite transformation temperatures after first heating cycle was attributed to the
increase in formation of dislocations [151].
5.4.4 Effect of thermal cycling on phase transformation temperatures of 15Cu
alloy
Fig. 5.17 represents the DSC heating and cooling curves for the evolution of
transformation temperatures during five thermal cycles for solution treated sample of
15Cu alloy. The calculated transformation temperatures were plotted in Fig. 5.18 to
compare the variation in transformation temperatures due to increasing thermal
cycles. All the DSC curves presented sharp peaks with well estimated phase
transformation temperatures. It can be observed from Fig. 5.18 that the
transformation temperatures dropped quickly in the second cycle and then decreased
slightly in the third cycle and become stable in the fourth cycle.
97
100 150 200 250
Hea
t Fl
ow
(W
/g)
exo
up
Temperature (oC)
0.3
5th Cycle
5th Cycle
1st Cycle
Cooling
Heating
Fig. 5.17 DSC curves representing the transformation temperatures during five
thermal cycles of 15Cu alloy
110
140
170
200
230
Tran
sfo
rmat
ion
Tem
per
atu
res
(oC
)
Cycle Number
Ms
Mf
As
Af
1 2 3 4 5
Fig. 5.18 Effect of thermal cycle on transformation temperatures of 15Cu alloy
98
The Ms dropped by 2°C from 203°C to 201°C, whereas the Af dropped by 4°C from
219°C to 215°C in the second cycle. Thermal hysteresis in the second cycle
decreased by 2°C from 16°C to 14°C. However the decrease in transformation
temperature and thermal hysteresis in third thermal cycle was lower as compared to
the second thermal cycle. The Ms decreased by 1°C from 201°C to 200°C and Af
decreased by 1°C from 215°C to 214°C in third cycle. Thermal hysteresis in third
cycle remained stable at 14°C. In fourth and fifth thermal cycles, the Ms and Af
remained stable at 200°C and 214°C respectively, while thermal hysteresis also
remained unchanged at 14°C. The faster decrease in transformation temperatures and
thermal hysteresis during second thermal cycle has been explained in section 5.4.1.
During five thermal cycles, the Ms and Af decreased by 3°C and 5°C, whereas
thermal hysteresis dropped by 2°C.
5.5 Summary
In this chapter the effect of 5%, 10% and 15% Cu addition in TiNiPd and aging at
400°C, 500°C, 600°C and 700°C for 3 hours of TiNiPdCu alloys on transformation
temperatures under stress free condition were presented. Transformation
temperatures of all four alloys in both conditions; solution treated and aged were
found by differential scanning calorimetry.
Increasing the Cu concentration in TiNiPd alloys, the transformation temperatures
and transformation heats increased significantly whereas thermal hysteresis
decreased in solution treated condition. Adding the Cu content by 5 ~ 15%, the
martensite start temperatures increased by 9 ~ 42% and thermal hysteresis decreased
by 6 ~ 35%. This behavior confirmed that addition of Cu in place of Ni in TiNiPd
alloys was very beneficial to improve its transformation temperatures and decrease
its thermal hysteresis.
When 0Cu alloy was aged, the transformation temperatures slightly increased as the
aging temperature increased from 400°C to 700°C. The martensite start temperature
was increased by 1% at aging temperature of 600°C as compared to solution treated
condition. By aging the 5Cu alloy, the transformation temperatures increased slightly
when the aging temperature was increased from 400°C to 600°C and then slightly
99
decreased at aging temperature of 700°C. At aging temperature of 600°C, the
martensite start temperature was increased by 12%. For 10Cu and 15Cu alloys, the
transformation temperatures decreased significantly as the aging temperature
increased from 400°C to 600°C and then increased comparable to that of solution
treated condition when the aging temperature was increased to 700°C. The
martensite start temperatures of 10Cu and 15Cu alloys were decreased by 18% and
19% respectively at aging temperature of 600°C. From these observations it can be
concluded that aging of 0Cu and 5Cu alloys was beneficial to improve their
transformation temperatures whereas in case of 10Cu and 15Cu alloys, the
transformation temperatures decreased remarkably.
In solution treated 0Cu alloy, after completion of five thermal cycles, the Ms and Af
decreased by 7.5°C and 14°C respectively, whereas thermal hysteresis decreased by
6.5°C. In case of 5Cu alloy, the Ms, Af and thermal hysteresis dropped by 5°C,
10°C and 5°C respectively. By thermal cycling the 10Cu alloy, the Ms and Af
lowered by 3.7°C and 8.7°C respectively, whereas thermal hysteresis lowered by
4.9°C. Similarly, for 15Cu alloy the Ms, Af and thermal hysteresis decreased by 3°C,
5°C and 2°C respectively. Moreover, during thermal cycling, the transformation
temperatures of solution treated 0Cu alloys decreased by increasing the number of
thermal cycles, however the drop was faster in initial cycles and then reduced in the
final cycles. Similar behavior was also observed during thermal cycling of 5Cu,
10Cu and 15Cu alloys, however the drop in transformation temperatures was less as
compared to 0Cu alloy. This behavior confirmed that, by increasing the Cu content in
TiNiPd alloys, thermal stability improved. Hence 5Cu, 10Cu and 15Cu alloys were
more thermally stable as compared to 0Cu alloy.
100
Chapter – 6
Effect of Copper Addition and Aging
on Mechanical Properties of TiNiPd
Alloys
6.1 Introduction
In this chapter, the mechanical properties i.e. Stress-strain relation, yield strength,
ultimate tensile strength (fracture strength) and hardness of 0Cu, 5Cu, 10Cu and
15Cu alloys for different heat treated conditions (solution treated and aged) have
been shown. It has been investigated that the mechanical properties can be improved
by applying two techniques; solid solution strengthening by 5%, 10% and 15% Cu
addition, and precipitation strengthening by aging the 0Cu, 5Cu, 10Cu and 15Cu
alloys at different temperatures ranging from 400°C to 700°C for 3 hours.
6.2 Effect of Cu addition on hardness
Fig. 6.1 represents the Vickers microhardness of 0Cu, 5Cu, 10Cu and 15Cu alloys
solution treated at 900°C for 1 hour. The hardness values of each alloy were
determined from the average of 5 measurements taken at different locations on each
sample. It was observed that the hardness increased by increasing the Cu content in
place of Ni. Hardness of 0Cu alloy; 245 Hv increased to 252 Hv when Ni was
replaced by 5% Cu. By increasing the Cu concentration from 5% to 10%, the
hardness further increased to 255 Hv. Similarly by further increasing Cu
concentration to 15%, the hardness further increased to 263 Hv. The behavior of
increasing hardness was identical to increase in transformation temperatures as
discussed in Chapter – 5. Thus it can be concluded that by increasing the substitution
of Ni by Cu, both the hardness and transformation temperatures were increased. The
increase in hardness by increasing the substitution of Ni by Cu was attributed to the
solid solution strengthening. As the atomic radius of Cu (0.128 nm) is relatively
greater than the atomic radius of Ni (0.125 nm), thus caused to increase the hardness.
101
6.3 Effect of aging on hardness
6.3.1 Effect of aging on hardness of 0Cu alloy
Fig. 6.2 represents the hardness of 0Cu alloy; solution treated and aged at given
temperatures for 3 hours. Hardness of solution treated sample; 245 Hv slightly
decreased to 236 Hv when the same alloy was aged at 400°C. The slight decrease in
hardness was attributed to the removal of thermal stresses which were produced
during quenching process due to higher temperature gradient. When the alloy was
aged at 500°C, the hardness of the alloy increased to 258 Hv, suggesting that
precipitation process has been initiated at this temperature. However due to very fine
size of precipitates, it could not be detected in backscattered SEM images given in
Fig. 4.6(a – d). By increasing the aging temperatures to 600°C, the hardness of the
alloy increased slightly to 263 Hv. This slight increase in hardness can be attributed
to the slight increase in volume fraction of the precipitates, however the size
remained the same and could not be detected. By increasing the aging temperature
further to 700°C, again the hardness of alloy reduced to 249 Hv. The decrease in
hardness showed that precipitation was not initiated at higher aging temperature.
200
220
240
260
280
300
Mic
roh
ard
nes
s (H
v)
Cu (at%)
0 5 10 15
Fig. 6.1 Microhardness of solution treated samples of 0Cu, 5Cu, 10Cu and 15Cu
alloys
102
6.3.2 Effect of aging on hardness of 5Cu alloy
Fig. 6.3 represents the hardness of 5Cu alloy; solution treated and aged at given
temperatures for 3 hours. Hardness of solution treated sample; 252 Hv slightly
decreased to 242 Hv when the same alloy was aged at 400°C. The slight decrease in
hardness is attributed to the removal of thermal stresses. Moreover it can also be
confirmed that the formation of precipitates was not started at the said temperature
150
250
350
Har
dn
ess
(Hv)
Aging Temperature (°C)
ST 400 500 600 700
Fig. 6.2 Microhardness of 0Cu alloy aged at different temperatures
150
250
350
Har
dn
ess
(H
v)
Aging Temperature (°C)
ST 400 500 600 700
Fig. 6.3 Microhardness of 5Cu alloy aged at different temperatures
103
and thus not detected in Fig. 4.7a. When the alloy was aged at 500°C, the hardness of
the alloy increased to 280 Hv, suggesting that precipitation process has been
initiated. It can be observed from backscattered SEM images of Fig. 4.7b, that
precipitates of brighter contrast have been formed along the grain boundaries caused
to increase the hardness. By further increasing the aging temperatures to 600°C, the
hardness of the alloy decreased slightly to 276Hv. By increasing the aging
temperature further to 700°C, again the hardness of alloy reduced to 246 Hv. This
result shows that, although precipitates have been formed as shown in Fig. 4.7d,
however due to incoherency, the hardness decreased.
6.3.3 Effect of aging on hardness of 10Cu alloy
Fig. 6.4 shows the hardness of the samples solution treated and aged at given
temperatures for 3 hours. Hardness of solution treated sample; 255 Hv slightly
increased to 260 Hv when the same alloy was aged at 400°C. It was observed
previously that, at aging temperature of 400°C, hardness of 0Cu and 5Cu alloys were
decreased due to removal of thermal stresses and absence of precipitation. However
in case of 10Cu alloy, the hardness slightly increased confirming the formation of
precipitates started at the said temperature. By increasing the aging temperatures to
Fig. 6.4 Microhardness of 10Cu alloy aged at different temperatures
150
250
350
450
Har
dn
ess
(H
v)
Aging Temperature (°C)
ST 400 500 600 700
104
500°C and 600°C, the hardness of the samples abruptly increased to 369 Hv and 408
Hv, respectively. The significant increase in the hardness values of the samples aged
at 500°C and 600°C were suggested to be the outcome of precipitation process. The
nucleation sites along the grain boundaries have been observed to be formed at
400°C as shown in Fig. 4.8a, resulted in slight increase in the hardness. By
increasing the aging temperatures to 500°C and 600°C, the sizes and densities of
both types of precipitates rapidly increased as shown in Fig. 4.8b and c , caused
further significant increase in the hardness. From these observations, it can be easily
derived that both type of precipitates formed at 400°C to 600°C were found to be
coherent.
However, by aging at 700°C, the hardness of the alloy abruptly decreased to 258 Hv
almost equal to the hardness of solution treated sample. Here the aging temperature
was higher than the recrystallization temperature of the TiNiPd alloys (~ 600°C)
[154], hence the precipitation hardening was not promoted. It can be observed from
Fig. 4.8d, that both types of precipitates were present in the alloy aged at 700°C,
however it could not cause to increase the hardness. The decrease in the hardness at
aging temperatures of 700°C can be attributed to the formation of incoherent
precipitates with comparatively low densities of relatively larger sizes.
From the above results, it can be summarized that at low aging temperatures (400°C
and below), no remarkable increase in hardness was observed. Aging at intermediate
temperatures (500°C and 600°C), increased the hardness of the alloy significantly.
However at high aging temperatures (700°C and above), the hardness of the alloy
was observed to be decreased significantly.
