Metastable Beta Titanium Alloys...Metastable Beta Titanium Alloys: Tuning Their Beta Phase Stability and Low-Temperature Martensitic Transformation by Le Xu to obtain the degree of

Metastable Beta TitaniumAlloys:Tuning Their Beta Phase Stability andLow-Temperature Martensitic Trans-formation

Le XuM.Sc. ThesisNovember, 2015

;

Metastable Beta Titanium Alloys:Tuning Their Beta Phase Stability and

Low-Temperature Martensitic Transformation

by

Le Xuto obtain the degree of Master of Science

at Delft University of Technology,

to be defended publicly on Friday November 6, 2015 at 11:00 AM.

Student number: 4323408

Project duration: October 2014 – November 2015

Thesis committee: Prof. dr. ir. S. van der Zwaag, TU Delft, Chairman of Thesis Committee

Prof. dr. W. Xu, TU Delft & Northeastern University, China

Dr. ir. M. J. M. Hermans, TU Delft

An electronic version of this thesis is available at http://repository.tudelft.nl/.

http://repository.tudelft.nl/

Preface

I am grinding my sample in the lab,

some sunny day.

I am sailing myself to somewhere,

somewhere we know,

free and easy.

I hear the bird chirping from somewhere,

somewhere you know,

peaceful and quiet.

I am not preparing the sample,

but preparing myself.

I know we will meet again,

some sunny day.

Le Xu

Delft, September 25, 2015

iii

Acknowledgements

During the past two years at TU Delft, I have learned and experienced way more than what

I could have anticipated and imagined at the time the wonderful journey began. I would

like to express my earnest gratitude to all who taught, helped and guided me throughout

the journey.

Many thanks to Prof. Sybrand van der Zwaag, my supervisor. You offered me the great

opportunity to do my master thesis project at the NovAM group. And many thanks to Prof.

Wei Xu and Dr. Qi Lu for co-supervision. I appreciate your unconditional help and inspira-

tional guidance throughout my thesis project. I learned a lot from you, especially the way of

thinking, the state of mind, and the qualities and methodologies to become an outstanding

professional.

Many thanks to my colleagues at the NovAM group and the office NB 1.20. I enjoyed the

fantastic year working with you all.

Many thanks to Lijing Xue, Frans Oostrum, Sander van Asperen, Dr. Maruti Hegde, Dr.

Richard Huizenga, and Kees Kwakernaak for technical support and kind help in the lab.

Special thanks to Adien and Jakob for your warm and unconditional hospitality during

the past two years. I will never forget the wonderful memories of hanging out with you

guys, the delicious food you cooked, and the board game we played.

Special thanks to all my friends around the world for unconditional support and en-

couragement from different aspects all the way regardless of the distance between us.

Lastly, my heartfelt gratitude to my family. You gave me everything. I do not know how

to put it into words. I love you.

Le Xu

Delft, October 2, 2015

v

Abstract

This research focuses on the stability of three metastableβ titanium alloys – Ti-10V-1Fe-3Al

(T i −1013), Ti-10V-2Cr-3Al (T i −102Cr 3) and Ti-10V-2Fe-3Al-B (T i −1023B).

In the first stage of the work, new β+ (α+β) heat treatment schemes were designed

on the basis of Thermo-Calc® simulations and previous research work and carried out to

such an extent that the β phase stability is tuned to the right level and a sufficient amount

of retained β phase is obtained at room temperature available for further transformation.

At the second stage, the effectiveness of the thermal driving force in triggering a marten-

sitic transformation in the retained β phase was investigated. Both in-situ and ex-situ

transformations were studied via DSC and further quenching in self-made cooling baths,

respectively.

The starting temperatures of low-temperature martensitic transformations were deter-

mined for each alloy. It was found that the low-temperature martensitic transformation

for the three alloys was of an isothermal character. Through XRD testing, the phases of

T i −1013 before and after low-temperature martensitic transformation were determined;

as in the commercial Ti-10V-2Fe-3Al (T i −1023) alloy, the type of martensite formed in re-

sponse to quenching from the two-phase region isα′′ in stead ofα′ martensite, due to over-

saturation of the martensite. EPMA testing proved the formation of Martensite-Free Zone

(MFZ) observed in the case of T i −102Cr 3 when quenched to a temperature just below the

starting temperature of the low-temperature martensitic transformation does not relate to

the diffusion behavior of alloying elements. Electrical Discharge Machining (EDM) was

found to be detrimental in this research as the machining process changed the designed

microstructure of the three metastable β titanium alloys.

