CITY UNIVERSITY OF HONG KONG

DEPARTMENT OF

PHYSICS AND MATERIALS SCIENCE

BACHELOR OF ENGINEERING (HONS) IN MATERIALS ENGINEERING

2007-2008

DISSERTATION

Ni-free Shape Memory Alloys (SMA) for biomedical application

by

Yeung Che Yan

March 2008

Ni-free Shape Memory Alloys (SMA) for biomedical application

By

Yeung Che Yan

Submitted in partial fulfilment of the

requirements for the degree of

BACHELOR OF ENGINEERING (HONS)

IN

MATERIALS ENGINEERING

from

City University of Hong Kong

March 2008

Project Supervisor : Dr. Jonathan C. Y. Chung

Acknowledgement

Firstly, I would like to express my sincere gratitude to my project supervisor, Dr. Jonathan

C.Y. Chung for his support and invaluable comments throughout the entire project. Through

his patient guidance, I have learnt much on how to make efficient project plans and better

experiment designs.

Secondly, I appreciate research students Mr. J.L. Zhang, Mr. Y.P. Zhang, Miss P.M. Chan,

Mr. Z.G. Lu and Mr. H. Li who help and encourage me in my experimental works.

Finally, I would like to thank all laboratory technicians and staffs in the Department. Without

their assistance, it would be extremely tough to carry out the works.

i

Table of Contents

Acknowledgement i

Table of Contents ii

List of Figures v

List of Tables vii

Abstract viii

Page

1) Introduction 1

2) Literature Review 2

2.1 Shape Memory Alloys (SMAs) 2

2.2 Crystal Structure of SMAs 2

2.3 Shape Memory Properties 3

2.3.1 Pseudoelasticity (PE) or Superelasticity (SE) 3

2.3.2 Shape Memory Effect (SME) 4

2.4 Nickel Titanium (NiTi) SMAs 4

2.4.1 Shape Memory Properties of NiTi 4

2.4.2 Biomedical Application 5

2.5 Nickel-free SMAs 5

2.5.1 Nickel-hypersensitivity 5

2.5.2 Surface Modification of NiTi and Development of Ni-free SMAs 5

2.5.3 Substitutes to Nickel in Titanium-based SMAs 5

2.6 Titanium Niobium (TiNb) SMAs 6

2.6.1 Development of TiNb Binary and TiNbZr Ternary SMAs 6

2.6.2 SME and PE of TiNbZr SMAs 7

2.6.3 Martensitic Transformation Temperature of TiNb-based Ternary SMAs 8

2.7 Thermal Treatment to Obtain Desired Shape Memory Properties 8

2.7.1 Effect of Thermal Treatment Parameters on Af of NiTi SMAs 8

2.7.2 Effect of Precipitation Hardening on PE of TiNb SMAs 9

3) Objectives 10

ii

4) Methodology 11

4.1 TiNbZr SMAs Fabrication 11

4.1.1 Etching and Batching 11

4.1.2 Arc Melting 11

4.2 Mechanical Treatment (Cold Rolling) 11

4.3 Thermal Treatment (Solution Treatment and Heat Treatment) 12

4.4 Material Characterization 12

4.4.1 Composition and Structural Analysis 13

(1) X-ray Fluorescence Test (XRF) 13

(2) X-Ray Diffractometry (XRD) 13

4.4.2 Mechanical Properties Study 14

(1) Tensile Test 14

(2) Vickers Hardness Measurement 14

4.4.3 Shape Memory Properties Study 14

(1) Strain Increment 5-cycle Tensile and Compression Tests 14

(2) Differential Scanning Calorimetry (DSC) 15

(3) Thermomechanical Analysis (TMA) 15

(4) Dynamic Mechanical Analysis (DMA) 15

5) Results and Discussions 16

5.1 Composition Analysis 16

5.2 Crystal Structure Analysis 17

5.2.1 Phase at Room Temperature and 400℃ 17

5.2.2 Changes in Grain Orientation 18

5.2.3 Grain Size Estimated from Full Width Half Maximum (FWHM) 19

5.3 Determination of Martensitic Transformation Temperatures 20

5.3.1 Differential Scanning Calorimetry (DSC) 21

5.3.2 Thermomechanical Analysis (TMA) 22

5.3.3 Dynamic Mechanical Analysis (DMA) 22

5.4 Effect of Thermo-mechanical Treatments on Vickers Hardness 25

5.5 Tensile Stress-strain Behaviour in Ti-22Nb-(2-10)-Zr (at.) Alloys 27

iii

5.6 Elastic and Plastic Deformation Behaviour in Strain Increment 5-cycle Tensile 28

and Compression Tests

5.6.1 Elastic and Plastic Deformation Behaviour 28

5.6.2 Recoverable Strain 30

5.6.3 Young’s Modulus in the Cyclic Tensile and Cyclic Compression Tests 30

5.7 Further Discussions and Recommendations 32

6) Conclusions 34

7) References 36

Appendix I – Experimental Details

Appendix II – XRD Results

Appendix III – Cyclic Tensile and Cyclic Compression Results

iv

List of Figures

Page

Fig. 2.3.1 Schematic diagram showing stress-induced martensitic transformation 4

Fig. 2.6.2 Schematic diagram showing the relationship of temperature with the critical 8

stress to induce martensite(σSIM) and the critical stress for slip deformation

(σs)

Fig. 4.4.2 Geometry of the tensile specimen 14

Fig. 5.1a XRF spectrum indicating the presence of Ti, Nb and Zr in the 93% CR 16

Specimen underwent ST, and HT at 550℃ for 30min with subsequent oil

quench

Fig. 5.1b Line scan XRF spectrum along the longitudinal direction of the 93% CR 16

specimen underwent ST

Fig. 5.2.1 XRD spectra at room temperature (left) and 400℃ (right) of the 93% CR 18

specimen underwent ST, and HT at 550℃ 120min with subsequent oil quench

Fig. 5.2.2 Peak intensities of the (110), (200), (211) and (222) planes of specimens with 19

different mechanical treatments (left) and thermal treatments (right)

Fig. 5.3.1 DSC curve of the 93% CR specimen underwent ST 21

Fig. 5.3.2 TMA curve of the as-cast rod-shaped specimen 22

Fig. 5.3.3a DMA results of the 93% CR specimen underwent ST, and HT at 550℃ for 23

30min with subsequent oil quench. 100 Hz, 20 Hz and 1 Hz were applied

Fig. 5.3.3b DMA results at 20 Hz of the 50% CR specimen 24

Fig. 5.4a Vicker’s hardness of specimens with different thermo-mechanical treatments. 25

Columns 1 to 3: increasing degree of CR; Column 4: 93% CR specimen with

ST; Columns 5 to 10: 93% CR specimens with ST, and HT at conditions stated

in the graph

Fig. 5.4b Percentage change in Vickers hardness against the degree of mechanical work 27

Fig. 5.4c Schematic diagram showing the dependence of hardness on ageing time at 27

different ageing temperatures

Fig. 5.5a Tensile stress-strain curve of Tensile stress-strain curve of the 93% CR 28

specimen underwent ST

Fig. 5.5b Mechanical properties variation of the Ti-22Nb-xZr (at.%) system 28

v

Fig. 5.6.1 Stress-strain curves for cyclic tensile test of the 93% CR specimen 29

underwent ST, with imposed tensile strains 2.5%, 3%, 3.5%, 4%, 5%, 6.5%,

8% and 10%

Fig. 5.6.2b Recoverable strain in cyclic tensile test of 93% CR specimens underwent ST, 30

and underwent ST plus HT at 300℃ for 30min with subsequent oil quench

Fig. 5.6.2c Recoverable strain in cyclic compression test of specimen without any heat 30

treatment, and specimen underwent HT at 300℃ for 30min with subsequent

oil quench

Fig. 5.6.3a Young’s modulus of cycles 1 to 5 in cyclic tensile test of 93% CR specimen 31

underwent ST

Fig. 5.6.3b Schematic diagram showing the determination of E in a non-linear stress- 32

strain curve found in cycles at high imposed strain

Fig. 5.6.3c Young’s modulus of cycles 1 to 5 in cyclic compression test of specimen 32

without any heat treatment

vi

List of Tables

Page

Table 4.1.1 Etching conditions to remove surface oxide for the pure elements of TiNbZr 11

Table 4.4 Thermo-mechanical treatment conditions for the samples characterized by 12

different techniques

Table 5.2.1 Position of all detected peaks and the corresponding hkl planes 17

vii

Abstract Button-shaped ingots of Ti-22Nb-10Zr (at.%) were prepared by arc melting method in an

argon atmosphere. Cylindrical rod specimens, used for cyclic compression test and

thermomechanical analysis (TMA), were prepared by applying vacuum suction in the copper

mould casting.

The ingots were cold rolled for mechanical treatment with 93% thickness reduction. Solution

treatment was applied on all cold rolled thin plates. Heat treatment at 300℃, 450℃ and 550

℃ for 30min and 120min with subsequent oil quench were carried out. Effects of thermo-

mechanical treatments on the crystal structure, transformation temperatures, hardness and

elastic deformation behaviour were investigated in this project.

Austenite phase was found in the X-ray diffraction (XRD) of all specimens. Changes in grain

orientation after thermo-mechanical treatments were observed. This was attributed to the

oriented nucleation and oriented growth of the deformed microstructure in the

recrystallization process.