The hardness of the same alloy aged at 500°C for 1 hour, processed and heat treated
at the same conditions was reported by Imran et al. [140]. It was observed that the
hardness was not increased and remained at the same level as that of solution treated
sample. In this research, the hardness of the sample aged at 500°C for 3 hours
increased significantly up to 369 Hv. From these results, it can be confirmed, that for
the proper age hardening process, aging temperature as well as proper time duration
are important. From the same reference, it was also noticed that by increasing the
annealing temperature from 350°C to 400°C, the hardness of the alloy decreased
105
significantly. However, annealing at 450°C resulted in remarkable increase in the
hardness from 352 Hv to 544 Hv and then the hardness decreased continuously by
increasing the annealing temperature to 700°C. It is important to note that the
decrease in the hardness at annealing temperature of 400°C was attributed to the
recovery process, whereas formation of nano-scaled precipitates was responsible for
the remarkable increase in the hardness at 450°C. On the other hand, by increasing
the aging temperature from 400°C to 600°C, the hardness of the alloy increased
continuously from 260 to 408 Hv and then decreased in the same manner when the
aging temperature was further increased to 700°C, in the present study. By
comparing these results, it can be concluded that the age hardening behavior in the
present alloy was observed to be significantly different from that of annealing after
cold working.
6.3.4 Effect of aging on hardness of 15Cu alloy
Fig. 6.5 shows the hardness of the samples solution treated and aged at given
temperatures for 3 hours. Hardness of solution treated sample; 263 Hv slightly
increased to 270 Hv when the same alloy was aged at 400°C. The increase in
hardness at aging temperature of 400°C indicated that precipitation process has been
started at the said temperature like 10Cu alloy. By increasing the aging temperatures
to 500°C and 600°C, the hardness of the samples abruptly increased to 442 Hv and
480 Hv respectively. The significant increase in the hardness values of the samples
aged at 500°C and 600°C were suggested to be the outcome of precipitation process.
The nucleation sites along the grain boundaries have been observed to be formed at
400°C as shown in Fig. 4.9a, resulted in slight increase in the hardness. By
increasing the aging temperatures to 500°C and 600°C, the sizes and densities of
both types of precipitates rapidly increased as shown in Fig. 4.9b and c , caused
further significant increase in the hardness. From these observations, it can be easily
derived that both type of precipitates formed at 400°C to 600°C were found to be
coherent.
106
However, by aging at 700°C, the hardness of the alloy abruptly decreased to 260 Hv
almost equal to the hardness of solution treated sample. Here the aging temperature
was higher than the recrystallization temperature of the TiNiPd alloys (~ 600°C)
[154], hence the precipitation hardening was not promoted. It can be observed from
Fig. 4.9d, that both types of precipitates were present in the alloy aged at 700°C,
however it could not increase the hardness of the alloy. The decrease in the hardness
values at aging temperatures of 700°C can be attributed to the formation of
incoherent precipitates with comparatively low densities of relatively larger sizes.
From the above results, it can be summarized that at low aging temperatures (400°C
and below), no remarkable increase in hardness was observed. Aging at intermediate
temperatures (500°C and 600°C), resulted significant increase in the hardness
whereas at high aging temperatures (700°C and above), the hardness of the alloy was
observed to be decreased significantly.
150
250
350
450
550
Har
dn
ess
(Hv)
Aging Temperature (°C)
ST 400 500 600 700
Fig. 6.5 Vickers microhardness of 15Cu alloy aged at different temperatures
107
6.4 Effect of Cu addition on mechanical strength
6.4.1 Effect of Cu addition on mechanical strength in martensite phase
It is important to investigate the mechanical properties of shape memory alloys in
both phases i.e. martensite and austenite, because the stress-strain relations in both
phases are different from each other. Other conventional structural materials like
steel, has maximum yield strength at room temperature and then decreases by
increasing the testing temperature above 100°C. Conversely, for high temperature
shape memory alloys, the yield strength at room temperature in martensite phase
(lower than Mf) must be less than the yield strength at higher temperature (greater
than Af) in austenite phase. Austenite yield stress represents the critical stress for
slip deformation while martensite yield stress represents critical stress for shear of
martensite twins. For feasible actuators, the critical stress for shear must be lower
than the critical stress for slip, so that when stress is applied, it results in shape
deformation by shear of martensite twins rather than via dislocation generation and
movement.
Tensile stress-strain curves of solution treated samples tested in the martensite phase,
50°C below Mf of 0Cu, 5Cu, 10Cu and 15Cu alloys are shown in Fig. 6.6. The yield
stress for all alloys was calculated by drawing a parallel line, 0.2% offset to the
elastic region of stress-strain curve as shown in Fig. 6.6. The yield stress (σy),
fracture stress (σf) and fracture strain (εf) calculated from these curves in the
martensitic condition for all alloys are shown in Table 6.1.
The martensite yield stress corresponds to the stress required for reorientation of
martensite twins (also called stress for detwinned martensite). From Table 6.1, it can
be observed that martensite yield stress of 5Cu alloy increased by 33 MPa from 290
to 323 MPa with respect to the baseline 0Cu alloy. Similarly, the stress at which the
fracture occurred in the alloy (σf) also increased by 60 MPa; 951.6 MPa of 0Cu alloy
increased to 1011.6 MPa for 5Cu alloy. However the strain at which the fracture
occurred in the alloy (εf) decreased by 0.75%; fracture strain of 10% for 0Cu alloy
decreased to 9.25% for 5Cu alloy.
108
In the same manner, when Ni was substituted by 10% Cu, the relevant mechanical
properties further increased. The martensite yield stress of 10Cu alloy increased by
56 MPa from 290 to 346 MPa, whereas the fracture stress increased by 88.4 MPa;
951.6 MPa of 0Cu alloy increased to 1040 MPa as compared to 0Cu alloy. However
the fracture strain decreased by 1.03%; fracture strain of 10% for 0Cu alloy
decreased to 8.97% for 10Cu alloy. The martensite yield stress of 15Cu alloy
increased by 76 MPa from 290 to 366 MPa with respect to the baseline 0Cu alloy.
Fig. 6.6 Tensile stress-strain curves of solution treated 0Cu, 5Cu, 10Cu and 15Cu alloys,
tested in martensite phase (Mf – 50°C)
0
300
600
900
1200
0 3 6 9 12
Stre
ss (
MP
a)
Strain (%)
Mf = 119°C
Testing Temperature = 69°C
0Cu alloy
0
300
600
900
1200
0 3 6 9 12Strain (%)
Mf = 133°C
Testing Temperature = 83°C
5Cu alloy
0
300
600
900
1200
0 3 6 9 12
Stre
ss (
MP
a)
Strain (%)
Mf = 151°C
Testing Temperature = 101°C
10Cu alloy
0
300
600
900
1200
0 3 6 9 12Strain (%)
Mf = 176°C
Testing Temperature = 126°C
15Cu alloy
109
Similarly the fracture stress increased by 114.4 MPa; 951.6 MPa of 0Cu alloy
increased to 1066 MPa. However the fracture strain decreased by 1.28%; fracture
strain of 10% for 0Cu alloy decreased to 8.72%.
Table 6.1 Yield stress, fracture stress and fracture strain calculated from stress-strain
curves of Fig. 6.6 for solution treated alloys tested in martensite phase (Mf – 50°C)
Alloy Testing Temp (°C) Yield Stress, σy
(MPa)
Fracture Stress,
σf (MPa)
Fracture
Strain, εf (%)
0Cu 69 290 951.6 10
5Cu 83 323 1011.6 9.25
10Cu 101 346 1040 8.97
15Cu 126 366 1066 8.72
6.4.2 Effect of Cu addition on mechanical strength in austenite phase
Fig. 6.7 represents the tensile stress-strain curves of solution treated samples tested in
the austenite phase, 50°C above Af of 0Cu, 5Cu, 10Cu and 15Cu alloys.
Measurement of austenite yield stress was carried out by the same procedure as
discussed earlier in section 6.4.1. σy, σf and εf calculated from these curves in the
austenite phase for all alloys are shown in Table 6.2.
The austenite yield stress represents the critical stress for slip; stress required for the
deformation of material through dislocations generation and their movement. It can
be observed from Table 6.2 that the austenite yield stress of 440 MPa for 0Cu alloy
increased to 502 MPa for 5Cu alloy resulted net increase of 62 MPa. The fracture
stress also increased by 40.9 MPa; 1121.3 MPa for 0Cu alloy increased to 1162.2
MPa for 5Cu alloy. Here the fracture strain was observed to be decreased by 1.2%;
fracture strain of 8.5% for 0Cu alloy decreased to 7.3% for 5Cu alloy.
Similarly the austenite yield stress for 10Cu alloy increased by 102 MPa to 542 MPa
as compared to 440 MPa of the baseline 0Cu alloy. The fracture stress also increased
by 80.7 MPa; 1121.3 MPa for 0Cu alloy increased to 1202 MPa for 10Cu alloy. The
fracture strain was observed to be decreased by 1.54%; fracture strain of 8.5% for
110
0Cu alloy decreased to 6.96% for 10Cu alloy. The austenite yield stress for 15Cu
alloy further increased by 124 MPa from 440 MPa to 564 MPa. The fracture stress
also increased by 98.7 MPa; 1121.3 MPa for 0Cu alloy increased to 1220 MPa for
15Cu alloy. The fracture strain was observed to be decreased further by 1.74%;
fracture strain of 8.5% for 0Cu alloy decreased to 6.76% for 15Cu alloy.
0
300
600
900
1200
0 3 6 9
Stre
ss (
MP
a)
Strain (%)
Af = 167°C
Testing Temperature = 217°C
0Cu alloy
0
300
600
900
1200
0 3 6 9Strain (%)
Af = 178°C
Testing Temperature = 228°C
5Cu alloy
0
300
600
900
1200
0 3 6 9
Stre
ss (
MP
a)
Strain (%)
Af = 192°C
Testing Temperature = 242°C
10Cu alloy
0
300
600
900
1200
0 3 6 9Strain (%)
Af = 219°C
Testing Temperature = 269°C
15Cu alloy
Fig. 6.7 Tensile stress-strain curves of solution treated 0Cu, 5Cu, 10Cu and 15Cu alloys,
tested in austenite phase (Af + 50°C)
111
Table 6.2 Yield stress, fracture stress and fracture strain calculated from stress-strain
curves of Fig. 6.7 for solution treated alloys tested in austenite phase (Af + 50°C)
Alloy Testing Temp (°C) Yield Stress, σy
(MPa)
Fracture Stress,
σf (MPa)
Fracture
Strain, εf (%)
0Cu 217 440 1121.3 8.5
5Cu 228 502 1162.2 7.3
10Cu 242 542 1202 6.96
15Cu 269 564 1220 6.76
6.5 Comparison between the mechanical properties of martensite
and austenite phases in solution treated condition
Fig. 6.8 represents the effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite yield stress, tested at (Mf – 50°C) and austenite yield stress,
tested at (Af + 50°C).
150
300
450
600
Yie
ld S
tre
ss (
MP
a)
Cu (at%)
Martensite yield stress
Austenite yield stress
0 5 10 15
Fig. 6.8 Effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite yield stress, tested at (Mf – 50°C) and
austenite yield stress, tested at (Af + 50°C)
112
It can be observed that by substitution of Ni by 5%, 10% and 15% Cu, the yield
stress in the martensite phase increased by 33, 56 and 76 MPa respectively, whereas
the austenite yield stress increased by 62, 102 and 124 MPa respectively.
Fig. 6.9 represents the effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite fracture stress, tested at (Mf – 50°C) and austenite fracture
stress, tested at (Af + 50°C). It can be observed that by substitution of Ni by 5%, 10%
and 15% Cu, the fracture stress in the martensite phase increased by 60, 88.4 and
114.4 MPa respectively, whereas the austenite fracture stress increased by 40.9, 80.7
and 98.7 MPa respectively.
Fig. 6.10 represents the effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite fracture strain, tested at (Mf – 50°C) and austenite fracture
strain, tested at (Af + 50°C). It can be observed that by substitution of Ni by 5%, 10%
and 15% Cu, the fracture strain in the martensite phase decreased by 0.75%, 1.03%
and 1.28% respectively, whereas the austenite fracture strain decreased by 1.2%,
1.54% and 1.74% respectively.
900
1000
1100
1200
1300
Frac
ture
Str
ess
(MP
a)
Cu (at%)
Martensite fracture stress
Austenite fracture stress
0 5 10 15
Fig. 6.9 Effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite fracture stress, tested at (Mf – 50°C) and
austenite fracture stress, tested at (Af + 50°C)
113
The increase in both yield stresses (stress for reorientation of martensite and stress
for slip deformation) and fracture stress (maximum stress) is attributed to the solid
solution strengthening due to partial substitution of Ni by Cu. Here the atomic radius
of Cu (0.128 nm) is relatively greater than the atomic radius of Ni (0.125 nm) and
therefore responsible for solid solution strengthening. By solution strengthening
effect, the ductility of 0Cu alloy lowered and resulted in relatively low fracture strain
in both phases.