Finally, a proposal is made to produce tailored α+β titanium metallic composite via

SPS.

vii

Contents

1 Introduction 11.1 Classification of titanium alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.2 β and metastable β titanium alloys . . . . . . . . . . . . . . . . . . . . . . . . . . 2

1.3 Metastable phases in metastable β titanium alloys . . . . . . . . . . . . . . . . . 4

1.3.1 α′ martensite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4

1.3.2 α′′ martensite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5

1.3.3 ω phase . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

1.4 Nucleation and diffusional growth process in metastable β titanium alloys . . 7

1.5 Stability of the β phase and its relation to the MS temperature . . . . . . . . . . 7

1.5.1 Alloying elements and β phase stability. . . . . . . . . . . . . . . . . . . . 8

1.5.2 Molybdenum equivalence. . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

1.5.3 MS temperature . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9

1.6 Martensitic transformation in metastable β titanium alloys. . . . . . . . . . . . 9

1.6.1 Martensite formed during quenching . . . . . . . . . . . . . . . . . . . . . 10

1.6.2 Athermal and isothermal martensitic transformation . . . . . . . . . . . 10

1.6.3 Low-temperature martensitic transformation . . . . . . . . . . . . . . . . 11

1.6.4 Stress-induced martensitic transformation . . . . . . . . . . . . . . . . . 11

1.7 Objectives of this work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11

2 Materials 132.1 Composition of as-received materials . . . . . . . . . . . . . . . . . . . . . . . . . 13

2.2 Initial heat treatment state and resulting microstructure . . . . . . . . . . . . . 15

2.3 Sample preparation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 15

3 Thermo-Calc® Simulation 173.1 Objectives of the simulation work . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

ix

x Contents

3.2 Simulation results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 18

3.3 Analysis and discussion of simulation results . . . . . . . . . . . . . . . . . . . . 18

3.3.1 Comparison of the simulation results . . . . . . . . . . . . . . . . . . . . . 21

3.3.2 Effects of the redistribution of solute elements . . . . . . . . . . . . . . . 21

3.3.3 Selection of the most relevant result. . . . . . . . . . . . . . . . . . . . . . 21

4 Heat Treatment andMetallography 23

4.1 Heat treatment design process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 23

4.1.1 Selection of heat treatment scheme . . . . . . . . . . . . . . . . . . . . . . 23

4.1.2 Selection of annealing temperature . . . . . . . . . . . . . . . . . . . . . . 24

4.2 Heat treatment experiments. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

4.3 Methods for metallographic specimen preparation and examination . . . . . . 25

4.3.1 Electropolishing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

4.3.2 Etching . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 26

4.3.3 Optical microscopy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

4.3.4 Metallographic image processing . . . . . . . . . . . . . . . . . . . . . . . 27

4.4 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27

4.4.1 Quenching from β single-phase region . . . . . . . . . . . . . . . . . . . . 27

4.4.2 Annealing for reaching equilibrium state . . . . . . . . . . . . . . . . . . . 28

4.4.3 Comparison between experiment and simulation results . . . . . . . . . 32

4.5 Conclusion of the heat treatment experiments . . . . . . . . . . . . . . . . . . . 33

5 Low-TemperatureMartensitic Transformation 35

5.1 In-situ method via Differential Scanning Calorimetry . . . . . . . . . . . . . . . 35

5.1.1 Reason of choosing DSC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

5.1.2 Sample preparation for DSC . . . . . . . . . . . . . . . . . . . . . . . . . . 36

5.1.3 DSC testing and results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

5.1.4 Discussion of DSC results . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36

5.1.5 Metallographic examination of DSC samples and discussion. . . . . . . 37

5.1.6 Ex-situ testing of DSC cooled samples . . . . . . . . . . . . . . . . . . . . 39

5.1.7 Final remarks . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

Contents xi

5.2 Ex-situ method via further quenching . . . . . . . . . . . . . . . . . . . . . . . . . 42

5.2.1 Further-quenching experiments . . . . . . . . . . . . . . . . . . . . . . . . 42

5.2.2 Results and discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42

5.3 X-ray diffraction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53

5.3.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53

5.3.2 Experimental details . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53

5.3.3 XRD results and discussions . . . . . . . . . . . . . . . . . . . . . . . . . . 54

5.4 Electron probe micro-analysis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59

5.4.1 Experimental details . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 59

5.4.2 EPMA results and discussions . . . . . . . . . . . . . . . . . . . . . . . . . 59

5.5 Conclusion of low-temperature experiments. . . . . . . . . . . . . . . . . . . . . 61

6 Conclusion 63

7 Recommendations for FutureWork 65

A Fabricating Novel TitaniumMetallic Composites by SPS 67A.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 67

A.2 Research work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

A.2.1 Project idea. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

A.2.2 Theoretical basis of applying SIM effect as reinforcement. . . . . . . . . 70

A.2.3 Selection of powder bed preheating temperature . . . . . . . . . . . . . . 73

A.2.4 Controlled parameters . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73

A.3 Expected research outcomes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74

A.4 Supplier information of titanium powders . . . . . . . . . . . . . . . . . . . . . . 75

Bibliography 77

1Introduction

Featuring a relatively low density, a high specific strength, a good corrosion resistance,

superior biocompatibility, and an outstanding high-temperature strength and creep be-

havior, titanium and its alloys are superior lightweight engineering materials and are play-

ing an increasingly important role in airframes, jet engines and other aerospace engineer-

ing applications, especially for supersonic aircraft due to its higher requirement for high-

temperature strength.