Martensitic transformation temperatures were investigated by differential scanning

calorimetry (DSC), thermomechanical analysis (TMA) and dynamic mechanical analysis

(DMA). Both DSC and TMA failed to reveal the transition temperatures. The peaks at 20℃

and 45℃ in the heating curve of DMA were related to the reverse martensitic transformation.

The Vickers hardness was increased by mechanical treatment. Thermal treatment also played

a role in the hardening, although it was not as effective as mechanical treatment. Increasing

the heat treatment temperature from 300℃ to 450℃ and 550℃ caused softening. However,

increasing the heat treatment time from 30min to 120min showed little effect on the hardness.

Tensile test was carried out on rectangular specimen with 20mm gage length. The Young’s

modulus determined was 18GPa, which is much lower than previous reports of similar alloys.

The critical stress for slip deformation also showed a dramatic increase compared with others’

published work. This was attributed to the difference in grain size and rolling and

recrystallization texture.

Elastic and plastic deformation behaviour in strain increment cyclic loading was investigated.

The first cycles of all specimens at imposed strains greater than 2.5% deformed plastically.

viii

ix

On the contrary, the second to fifth cycles superimposed themselves and deformed elastically.

A maximum recoverable strain of 5% was achieved in the cyclic tensile test of the specimen

with 93% CR and underwent solution treatment.

1 Introduction Materials for biomedical application should possess good biocompatibility, low cytotoxicity,

high corrosion resistance, favourable for cell attachment and cell growth. The mechanical

properties such as strength, wear resistance, fatigue life and Young’s modulus are considered

important in biomedical application whereby Young’s modulus close to that of human hard

tissues, say bones, are preferred.

In the development of biomedical SMAs, the Ni-free ones are preferred as Ni is suspected to

cause hypersensitive allergic reaction. It is a must for biomedical SMAs to exhibit SM

properties such as SME and/or PE properties (also known as superelastic SE properties) at

human body temperature. Stable SM properties over many cycles, large recoverable strain and

high stress for slip deformation are required. There are three major factors affecting the SME

transformation characteristics and SE behaviour. They are composition, thermal treatments

and mechanical treatments.

Ti-22Nb-10Zr (at.%) was chosen in this study because the TiNb binary alloy system is one of

the most promising Ni-free SMAs. The addition of Zr, which is a β-stabilizer, can help to

stabilize the β phase and to augment the SM properties. This 10% Zr system was expected to

demonstrate different mechanical properties and SM properties compared with the 2% to 8%

added systems [1].

CR was the most popular mechanical treatment for SMAs [1-5]. Solution treatment (ST) was

carried out at 900℃ for 30min within the β field. This temperature was chosen following the

work of previous study on similar composition [1]. Comparison between the results obtained

in this project and those in the previous study is another reason for choosing 900℃ 30min to

be the ST conditions.

In Kim’s study [2], the heat treatment at 300℃ was effective to increase the critical stress for

slip deformation in the Ti-(25-27)Nb (at.%) alloys. Therefore, heat treatment at 300℃, 450℃

and 550℃ for 30min and 120min after ST was selected in this project. The effect of different

heat treatment temperatures and time was investigated. Fast cooling was adopted in the heat

treatments as a higher cooling rate would give rise to a smaller Young’s modulus [6].

1

2 Literature Review

2.1 Shape Memory Alloys (SMAs) Shape memory alloys (SMAs) are metals that exhibit two unique properties, namely shape

memory effect (SME) and pseudoelasticity (PE, also known as superelasticity SE). Today, the

most frequently used types are the NiTi, CuZnAl and CuAlNi, all possess good SM

properties.

NiTi SMAs are widely used in the biomedical field such as blood vessel stents, orthodontic

braces and surgical equipments. The copper-based ones are often used in industrial

application due to their cheap prices, but they cannot replace the NiTi in biomedical

application as they lack biocompatibility [7-9].

2.2 Crystal Structure of SMAs Both the SME and PE of SMAs originate from the reversible martensitic transformation ─ a

solid-to-solid phase transformation between the austenite and martensite phases. The stronger

phase austenite exists at higher temperature and the softer martensite which is easier to

deform exists at lower temperature. Martensite starts to form when a SMA is cooled from the

high temperature state to the martensitic start temperature Ms. When the temperature is further

reduced, more and more martensite crystals are formed. Eventually the transformation

completes and all austenite crystals are transformed to martensite at Mf. When the

temperature is raised, reverse martensitic transformation occurs where austenite crystals start

to form at As and the transformation completes at Af. The four transformation temperatures

are characteristics of a particular SMA, which depend on the composition, thermal history and

mechanical treatments.

The shape change of the SMA during martensitic transformation cannot be observed by naked

eyes unless deformation is applied. Macroscopically, the crystal lattice of the SMA has been

twinned from the original cubic structure austenite [8, Figure 2]. The crystal is softened and

therefore shape change is easy to be attained when deformation is applied [7-8].

Martensitic transformation is defined as the lattice transformation involving shear

deformation and resulting from cooperative atomic movement. It can be achieved by

supercooling to below the martensitic start temperature Ms or loading beyond the critical

stress leading to SME and PE respectively. Non-thermoelastic martensitic transformation

proceeds by creating new martensite crystals. On the other hand, the martensite crystals grow

2

at a speed proportional to the cooling rate once they are nucleated in thermoelastic

transformation. After they attain the thermal equilibrium, any change in temperature or

applied stress would lead to grow or shrink of the martensite crystals.

In most cases, SMAs have superlattice structure with body-centered cubic (BCC) parent

phases, with a few exceptions possessing face-centered cubic (FCC) parent phases.

Regardless of the alloy type, SMAs having BCC parent phases are denoted by β phase alloys.

They are categorized into β2 and β1 by their composition ratio and lattice structure. β2 refers to

the lattices order like CsCl with 50:50 elemental ratio, while the lattices order like Fe3Al with

75:50 composition ratio are denoted by β1. The martensitic phases obtained from the β2 are

labelled as β’2, α’2 or γ’2. Similar notation also applies to β1 alloys.

The cubic structure of a β2 unit cell can be viewed as the result of alternately stacking the

(110) planes [7, p. 15, Figure 1.8]. Supercooling (in SME) or applying certain amount of

stress (in PE) leads martensitic transformation to occur by shearing the crystal in the [110]

direction along the (100) planes. There are three types of close packed planes [7, p. 16, Figure

1.9] and the complete martensitic structure is constructed by regularly stacking these close

packed planes [7-8].

2.3 Shape Memory Properties 2.3.1 Pseudoelasticity (PE) or Superelasticity (SE)

Pseudoelasticity (PE), also named as superelasticity (SE), is demonstrated by applying a load

at temperature above Af. This can be attained because transformation temperatures increase

with applied loading (Fig. 2.3.1). The SMA transforms from the austenite phase to the

martensite phase because the martensite phase is thermodynamically more stable at the

stressed state. The applied mechanical energy is absorbed by the soft martensite. Upon

unloading, the stored energy is released and the martensite transforms back to austenite since

austenite has lower energy at temperature above Af. When we talk about PE in SMAs, both

the martensitic transformation and the reverse transformation are driven by stress and no

thermal energy is involved in the shape recovery [7-8]. The critical stress required to induce

martensite is represented by the dotted line in the Fig. 2.3.1. Complete strain recovery by SE

can only be achieved if the critical stress for slip deformation (plastic deformation) is higher

than it [7].

3

Fig. 2.3.1 Schematic diagram showing stress-induced martensitic transformation

2.3.2 Shape Memory Effect (SME)

On the other hand, when deformation is performed at temperature below Af, residual strain

can be observed after unloading. If the applied stress has not exceeded the critical stress for

slip deformation (due to dislocation or twinning), strain recovery can be achieved by heating

the SMA above Af. During the heating process, the deformed martensite is driven to undergo

reverse transformation to become austenite, which is configured of the original shape. The

phenomenon to obtain shape recovery by heating is known as SME [7, 8].

2.4 Nickel Titanium (NiTi) SMAs

2.4.1 Shape Memory Properties of NiTi

The parent phase of nickel titanium (NiTi or nitinol) SMA has a BCC B2 structure while the

martensite phase has a monoclinic B19’ crystal structure. In an intermediate temperature

range, a rhombohedral R phase may exist [7, 10-11].

Near-equiatomic NiTi SMAs have been studied for many years. They demonstrate excellent

SME and PE. In a typical Ni-49Ti (at.%) SMA, the transformation temperatures are Mf = -

150℃, Ms = -114℃, As = -89℃ and Af = -40℃. Deformation conducted in the range -74℃ to

-49℃ demonstrates perfect SME, the shape is almost completely retained after unloading and

perfect shape recovery is demonstrated when it is heated above Af. PE is demonstrated in the

range Af ± 15℃ in which a characteristic plateau is observed in the stress-strain curve. Upon

unloading, there is no residual strain. When deformation is made above Af + 15℃, imperfect

PE is found where plastic strain remains after uploading [7]. In application we manipulate the

transformation temperatures by varying the composition and thermo-mechanical treatments,

such that the desire SM behaviour is demonstrated at specific temperature range. For example,

4

a deformed vena cava filter recovers its original shape at human body temperature

demonstrating SME.

2.4.2 Biomedical Application

NiTi alloys are the most widely used SMAs in biomedical application. Implants and medical

instruments are made in vitae of PE and SME. These include arch wires in dental application,

guide wires in endoscopic instrument, vena cava filters, plates for fixing fractured bones and

spinal correction instrumentations in scoliosis surgery [9-10].