It can also be observed that the increase in stress for reorientation of martensite
variants; martensite yield stress (although it was not required for better shape
memory alloys) was observed to be 33, 56 and 76 MPa for 5Cu, 10Cu and 15Cu
alloys respectively. On the other hand, the critical stress for slip deformation;
austenite yield stress (essentially required for better shape memory alloys) was
increased by 62, 102 and 124 MPa for 5Cu, 10Cu and 15Cu alloys respectively. As
the increase in critical stress for slip was more beneficial than the increase in stress
for reorientation of martensite variants, thus it is suggested that by partial substitution
of Ni with 5%, 10% and 15% Cu, the critical stress for slip can be increased by 29,
46 and 48 MPa (difference in net increase in critical stress for slip and net increase in
the stress for reorientation of martensite variants).
5
7
9
11
Frac
ture
Str
ain
(%
)
Cu (at%)
Martensite fracture strain
Austenite fracture strain
0 5 10 15
Fig. 6.10 Effect of partial substitution of Ni by Cu in solution treated
0Cu alloy on martensite fracture strain, tested at (Mf – 50°C) and
austenite fracture strain, tested at (Af + 50°C)
114
6.6 Effect of aging on mechanical strength
6.6.1 Effect of aging on mechanical strength in martensite phase
Tensile stress-strain curves of 600°C-aged samples tested in the martensite phase,
50°C below Mf for 0Cu, 5Cu, 10Cu and 15Cu alloys are shown in Fig. 6.11. The
yield stress, fracture stress and fracture strain calculated from these curves in the
martensitic condition for all alloys are shown in Table 6.3. From Table 6.3, it can be
observed that martensite yield stress of 5Cu alloy increased by 32 MPa from 300 to
332 MPa with respect to the baseline 0Cu alloy. Similarly, the fracture stress also
increased by 65 MPa; 960 MPa of 0Cu alloy increased to 1025 MPa for 5Cu alloy.
However the fracture strain decreased by 0.6%; fracture strain of 9.5% for 0Cu alloy
decreased to 8.9% for 5Cu alloy. In the same manner, when Ni was substituted by
10% Cu, the relevant mechanical properties further increased. The martensite yield
stress of 10Cu alloy increased by 55 MPa from 300 to 355 MPa with respect to the
baseline 0Cu alloy. The fracture stress increased by 195 MPa; 960 MPa of 0Cu alloy
increased to 1155 MPa for 10Cu alloy. However the fracture strain decreased by
3.2%; fracture strain of 9.5% for 0Cu alloy decreased to 6.3% for 10Cu alloy. The
martensite yield stress of 15Cu alloy increased by 85 MPa from 300 to 385 MPa
whereas the fracture stress increased by 270 MPa; 960 MPa of 0Cu alloy increased to
1230 MPa for 15Cu alloy. However the fracture strain decreased by 3.3% from 9.5%
to 6.2% for 15Cu alloy.
115
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Stre
ss (
MP
a)
Strain (%)
0Cu alloy
Mf = 121°C Testing Temperature = 71°C
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Strain (%)
5Cu alloy
Mf = 144°C
Testing Temperature = 94°C
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Stre
ss (
MP
a)
Strain (%)
10Cu alloy
Mf = 130°C Testing Temperature = 80°C
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Strain (%)
15Cu alloy
Mf = 140°C Testing Temperature = 90°C
Fig. 6.11 Tensile stress-strain curves of 600°C-aged 0Cu, 5Cu, 10Cu and 15Cu alloys,
tested in martensite phase (Mf – 50°C)
116
Table 6.3 Yield stress, fracture stress and fracture strain calculated from stress-strain
curves of Fig. 6.11 for 600°C-aged alloys tested in martensite phase (Mf – 50°C)
Alloy Testing Temp (°C) Yield Stress, σy
(MPa)
Fracture Stress,
σf (MPa)
Fracture
Strain, εf (%)
0Cu 71 300 960 9.5
5Cu 94 332 1025 8.9
10Cu 80 355 1155 6.3
15Cu 90 385 1230 6.2
6.6.2 Effect of aging on mechanical strength in austenite phase
Tensile stress-strain curves of 600°C-aged samples tested in the austenite phase,
50°C above Af for 0Cu, 5Cu, 10Cu and 15Cu alloys are shown in Fig. 6.12. The yield
stress, fracture stress and fracture strain calculated from these curves in the austenite
phase for all alloys are shown in Table 6.4. From Table 6.4, it can be observed that
austenite yield stress of 5Cu alloy increased by 95 MPa from 410 to 505 MPa
whereas the fracture stress increased by 100 MPa; 1070 MPa of 0Cu alloy increased
to 1170 MPa for 5Cu alloy. However the fracture strain decreased by 1.5% from
8.3% for 0Cu alloy to 6.8% for 5Cu alloy. In the same manner, when Ni was
substituted by 10% Cu, the relevant mechanical properties further increased. The
austenite yield stress of 10Cu alloy increased by 135 MPa from 410 to 545 MPa
whereas the fracture stress increased by 190 MPa; 1070 MPa of 0Cu alloy increased
to 1260 MPa for 10Cu alloy. However the fracture strain decreased by 1.8%; fracture
strain of 8.3% for 0Cu alloy decreased to 6.5% for 10Cu alloy.
117
When Ni was substituted by 15% Cu, the mechanical properties further increased.
The austenite yield stress of 15Cu alloy increased by 180 MPa from 410 to 590 MPa
whereas the fracture stress increased by 250 MPa; 1070 MPa of 0Cu alloy increased
to 1320 MPa for 15Cu alloy. On the hand, the fracture strain decreased by 3.05%
from 8.3% to 5.25% for 15Cu alloy.
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Stre
ss (
MP
a)
Strain (%)
Af = 169°C Testing Temperature = 219°C
0Cu alloy
0
300
600
900
1200
1500
0 2.5 5 7.5 10Strain (%)
Af = 194°C Testing Temperature = 244°C
5Cu alloy
0
300
600
900
1200
1500
0 2.5 5 7.5 10
Stre
ss (
MP
a)
Strain (%)
Af = 166°C Testing Temperature = 216°C
10Cu alloy
0
300
600
900
1200
1500
0 2.5 5 7.5 10Strain (%)
Af = 189°C Testing Temperature = 239°C
15Cu alloy
Fig. 6.12 Tensile stress-strain curves of 600°C-aged 0Cu, 5Cu, 10Cu and 15Cu alloys, tested
in austenite phase (Af + 50°C)
118
Table 6.4 Yield stress, fracture stress and fracture strain calculated from stress-strain
curves of Fig. 6.12 for 600°C-aged alloys tested in austenite phase (Af + 50°C)
Alloy Testing Temp (°C) Yield Stress, σy
(MPa)
Fracture Stress,
σf (MPa)
Fracture
Strain, εf (%)
0Cu 219 410 1070 8.3
5Cu 244 505 1170 6.8
10Cu 216 545 1260 6.5
15Cu 239 590 1320 5.25
6.7 Comparison between the mechanical properties of martensite
and austenite phases in 600°C-aged condition
Fig. 6.13 represents the effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite yield stress, tested at (Mf – 50°C) and austenite yield stress,
tested at (Af + 50°C). It can be observed that by substitution of Ni by 5%, 10% and
15% Cu, the yield stress in the martensite phase increased by 32, 55 and 85 MPa
respectively, whereas the austenite yield stress increased by 95, 135 and 180 MPa
respectively.
Fig. 6.14 represents the effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture stress, tested at (Mf – 50°C) and austenite fracture stress,
tested at (Af + 50°C). It can be observed that by substitution of Ni by 5%, 10% and
15% Cu, the fracture stress in the martensite phase increased by 65, 195 and 270
MPa respectively, whereas the austenite fracture stress increased by 100, 190 and
250 MPa respectively.
119
150
300
450
600
Yie
ld S
tres
s (M
Pa)
Cu (at%)
Martensite yield stress
Austenite yield stress
0 5 10 15
Fig. 6.13 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite yield stress, tested at (Mf – 50°C) and austenite
yield stress, tested at (Af + 50°C)
900
1000
1100
1200
1300
1400
Frac
ture
Str
ess
(MP
a)
Cu (at%)
Martensite fracture stress
Austenite fracture stress
0 5 10 15
Fig. 6.14 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture stress, tested at (Mf – 50°C) and austenite
fracture stress, tested at (Af + 50°C)
120
Fig. 6.15 represents the effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture strain, tested at (Mf – 50°C) and austenite fracture strain,
tested at (Af + 50°C). It can be observed that by substitution of Ni by 5%, 10% and
15% Cu, the fracture strain in the martensite phase decreased by 0.6%, 2.2% and
3.3% respectively, whereas the austenite fracture strain decreased by 1.5%, 1.8% and
3.05% respectively. The increase in both yield stresses of 600°C-aged alloys was
attributed to the solid solution strengthening due to partial substitution of Ni by Cu.
It can also be observed that the increase in stress for reorientation of martensite
variants was observed to be 32, 55 and 85 MPa for 5Cu, 10Cu and 15Cu alloys
respectively. On the other hand, the critical stress for slip was increased by 95, 135
and 180 MPa for 5Cu, 10Cu and 15Cu alloys respectively. As the increase in critical
stress for slip is more beneficial than the increase in stress for reorientation of
martensite variants, thus it is suggested that by partial substitution of Ni with 5%,
10% and 15% Cu, the critical stress for slip can be increased by 63, 80 and 95 MPa
(difference in net increase in critical stress for slip and net increase in the stress for
reorientation of martensite variants).
4
6
8
10
Frac
ture
Str
ain
(%
)
Cu (at%)
Martensite fracture strain
Austenite fracture strain
0 5 10 15
Fig. 6.15 Effect of partial substitution of Ni by Cu in 600°C-aged 0Cu
alloy on martensite fracture strain, tested at (Mf – 50°C) and austenite
fracture strain, tested at (Af + 50°C)
121
6.8 Summary
In this chapter the effect of 5%, 10% and 15% Cu addition and aging at 400°C,
500°C, 600°C and 700°C for 3 hours of TiNiPd alloys on mechanical properties was
presented. The hardness of all four alloys in both conditions; solution treated and
aged were found by microhardness tester whereas the yield stress, stress at fracture
and strain at fracture were found by using the tensile testing machine.
Increasing the Cu concentration in TiNiPd alloys, the hardness of the resultant alloys
increased by 3%, 4% and 7% for 5Cu, 10Cu and 15Cu alloys respectively, as
compared to 0Cu alloy, in solution treated condition. By aging the 0Cu and 5Cu
alloys at 400°C, 500°C, 600°C and 700°C for 3 hours, the hardness was slightly
increased till aging temperature of 600°C and then slightly decreased at 700°C. The
hardness of 0Cu and 5Cu alloys increased by 7% and 10% respectively when the
alloys were aged at 600°C. The hardness of 10Cu and 15Cu alloys was remarkably
increased when its aging temperature was increased from 400°C to 600°C and then
decreased remarkably almost equal to solution treated condition. It was noted that the
hardness of 10Cu and 15Cu alloys was increased by 60% and 83% respectively.
In solution treated condition, the yield stress and fracture stress were significantly
increased by increasing the Cu concentration in TiNiPd alloys in both martensite and
austenite phases, however the fracture strain was decreased. Yield stress, fracture
stress and fracture strain were found only for 600°C-aged TiNiPdCu alloys. It was
observed that the yield and fracture stresses were significantly increased by
increasing the Cu concentration in TiNiPd alloys, however fracture strain was
decreased. The yield stress of all four alloys was increased by 3 ~ 5% by aging them
at 600°C as compared to solution treated condition. From these results it can be
concluded that addition of Cu in place of Ni in TiNiPd alloys was beneficial to
improve its hardness, yield and fracture stresses.
By comparing the mechanical properties in martensite and austenite phases, it was
observed that yield and fracture stresses in austenite phase were significantly higher
than that of martensite phase. In austenite phase, the yield and fracture stresses of all
four alloys were found to be greater by 52% ~ 57% and 14% ~ 18% respectively, as
122
compared to martensite phase. However the fracture strain in austenite phase was
less as compared to martensite phase.
123
Chapter – 7
Effect of Copper Addition on Shape
Memory Properties of TiNiPd Alloys
7.1 Introduction
In this chapter, the shape memory properties; transformation temperatures under
biased load, recovered and irrecoverable strains, recovery ratio and work output of
0Cu, 5Cu, 10Cu and 15Cu alloys were determined in solution treated condition. The
detailed results and discussions are presented in the subsequent sections.
7.2 Shape memory properties of TiNiPd alloys with varying Cu
percentage
To investigate the effect of Cu addition on shape memory properties of TiNiPd high
temperature shape memory alloys, constant stress thermal cycling experiments were
conducted at various stress levels in solution treated condition of 0Cu, 5Cu, 10Cu
and 15Cu alloys.