From the advent of first commercial titanium produced by DuPont Company in 1948,

over 100 titanium alloys have been developed up to now. Only have 20 to 30 of them did

realize commercial application. Half the commercial titanium market share is occupied by

the classic Ti-6Al-4V alloy and roughly 25% is taken up by unalloyed titanium [1].

1.1. Classification of titanium alloysTitanium alloys are generally classified as α, β and α+β alloys, based on their crystal-

lographic phases at room temperature [2]. The temperature where the last remaining α-

phase transforms into the β phase upon heating is referred to as the β-transus. Pure tita-

nium mainly contains two phases: theα phase – with hexagonal close-packed (hcp) crystal

structure and existing below 882◦C – and the β phase – with body-centered cubic (bcc)

1

2 1. Introduction

Figure 1.1: Effects of different groups of alloying elements on schematic titanium phase diagrams [1].

structure and existing from 882◦C up to melting point 1668◦C [3]. Alloying elements in

titanium can affect the β-transus temperature. Therefore, depending on their respective

influences, alloying elements are classified as neutral, α-stabilizers and β-stabilizers, the

schematics of the influence of representative alloys elements on titanium phase diagrams

are shown in Figure 1.1.

By tailoring the alloying element concentration and heat treatments, different types of

phases are obtained at room temperature. Depending on the type of phases and volume

fraction, morphology and arrangement, titanium alloys will exhibit different properties. A

summary of microstructural features, β-transus temperature, properties and applications

for the major categories of titanium alloys is given in Table 1.1.

β and metastable β titanium alloys are the most versatile category among titanium al-

loys, featuring the most complex and tunable microstructure. They possess the highest

strength-to-weight ratio and an outstanding combinations of strength, toughness and fa-

tigue resistance [1]. Alloys of this type are represented by the β-isomorphous phase dia-

gram (shown in Figure 1.1), of which vanadium, niobium, tantalum and molybdenum are

the representative β-isomorphous-stabilizers which, by definition, suppress the formation

of the α phase at room temperature.

1.2. β and metastable β titanium alloysβ titanium alloys are capable of being hardened to extremely high strength levels, for in-

stance, more than 1400 MPa (tensile strength) for the Beta C titanium alloy (alloy com-

position: Ti-3Al-8V-6Cr-4Mo-4Zr, metastable β alloy) [6]. Their complex microstructure

1.2. β and metastable β titanium alloys 3

Tab

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Pro

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120-

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890-

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300-

400

110-

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800-

1200

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8-19

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750-

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res

4 1. Introduction

enables physical metallurgists to tune the microstructure toward better combinations of

strength and toughness.

Although β alloys inevitably have some drawbacks such as a high density because of

alloying with heavy elements, unsatisfactory low temperature and high temperature prop-

erties, higher cost and microstructural instabilities, they have a place in the extremely com-

petitive engineering materials market owing to their high strength-to-density ratio, low

Young’s modulus, high toughness and heat-treatability, and improved formability at room

temperature [7].

A schematic phase diagram of β titanium alloy is given in Figure 1.2, in which the re-

gion of metastable β alloys is that between β stabilizing elements composition βc and βs .

Furthermore, a small section inside metastable β region between βc and βω called me-

chanically instable region exists.

Commonly used β alloys include Ti-35V-15Cr (Alloy C), Ti-40Mo, Ti-30Mo. The most

commonly used metastable β alloys are Ti-10V-2Fe-3Al (T i −1023), Ti-5Al-2Sn-2Cr-4Mo-

4Zr-1Fe (β−C E Z ), Ti-15Mo-2.6Nb-3Al-0.2Si (TIMETAL 21S), Ti-3Al-8V-6Cr-4Mo-4Zr (Beta

C), Ti-15Mo (IMI 205), Ti-8V-8Mo-2Fe-3Al (T i − 8823), Ti-15V-3Cr-1Mo-0.5Nb- 3Al-3Sn-

0.5Zr (VT 35), Ti-5V-5Mo-1Cr-1Fe-5Al (VT 22), Ti-11.5Mo-6Zr-4.5Sn (Beta III) [1][5][8].

1.3. Metastable phases in metastable β titanium alloys

In general, decomposition products of metastable β alloys areα (hcp), α′ martensite (hcp),

α′′ martensite (orthorhombic), and the ω phase (complex structure) [9][10][11][12]. Of our

primary interest, martensitic transformation can be tailored primarily within the mechan-

ically instable region (between βc and βw ) as shown in Figure 1.2. The transformation can

be tuned by heat treatments that control the β-stabilizers’ contents in the β phase [5][13].