2.5 Nickel-free SMAs

2.5.1 Nickel-hypersensitivity

The intermetallic compound NiTi is the best known SM system in biomedical applications

due to its high corrosion resistance and large recovery strain and stress. However, the safety

of use for implantation was aroused by medical doctors and dentists.

Nickel is an allergen and consequently may induce a delayed hypersensitivity reaction (type

IV immune response). Up to present, there is no solid evidence that patients are at a

significant risk of developing hypersensitivity solely due to contact with Ni-containing

implants. But there are a few cases that prior sensitization from non-medical contact lead to

Ni-sensitivity [12-13].

2.5.2 Surface Modification of NiTi and Development of Ni-free SMAs

For the above reason, researchers have worked to minimize the release of Ni ions recently.

One effective way to impede the leaching out of Ni ions is to produce a barrier layer on the

NiTi SMA by surface modification. Sol-gel technique, plasma immersion ion implantation

(PIII), chemical vapor deposition (CVD) and biomimetic deposition are some surface

modification technologies that researchers are now working on [14]. On the other hand, some

researchers are interested in developing Ni-free Ti-based SMAs, which are absolutely safe

from the exposure of the toxic Ni ions [1-6, 15-19].

2.5.3 Substitutes to Nickel in Titanium-based SMAs

Other Ti-based SMAs are being investigated to substitute the NiTi SMAs. Ti alloys are

promising biomaterials due to their reduced elastic modulus (so as to ease the stress shielding

5

effect), high strength, light weight, superior biocompatibility and good corrosion resistance

compared with conventional implant materials stainless steel and CoCr alloys [3, 6].

β type Ti-based alloys exhibit two stable phases, the austenite β phase (B2, disordered BCC)

at higher temperature and the α phase (HCP) at lower temperature. It also exhibits three

metastable phases α’, α’’ and ω. Transformation from the β phase at high cooling rate gives

rise to hexagonal martensite α’ or orthorhombic martensite α’’. The reverse transformation

from the orthorhombic martensite α’’ phase to the β phase is important in this study, as this is

the origin of shape recovery. The transformation temperature can be controlled by adjusting

the alloying composition, and optimizing the thermal and mechanical treatments. The

metastable phase formed when quenching from high temperature is known as athermal ω,

while that formed by heat treatment at intermediate temperature is known as thermal ω [1-2, 4,

6, 15]. The formation of these two ω phases is unavoidable during heat treatment processes

and they are sometimes utilized to increase the critical stress for slip deformation. This will be

discussed in detail in the next section.

Within the many types of Ni-free SMAs, TiV-based alloys, TiMo-based alloys, and TiNb-

based alloys were found to demonstrate good SME and PE. TiV-based SMAs are not suitable

for biomedical application because of vanadium’s cytotoxicity [1-2]. In binary alloys, Nb and

Mo act as β-stabilizers to stabilize the high temperature phase of Ti-based SMAs. It was

found that mechanical properties could be improved by the addition of a third element, such

as Sn, Al, Ta, Pd and Zr [15].

2.6 Titanium Niobium (TiNb) SMAs

2.6.1 Development of TiNb Binary and TiNbZr Ternary SMAs

Baker first reported that Ti-35Nb (wt.%) exhibits SME. Miyazaki et al. puts great effort on

the development of TiNb binary SMAs. They reported the SME in Ti-(22-25)Nb (at.%) and

SE in Ti-(25.5-27)Nb (at.%) at room temperature. The maximum SE strain recovered in

solution treated Ti-26Nb (at.%) alloys is about 3%, much smaller than that in NiTi SMAs.

The small SE strain is attributed to the low critical stress for slip deformation σs. Thus, the

strategy to enhance SE is to increase σs. Adding a third element as β-stabilizer and

introducing ω precipitates was suggested to make improvement [15].

6

The aforementioned group also investigated Ti-based ternary SMAs. Biocompatibility is the

first consideration for implant materials. Zr is chosen to be the third element because of its

low cytotoxicity. It has been proved that the addition of Zr to the TiNb SMAs can improve the

mechanical properties and it works together with Nb as β-stabilizers to stabilize the SME and

SE. Because of the solution hardening effect, σs increases with Zr content. Thus, the

recoverable strain increases with Zr content. A maximum recoverable strain of 4.3% is

obtained in the Ti-22Nb-6Zr (at.%) SMA [1], which is 1.3% above that achieved in TiNb

binary alloy [2]. This means that adding Zr is an effective method to increase σs and

ultimately increases the recoverable strain [1, 15].

2.6.2 SME and PE of TiNbZr SMAs

Kim et al. reported that Ti-22Nb-(2-4)Zr (at.%) and Ti-22Nb-6Zr (at.%) alloys exhibit stable

SME and PE respectively at room temperature [1]. To investigate the stability of SME and

PE, cyclic tensile tests were performed on these two SMAs, in which the specimens were

loaded to 2.5% strain in each cycle.

Incomplete SME is exhibited after heating in the first cycle while nearly complete SME was

achieved in the second cycle onwards [1, p. 338, Figure 12]. This shows that SME of Ti-

22Nb-4Zr (at.%) becomes stable at room temperature after the first cycle. The stress for slip

deformation marked by the white-headed arrow drops from 140MPa in the first cycle to

90MPa in the second cycle and then remains almost unchanged. On the other hand, the stress

required to attain 2.5% strain increases with cycles, with the exception of the first cycle.

Similar findings on stability of SE is also presented in Ti-22Nb-6Zr (at.%). Incomplete SE is

exhibited in the first cycle while almost complete SE is observed after the fourth cycle [1, p.

338, Figure 12].

In the investigation of SE at different temperatures on the Ti-22Nb-6Zr (at.%) alloys, it was

found that the residual strain increases with testing temperature [1]. As shown schematically

in Fig. 2.6.2, the critical stress to induce martensite (σSIM) increases with temperature while

the critical stress for slip deformation (σs) decreases with temperature [1-2]. When

deformation is made at temperature TA (indicates by the arrows), slip deformation occurs prior

to martensitic transformation as σs σSIM. This explains the increase of residual strain – the

result of slip deformation – at higher temperature.

⟨

7

Fig. 2.6.2 Schematic diagram showing the relationship of temperature with the critical stress to induce martensite (σSIM) and the critical stress for slip deformation (σs) 2.6.3 Martensitic Transformation Temperature of TiNb-based Ternary SMAs

It has been established that Ms decreases with the amount of alloying β-stabilizer. The study

on Ti-22Nb-xZr (at.%) showed that every 1 at.% addition of this Zr lowers the transformation

temperature by 35℃. The reduction of Ms also follows a linearly relationship with the amount

of Ta or O added to the Ti-22Nb alloys [15].

The relationship between Ms and amount of alloying element is important for developing a

biomedical SMA as it has to exhibit the desired SM properties at human body temperature.

Perfect SE is expected to be found in the range Af ± 15℃. As a consequence, it would be

desirable to have the human body temperature fell into this temperature range.

2.7 Thermal Treatment to Obtain Desired Shape Memory Properties SM properties in SMAs are often exhibited after optimal thermo-mechanical treatments. In

fact, changing the processing conditions can “tune” the transformation temperature by 50℃

without altering the composition [10]. As such, many researchers investigate the behaviour of

SMAs under different thermo-mechanical treatment conditions. Thermal treatment strategy

such as thermal cycling and series of heat treatment are often adopted to manipulate the

transformation temperatures. Besides, the processing-property interplay on microstructures

and SE properties are also investigated [4, 6, 10-11, 20-21].

2.7.1 Effect of Thermal Treatment Parameters on Af of NiTi SMAs

There are many reports that the transformation temperatures of NiTi are affected by thermal

treatment. For example, Paula et al. reported that during thermal cycling, Ms and Mf

temperatures increase, As remains constant and Af decreases. The transformation temperature

8

hysteresis is also reduced by fast cooling after annealing [11]. Liu et al. reported the increase

of austenitic phase transformation temperature by ageing at 325 - 375℃. This is caused by the

internal elastic stress fields inside the parent phase, which is created during the ageing

process. Yeung et al. also reported that proper ageing can increase the austenitic phase

transformation temperature from a few degree Celsius to human body temperature in

biomedical application [10].

2.7.2 Effect of Precipitation Hardening on PE of TiNb SMAs

In the development of TiNbZr SMAs, the main obstacle is the insufficient SE strain resulted

from low σs. There are two approaches to enhance σs. One is precipitation hardening by

ageing via the formation of fine ω and α precipitates. The other approach is work hardening

with or without subsequent annealing.

Two metastable ω phases exist in Ti-based alloys: athermal ω and thermal ω. The former one

is formed during quenching in the process of solution treatment or annealing, while the later

one is formed during ageing treatment.

The formation of fine and dense thermal ω precipitates has been proved to enhance SE in

binary TiNb SMAs. As the ω phase is Ti-rich, forming the precipitates would increase the Nb

content in the matrix. Thus, the martensitic transformation temperature Ms is reduced.

Moreover, the dispersed ω particles can suppress martensitic transformation mechanically. As

a result, σSIM is increased by forming thermal ω phase. On the contrary, it has been proved

experimentally that σs increases with increasing ageing time. In other words, the formation of

thermal ω phase increases σs. Combining the two effects – increase in both σSIM and σs – the

maximum stress for perfect SE deformation can be enhanced by forming thermal ω phase.

The thermal ω phase does not solely improve SE, it also improves the tensile strength, yield

stress and hardness. However, it is noted that ductility is severely sacrificed at the same time.