7.2.1 Shape memory properties of 0Cu alloy
Fig. 7.1 shows the strain-temperature curves for solution treated 0Cu alloy at stress
levels of 100 – 500 MPa. Thermal cycling experiments (loading, heating and
cooling) at various stress levels were carried out by the same manner as explained in
section 3.6.2. Transformation temperatures (Ms, Mf, As, Af) at different stress levels
were measured by the tangent intersection method as shown in Fig. 3.7.
Fig. 7.2 represents the change in transformation temperatures (measured from Fig.
7.1) with respect to the applied stress levels for solution treated 0Cu alloy. It can be
observed from Fig. 7.2 that all the transformation temperatures increased as the stress
level increased from 100 MPa to 500 MPa. The Ms and Af temperatures; 150°C and
200°C at 100 MPa, increased to 163°C and 220°C respectively when the applied
stress was increased to 300 MPa. Similarly when the stress was increased to 500
MPa, the Ms and Af temperatures also increased to 177°C and 240°C respectively.
124
Moreover, a linear relationship between the transformation temperatures and applied
stress was observed for all stress levels, satisfied the Clausius-Clapeyron (Cs-Cl)
equation (dσ/dT = constant) [4].
The recovered strain [equal to the difference in strain between the Af + 50°C and
50°C of heating cycle] and irrecoverable strain [equal to the strain between heating
and cooling cycle at Af + 50°C] were calculated by the measurement scheme shown
in Fig. 4.7. The calculated transformation strains (εrec, εirr) for the solution treated
0Cu alloy were re-plotted in Fig. 7.3.
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 7.1 Strain-temperature curves representing the shape memory
properties of solution treated 0Cu alloy at stress levels of 100 – 500 MPa
125
It can be observed from Fig. 7.3, that by increasing the applied stress, both the
recovered and irrecoverable strains also increased. This increase in transformation
strain can be attributed to the increase in number of favored martensite variants in
direction of applied stress [155]. At 100 MPa, the recovered strain of 2.61%
increased to 4.09% when the applied stress was increased to 300 MPa. By further
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 7.2 Change in transformation temperatures of solution treated
0Cu alloy at stress levels of 100 – 500 MPa
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 7.3 Recovered and irrecoverable strains of solution treated 0Cu
alloy at stress levels of 100 – 500 MPa
126
increasing the applied stress to 500 MPa, the recovered strain further increased to
4.8%. Similarly at 100 MPa, the irrecoverable strain of 0.82% increased to 1.51%
when the applied stress was increased to 300 MPa. By further increasing the applied
stress to 500 MPa, the irrecoverable strain further increased to 2.2%.
Fig. 7.4 represents the recovery ratio and work output for solution treated 0Cu alloy,
calculated from the transformation strains and their corresponding stress levels. It
can be noted, that the recovery ratio decreased as the applied stress was increased
from 100 MPa to 500 MPa. The recovery ratio of 76% resulted at 100 MPa
decreased to 73% at 300 MPa and then further decreased to 68% when the applied
stress was further increased to 500 MPa. As shown in Fig. 7.4, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 2.61 J/cm3 at 100 MPa increased to 12.27 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 24 J/cm3.
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 7.4 Recovery ratio and work output of solution treated 0Cu alloy
at stress levels of 100 – 500 MPa
127
7.2.2 Shape memory properties of 5Cu alloy
Fig. 7.5 shows the strain-temperature curves for solution treated 5Cu alloy at stress
levels of 100 – 500 MPa. The transformation temperatures at various stress levels
were calculated by the same manner as explained in section 7.2.1.
Fig. 7.6 represents the variation of transformation temperatures (measured from Fig.
7.5) with respect to the applied stress levels for solution treated 5Cu alloy. It can be
observed from Fig. 7.6 that all the transformation temperatures increased as the stress
level increased from 100 MPa to 500 MPa. The Ms and Af temperatures; 165°C and
215°C at 100 MPa, increased to 180°C and 235°C respectively when the applied
stress was increased to 300 MPa. Similarly when the stress was increased to 500
MPa, the Ms and Af temperatures also increased to 195°C and 253°C respectively.
Moreover, a linear relationship between the transformation temperatures and applied
stress was observed for all stress levels, satisfied the Clausius-Clapeyron (Cs-Cl)
equation.
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 7.5 Strain-temperature curves representing the shape memory
properties of solution treated 5Cu alloy at stress levels of 100 – 500 MPa
128
The recovered and irrecoverable strains were calculated from Fig. 7.5 for the solution
treated 5Cu alloy and re-plotted in Fig. 7.7. It can be observed, that by increasing the
applied stress, both the recovered and irrecoverable strains also increased. At 100
MPa, the recovered strain of 2.56% increased to 4.15% when the applied stress was
increased to 300 MPa. By further increasing the applied stress to 500 MPa, the
recovered strain further increased to 4.9%. Similarly at 100 MPa, the irrecoverable
strain of 0.6% increased to 1.2% when the applied stress was increased to 300 MPa.
By further increasing the applied stress to 500 MPa, the irrecoverable strain further
increased to 1.7%.
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 7.6 Change in transformation temperatures of solution treated
5Cu alloy at stress levels of 100 – 500 MPa
129
Fig. 7.8 represents the recovery ratio and work output for solution treated 5Cu alloy,
calculated from the transformation strains and their corresponding stress levels. It
can be noted, that the recovery ratio decreased as the applied stress was increased
from 100 MPa to 500 MPa. The recovery ratio of 81% resulted at 100 MPa
decreased to 77% at 300 MPa and then further decreased to 74% when the applied
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 7.7 Recovered and irrecoverable strains of solution treated 5Cu
alloy at stress levels of 100 – 500 MPa
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 7.8 Recovery ratio and work output of solution treated 5Cu alloy
at stress levels of 100 – 500 MPa
130
stress was further increased to 500 MPa. As shown in Fig. 7.8, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 2.56 J/cm3 at 100 MPa increased to 12.45 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 24.5 J/cm3.
7.2.3 Shape memory properties of 10Cu alloy
Fig. 7.9 shows the strain-temperature curves for solution treated 10Cu alloy at stress
levels of 100 – 500 MPa.
Fig. 7.10 represents the variation of transformation temperatures (measured from Fig.
7.9) with respect to the applied stress levels for solution treated 10Cu alloy. It can be
observed from Fig. 7.10 that all the transformation temperatures increased as the
stress level increased from 100 MPa to 500 MPa. The Ms and Af temperatures;
198°C and 245°C at 100 MPa, increased to 207°C and 263°C respectively when the
applied stress was increased to 300 MPa. Similarly when the stress was increased to
500 MPa, the Ms and Af temperatures also increased to 234°C and 283°C
respectively. Moreover, a linear relationship between the transformation
temperatures and applied stress was observed for all stress levels, satisfied the
Clausius-Clapeyron (Cs-Cl) equation.
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 7.9 Strain-temperature curves representing the shape memory
properties of solution treated 10Cu alloy at stress levels of 100 – 500 MPa
131
The recovered and irrecoverable strains were calculated from Fig. 7.9 for the solution
treated 10Cu alloy and re-plotted in Fig. 7.11. It can be observed, that by increasing
the applied stress, both the recovered and irrecoverable strains also increased. At 100
MPa, the recovered strain of 3.95% increased to 5.2% when the applied stress was
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 7.10 Change in transformation temperatures of solution treated
10Cu alloy at stress levels of 100 – 500 MPa
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 7.11 Recovered and irrecoverable strains of solution treated 10Cu
alloy at stress levels of 100 – 500 MPa
132
increased to 300 MPa. By further increasing the applied stress to 500 MPa, the
recovered strain remained stable at 5.2%.
Similarly at 100 MPa, the irrecoverable strain of 0.1% increased to 0.6% when the
applied stress was increased to 300 MPa. By further increasing the applied stress to
500 MPa, the irrecoverable strain further increased to 1.6%.
Fig. 7.12 represents the recovery ratio and work output for solution treated 10Cu
alloy, calculated from the transformation strains and their corresponding stress
levels. It can be noted that the recovery ratio decreased as the applied stress was
increased from 100 MPa to 500 MPa. The recovery ratio of 97% resulted at 100 MPa
decreased to 90% at 300 MPa and then further decreased to 76% when the applied
stress was further increased to 500 MPa. As shown in Fig. 7.12, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 3.95 J/cm3 at 100 MPa increased to 15.6 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 26 J/cm3.
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600W
ork
Ou
tpu
t (J
/cm
3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 7.12 Recovery ratio and work output of solution treated 10Cu
alloy at stress levels of 100 – 500 MPa
133
7.2.4 Shape memory properties of 15Cu alloy
Fig. 7.13 shows the strain-temperature curves for solution treated 15Cu alloy at stress
levels of 100 – 500 MPa. Fig. 7.14 represents the variation of transformation
temperatures (measured from Fig. 7.13) with respect to the applied stress levels for
solution treated 15Cu alloy. It can be observed from Fig. 7.14 that all the
transformation temperatures increased as the stress level increased from 100 MPa to
500 MPa. The Ms and Af temperatures; 214°C and 261°C at 100 MPa, increased to
226°C and 277°C respectively when the applied stress was increased to 300 MPa.
Similarly when the stress was increased to 500 MPa, the Ms and Af temperatures also
increased to 244°C and 295°C respectively. Moreover, a linear relationship between
the transformation temperatures and applied stress was observed for all stress levels,
satisfied the Clausius-Clapeyron (Cs-Cl) equation.
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 7.13 Strain-temperature curves representing the shape memory
properties of solution treated 15Cu alloy at stress levels of 100 – 500 MPa
134
The recovered and irrecoverable strains were calculated from Fig. 7.13 for the
solution treated 15Cu alloy and re-plotted in Fig. 7.15. It can be observed, that by
increasing the applied stress, both the recovered and irrecoverable strains also
increased. At 100 MPa, the recovered strain of 3.65% increased to 4.8% when the
applied stress was increased to 300 MPa. By further increasing the applied stress to
500 MPa, the recovered strain further increased to 5.1%. Similarly at 100 MPa, the
irrecoverable strain of 0.2% increased to 0.62% when the applied stress was
increased to 300 MPa. By further increasing the applied stress to 500 MPa, the
irrecoverable strain further increased to 1.7%.
Fig. 7.16 represents the recovery ratio and work output for solution treated 15Cu
alloy, calculated from the transformation strains and their corresponding stress
levels. It can be noted, that the recovery ratio decreased as the applied stress was
increased from 100 MPa to 500 MPa. The recovery ratio of 95% resulted at 100 MPa
decreased to 89% at 300 MPa and then further decreased to 75% when the applied
stress was further increased to 500 MPa. As shown in Fig. 7.16, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 3.65 J/cm3 at 100 MPa increased to 14.4 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 25.5 J/cm3.
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 7.14 Change in transformation temperatures of solution treated
15Cu alloy at stress levels of 100 – 500 MPa
135
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 7.15 Recovered and irrecoverable strains of solution treated 15Cu
alloy at stress levels of 100 – 500 MPa
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 7.16 Recovery ratio and work output of solution treated 15Cu
alloy at stress levels of 100 – 500 MPa
136
7.3 Effect of Cu addition on transformation temperatures
Fig. 7.17 and 7.18 represent the martensite start and austenite finish temperatures of
solution treated 0Cu, 5Cu, 10Cu and 15Cu alloys under stress levels of 100 – 500
MPa, respectively. The Ms and Af temperatures of 0Cu alloy increased each by 25°C
upon addition of 5% Cu in place of Ni at stress level of 100 MPa. The Ms and Af
temperatures of 0Cu alloy further increased by 48°C and 45°C respectively when Ni
was replaced by 10% Cu. When Ni was replaced by 15% Cu, the Ms and Af
temperatures further increased by 64°C and 61°C respectively as compared to
baseline 0Cu alloy.
Similarly in the same manner, it can be observed from Fig. 7.17 that the martensite
start temperatures at different stress levels increased by increasing the Cu-content in
place of Ni. The increase in transformation temperatures can be attributed to the
decrease in Ni-content and Ni/Pd ratio, because Cu content does not affect the
transformation temperatures [146]. It has been reported that according to the
TiNiTiPd pseudo-binary phase diagram, the increase in Pd-content and decrease in
Ni-content, increases the transformation temperatures [147]. Thus the decrease in
Ni/Pd ratio (as Ni content decreased and Pd content remained constant) resulted in
increase in transformation temperatures.
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Te
mp
erat
ure
(°C
)
Stress (MPa)
Ms
0Cu
5Cu
10Cu
15Cu
Fig. 7.17 Comparison of martensite start temperatures of solution treated
0Cu, 5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500 MPa
137
7.4 Effect of Cu addition on transformation strains
Fig. 7.19 and 7.20 show the transformation strain (εrec) and irrecoverable strain (εirr)
respectively of solution treated 0Cu, 5Cu, 10Cu and 15Cu alloys under stress levels
of 100 – 500 MPa. It can be observed from Fig. 7.19 that the εrec of 0Cu and 5Cu
alloys were remained almost at the same level whereas it was increased by 1.34%
when Ni was replaced by 10% Cu at stress level of 100 MPa. However, the εrec of
15Cu alloy was slightly decreased as compared to 10Cu alloy. Similarly in the same
manner, the εrec of 0Cu and 5Cu alloys were remained at the same level and then it
increased for 10Cu alloy and then decreased for 15Cu alloy at all stress levels.
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Stress (MPa)
Af
0Cu
5Cu
10Cu
15Cu
Fig. 7.18 Comparison of austenite finish temperatures of solution treated
0Cu, 5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500 MPa
138
It can be observed from Fig. 7.20 that εirr of 5Cu and 10Cu alloys was decreased by
0.22% and 0.72% respectively as compared to 0Cu alloy at stress level of 100 MPa.
However, εirr of 15Cu was remained almost at the same level as of 10Cu alloy.
Similarly in the same manner, the εirr of 5Cu and 10Cu alloys was decreased as
compared to 0Cu alloy at all stress levels. Moreover it can also be observed from Fig.
2
3
4
5
6
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
εrec
0Cu
5Cu
10Cu
15Cu
Fig. 7.19 Comparison of recovered strain of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys under stress level of 100 – 500 MPa
0
1
2
3
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
εirr
0Cu
5Cu
10Cu
15Cu
Fig. 7.20 Comparison of irrecoverable strain of solution treated 0Cu,
5Cu, 10Cu and 15Cu alloys under stress level of 100 – 500 MPa
139
7.19 and 7.20 that the transformation strains increased by increasing the stress level
from 100 MPa to 500 MPa.
7.5 Effect of Cu addition on recovery ratio and work output
Fig. 7.21 and 7.22 show the recovery ratio and work output of solution treated 0Cu,
5Cu, 10Cu and 15Cu alloys under stress levels of 100 – 500 MPa. It can be observed
from Fig. 7.21 that the recovery ratio of 5Cu and 10Cu alloys was increased by 5%
and 21% as compared to 0Cu alloy at stress level of 100 MPa. However, the recovery
ratio of 15Cu was slightly decreased by 2% as compared to 10Cu alloy and the same
trend remained till higher stress levels. Moreover it can also be observed that for
each alloy the recovery ratio was decreased as the applied stress level was increased
from 100 MPa to 500 MPa.
Work output of 5Cu alloy was observed to be remained almost the same as of 0Cu
alloy, however it was increased by 1.34 J/cm3
for 10Cu alloy as compared to 0Cu
alloy. Again for 15Cu alloy the work output was remained almost the same as of
10Cu alloy. Similarly in the same manner, the work output of 0Cu and 5Cu alloys
was observed to be remained at the same level at higher stress levels, whereas it was
increased for 10Cu alloy. For 15Cu alloy the work output was again slightly
decreased as compared to 10Cu alloy. Moreover it can also be observed that for each
alloy the work output was increased as the applied stress level was increased from
100 MPa to 500 MPa.
140
0
10
20
30
0 100 200 300 400 500 600 700
Wo
rk O
utp
ut
(J/c
m3)
Stress (MPa)
0Cu
5Cu
10Cu
15Cu
Fig. 7.22 Comparison of work output of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys under stress level of 100 – 500 MPa
60
70
80
90
100
0 100 200 300 400 500 600 700
Re
cove
ry R
atio
(%
)
Stress (MPa)
0Cu
5Cu
10Cu
15Cu
Fig. 7.21 Comparison of recovery ratio of solution treated 0Cu, 5Cu,
10Cu and 15Cu alloys under stress level of 100 – 500 MPa
141
7.6 Summary
To study the effect of Cu addition on shape memory properties of TiNiPd high
temperature shape memory alloys, constant stress thermomechanical cycling
experiments in solution treated condition were conducted under various stress levels
of 100 – 500 MPa. By substitution of Cu in place of Ni, the shape memory
properties; transformation temperatures, recovered and irrecoverable strains,
recovery ratio and work output of the alloy were changed. The transformation
temperatures of baseline 0Cu alloy increased significantly when Ni was replaced by
5% Cu. Increasing the Cu content to 10% and 15%, the transformation temperatures
of the alloy were further increased and got maximum values at 15% Cu. The
martensite start temperatures of 5Cu, 10Cu and 15Cu alloys increased by 16%, 32%
and 38% respectively as compared to 0Cu alloy.
The recovered and irrecoverable strains of alloys were increased up to addition of
10% Cu and then slightly decreased when 15% Cu was added. The strain recovery of
5Cu, 10Cu and 15Cu alloys were observed to be increased by 2%, 8% and 6%
respectively as compared to 0Cu alloy. In the same manner, the irrecoverable strain
was decreased up to addition of 10% Cu and then slightly increased by addition of
15% Cu. It was observed that the irrecoverable strain was decreased by 23%, 27%
and 23% for 5Cu, 10Cu and 15Cu alloys respectively. Like recovered strain, the
recovery ratio was also increased by increasing the Cu content up to 10% and then
slightly decreased by addition of 15% Cu. Recovery ratio of 5Cu, 10Cu and 15Cu
alloys was increased by 6%, 27% and 25% respectively as compared to 0Cu alloy.
Similarly the work output of the alloy was increased up to addition of 10% Cu and
then slightly decreased when 15% Cu was added. The work output of 5Cu, 10Cu and
15Cu alloys was increased by 2%, 8% and 6% respectively as compared to the
baseline 0Cu alloy. From these observations it can be firmly concluded that addition
of 10% Cu in place of Ni in TiNiPd alloy resulted an improved shape memory
properties in terms of recovered and irrecoverable strains, recovery ratio and work
output.
142
Chapter – 8
Effect of Aging on Shape Memory
Properties of TiNiPd Alloys
8.1 Introduction
In this chapter, the shape memory properties; transformation temperatures under
biased load, recovered and irrecoverable strains, recovery ratio and work output were
investigated for of 0Cu, 5Cu, 10Cu and 15Cu alloys, aged at aging temperatures of
600°C for 3 hours. The detailed results and discussion are given in the following
sections.
8.2 Shape memory properties of 600°C-aged TiNiPd alloys with
varying Cu percentage
To investigate the effect of aging on shape memory properties, constant stress
thermal cycling experiments were performed. Later on the transformation
temperatures, recovered and irrecoverable strains, recovery ratio and work output of
600°C-aged 0Cu, 5Cu, 10Cu and 15Cu alloys were compared to that of solution
treated samples presented in Chapter – 7.
8.2.1 Shape memory properties of 600°C-aged 0Cu alloy
Fig. 8.1 shows the strain-temperature curves of 600°C-aged 0Cu alloy at stress levels
of 100 – 500 MPa. Thermal cycling experiments (loading, heating and cooling) at
various stress levels were carried out by the same manner as explained in section
3.6.2. Transformation temperatures (Ms, Mf, As, Af) at different stress levels were
measured by the tangent intersection method as shown in Fig. 3.7.
Fig. 8.2 represents the variation of transformation temperatures (measured from Fig.
8.1) with respect to the applied stress levels for 600°C-aged 0Cu alloy. It can be
observed from Fig. 8.2 that all the transformation temperatures increased as the stress
level increased from 100 MPa to 500 MPa. The Ms and Af temperatures; 150°C and
200°C at 100 MPa, increased to 165°C and 224°C respectively when the applied
143
stress was increased to 300 MPa. Similarly when the stress was increased to 500
MPa, the Ms and Af temperatures also increased to 180°C and 245°C respectively.
Moreover, a linear relationship between the transformation temperatures and applied
stress was observed for all stress levels, satisfied the Clausius-Clapeyron (Cs-Cl)
equation [4].
The recovered and irrecoverable strains were calculated by the measurement scheme
shown in Fig. 3.7. The calculated transformation strains (εrec, εirr) for the 600°C-aged
0Cu alloy were re-plotted in Fig. 8.3.
It can be observed, that by increasing the applied stress, both the recovered and
irrecoverable strains also increased. This increase in transformation strains can be
attributed to the increase in number of favored martensite variants in direction of
applied stress [155].
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 8.1 Strain-temperature curves representing the shape memory
properties of 600°C-aged 0Cu alloy at stress levels of 100 – 500 MPa
144
At 100 MPa, the recovered strain of 2.59% increased to 4.00% when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the recovered strain further increased to 4.7%. Similarly at 100 MPa, the
irrecoverable strain of 0.8% increased to 1.4% when the applied stress was increased
to 300 MPa. By further increasing the applied stress to 500 MPa, the irrecoverable
strain further increased to 2.05%.
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 8.2 Change in transformation temperatures of 600°C-aged
0Cu alloy at stress levels of 100 – 500 MPa
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 8.3 Recovered and irrecoverable strains of 600°C-aged 0Cu alloy
at stress levels of 100 – 500 MPa
145
Fig. 8.4 represents the recovery ratio and work output of 600°C-aged 0Cu alloy,
calculated from the transformation strains and their corresponding stress levels. It
can be noted that the recovery ratio decreased as the applied stress was increased
from 100 MPa to 500 MPa. The recovery ratio of 76% resulted at 100 MPa
decreased to 74% at 300 MPa and then further decreased to 70% when the applied
stress was further increased to 500 MPa. As shown in Fig. 8.4, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 2.59 J/cm3 at 100 MPa increased to 12.00 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 23.5 J/cm3.
8.2.2 Shape memory properties of 600°C-aged 5Cu alloy
The strain-temperature curves of 600°C-aged 5Cu alloy at stress levels of 100 – 500
MPa is shown in Fig. 8.5.
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 8.4 Recovery ratio and work output of 600°C-aged 0Cu alloy at
stress levels of 100 – 500 MPa
146
The variation of transformation temperatures (measured from Fig. 8.5) with respect
to the applied stress levels for 600°C-aged 5Cu alloy is shown in Fig. 8.6. It can be
observed from Fig. 8.6 that all the transformation temperatures increased as the stress
level increased from 100 MPa to 500 MPa. The Ms and Af temperatures; 170°C and
220°C at 100 MPa, increased to 184°C and 237°C respectively when the applied
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
100 MPa
200 MPa
300 MPa
400 MPa
500 MPa
Heating
Cooling 10
%
Fig. 8.5 Strain-temperature curves representing the shape memory
properties of 600°C-aged 5Cu alloy at stress levels of 100 – 500 MPa
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Te
mp
erat
ure
(oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 8.6 Change in transformation temperatures of 600°C-aged 5Cu
alloy at stress levels of 100 – 500 MPa
147
stress was increased to 300 MPa. Similarly when the stress was increased to 500
MPa, the Ms and Af temperatures also increased to 208°C and 265°C respectively.
Moreover, a linear relationship between the transformation temperatures and applied
stress was observed for all stress levels, satisfied the Clausius-Clapeyron (Cs-Cl)
equation.
The recovered and irrecoverable strains were calculated from Fig. 8.5 for the 600°C-
aged 5Cu alloy and then it were re-plotted in Fig. 8.7.
It can be observed, that by increasing the applied stress, both the recovered and
irrecoverable strains also increased. At 100 MPa, the recovered strain of 2.5%
increased to 4.1% when the applied stress was increased to 300 MPa. By further
increasing the applied stress to 500 MPa, the recovered strain further increased to
4.8%. Similarly at 100 MPa, the irrecoverable strain of 0.5% increased to 1.0% when
the applied stress was increased to 300 MPa. By further increasing the applied stress
to 500 MPa, the irrecoverable strain further increased to 1.5%.
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 8.7 Recovered and irrecoverable strains of 600°C-aged 5Cu alloy
at stress levels of 100 – 500 MPa
148
Fig. 8.8 represents the recovery ratio and work output of 600°C-aged 5Cu alloy,
calculated from the transformation strains at their corresponding stress levels. It can
be noted, that the recovery ratio decreased as the applied stress was increased from
100 MPa to 500 MPa. The recovery ratio of 83% resulted at 100 MPa decreased to
80% at 300 MPa and then further decreased to 76% when the applied stress was
further increased to 500 MPa. As shown in Fig. 8.8, the work output increased
continuously by increasing the applied stress from 100 MPa to 500 MPa. The work
output of 2.5 J/cm3 at 100 MPa increased to 12.3 J/cm
3 when the applied stress was
increased to 300 MPa. By further increasing the applied stress to 500 MPa, the work
output further increased to 24.0 J/cm3.