In the particular case of Ti-26Nb (at.%), ageing at 300℃ for 1 hour effectively improves the

superelastic behaviour and perfect SE is exhibited until the fourth cycle. On the contrary,

prolong ageing at 300℃ for 10 hours has led to embrittlement. Fracture occurs after three

cycles of loading [2-3].

9

3 Objectives NiTi SMAs are one of the most promising biomaterials due to their light weight, good

biocompatibility and high corrosion resistance. Moreover, the SME and PE properties give

NiTi SMAs extra functionalities in their biomedical application compared with the traditional

metallic biomaterials, such as stainless steel and cobalt-chromium alloys. Unfortunately, NiTi

SMAs may cause Ni-hypersensitive allergic reaction when implanted in human body.

Researchers have put a lot of efforts to reduce the Ni-release by forming a surface passive

layer. On the other hand, replacing the NiTi with Ni-free Ti-based SMAs solves the problem

more thoroughly as there is no Ni at all.

There are reports on the microstructure, transformation behaviour and mechanical properties

on the TiNb binary SMAs. But the previous studies on TiNbZr mainly focused on the

mechanical properties at different temperatures and the transformation temperatures were

derived from the stress-strain relationship. The main obstacle encountered so far is the low

recoverable strain and therefore imperfect SME and PE behaviour.

In this project, the main objectives were to find out the effect of thermo-mechanical

treatments on both the crystal structure and mechanical properties. It is thought that

introducing thermo-mechanical treatments would increase the recoverable strain and therefore

the SME and PE behaviour. On top of deriving transformation temperatures from the stress-

strain relationship, various characterization techniques were applied in this project, aiming to

give additional information on the transformation behaviour. The ultimate goal of the project

was to obtain a Ni-free SMA with good SM properties.

10

4 Methodology

4.1 TiNbZr SMAs Fabrication The Ti-22Nb-10Zr (at.%) ternary SMAs were prepared by arc melting method from the pure

elements (Ti: 99.8%, Element Titanium Metal Manufacture Co., Ltd., Nb and Zr: 99.8%,

Oryx Sputtering Targets (Shenzhen) Co., Ltd.). The abbreviation TiNbZr is used hereafter as

this is the only chemical composition in the entire project.

4.1.1 Etching and Batching

To remove the surface oxide, the pure elements in the form of chips were acid etched

according to Table 4.1.1. They were then rinsed by distilled water and ethanol. Batches of 8 g

were weighed according to the alloying composition immediately before arc melting.

Element Etching medium Etching time Ti HNO3:H2O, in the ratio of 13:7 15 min

Nb HF:HNO3:H2O, in the ratio of 1:4:25 3 min

Zr HF:HNO3:H2O, in the ratio of 1:4:25 1 min

Table 4.1.1 Etching conditions to remove surface oxide for the pure elements of TiNbZr

4.1.2 Arc Melting

TiNbZr ingots with 16.5mm diameter and 10mm thickness were fabricated from the pure

elements by the arc melting method in a water-cooling copper mould. Vacuum was reached in

two stages, 5 x 10-2 mbar by rotary pump and 9 x 10-5 mbar by turbo pump. 950 mbar argon

was pumped before melting to provide an inert atmosphere. Due to the large difference in

melting points (Ti: 1668℃, Nb: 2477℃, Zr: 1855℃) among the pure elements, the ingots

were inverted and re-melted for two more times to promote chemical homogeneity.

Cylindrical rod specimens, used for cyclic compression test and TMA, were obtained by

copper mould casting. 1.7g TiNbZr alloy was put in the copper mould and vacuum suction

was applied when the alloy was melted.

4.2 Mechanical Treatment (Cold Rolling)

Thickness reduction of 60%, 70% and 93% was achieved by rolling at room temperature (W.

Durston Rolling Mill). To obtain an elongated thin plate from the button-shaped ingot, rolling

was performed along one direction.

11

4.3 Thermal Treatment (Solution Treatment and Heat Treatment) The specimens were first cut to the desired shape before solution treatment (ST) to minimize

the effect of internal stress on the transformation temperatures. At the first stage, all

specimens underwent ST at 900℃ for 30min in the argon-filled tube furnace (Thermolyne),

followed by quenching in oil. After ST, heat treatment (HT) were conducted in the same

furnace at 300℃, 450℃ and 550℃ for 30min and 120min. Oil quench were applied.

4.4 Material Characterization Different material characterization techniques were performed to study the crystal structure,

general mechanical behaviour and SM properties of TiNbZr. Table 4.4 below shows the

thermo-mechanical treatment conditions for the samples characterized by different techniques.

Characterization

techniques

Mechanical treatment Thermal treatment CR

(Percentage of CR) ST

HT

(Temp, Time, Cooling method)

93% Yes No

93% Yes 300℃, 120min, OQ XRF

93% Yes 550℃, 120min, OQ

No, 60%, 70%, 93% No No

93% Yes No

93% Yes 300℃, 30min, OQ

93% Yes 300℃, 120min, OQ XRD

93% Yes 450℃, 30min, OQ

93% Yes 450℃, 120min, OQ

93% Yes 550℃, 30min, OQ

93% Yes 550℃, 120min, OQ

In-situ XRD at 400℃ 93% Yes 550℃, 120min, OQ

93% No No

93% Yes No DSC

93% Yes 450℃, 30min, OQ

93% Yes 450℃, 120min, OQ

TMA Cylindrical specimen No No

93% No No

93% Yes No DMA

93% Yes 300℃, 30min, OQ

93% Yes 300℃, 120min, OQ

12

93% Yes 450℃, 30min, OQ

93% Yes 450℃, 120min, OQ DMA

93% Yes 550℃, 30min, OQ

93% Yes 550℃, 120min, OQ

No, 70%, 93% No No

93% Yes No

93% Yes 300℃, 30min, OQ

93% Yes 300℃, 120min, OQ Vickers hardness

93% Yes 450℃, 30min, OQ

93% Yes 450℃, 120min, OQ

93% Yes 550℃, 30min, OQ

93% Yes 550℃, 120min, OQ

Tensile test 93% Yes No

93% Yes No Cyclic tensile test

93% Yes 300℃, 30min, OQ

Cylindrical specimen No No Cyclic compression test

No 300℃, 30min, OQ Cylindrical specimen

Table 4.4 Thermo-mechanical treatment conditions for the samples characterized by different techniques

(Remarks: ST refers to solution treatment at 900℃ for 30min with subsequent oil quench. OQ refers to oil quench)

4.4.1 Composition and Structural Analysis

(1) X-ray Fluorescence Test (XRF)

The chemical composition of specimens before and after heat treatment were checked by the

XRF (Micro EDXRF Eagle III X-ray Fluorescence Machine). Two scanning modes were

adopted: (1) average composition of 32 points evenly distributed on the specimen; (2)

composition distribution along a straight line on the specimen (line scan).

(2) X-Ray Diffractometry (XRD)

A θ-2θ XRD analysis (Philips XPert MRD X-ray Diffractometer) was conducted to determine

the phase distribution at room temperature using Cu Kα radiation. Scanning range of 20˚ to

150˚ was selected at a scanning rate 0.3˚/s. The average grain size was calculated by the full

width half maximum (FWHM) using the Scherrer equation. In-situ XRD at 400℃ (Siemens

D500 X-ray Diffractometer) was used to investigate the phase existing at high temperature.

13

4.4.2 Mechanical Properties Study

(1) Tensile Test

The tensile test was carried out in a MTS machine (Instron 5567) at room temperature to

determine the Young’s modulus, stress for slip deformation, ultimate tensile strength and

fracture strain.

Fig. 4.4.2 shows the geometry of the tensile specimen cut by electro-discharge machine

(EDM) from a cold rolled thin plate. The gage size was 20mm length, 6mm width, and 0.7mm

thick. The specimens were machined such that the loading axis aligned with the rolling

direction. The tensile loading rate was set at 0.25 mm/min.

Fig. 4.4.2 Geometry of the tensile specimen

(2) Vickers Hardness Measurement

Vickers hardness test was performed by a 1 kg load (corresponding to indentation of about

one-tenth of sample thickness). The dwell time was set at 10s. The surface layer formed due

to high temperature oxidation in thermal treatments was grinded before conducting the

hardness test.

4.4.3 Shape Memory Properties Study

(1) Strain Increment 5-cycle Tensile and Compression Tests

The elastic and plastic deformation behavior of the TiNbZr SMAs was evaluated by the strain

increment 5-cycle tensile and compression tests. Two loading modes tensile and compression

were adopted in the MTS machine stated in Section 3.4.2. In both modes, loading and

unloading were performed on the same piece of specimen for 5 cycles at imposed strains

2.5%, 3%, 3.5%, 4%, 5%, 6.5%, 8% and finally 10%. The crosshead speeds for the cyclic

tensile and cyclic compression tests were 0.25 mm/min and 0.08 mm/min respectively.

The specimens for cyclic tensile test had the same geometry as the previous tensile specimen

(Fig. 4.4.2). The test was conducted on a specimen underwent ST and a specimen underwent

ST, and HT at 300℃ for 30min with subsequent oil quench.

14

The cyclic compression specimens were cylindrical rods with 6.9mm height and 4.7mm

diameter. The test was applied on specimens before and after HT at 300℃ for 30min with

subsequent oil quench.

(2) Differential Scanning Calorimetry (DSC)

The transformation temperatures were examined by a DSC (MDSC 2910 TA Instruments).