8.2.3 Shape memory properties of 600°C-aged 10Cu alloy
Fig. 8.9 shows the strain-temperature curves of 600°C-aged 10Cu alloy at stress
levels of 100 – 500 MPa.
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 8.8 Recovery ratio and work output of 600°C-aged 5Cu alloy at
stress levels of 100 – 500 MPa
149
Fig. 8.10 represents the variation of transformation temperatures (measured from Fig.
8.9) with respect to the applied stress levels for 600°C-aged 10Cu alloy. It can be
observed from Fig. 8.10 that all the transformation temperatures increased as the
stress level increased from 100 MPa to 500 MPa. The Ms and Af temperatures;
162°C and 225°C at 100 MPa, increased to 175°C and 230°C respectively when the
applied stress was increased to 300 MPa. Similarly when the stress was increased to
500 MPa, the Ms and Af temperatures also increased to 205°C and 233°C
respectively. Moreover, a linear relationship between the transformation
temperatures and applied stress was observed for all stress levels, satisfied the
Clausius-Clapeyron (Cs-Cl) equation.
The recovered and irrecoverable strains were calculated from Fig. 8.9 were re-plotted
in Fig. 8.11. It can be observed, that by increasing the applied stress, both the
recovered and irrecoverable strains also increased. At 100 MPa, the recovered strain
of 0.45% increased to 0.96% when the applied stress was increased to 300 MPa. By
further increasing the applied stress to 500 MPa, the recovered strain further
increased to 1.4%. Similarly at 100 MPa, the irrecoverable strain of 0% increased to
0.03% when the applied stress was increased to 300 MPa. By further increasing the
applied stress to 500 MPa, the irrecoverable strain further increased to 0.09%.
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
2.5
%
300 MPa
100 MPa
500 MPa
400 MPa
200 MPa
Heating
Cooling
Fig. 8.9 Strain-temperature curves representing the shape memory
properties of 600°C-aged 10Cu alloy at stress levels of 100 – 500 MPa
150
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 8.10 Change in transformation temperatures of 600°C-aged 10Cu
alloy at stress levels of 100 – 500 MPa
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 8.11 Recovered and irrecoverable strains of 600°C-aged 10Cu
alloy at stress levels of 100 – 500 MPa
151
Fig. 8.12 represents the recovery ratio and work output of 600°C-aged 10Cu alloy. It
can be noted that the recovery ratio decreased as the applied stress was increased
from 100 MPa to 500 MPa. The recovery ratio of 100% resulted at 100 MPa
decreased to 97% at 300 MPa and then further decreased to 94% when the applied
stress was further increased to 500 MPa. As shown in Fig. 8.12, the work output
increased continuously by increasing the applied stress from 100 MPa to 500 MPa.
The work output of 0.45 J/cm3 at 100 MPa increased to 2.88 J/cm
3 when the applied
stress was increased to 300 MPa. By further increasing the applied stress to 500
MPa, the work output further increased to 7.0 J/cm3.
8.2.4 Shape memory properties of 600°C-aged 15Cu alloy
Fig. 8.13 shows the strain-temperature curves of 600°C-aged 15Cu alloy at stress
levels of 100 – 500 MPa. Fig. 8.14 represents the variation of transformation
temperatures (measured from Fig. 8.13) with respect to the applied stress levels for
600°C-aged 15Cu alloy. It can be observed from Fig. 8.14 that all the transformation
temperatures increased as the stress level increased from 100 MPa to 500 MPa. The
Ms and Af temperatures; 178°C and 241°C at 100 MPa, increased to 187°C and
251°C respectively when the applied stress was increased to 300 MPa.
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 8.12 Recovery ratio and work output of 600°C-aged 10Cu alloy at
stress levels of 100 – 500 MPa
152
-50 50 150 250 350
Stra
in (
%)
Temperature (°C)
2.5
%
300 MPa
100 MPa
500 MPa
400 MPa
200 MPa
Heating
Cooling
Fig. 8.13 Strain-temperature curves representing the shape memory
properties of 600°C-aged 15Cu alloy at stress levels of 100 – 500 MPa
50
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Te
mp
erat
ure
(oC
)
Stress (MPa)
Ms
Mf
As
Af
Fig. 8.14 Change in transformation temperatures of 600°C-aged 15Cu
alloy at stress levels of 100 – 500 MPa
153
Similarly when the stress was increased to 500 MPa, the Ms and Af temperatures also
increased to 211°C and 258°C respectively. Moreover, a linear relationship between
the transformation temperatures and applied stress was observed for all stress levels,
satisfied the Clausius-Clapeyron (Cs-Cl) equation.
0.0
0.5
1.0
1.5
2.0
2.5
0
1
2
3
4
5
6
0 100 200 300 400 500 600
Irre
cove
rab
le S
trai
n (
%)
Re
cove
red
Str
ain
(%
)
Stress (MPa)
Ɛrec
Ɛirr
Fig. 8.15 Recovered and irrecoverable strains of 600°C-aged 15Cu
alloy at stress levels of 100 – 500 MPa
0
5
10
15
20
25
30
50
60
70
80
90
100
110
0 100 200 300 400 500 600
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Recovery Ratio
Work Output
Fig. 8.16 Recovery ratio and work output of 600°C-aged 15Cu alloy at
stress levels of 100 – 500 MPa
154
The transformation strains were calculated from Fig. 8.13 for the 600°C-aged 15Cu
alloy and their values were re-plotted in Fig. 8.15.
It can be observed, that by increasing the applied stress, both the recovered and
irrecoverable strains also increased. At 100 MPa, the recovered strain of 0.43%
increased to 0.91% when the applied stress was increased to 300 MPa. By further
increasing the applied stress to 500 MPa, the recovered strain further increased to
1.34%. Similarly at 100 MPa, the irrecoverable strain of 0% remained at 0% when
the applied stress was increased to 300 MPa. By further increasing the applied stress
to 500 MPa, the irrecoverable strain increased to only 0.07%.
Fig. 8.16 represents the recovery ratio and work output for 600°C-aged 15Cu alloy,
calculated from the transformation strains at their corresponding stress levels. It can
be noted that the recovery ratio decreased as the applied stress was increased from
100 MPa to 500 MPa. However the recovery ratio remained stable till the stress level
of 300 MPa. The recovery ratio of 100% resulted at 100 MPa remained at the same
level at 300 MPa and then decreased to 94% when the applied stress was increased to
500 MPa. As shown in Fig. 8.16, the work output increased continuously by
increasing the applied stress from 100 MPa to 500 MPa. The work output of 0.43
J/cm3 at 100 MPa increased to 2.73 J/cm
3 when the applied stress was increased to
300 MPa. By further increasing the applied stress to 500 MPa, the work output
further increased to 6.7 J/cm3.
8.3 Comparison of shape memory properties between solution
treated and 600°C-aged 0Cu alloys
8.3.1 Comparison of transformation temperatures
Fig. 8.17 shows the comparison of transformation temperatures of solution treated
and 600°C-aged 0Cu alloys under stress levels of 100 – 500 MPa. It can be observed
that the Ms and Af temperatures of solution treated and 600°C-aged samples
remained unchanged at lower stress levels. However at higher stress levels the Ms
and Af temperatures of 600°C-aged sample were slightly increased as compared to
solution treated condition. At 500 MPa, the Ms and Af temperatures of 600°C-aged
sample were increased by 3°C and 5°C respectively. This result showed that in case
155
of 0Cu alloy, there was no effect on transformation temperatures when the alloy aged
at 600°C for 3 hours. The slight increase in transformation temperatures was due to
the grain growth as discussed in section 5.3.1.
8.3.2 Comparison of transformation strains
The comparison of transformation strains between solution treated and 600°C-aged
0Cu alloys under stress levels of 100 – 500 MPa is shown in Fig. 8.18. At lower
stress levels, the transformation strains of solution treated and 600°C-aged samples
remained same whereas at higher stress levels the transformation strains of 600°C-
aged sample were slightly decreased as compared to solution treated condition. At
500 MPa, the εrec and εirr of 600°C-aged sample were decreased by 0.1% and 0.15%
respectively. This result indicated that there was no effect on transformation strains
by aging the 0Cu alloy at 600°C for 3 hours.
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Te
mp
erat
ure
(°C
)
Stress (MPa)
Soln Treated
Ms
Af
600 °C-aged Ms
Af
Fig. 8.17 Comparison of transformation temperatures of solution
treated and 600°C-aged 0Cu alloy under stress level of 100 – 500
MPa
156
8.3.3 Comparison of recovery ratio and work output
Fig. 8.19 represents the comparison of recovery ratio and work output between
solution treated and 600°C-aged 0Cu alloys under stress levels of 100 – 500 MPa.
Here again it can be observed that at lower stress levels, the recovery ratio and work
output of solution treated and 600°C-aged samples remained same whereas at higher
stress levels the work output of 600°C-aged sample was slightly decreased as
compared to solution treated condition. However the recovery ratio of 600°C-aged
sample was slightly improved. At 500 MPa, the work output was decreased by 0.02
J/cm3 and recovery ratio was increased by 2% when the alloy was aged at 600°C for
3 hours. This result suggested that there was no significant effect on recovery ratio
and work output by aging the 0Cu alloy at 600°C for 3 hours.
0
2
4
6
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
Soln Treated εrec
εirr
600 °C-aged εrec
εirr
Fig. 8.18 Comparison of transformation strains of solution treated and
600°C-aged 0Cu alloy under stress level of 100 – 500 MPa
157
8.4 Comparison of shape memory properties between solution
treated and 600°C-aged 5Cu alloys
8.4.1 Comparison of transformation temperatures
The comparison of transformation temperatures of solution treated and 600°C-aged
samples under stress levels of 100 – 500 MPa for 5Cu alloys is shown in Fig. 8.20.
By comparing their transformation temperatures, it can be observed that the Ms and
Af temperatures of 600°C-aged sample were slightly increased than that of solution
treated sample at lower stress levels. Upon further increasing the stress level to 500
MPa, the Ms and Af temperatures of 600°C-aged sample were further increased as
compared to solution treated condition. At 100 MPa, the Ms and Af temperatures of
600°C-aged sample were increased by 5°C each whereas at 500 MPa, the Ms and Af
temperatures of 600°C-aged sample were further increased by 13°C and 12°C
respectively. This result showed that by aging the 5Cu alloy at 600°C for 3 hours,
slight increase occurred in transformation temperatures at lower as well as at higher
stress levels. The increase in transformation temperatures of 600°C-aged 5Cu alloy
can be attributed to the grain growth as discussed in section 5.3.2.
0
10
20
30
60
70
80
90
100
110
0 100 200 300 400 500 600 700
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Soln Treated
600 °C-aged
Rec Ratio
Work Output
Rec Ratio
Work Output
Fig. 8.19 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 0Cu alloy under stress level of 100 – 500 MPa
158
8.4.2 Comparison of transformation strains
Fig. 8.21 shows the comparison of transformation strains between solution treated
and 600°C-aged 5Cu alloys under stress levels of 100 – 500 MPa. It can be observed
that at lower stress levels, the recovered and irrecoverable strains of solution treated
and 600°C-aged samples were remained same whereas at higher stress levels the
recovered strain of 600°C-aged sample were slightly decreased as compared to
solution treated condition. At 500 MPa, the εrec and εirr of 600°C-aged sample were
decreased by 0.1% and 0.2% respectively. This result indicated that there was no
significant effect on transformation strains when the 5Cu alloy was aged at 600°C for
3 hours.
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Te
mp
erat
ure
(°C
)
Stress (MPa)
Soln Treated
Ms
Af
600 °C-aged Ms
Af
Fig. 8.20 Comparison of transformation temperatures of solution treated
and 600°C-aged 5Cu alloy under stress level of 100 – 500 MPa
159
8.4.3 Comparison of recovery ratio and work output
Fig. 8.22 represents the comparison of recovery ratio and work output between
solution treated and 600°C-aged 5Cu alloys under stress levels of 100 – 500 MPa.
Work output of solution treated and 600°C-aged samples were remained same at
lower stress levels whereas at higher stress levels the work output of 600°C-aged
sample was slightly decreased as compared to solution treated condition. However
the recovery ratio of 600°C-aged sample was slightly improved at all stress levels. At
500 MPa, the work output of 600°C-aged sample was decreased by 0.06 J/cm3. This
result suggests that there was no significant effect on recovery ratio and work output
by aging the 5Cu alloy at 600°C for 3 hours.