The test was performed by heating from -150℃ to 350℃ at 5 ℃/min and then cooling from

350℃ to 80℃ at the same rate. The weight of specimens was between 10mg and 15mg.

(3) Thermomechanical Analysis (TMA)

Transformation temperatures of a cylindrical specimen 4.4mm height and 2.9mm diameter

was investigated by a TMA (Setaram TMA). The scanning range was 30℃ to 350℃, the

weight of the measuring probe was balanced to give a consistent measuring force of 2g, and

heating and cooling were performed at scanning rate 2 ℃/min.

(4) Dynamic Mechanical Analysis (DMA)

The DMA damping test was performed with a DMA (DMA 2980 TA Instruments) in the

single cantilever mode. Rectangular specimens of 30mm length, 4mm width and 0.7mm

thickness were pre-heated to 350℃ in the DMA. The test was performed by cooling to -110℃

and then heating to 550℃ at a rate of 5 ℃/min. The frequency and amplitude were set at 20

Hz and 10μm respectively.

15

5 Results and Discussions

5.1 Composition Analysis Kα and Kβ lines of Ti, Nb and Zr and small amount of Lα line of Nb were detected in the XRF.

The signal of 32 points evenly distributed on the surface of a specimen gave an average

composition and the average composition of three specimens was Ti-22.97Nb-10.17Zr (at.%).

The chemical composition of specimens before and after HT showed little differences and one

of them is plotted in the below figure.

XRF spectrum

Energy (keV)

Inte

nsit

y (a

.u.)

Nb Kα

Ti Kα

Zr Kα

Zr Kβ Nb Kβ Nb Lα Ti Kβ

Fig. 5.1a XRF spectrum indicating the presence of Ti, Nb and Zr in the 93% CR specimen

underwent ST, and HT at 550℃ for 30min with subsequent oil quench

The composition distribution along a straight line on the specimen was investigated by using

the line scanning mode. Similar composition was observed in both longitudinal and transverse

directions of the 93% CR specimen underwent ST.

Line scan along longitudinal direction

0

20

40

60

0 2 4 6 8 10 12 14Distance (mm)

Wei

ght

(%) Ti

Nb

Zr

Fig. 5.1b Line scan XRF spectrum along the longitudinal direction of the 93% CR specimen underwent ST

16

5.2 Crystal Structure Analysis 5.2.1 Phase at Room Temperature and 400℃

From the previous studies on Ti-22Nb-(2-8)Zr (at.%) systems, peaks for both HCP and BCC

structure are found in the 2% Zr system and peaks for solely BCC structure are found at 4% to

8% Zr systems [1]. Addition of Zr content by 1 at.% lowers the Ms by 35℃ and Ti-22Nb-4Zr

(at.%) demonstrates Ms at near room temperature [15].

With higher content of beta-stabilizers, the present system, Ti-22Nb-10Zr (at.%), showed

agreement with the previous studies. Only peaks corresponding to the BCC structure (beta)

were found in all specimens and no peaks corresponding to HCP structure (alpha) were found.

This implied the existence of high temperature phase austenite at room temperature.

The XRD spectra of specimens underwent different thermo-mechanical treatments were

shown in Appendix II. The peak positions for all specimens and the corresponding hkl planes

were extracted and listed in Table 5.2.1. Within the nine peaks listed here, some existed only

in specimens with particular thermo-mechanical treatments. The peak at 2 Theta 65.266˚

could not be matched with any existing JCPDS database (1997 version). This peak was found

in specimens with heavy deformation (minimum 80% CR), underwent both ST and HT [22].

It is thought to be due to a new phase formed after such treatment.

d-value α1 (Å) hkl Crystal structure 2 Theta (degree) 38.481 2.3375 110 BCC 55.541 1.6532 200 BCC 65.266 1.4284 Unknown Unknown 69.605 1.3496 211 BCC 82.444 1.1689 220 BCC 94.924 1.0454 310 BCC 107.624 0.9544 222 BCC 121.303 0.8837 321 BCC 137.455 0.8266 400 BCC

Table 5.2.1 Position of all detected peaks and the corresponding hkl planes

In-situ XRD spectrum at 400℃ was obtained by using a hot stage. The spectrum was similar

to that obtained at room temperature. Therefore, we conclude that the austenite phase existed

at both room temperature and 400℃. It could be deduced that all transformation temperatures

Ms, Mf, As and Af lied below room temperature. Previous studies showed that martensitic

17

transformation in Ti-22Nb-4Zr (at.%) started at about room temperature and additional Zr

content lowers the Ms [1, 15]. The martensite phase of the present system, Ti-22Nb-10Zr

(at.%), ought to exist below room temperature. The above results concur to the expected trend.

In this study, the intensity percentage of a particular peak was calculated by dividing its signal

counts by the summation of signal counts of all peaks in that specimen. It was observed that

the intensity of the (222) plane, being parallel to the rolling plane, lowered when the specimen

was characterized at 400℃. The intensity of the (211) peak increased simultaneously. This

was attributed to the thermal effect on the specimen leading to oriented grain growth. Further

explanation is provided in the following section.

Spectrum at room temperature

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Spectrum at 400℃

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

211

211 222100 100 222

Fig. 5.2.1 XRD spectra at room temperature (left) and 400℃ (right) of the 93% CR specimen underwent ST, and HT at 550℃ 120min with subsequent oil quench

5.2.2 Changes in Grain Orientation

The change in intensities of the eight recognized peaks corresponding to the BCC structure

(stated in Table 5.2.1) will be discussed in this section. The intensities of the peaks at 2 Theta

65.266˚ in all specimens were low (spectra shown in Appendix II). As such, changes would be

difficult to recognize. So this peak was not used in the analysis.

The as-cast specimen possessed eight BCC peaks. Within these eight peaks, the (220), (310),

(321) and (400) planes were minorities in this crystal structure and they nearly disappeared

after ST. On the other hand, the peaks for the (110), (200), (211) and (222) planes showed

different intensities under different thermo-mechanical treatments. This could be explained by

the change in grain orientation due to formation of deformation and recrystallization texture.

The intensity of the (110) planes was substantially lowered when the sample thickness

reduced from 3mm (70% CR) to 0.7mm (93% CR). At the same time, the intensity of the (200)

18

planes increased by nearly the same amount. This could be thought that part of the (110)

planes, the slip plane for the BCC structure, re-arranged themselves into (222) planes when

plastically deformed during CR.

Effect of mechanical treatments

0

20

40

60

80

0% 20% 40% 60% 80% 100%

Thickness Reduction

Inte

nsit

y (%

)

110

200

211

222

Effect of thermal treatments

0

20

40

60

80

Nil ST 300℃

120min

450℃

120min

550℃

120min

Inte

nsit

y (%

)

Fig. 5.2.2 Peak intensities of the (110), (200), (211) and (222) planes of specimens with different mechanical treatments (left) and thermal treatments (right)

The graph titled “Effect of thermal treatments” in Fig. 5.2.2 compares the effect of different

thermal treatments on the peak intensities of the 93% CR specimens. The peak intensities

before thermal treatment (indicated by Nil), after ST, and ST plus HT at 300℃, 450℃ and

550℃ for 120 minutes were shown. The drastic increase of (211) and (222) intensities after

ST could be explained by the creation of pre-dominant sites for crystal growth during the

mechanical deformation. Dislocations and defects were introduced when a specimen was cold

rolled. These were the pre-dominant sites for recrystallization which took place prior to ST at

lower temperature. The rearrangement (re-orientation) of crystal planes, dislocations and

defects of the as-rolled specimen determined the location of nucleation and direction of

crystal growth. Therefore, the (211) and (222) intensities were enhanced at the expense of

(200) planes.

5.2.3 Grain Size Estimated from Full Width Half Maximum (FWHM)

Direct measurement of micro-sized grains can be made in optical microscope (OM) and

scanning electron microscope (SEM) after polishing and etching. However, nano-sized grains

may not be revealed by these two techniques because of limited resolution and sample

preparation techniques. In this project, the grain size could not be determined using OM [22].

In order to clarify whether the grain size was in the nanometer range, the grain size was

further investigated by using the full width half maximum (FWHW) in the XRD diffraction

peaks.

19

Grain boundary can be considered a kind of imperfect crystal structure in the 3D-ordered

crystals leading to diffraction peak broadening. The Scherrer equation can be used to estimate

the grain size utilizing this phenomenon. The mean crystallite thickness (i.e. the grain size) is

determined by θ

λcosBKD = , where D is the average grain size, K is the shape factor taken as

1 (assuming spherical grains), λ is the wavelength, B is the FWHW in radian and θ is the

position of the peak [23].

The grain size was evaluated from the (110) and (211) planes and it was found that they were

in nanometer range between 4nm and 30nm. This explained why we could not obtain the

image of grains in OM.

The Scherrer equation can only be served as a preliminary evaluation to calculate the grain

size. Further investigation using transmission electron microscopy (TEM) should be carried

out for verification. Firstly, the power of the XRD machine may not be sufficient and this

would lead to broadening of peak. Secondly, the grains were assumed to be spherical in the

calculation (shape factor K taken as 1). However, the actual geometry and the orientation of

the grains were still unknown. The grains may be elongated in shape and the grain orientation

may not be random. In this way, K should be taken smaller than 1. Thirdly, the diffraction

peaks obtained in this study showed fluctuation so that inaccuracy existed in FWHM

determination. Combining the three factors, high degree of inaccuracy is expected in the grain

size determination. Nevertheless, there is a high possibility that the specimens consisted of

nano-sized grains.

5.3 Determination of Martensitic Transformation Temperatures The appearance of SME and PE which distinguish SMAs from normal metals are attributed to

the phase transformation between the martensite phase and austenite phase. Therefore, it is of

empirical importance to find the transformation temperatures when we develop a new SMA

system.

DSC is the most commonly used technique to determine the martensitic transformation

temperatures in all kinds of SMAs. Apart from the DSC thermal analysis, there are other

methods to determine these temperatures. SMAs in martensite phase exhibit mechanical

behaviour different from the austenite phase. It has been reported that the thermal strain [24-

25] and damping properties [26-30] of SMAs have turning points across the phase

transformation temperatures. In this connection, DSC, TMA and DMA were carried out in

20

this study to determine the transformation temperatures.

Before carrying out the experiments, the Ms of the current alloy system, Ti-22Nb-10Zr (at.%),

was estimated from the previous study. In the study of Kim et al. [1] on the Ti-22Nb-(2-8)Zr

(at.%) SMAs, Ms drops by 35℃ for 1 at.% addition of Zr. They conducted tensile test at

various temperatures and determined the Ms at the temperature which minimum stress for slip

deformation is obtained. The Ms of Ti-22Nb-8Zr (at.%) so determined is around -90℃. As a

result, the Ms of the our alloy is estimated to be -160℃ by extrapolation.

5.3.1 Differential Scanning Calorimetry (DSC)

In NiTi SMAs, transformation from the austenite phase to the martensite phase is usually an

exothermic reaction while the reverse martensitic transformation is endothermic. However, it

was found that the TiNbZr alloy underwent ST and 450℃ heat treatment did not have

significant changes in the heat flow of the DSC analysis. There may be two different

explantions. One is that TiNbZr alloy was not a SMA. Another reason is that the TiNbZr alloy

was a SMA but the transformation was not revealed in the test due to inaccurate scanning

range or insufficient enthalpy of reaction.

Generally a specimen was not a SMA when no peak could be observed in the DSC thermal

analysis. However, we believe that the TiNbZr in this study was a SMA as it exhibited an

abnormal 4% to 5% recoverable strain in the cyclic tensile and cyclic compression tests

(Section 5.6). From the extrapolation of the Ti-22Nb-(2-8)Zr (at.) systems [1], the Ms of the

current system is at around -160℃, which was below the scanning range of the DSC. The

transformation peak may be revealed if we did further cooling. Another possibility of not

showing a peak is the insufficient variation of heat flow due to a small enthalpy of reaction.

-1.5

-1

-0.5

0

0.5

1

-200 -100 0 100 200 300 400

Temperature (℃)

Hea

t flo

w (m

W)

HeatingCooling

Fig. 5.3.1 DSC curve of the 93% CR specimen underwent ST

21

5.3.2 Thermomechanical Analysis (TMA)

TMA was carried out as an alternative to DSC in determining the transformation temperatures.

In recent studies, changes in thermal coefficient of expansion across phase transformation

temperatures were clearly observed in near-equiatomic NiTi [24-25]. On the contrary, no

changes were observed in the TiNbZr alloy. It is proposed that the phase transformation

temperatures were below the scanning range 20℃ to 350℃. The flatten region in Fig. 5.3.2

from 250℃ to 350℃ was due to insufficient power of the TMA machine in reaching the

temperature changing rate.

-5

0

5

10

15

20

25

0 100 200 300 400

Temperature (℃)

Am

plitu

de (u

m)

HeatingCooling

Fig. 5.3.2 TMA curve of the as-cast rod-shaped specimen

5.3.3 Dynamic Mechanical Analysis (DMA)

It has been reported that the internal friction of the austenite phase and martensite phase are

nearly constant against temperature. The change in the internal friction originates from the

damping due to the mobile twin boundaries and/or phase boundaries. As a consequence, a

sharp peak is shown at transformation as there is a change in volume fraction of martensite.

[29-30].

Fig. 5.3.3a shows the DMA results obtained at different frequencies. At 100 Hz, the result

fluctuated and it is difficult to perform conclusive analysis unambiguously. At 20 Hz and 1

Hz, both cooling and heating curves give a broad symmetric peak, with a shift to the higher

temperature during the heating process. In the 20 Hz scan, the peak ended at around 150℃

during the cooling process. The peak in the heating curve also started at around 150℃. In the

1 Hz scan, the peaks ended and started at around 110℃ during the cooling and heating

process, respectively.

22

We believe that these peaks are not transformation peaks. Firstly, the peaks were broad and

spread over a wide range of temperature which was different from typical martensitic

transformation behaviour. Secondly, the “Mf ”thus determined (end point of peak in the

cooling curve) should not coincide with “As“ (starting point of peak in the heating curve)

because hysteresis is usually observed in reversible martensitic transformation. Thirdly,

substantial peak shift was observed in different scanning frequencies of 20 Hz and 1 Hz, the

result even became fluctuated at 100 Hz. This should not happen to typical reversible

martensitic transformation.

The peaks observed in the scan of 20 Hz and 1 Hz were believed related to recrystallization.

During recrystallization, the grain boundaries become more mobile. Viscoplastic behaviour

was expected and the internal friction would also change with respect to temperature. Unlike

martensitic transformation, recrystallization and grain growth continued through a wide

temperature range. This explained the broadening of peak in the DMA results. Moreover,

recrystallization took place whenever the temperature was sufficient regardless of cooling or

heating. So in each frequency, the starting points of both the cooling and heating curves were

at similar position.

DMA results at different applying frequencies

-0.03

-0.02

-0.01

0

0.01

0.02

0 100 200 300 400

Temperature (℃)

Tan

Del

ta

100 Hz Cooling100 Hz Heating20 Hz Cooling20 Hz Heating1 Hz Cooling1 Hz Heating

Fig. 5.3.3a DMA results of the 93% CR specimen underwent ST, and HT at 550℃ for 30min with subsequent oil quench. 100 Hz, 20 Hz and 1 Hz were applied

23

Similar curves were obtained for all 93% CR specimens which had different thermal

treatments. Those specimens were scanned in the range -110℃ and 350℃. The starting

position of the peaks ranged from 140℃ to 160℃ . Similar explanations applied. It is

suggested to have lowered the upper limit of the DMA scanning to avoid recrystallization.

The DMA result shown below is obtained from another final year project student Wong [20],

which the composition was the same as the aforementioned specimens. The specimen was

cold rolled to reduce its thickness by 50%. The frequency applied in the DMA test was 20 Hz

while the amplitude was 20μm. After pre-cooling to -110℃, the test was carried out by

heating to 200℃and then cooling to -110℃at 5 ℃/min.

From the predicted Ms temperature (around -160℃), martensitic transformation should not be

revealed in the test as the scanning was conducted in the range -110℃ to 200℃ . The

prediction was in agreement with the cooling curve in Fig. 5.3.3b. As and Af should be above

-135℃ and fall within the scanning range. In fact, the two peaks at 20℃ and 45℃ in the

DMA heating curve were thought to be relating to the reverse martensitic transformation.

The peaks found in the DMA results may not be at the same position as those predicted from

the tensile test results. In DMA, the transformation was characterized by the change of

internal friction. In the tensile test that Kim et al. conducted, the transformation was

characterized by the temperature dependence of σs and σSIM. As different mechanisms were

involved, the two techniques may direct to different transformation temperatures. The

temperatures at which damping peaks were found may shift in position if we determine the

transformation temperatures by the tensile test.

Specimen with 50% thickness reduction

0.0020

0.0025

0.0030

0.0035

0.0040

0.0045

0.0050

-150 -50 50 150 250Temperature (℃)

Tan

Del

ta

Heating

Cooling

Fig. 5.3.3b DMA results at 20 Hz of the 50% CR specimen

24

5.4 Effect of Thermo-mechanical Treatments on Vickers Hardness Average reading of 10 indentations was used to find out the effect on hardening due to

thermo-mechanical treatments. In Fig. 5.4a, the first three columns indicated the Vickers

hardness of specimens with increasing amount of CR, which is proportional to the percentage

of thickness reduction. It is reasonable to use the conventional dislocation mechanism to

explain the enhancement of hardness.

Thermal treatments were performed on the 93% CR specimens. Their hardness values were

indicated by columns 4 to 10. It could be easily seen from the figure that the error bars were

enlarged significantly after thermal treatment(s). In fact, during the hardness measurement,

the hardness of the aged specimens fluctuated from region to region. During CR a specimen

was deformed to a thin plate with a small amount of curvature owing to the uneven

distribution of applying force in the rollers. In the thermal treatment, energy was provided to

overcome the kinetic barrier. Residual stress field from cold working was relieved leading to

recrystallization and grain growth. The CR deformation might be non-uniform and the

resulting recrystallization microstructure might be inhomogeneous. Hence, the variation of

hardness and thus the error bars became larger after heat treatments.

Fig. 5.4a Vicker’s hardness of specimens with different thermo-mechanical treatments. Columns 1 to 3: increasing degree of CR; Column 4: 93% CR specimen with ST; Columns 5 to 10: 93% CR specimens with ST, and HT at conditions stated in the graph

25

Mechanical treatment was found to be more effective in increasing the hardness than thermal

treatment(s). When 70% and 93% CR were imposed, the hardness increased by 18% and 38%

respectively. However, the hardness changed by only 5% before and after ST for the 93% CR

specimens.