0
2
4
6
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
Soln Treated εrec
εirr
600 °C-aged εrec
εirr
Fig. 8.21 Comparison of transformation strains of solution treated and
600°C-aged 5Cu alloy under stress level of 100 – 500 MPa
160
8.5 Comparison of shape memory properties between solution
treated and 600°C-aged 10Cu alloys
8.5.1 Comparison of transformation temperatures
Fig. 8.23 shows the comparison of transformation temperatures of solution treated
and 600°C-aged samples under stress levels of 100 – 500 MPa for 10Cu alloys. It can
be noted that the transformation temperatures of 600°C-aged sample were observed
to be significantly lower than that of solution treated sample at each stress levels. At
stress level of 100 MPa, the Ms and Af temperatures of 600°C-aged sample decreased
by 36°C and 20°C respectively as compared to solution treated condition and the
same trend remained till higher stress levels. This result suggested that due to aging
at 600°C for 3 hours, precipitates were formed, caused to decrease their
transformation temperatures. The same result can be supported from back-scattered
SEM images shown in Fig. 4.8, where both types of precipitates can be observed on
the grain boundaries as well as in the grain interiors.
0
10
20
30
60
70
80
90
100
110
0 100 200 300 400 500 600 700
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Soln Treated
600 °C-aged
Rec Ratio
Work Output
Rec Ratio
Work Output
Fig. 8.22 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 5Cu alloy under stress level of 100 – 500 MPa
161
8.5.2 Comparison of transformation strains
Fig. 8.24 shows the comparison of transformation strains between solution treated
and 600°C-aged 10Cu alloys under stress levels of 100 – 500 MPa. From Fig. 8.24 it
is evident that at all stress levels, the recoverable strain of 600°C-aged sample
remarkably decreased as compared to solution treated one. At 100 MPa, the εrec of
600°C-aged sample was decreased by 3.5% whereas at 500 MPa, the same value
further decreased to 3.8%. Similarly the irrecoverable strain of 600°C-aged sample
was decreased slightly at lower stress level whereas at higher stress levels the same
value further decreased. At 100 MPa, the εirr of 600°C-aged sample was decreased
by 0.1% whereas at 500 MPa, the same value was further decreased by 1.51% as
compared to that of solution treated sample. The remarkable change in
transformation strains confirms the formation of both types of precipitates when the
10Cu alloy was aged at 600°C for 3 hours.
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Stress (MPa)
Soln Treated
Ms
Af
600 °C-aged Ms
Af
Fig. 8.23 Comparison of transformation temperatures of solution treated
and 600°C-aged 10Cu alloy under stress level of 100 – 500 MPa
162
8.5.3 Comparison of recovery ratio and work output
Fig. 8.25 represents the comparison of recovery ratio and work output between
solution treated and 600°C-aged 10Cu alloys under stress levels of 100 – 500 MPa. It
can be observed that recovery ratio of 600°C-aged sample was found to be greater at
all stress levels. However at lower stress levels the difference in their recovery ratio
was smaller and then become greater when the applied stress level was increased to
500 MPa. At 100 MPa the recovery ratio of 600°C-aged sample was increased by 3%
whereas at 500 MPa the same value further increased by 18% as compared to
solution treated sample. Conversely the work output of 600°C-aged sample was
significantly decreased at all stress levels as compared to solution treated condition.
At 100 MPa, the work output of 600°C-aged sample was significantly decreased by
3.5 J/cm3 and then further decreased by 19 J/cm
3 at 500 MPa as compared to that of
solution treated sample.
0
2
4
6
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
Soln Treated εrec
εirr
600 °C-aged εrec
εirr
Fig. 8.24 Comparison of transformation strains of solution treated and
600°C-aged 10Cu alloy under stress level of 100 – 500 MPa
163
8.6 Comparison of shape memory properties between solution
treated and 600°C-aged 15Cu alloys
8.6.1 Comparison of transformation temperatures
The comparison of transformation temperatures of solution treated and 600°C-aged
samples under stress levels of 100 – 500 MPa for 15Cu alloys is shown in Fig. 8.26.
By comparing their transformation temperatures, it can be observed that the Ms and
Af temperatures were significantly decreased when the alloy was aged at 600°C for 3
hours as compared to solution treated sample at each stress levels. At stress level of
100 MPa, the Ms and Af temperatures of 600°C-aged sample decreased by 36°C and
20°C respectively as compared to solution treated condition. Similarly the decrease
in transformation temperatures of 15Cu alloy due to aging at 600°C for hours was
remained till higher stress levels. This result indicated that by aging the 15Cu alloy at
600°C for 3 hours, significant decrease in transformation temperatures was observed
at each stress levels. This result confirmed that due to aging at 600°C for 3 hours,
precipitates were formed, caused to lower their transformation temperatures. The
same result can be supported from back-scattered SEM images shown in Fig. 4.9,
where both types of precipitates can be observed on the grain boundaries as well as
in the grain interiors.
0
10
20
30
60
70
80
90
100
110
0 100 200 300 400 500 600 700
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Soln Treated
600 °C-aged
Rec Ratio
Work Output
Rec Ratio
Work Output
Fig. 8.25 Comparison of recovery ratio and work output of solution
treated and 600°C-aged 10Cu alloy under stress level of 100 – 500 MPa
164
8.6.2 Comparison of transformation strains
The comparison of transformation strains between solution treated and 600°C-aged
15Cu alloys under stress levels of 100 – 500 MPa is shown in Fig. 8.27. It can be
observed that at all stress levels, the recoverable strain of 600°C-aged sample
remarkably decreased as compared to solution treated one. At 100 MPa, the εrec of
600°C-aged sample was decreased by 3.21% whereas at 500 MPa, the same value
further decreased to 3.76%. Similarly the irrecoverable strain of 600°C-aged sample
was decreased slightly at lower stress level whereas at higher stress levels the same
value further decreased. At 100 MPa, the εirr of 600°C-aged sample was decreased
by 0.2% whereas at 500 MPa, the same value was further decreased by 1.63% as
compared to that of solution treated sample.
100
150
200
250
300
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Tem
per
atu
re (
°C)
Stress (MPa)
Soln Treated
Ms
Af
600 °C-aged Ms
Af
Fig. 8.26 Comparison of transformation temperatures of solution treated
and 600°C-aged 15Cu alloy under stress level of 100 – 500 MPa
165
8.6.3 Comparison of recovery ratio and work output
Fig. 8.28 represents the comparison of recovery ratio and work output between
solution treated and 600°C-aged 0Cu alloys under stress levels of 100 – 500 MPa. It
can be observed that recovery ratio of 600°C-aged sample was increased at all stress
levels. However at lower stress levels the difference in their recovery ratio was
smaller and then become greater when the applied stress level was increased to 500
MPa. At 100 MPa the recovery ratio of 600°C-aged sample was increased by 5%
whereas at 500 MPa the same value further increased by 20% as compared to
solution treated sample. Conversely the work output of 600°C-aged sample was
significantly decreased at all stress levels as compared to solution treated condition.
At 100 MPa, the work output of 600°C-aged sample was significantly decreased by
3.22 J/cm3 and then further decreased by 18.8 J/cm
3 at 500 MPa as compared to that
of solution treated sample.
0
2
4
6
0 100 200 300 400 500 600 700
Tran
sfo
rmat
ion
Str
ain
(%
)
Stress (MPa)
Soln Treated εrec
εirr
600 °C-aged εrec
εirr
Fig. 8.27 Comparison of transformation strains of solution treated and
600°C-aged 15Cu alloy under stress level of 100 – 500 MPa
166
0
10
20
30
60
70
80
90
100
110
0 100 200 300 400 500 600 700
Wo
rk O
utp
ut
(J/c
m3)
Re
cove
ry R
atio
(%
)
Stress (MPa)
Soln Treated
600 °C-aged
Rec Ratio
Work Output
Rec Ratio
Work Output
Fig. 8.28 Comparison of recovery ratio and work output of solution treated
and 600°C-aged 15Cu alloy under stress level of 100 – 500 MPa
167
8.7 Summary
To investigate the effect of aging at 600°C for 3 hours on transformation
temperatures, transformation strains, recovery ratio and work output of TiNiPdCu
alloys, thermal cycling experiments were performed. The shape memory properties
of 600°C-aged alloys were then compared with properties of solution treated alloys
given in Chapter – 7.
By aging 0Cu alloy at 600°C for 3 hours, no change was observed on shape memory
properties as compared to solution treated samples. In case of 5Cu alloy, the
transformation strains and work output were remained unchanged. However
transformation temperatures and recovery ratio of 600°C-aged 5Cu alloy were
slightly increased as compared to that of solution treated one. The martensite start
temperature and recovery ratio increased by 12% and 6% respectively. In view of
these results it can be concluded that aging of 5Cu alloy was beneficial to enhance its
shape memory properties.
Comparing the shape memory properties of 10Cu alloy, it was noted that by aging
the alloy at 600°C for 3 hours, their properties were significantly changed.
Transformation temperatures, transformation strains and work output of 600°C-aged
samples were significantly decreased whereas recovery ratio was slightly increased
as compared to solution treated samples. The recovered strain and work output of the
alloy decreased each by 73% whereas recovery ratio increased by 23%. For 15Cu
alloy the same trend was observed as found in case of 10Cu alloy. Hence for 15Cu
alloy the recovered strain and work output of the alloy decreased each by 74%
whereas recovery ratio increased by 27%. These observations confirmed that aging
of 10Cu and 15Cu alloys was not beneficial in terms of transformation temperatures,
recovered strains and work output, however it was beneficial in terms of
irrecoverable strains and recovery ratio.
168
Chapter – 9
Summary of Results and Discussion
9.1 Introduction
In this chapter, the results and discussion already presented in Chapters 4, 5, 6, 7 and
8 are summarized. In Chapter – 4, the effect of Cu addition and aging on
microstructure has been discussed. The effect of Cu addition and aging on stress free
transformation temperatures has been presented in Chapter – 5. Dependence of
mechanical properties on Cu addition and aging has been discussed in Chapter – 6.
The effect of Cu addition on shape memory properties in solution treated condition
has been shown in Chapter – 7. Lastly, in Chapter – 8, the effect of aging at 600°C
temperature on shape memory properties has been presented.
9.2 Effect of Cu addition on microstructure
To investigate the effect of 5%, 10% and 15% Cu addition on microstructure of
TiNiPd alloys, optical microscopy, scanning electron microscopy, X-ray
diffractometry and energy dispersive spectroscopy were carried out. It was observed
that all the four alloys; 0Cu, 5Cu, 10Cu and 15Cu were consisted of single phase
having typical twinned martensite structure, B19 (orthorhombic) with clearly visible
grain boundaries. However by increasing the Cu concentration in TiNiPd alloys, the
grain size of the resultant TiNiPdCu alloys increased up to 35%. Moreover,
consistent shift of B19 phase was observed towards lower 2Ɵ angle in XRD patterns.
Second phase precipitates were formed in all four alloys at its grain boundaries with
same size having circular or elliptical shapes. This observation proved that addition
of Cu had no effect on the precipitate size and shape, however its density decreased
as the Cu concentration increased in TiNiPd alloys. The chemical composition of
second phase precipitates clearly indicated that the precipitates formed at the grain
boundaries were Ti2Ni which were formed during solidification.
169
9.3 Effect of Cu addition on transformation temperatures
It was observed that by increasing the Cu concentration in TiNiPd alloys, the
transformation temperatures and transformation heats were increased significantly
whereas thermal hysteresis was decreased. By adding 5 ~ 15% Cu in TiNiPd alloys,
the martensite start temperature increased by 9 ~ 42% and thermal hysteresis
decreased by 6 ~ 35%. From this result it can be confirmed that by addition of Cu in
place of Ni in TiNiPd alloys, the transformation temperatures could be significantly
increased whereas its thermal hysteresis could be decreased.
During thermal cycling, the transformation temperatures of solution treated 0Cu and
5Cu alloys decreased by increasing the number of thermal cycles, however the drop
was faster in initial cycles and then reduced in the final cycles. In 0Cu and 5Cu
alloys the drop in martensite start temperature after 5 thermal cycles was 7.5°C and
5°C respectively. Similar behavior was also observed during thermal cycling of
solution treated 10Cu and 15Cu alloys; however the drop in transformation
temperatures was less as compared to 0Cu and 5Cu alloys. In 10Cu and 15Cu alloys
the drop in martensite start temperature was observed to be 3.7°C and 3°C
respectively. This behavior showed that 10Cu and 15Cu alloys were more
dimensionally stable as compared to 0Cu and 5Cu alloys.
9.4 Effect of Cu addition on mechanical and shape memory
properties
During microhardness testing, it was observed that the hardness increased by 7%
when Cu concentration was increased up to 15% in TiNiPd alloys. Similarly, the
yield stress and fracture stress were increased by 26% and 12% respectively, in
martensite phase whereas the same values increased by 28% and 8% in austenite
phase, however the fracture strain decreased. From these results it can be concluded
that addition of Cu in place of Ni in TiNiPd alloys was beneficial to improve its
hardness, yield and fracture stress.
By comparing the mechanical properties in martensite and austenite phases, it was
observed that yield and fracture stresses in austenite phase were significantly higher
than that of martensite phase by 52% ~ 57% and 14% ~ 18% respectively, in all four
170
alloys. By substitution of Cu in place of Ni, the shape memory properties; recovered
and irrecoverable strains, recovery ratio and work output of the alloy were
significantly changed. It was observed that the recovered strain of alloy was
increased up to addition of 10% Cu and then slightly decreased when 15% Cu was
added. The strain recovery of 5Cu, 10Cu and 15Cu alloys was increased by 2%, 8%
and 6% respectively, with respect to the baseline 0Cu alloy. Similarly the
irrecoverable strain was decreased up to addition of 10% Cu and then slightly
increased when 15% Cu was added. The irrecoverable strain of 5Cu, 10Cu and 15Cu
alloys was decreased by 23%, 27% and 23% respectively. The recovery ratio was
observed to be increased by increasing the Cu content up to 10% and then slightly
decreased by addition of 15% Cu. Recovery ratio of 5Cu, 10Cu and 15Cu alloys was
improved by 6%, 27% and 25% respectively as compared to 0Cu alloy. Similarly the
work output of the alloy was increased up to addition of 10% Cu and then slightly
decreased when 15% Cu was added. The work output of 5Cu, 10Cu and 15Cu alloys
was increased by 2%, 8% and 6% respectively. From these observations it can be
firmly concluded that addition of 10% Cu in place of Ni in TiNiPd alloy resulted an
improved shape memory properties in terms of recovered and irrecoverable strains,
recovery ratio and work output.
9.5 Effect of aging on microstructure and transformation
temperatures
Aging at 400°C, 500°C, 600°C and 700°C for 3 hours had no significant effect on
microstructure of 0Cu and 5Cu alloys. However when 10Cu and 15Cu alloys were
aged at above mentioned temperatures, remarkable change in microstructure was
observed. By aging 10Cu and 15Cu alloys, two types of (darker and brighter
contrast) precipitates Ti2Pd and TiPdCu were formed at the grain boundaries and
grain interiors. The density of the precipitates were remained the same, however, the
size in 15Cu alloy was bigger as compared to that of 10Cu alloy.
When 0Cu alloy was aged, the transformation temperatures slightly increased as the
aging temperature increased from 400°C to 700°C. The martensite start temperature
was increased by 1% at aging temperature of 600°C as compared to solution treated
condition. By aging the 5Cu alloy, the transformation temperatures were slightly
171
increased when the aging temperature was increased from 400°C to 600°C and then
slightly decreased at aging temperature of 700°C. The martensite start temperature
increased by 12% when it was aged at 600°C. Conversely, for 10Cu and 15Cu alloys,
the transformation temperatures were significantly decreased as the aging
temperature was increased from 400°C to 600°C and then increased comparable to
that of solution treated condition when the aging temperature was increased to
700°C. The martensite start temperatures of 10Cu and 15Cu alloys were decreased
by 18% and 19% respectively at aging temperature of 600°C. These observations
indicated that aging of 0Cu and 5Cu alloys at 600°C was beneficial to improve its
transformation temperatures whereas aging of 10Cu and 15Cu alloys remarkably
decreased their transformation temperatures.
9.6 Effect of aging on mechanical and shape memory properties
By aging the 0Cu and 5Cu alloys at 400°C, 500°C, 600°C and 700°C for 3 hours, the
hardness was slightly increased till aging temperature of 600°C and then slightly
decreased at 700°C aging temperature. The hardness of 0Cu and 5Cu alloys
increased by 7% and 10% respectively when the alloys were aged at 600°C as
compared to solution treated condition. The hardness of 10Cu and 15Cu alloys was
remarkably increased when its aging temperature was increased from 400°C to
600°C and then decreased remarkably almost equal to solution treated condition. The
hardness of 10Cu and 15Cu alloys increased by 60% and 83% respectively when the
alloys were aged at 600°C. This behavior showed that by aging of TiNiPdCu alloys
having Cu content up to 5%, their hardness increased slightly whereas alloys having
higher Cu content, their hardness increased remarkably.
Yield stress, fracture stress and fracture strain were investigated only for 600°C-aged
TiNiPdCu alloys. It was observed that yield stress was slightly increased by aging
0Cu, 5Cu, 10Cu and 15Cu alloys at 600°C, however fracture strain was decreased as
compared to solution treated condition. The yield stress of all four alloys was
increased by 3 ~ 5% by aging them at 600°C as compared to solution condition.
When 0Cu alloy was aged at 600°C for 3 hours, no change in shape memory
properties was observed between solution treated and 600°C-aged samples. Similarly
172
for 5Cu alloy, the shape memory properties like transformation strains and work
output were remained unchanged. However transformation temperatures and
recovery ratio of 600°C-aged 5Cu alloy were slightly increased as compared to that
of solution treated one. The martensite start temperature and recovery ratio increased
by 12% and 6% respectively. From this discussion it can be concluded that aging of
5Cu alloy was beneficial to enhance its shape memory properties. Comparing the
shape memory properties of 10Cu alloy, it was noted that by aging the alloy at 600°C
for 3 hours, their properties were significantly changed. Transformation
temperatures, transformation strains and work output of 600°C-aged samples were
significantly decreased whereas recovery ratio was slightly increased as compared to
solution treated samples. The recovered strain and work output of the alloy decreased
by 73% whereas recovery ratio increased by 23%. For 15Cu alloy the same trend was
observed as found in case of 10Cu alloy. Hence for 15Cu alloy the recovered strain
and work output of the alloy decreased by 74% whereas recovery ratio increased by
27%. These observations confirmed that aging of 10Cu and 15Cu alloys was not
beneficial in terms of transformation temperatures, recovered strains and work
output, however it was beneficial in terms of irrecoverable strains and recovery ratio.
173
Chapter – 10
Conclusions and Recommendations for
Future Work
10.1 Summary of experimentation
The aim of this research was to improve the transformation temperatures, mechanical
and shape memory properties of TiNiPd high temperature shape memory alloys by
varying percentage of Cu and Ni in the alloy. The effect of aging (at different
temperatures of 400°C, 500°C, 600°C and 700°C for 3 hours) on the microstructure,
mechanical and shape memory properties of the alloys were also investigated. To
investigate the effect of Cu addition and aging on transformation temperatures,
mechanical and shape memory properties, a comprehensive characterization and
thermomechanical cycling experiments were conducted on solution treated and aged
alloys. Experimentation is summarized as below:
Four types of NiTi-based high temperature shape memory alloys i.e.
Ti50Ni25Pd25, Ti50Ni20Pd25Cu5, Ti50Ni15Pd25Cu10 and Ti50Ni10Pd25Cu15 (all in
atomic percent) were melted in vacuum arc melting furnace in desired
composition.
The required heat treatment i.e. homogenization at 950°C for 2 hours was
carried out in sealed quartz tubes.
Button of each alloy (20 g) was sliced in 0.4 mm thick strips using wire
EDM.
Cold rolling of 0.4 mm thick strips was done by 25% and its thickness was
reduced to 0.3 mm.
Solution treatment at 900°C for 1 hour of all samples and aging at 400°C,
500°C, 600°C and 700°C for 3 hours of selected samples were done in sealed
quartz tubes and then water quenched without breaking the quartz tubes.
Size, structure of grain and second phase precipitates was studied using
optical microscope and SEM. EDS was used for compositional analysis.
XRD analysis was carried out to investigate the presence of different phases.
174
Microhardness Tester and Mechanical Testing Machine were used to
investigate the hardness and yield strength.
Tensile Rupture and Creep Testing Machine was used to evaluate the shape
memory properties.
10.2 Conclusions
Four NiTi-based; Ti50Ni25-xPd25Cux (x=0, 5, 10 and 15) high temperature shape
memory alloys were developed. From the microstructural, mechanical
characterization and thermomechanical cycling experiments, the following
conclusions were drawn:
10.2.1 Effect of Cu addition
In solution treated condition, all four alloys consisted of single phase having
typical martensite B19 (orthorhombic) structure with clearly visible grain
boundaries. However by increasing the Cu content till 15%, the grain size
also increased up to 35%. Second phase precipitates of the same size formed
in all alloys, however the density of precipitates decreased with increasing
percentage of Cu. Ti2Ni type precipitates were observed to be formed along
the grain boundaries. Comparing the XRD patterns of all four alloys, it was
observed that the martensite phases consistently shifted towards lower 2Ɵ
angle with increasing content of Cu.
By increasing the Cu content up to 15%, the transformation temperatures
increased by 42%, whereas thermal hysteresis decreased by 35%. During
thermal cycling, the drop in transformation temperatures of 10Cu and 15Cu
alloys were less as compared to 0Cu and 5Cu alloys, hence increasing the Cu
content, the dimensional stability increased.
It was observed that addition of 15% Cu in the alloy, the hardness and yield
strength increased by 7% and 26% respectively in martensite phase, whereas
in austenite phase the yield strength increased by 28%. By comparing, it was
observed that the yield strength in austenite phase of all alloys was higher
than that of martensite phase by 52% ~ 57%.
The recovered strain, recovery ratio and work output increased by 2%, 6%
and 2% respectively when 5% Cu was added in TiNiPd alloy. For 10Cu alloy
175
the same values increased by 8%, 27% and 8% and similarly for 15Cu alloy
increased by 6%, 25% and 6% as compared to 0Cu alloy.
10.2.2 Effect of aging
By aging 0Cu and 5Cu alloys at 400°C, 500°C, 600°C and 700°C for 3 hours,
there was no significant effect on their microstructures in terms of martensite
phase, second phase precipitates and grain size. However, two types of
precipitates; Ti2Pd and TiPdCu (darker and brighter contrast respectively)
produced after aging of 10Cu and 15Cu alloys at mentioned aging
temperatures.
For 0Cu and 5Cu alloys, the martensite start temperature increased by 1% and
12% respectively, however in 10Cu and 15Cu alloys, the same temperature
decreased by 18% and 19% respectively by aging at 600°C. These results
indicated that aging of 0Cu and 5Cu alloys was beneficial to increase their
transformation temperatures, however the same temperatures were
significantly reduced in case of 10Cu and 15Cu alloys.
Hardness of the four alloys increased with increasing Cu content; however,
the increase was more pronounced in case of 10Cu and 15Cu alloys. For
10Cu and 15Cu alloys the hardness increased by 60% and 82% respectively
at aging temperature of 600°C. Similarly the yield strength of all alloys was
also increased by 3 ~ 5% by aging them at 600°C aging temperature.
The shape memory properties remained unchanged after aging 0Cu alloy at
600°C aging temperature. The recovered strain and work output remained
stable whereas the martensite start temperature and recovery ratio increased
by 12% and 6% respectively after aging the 5Cu alloy. By aging the 10Cu
alloy at 600°C, recovered strain and work output decreased by 73% whereas
recovery ratio increased by 23%. Similarly for 15Cu alloy, the same values
decreased by 74% and recovery ratio increased by 27% when it was aged at
600°C.
At the end, it can be concluded that addition of 5%, 10% and 15% Cu in place of Ni
in TiNiPd high temperature shape memory alloys was very useful to improve their
transformation temperatures, dimensional stability, mechanical and shape memory
176
properties. Moreover, it can also be concluded that aging of 0Cu and 5Cu alloy was
beneficial to improve their transformation temperatures and shape memory
properties. However, aging of 10Cu and 15Cu alloys produced an adverse effect on
their properties in terms of shape memory and transformation temperatures.
10.3 Recommendations for future work
Based upon these results, certain recommendations for future endeavors are detailed
below:
This research can be extended by varying the composition of Ni, Pd and Cu
in TiNiPdCu high temperature shape memory alloys.
In the recent research, mechanical properties like stress-strain curves have
been obtained till fracture of samples. If the samples are strained by pre-
defined limit before fracture and thermomechanical cycling (loading and
unloading) are performed for many cycles, it will be helpful to determine the
stress hysteresis and dimensional stability under biased load.
In the present research work, aging at different aging temperatures for
constant time duration has been studied. However it will be of significant
importance if aging at constant aging temperature for variable time duration
is studied.
To investigating the stresses produced due to addition of Cu in TiNiPd alloys
Application of TiNiPdCu alloys in actuators
177
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