A non-linear relationship between hardness and the degree of mechanical work was shown in

Fig. 5.4b. CR deformation introduced dislocations to the specimen and strain field was

created as a consequence. When the dislocation density was high, strain fields in

neighbouring regions superimposed each other. This led to a sudden increase in hardness

value when the thickness of specimen was reduced from 70% to 93%.

Ageing at 300℃ for 30min gave the maximum hardness value at 320HV. Increasing the

ageing temperature from 300℃ to 450℃ or 550℃ resulted in hardness reduction. This

phenomenon could be explained schematically by Fig. 5.4c. At a particular ageing

temperature (indicated by either one of the solid lines), there should be a peak time when

precipitation hardening gave the maximum hardness. Prolong ageing starting from this point

gradually reduced the hardness as grain growth was allowed. Consider at a fixed ageing time

which cut the curve of 300℃, the hardness would decrease with ageing temperature from

300℃ to 550℃. This was exactly what we observed in the hardness test.

When the ageing time was increased from 30min to 120min, little change in hardness could

be observed in specimens aged at 300℃, 450℃ and 550℃. It could be suggested that the

hardness value had already been saturated (represented by the dotted line in Fig. 5.4c) for the

ageing temperatures under concern. Further prolonging the ageing time is required to prove

this hypothesis.

In conclusion, the thermo-mechanical treatments successfully hardened the TiNbZr alloys.

Although ageing at 450℃ and 550℃ reduced the hardness, and thermal treatment was not as

effective as mechanical treatment, all thermal treatments carried out in this experiment gave

more superior hardness than having mechanical treatment alone.

26

Effect of mechanical treatment

0

10

20

30

40

50

0 20 40 60 80 100Percentage of thickness reduction (%)

Cha

nge

in V

icke

rsha

rdne

ss (%

)

Fig. 5.4b Percentage change in Vickers hardness against the degree of mechanical work

Fig. 5.4c Schematic diagram showing the dependence of hardness on ageing time at different ageing temperatures

5.5 Tensile Stress-strain Behaviour in Ti-22Nb-(2-10)-Zr (at.) Alloys

Mechanical properties like the Young’s modulus, elastic strain, and yield stress are essential

material constants in engineering applications. For biomedical implant application, the ideal

mechanical properties are those close to hard tissues. A weak implant (such as polymeric

materials) is not able to take up the loading sufficiently while an implant with mechanical

properties higher than the hard tissues would cause stress shielding effect (like the case of

stainless steel).

Fig. 5.5a shows the tensile stress-strain curve of a solution treated Ti-22Nb-10Zr (at.%)

specimen and Fig. 5.5b shows the mechanical properties variation with increasing Zr content

(information for Ti-22Nb-(2-8)Zr (at.%) obtained from previous study by Kim et al. [1]). The

fracture strain shown in Fig. 5.5a was 28%. In fact, the fracture strain decreased with the

addition of Zr because ductility decreased with alloying. Nevertheless, Ti-22Nb-10Zr (at.%)

was a ductile metal.

The stress for slip deformation σs was found at the cross section point of the two dotted lines

in Fig. 5.5a. It was found in Fig. 5.5b that σs and ultimate tensile stress (UTS) increased with

Zr content due to solution hardening effect. The drastic increase of σs and UTS from 8% to

10% Zr was due to different fabrication process in previous study by Kim et al. and the

current study. This led to different grain size. Nano-sized grains were found in the current

study (discussed in section 5.2.3 concerning XRD analysis), therefore the yield stress should

be enhanced by the nano-sized grains from the Hall Patch’s equation. The effect of solution

hardening together with different grain size explained the dramatic increase of σs and UTS

from 8% to 10% Zr.

27

0

200

400

600

800

0 10 20 30Tensile strain (%)

Stre

ss (M

Pa)

Mechanical properties variation

0

200

400

600

800

0 2 4 6 8 10 12

Zr content (at%)

Stre

ss (M

Pa)

20

30

40

50

60

Frac

ture

str

ain

(%)

σs(MPa)

UTS(MPa)

Fracturestrain(%)

Fig. 5.5b Mechanical properties variation of the Ti-22Nb-xZr (at.%) system [1]

Fig. 5.5a Tensile stress-strain curve of the 93% CR specimen underwent ST

The Young’s modulus of the current study was found to be 18GPa, much lower than 30GPa

found in Ti-22Nb-(4-6)Zr (at.%) [1]. This could also be explained by the difference in grain

size between the previous published work and the current study. It is noted that the Young’s

modulus of the current system was much closer to that of human hard tissues. This made the

system more appropriate for biomedical application as stress shielding effect found in

common metallic systems could be alleviated.

5.6 Elastic and Plastic Deformation Behaviour in Strain Increment 5-cycle

Tensile and Compression Tests 5.6.1 Elastic and Plastic Deformation Behaviour

In NiTi SMAs, 7% to 8% of the total elastic strain can be recovered. This includes the

conventional elastic strain recovered upon unloading and SE strain recovered due to the

reverse martensitic transformation. On the other hand, the recoverable strains that could be

observed in TiNb-based Ni-free SMAs were much lower. Ti-22Nb-(4-6)Zr (at.%) of Kim et al.

exhibited 4.3% total recoverable strain at room temperature [1].

Fig. 5.6.1 shows the cyclic tensile stress-strain curves of the 93% CR specimen underwent ST.

Tensile stress was applied until the strain reached about 2.5%, and then the stress was

removed. The loading and unloading process with maximum strain 2.5% was repeated 5 times.

Similar loading strategy with maximum strains 3%, 3.5%, 4%, 5%, 6.5%, 8% and 10% were

then imposed on the same piece of specimen. The stress-strain curves were shifted to the right

for the ease of analysis.

28

Fig. 5.6.1 Stress-strain curves for cyclic tensile test of the 93% CR specimen underwent ST, with imposed tensile strains 2.5%, 3%, 3.5%, 4%, 5%, 6.5%, 8% and 10%

100% shape recovery was observed when 2.5% strain was applied. As the applied strain

increased above the elastic limit, shape recovery could not be completed and residual strain

was observed from 3% imposed strains onwards for the first cycle. The first cycles of the

cycling sets 3% to 10% imposed strains deformed plastically. As residual strain could be

observed after unloading, the specimen was heated to 300℃ in an argon-filled tube furnace.

No shape recovery could be seen which deduced to no SME. Although permanent plastic

deformation occurred, the strain hardening effect of the first cycle happened to give higher

elastic deformation in the second to fifth cycles. In addition, the elastic deformation curves of

the second to fifth cycles superimposed themselves. This implied that the alloy was

completely elastic in the second to fifth cycles.

In typical PE deformation of SMAs, a characteristic plateau can be observed. The stress first

increases with strain linearly in the elastic region. When the critical stress is reached, selection

of martensite variances takes place where the stress is kept constant as strain increases.

However in the present alloy system, the stress always increased with strain in the second to

fifth cycles. The selection of martensite variances may not follow the same manner as those in

typical NiTi SMAs.

29

5.6.2 Recoverable Strain

Fig. 5.6.2b and c show the recoverable strain obtained after each set of cycling for the cyclic

tensile and cyclic compression tests respectively. It could be seen that the recoverable strain

increased with imposed strain and saturated at about 5%. The great amount of recoverable

strain distinguished TiNbZr from common metals like stainless steel and titanium alloys,

which exhibit small amount of elastic strain.

Recoverable strain in cyclictensile test

0

2

4

6

0 2 4 6 8 10 12Applied tensile strain (%)

Rec

over

able

str

ain

(%)

ST

ST + HT (300℃

30min OQ)

Recoverable strain in cycliccompression test

0

2

4

6

0 2 4 6 8 10 12Applied compression strain (%)

Rec

over

able

str

ain

(%)

Without HT

HT (300℃30min OQ)

Fig. 5.6.2b Recoverable strain in cyclic tensile test of 93% CR specimens underwent ST, and underwent ST plus HT at 300℃ for 30min with subsequent oil quench

Fig. 5.6.2c Recoverable strain in cyclic compression test of specimen without any heat treatment, and specimen underwent HT at 300℃ for 30min with subsequent oil quench

Though the first cycle of each set of cycling deformed plastically, it could be seen as another

form of mechanical treatment similar to strain hardening. The first cycle may have provided

strain hardening effect and resulted in higher recoverable strain in the subsequent cycles. Note

that the term “recoverable strain” was used here instead of “SE strain” or “elastic strain”,

because the mechanism was not clear.

5.6.3 Young’s Modulus in the Cyclic Tensile and Cyclic Compression Tests

The Young’s modulus (E) in the loading curves of cyclic tensile test from cycle 1 to cycle 5 is

shown in Fig. 5.6.3a. In the cycles of imposed strains 4% and onwards, the loading curves did

not follow a linear relationship. E was determined by drawing a straight line from the starting

point to the end point of an elastic loading curve (schematically shown in Fig. 5.6.3b). In fact,

this non-linear relationship is often observed in TiNb-based SMAs exhibiting PE deformation.

In the cyclic tensile test of the specimen underwent ST, the first cycle of 2.5% imposed strain

hardened the material such that E increased after the first cycle. In the case of 3% to 8%

30

imposed strains, the first cycles softened the material and E decreased as a consequence. The

average Young’s modulus over five cycles ( E ) decreased when the imposed strain decreased

from 2.5% to 6.5% because of strain softening. When the imposed strain was increased from

6.5% to 10%, E increased showing a strain hardening phenomenon. There existed a critical

imposed strain between 6.5% and 8% where the mechanism changed from strain softening to

strain hardening. E over five cycles of the tensile specimen underwent ST and then HT at

300℃ for 30min with subsequent oil quench gave the same trend as the aforementioned,

except that the various of E became smaller (from 12 - 20GPa to 14 - 20GPa)

Cyclic compression tests were performed on specimens before and after HT at 300℃ for

30min with subsequent oil quench. As shown in Fig. 5.6.3c, the value of the compression

modulus in the first cycle of 2.5% imposed strain was similar to that found in the cyclic

tensile test of the specimen underwent ST (around 18GPa). The material was softened by the

first cycle in 3% to 4% imposed strain and hardened by the first cycle in 6.5% to 10%

imposed strain. The various of E due to different imposed strains was narrowed from 12 -

20GPa observed in the cyclic tensile test, to 15 - 19GPa. The change in E in the cyclic

compression test of the specimens before and after HT followed the same manner in the

cyclic tensile test. The proposed mechanisms were similar. The critical imposed strain at

which the mechanism changed from strain softening to strain hardening fell between 4% to

6.5% in the two compression specimens.

E of cycles 1-5 in cyclic tensile test

10

12

14

16

18

20

22

Cycle 1 Cycle 2 Cycle 3 Cycle 4 Cycle 5

E (G

Pa)

2.5% (appliedstrain)3%

3.5%

6.5%

8%

10%

Fig. 5.6.3a Young’s modulus of cycles 1 to 5 in cyclic tensile test of 93% CR specimen underwent ST

31

Fig. 5.6.3b Schematic diagram showing the determination of E in a non-linear stress-strain curve found in cycles at high imposed strain

E of cycles 1-5 in cyclic compression test

10

12

14

16

18

20

22

Cycle 1 Cycle 2 Cycle 3 Cycle 4 Cycle 5

E (G

Pa)

2.5% (appliedstrain)3%

3.5%

4%

6.5%

10%

Fig. 5.6.3c Young’s modulus of cycles 1 to 5 in cyclic compression test of specimen without any heat treatment

5.7 Further Discussions and Recommendations (1) From the austenite phase found in the XRD spectra, it was deduced that Ms lied below

room temperature. In-situ XRD below the room temperature is suggested to confirm this

deduction and to find the Ms temperature at which martensite phase starts to form.

(2) From the FWHM calculation, it was believed that the TiNbZr consisted of nano-sized

grains. Transmission electron microscopy (TEM) with selected area electron diffraction

(SAED) is recommended for verification.

32

(3) Transformation temperatures were not detected in DSC and TMA. The Tan Delta peaks

in DMA of the 50% CR specimen were thought to be related to phase transformation.

Other material properties which depend on transformation temperatures are suggested for

characterization. These include the temperature dependence of electrical resistivity [31-

32], hardness and critical stress for slip deformation.

(4) Increase the ageing time from 30min to 120min was found not affecting the hardness

while increase the ageing temperature from 300℃ to 450℃ and 550℃ caused softening.

Short time heat treatment and prolong ageing at different temperatures are proposed to

find out the hardness and strength evolution for different ageing time.

(5) The Young’s modulus determined in the tensile test was much lower than the previous

reports of similar alloy system. It is believed to be the effect of ultra fine grain size in the

nanometer range. Microstructural analysis by TEM is necessary to provide evidence for

the presence of nano-crystalline structure affecting the Young’ modulus.

(6) In cyclic tensile and cyclic compression tests, strain hardening in the first cycle of each

set of cycling was used to explain the superior elastic deformation in the subsequent

second to fifth cycles. Hardness test between each cycle provides an easy route of

verification. The strain softening and hardening effect on the change in average Young’s

modulus over 5 cycles can also be proofed by the hardness test.

33

6 Conclusions (1) Only austenite phase was found in the XRD spectra of all specimens. Spectra taken at

room temperature and 400℃ had identical peak positions. This implied that the

transformation temperatures should lie below room temperature.

(2) Deformation texture was observed when deformation was severe (at 93% CR). The peak

intensity of the (110) planes decreased but at the same time the peak intensity of the (200)

planes increased.

(3) The uneven distribution of dislocation density formed during CR gave rise to oriented

nucleation and oriented grain growth in the recrystallization process. Preferential re-

orientation of planes was observed. The number of grains with (211) and (222) planes

which were parallel to the rolling sheet surface increased at the expense of the (200)

planes.

(4) Transformations were not reflected in both DSC and TMA. These were attributed to the

small enthalpy of reaction and inadequate range of scanning temperature constrained by

limitation of the available facilities.

(5) Ageing at 300℃ for 30min gave the maximum hardness value at 320HV. Increase the

ageing time from 30min to 120min did not affect the hardness value while increase the

ageing temperature from 300℃ to 450℃ and 550℃ caused softening.

(6) The Young’s modulus determined in the tensile test was 18GPa, much closer to human

hard tissues compared with previous studies of similar composition. This made the

system more suitable for biomedical application because less stress shielding effect is

expected.

(7) No typical plateau found in superelastic NiTi SMAs was observed in both cyclic tensile

and cyclic compression tests. Plastic deformation occurred in the first cycles of all

specimens at imposed strains greater than 2.5%. The second to fifth cycles deformed in

an elastic manner and maximum recoverable strain of 5% was observed in the specimen

underwent ST. As a result, the first cycles could be treated as mechanical treatments

which have led to strain hardening giving rise to higher recoverable strain in the

subsequent cycles.

34

(8) Specimens after HT at 300℃ for 30min with subsequent oil quench gave similar elastic

and plastic deformation behaviour as aforementioned. The recoverable strain after HT

increased in the cyclic compression test but decreased in the cyclic tensile test when

compared to the respective counterpart before HT.

(9) Changes in E over 5 cycles for increasing imposed strains were observed. E decreased

from 2.5% to 6.5% imposed tensile strains as an effect of strain softening and it increased

from 6.5% to 10% imposed tensile strains due to strain hardening. Similar explanation

imposed to the cyclic compression test where E decreased from 2.5% to 4% imposed

strains and increased from 4% to 10% imposed strains.

35

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Appendix I – Experimental Details

Fig. Ia The raw materials as received: titanium (left), niobium and zirconium (right)

Fig. Ib The arc melting furnace Fig. Ic The copper mould in the arc melting

furnace

Fig. Ie The as-rolled 93% CR thin plate Fig. Id The as-cast button-shaped

ingot

Fig. If The 93% CR tensile specimen (left), and the cylindrical compression specimens (right)

Fig. Ih Iron plates put inside the tube furnace to reduce high temperature oxidation of the samples

Fig. Ig The argon-filled tube furnace

Fig. Ii Crack formed during CR (top), the unflat thin plate after CR (bottom)

Fig. Ij Specimen broken into pieces when thermal treatment was carried out in air furnace

Appendix II – XRD Results The below figures show the XRD spectra of specimens underwent different thermo-

mechanical treatments.

0% thickness reduction

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Spectrum(110)(200)Unknown structure(211)(220)(310)(222)(321)(400)

Fig. IIa XRD spectrum of 0% CR specimen

60% thickness reduction

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

70% thickness reduction

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Fig. IIb XRD spectrum of 60% CR specimen Fig. IIc XRD spectrum of 70% CR specimen

93% thickness reduction

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

93%, ST

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Fig. IId XRD spectrum of 93% CR specimen Fig. IIe XRD spectrum of 93% CR specimen underwent ST

93%, ST, HT (300℃ 30min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

93%, ST, HT (300℃ 120min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Fig. IIf XRD spectrum of 93% CR specimen underwent ST, and HT at 300℃ for 30min with subsequent oil quench

Fig. IIg XRD spectrum of 93% CR specimen underwent ST, and HT at 300℃ for 120min with subsequent oil quench

93%, ST, HT (450℃ 30min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

93%, ST, HT (450℃ 120min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Fig. IIh XRD spectrum of 93% CR specimen underwent ST, and HT at 450℃ for 30min with subsequent oil quench

Fig. IIi XRD spectrum of 93% CR specimen underwent ST, and HT at 450℃ for 120min with subsequent oil quench

93%, ST, HT (550℃ 30min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

93%, ST, HT (550℃ 120min OQ)

0

20

40

60

80

20 40 60 80 100 120 1402 Theta (degree)

Inte

nsity

(%)

Fig. IIj XRD spectrum of 93% CR specimen underwent ST, and HT at 550℃ for 30min with subsequent oil quench

Fig. IIk XRD spectrum of 93% CR specimen underwent ST, and HT at 550℃ for 120min with subsequent oil quench

Appendix III – Cyclic Tensile and Cyclic Compression Results The below figures show the stress-strain curves of the strain increment cyclic tensile and

cyclic compression tests. The same specimen was loaded and unloaded for 5 cycles at

imposed tensile strain (or compression strain) 2.5%, 3%, 3.5%, 4%, 5%, 6.5%, 8% and finally

10%. The stress-strain curves were shifted for the ease of observation.

Fig. IIIa Cyclic tensile stress-strain curves of the 93% CR specimen underwent ST

Fig. IIIb Cyclic tensile stress-strain curves of the 93% CR specimen underwent ST, and HT at 300℃ for 30min with subsequent oil quench

Fig. IIIc Cyclic compression stress-strain curves of the as-cast cylindrical rod specimen

Fig. IIId Cyclic compression stress-strain curves of the cylindrical rod specimen with HT at 300℃ for 30min with subsequent oil quench