1.3.1. α′ martensite

Hexagonal α′ martensite exists in two forms: massive martensite (lath or packet marten-

site) and acicular martensite [4][14]. The former only exists in engineering-grade pure tita-

nium or alloys with high martensitic transformation temperature (MS); acicular martensite

exists in alloys that are heavily alloyed [4]. When increasing the solute content, the hcp α′

1.3. Metastable phases in metastable β titanium alloys 5

Figure 1.2: Generalized titanium-β-stabilizing-elements phase diagram with highlighted regions of various

alloy categories and β phase decomposition products, valid for isomorphic β-stabilizers that forms no chem-

ical compound with titanium.

martensite will be distorted. From a crystallographic perspective, the α′ martensite will

lose its hexagonal symmetry and transform into orthorhombic martensite that is referred

to as α′′ martensite [15][16][17].

The α′ phase is not an equilibrium phase, and can be oversaturated by alloying ele-

ments [8][9]. In other words, more alloying elements can be maintained in α′ than those in

α phase.

1.3.2. α′′ martensite

The orthorhombicα′′ phase can be regarded as an intermediate stage in between the trans-

formation from bcc to hcp. The α′′ phase is acicular/needle-like martensitic structure [8].

In titanium alloys, different from the situation in carbon-containing steels, martensite is

softer than the β phase [13]. Apart from β → α′′ one-step transformation, β+ωath →α′′ two-step stress-assisted transformation has also been observed in metastable β alloys

such as Ti-1023 [18]. On the other hand, the stress-induced martensite is generally the or-

thorhombic α′′ martensite [5][18][19].

6 1. Introduction

(a) bcc β and orthorhombic α′′ crystal structures (b) hcp α′ and orthorhombic α′′ crystal structures

Figure 1.3: Equivalent demonstration of the relations of α′′ with β, α′ crystal structures by using an or-

thorhombic cell [21].

XRD result shows that the α′′ martensite is characterized by cleavage of interference

lines that only occur in hexagonal structure; with increasing the content of β-stabilizer, the

cleavage grows [8]. Compared with α′ phase, α′′ phase can be even more oversaturated

with alloying elements [8][20].

With regard to shape memory effect, the α′′ → β reversion is of primary importance.

With the transformation from highly oriented α′′ after deformation back to β by heating,

the alloy will return back to its original shape. This process is believed to be an athermal

shearing process [18].

The relation between orthorhombic α′′ phase and β and α′ phases is illustrated by an

orthorhombic cell in Figure 1.3(a) and Figure 1.3(b), respectively. As a diffusionless trans-

formation, the homogeneous lattice-distortive strain (or Bain strain) of the transformation

from β to α′ is made up of a shrinkage of a-axis and expansions of b and c axes, along

with shuffles of atoms. In comparison, the β→ α′′ transformation is halfway the β→ α′

transformation with moderate lattice distortion [8][13][21].

1.3.3.ω phase

In published work, three types of ω phase – athermal ω phase (ωath), isothermal ω phase

(ωi so) and stress-inducedωphase (ωD ) – with complex crystal structure have been reported

for titanium alloys. From a crystallographic point of view, both ωath and ωi so have trigonal

1.4. Nucleation and diffusional growth process in metastable β titanium alloys 7

symmetry in heavily β-stabilized alloys and have a hexagonal crystal structure in solute-

leaner alloys [4][22]. It is difficult to avoid ωath and ωi so during quenching or during initial

stage of aging or reheating [9][18]; and the transformation between hexagonal and trigonal

structures is continuous with the variation of solute content.

On the other hand, reported in recent work [23][24], stress-induced ω phase (ωD ) has

been observed in Ti-9Mo-6W alloy that shows combined TRIP and TWIP effects. At the

same time, it has been found that in Ti-10V-3Fe-3Al alloy, stress-induced ωD phase exist,

leading to the stress plateau phenomena [25]. The crystal structure of ωD is hexagonal,

which has been proved for alloys having stress-induced bcc matrix to ωD phase trans-

formation not only in metastable β titanium alloys but also in metals such as tantalum

[24][26][27][28].

1.4. Nucleation and diffusional growth process in metastable

β titanium alloys

As can be seen from Figure 1.2, the phase transformation from β phase (bcc) to α phase

(hcp) while cooling through β-transus may take place martensitically or by conventional

diffusion-controlled nucleation and growth process. Which mechanism will prevail mainly

depends on applied cooling rate and alloy composition.

If the cooling rate from β single-phase field to α+β two-phase field through β-transus

is sufficiently low, α (also α′ [5]) phase will preferentially nucleate at the grain boundary

of β phase, microscopically resulting in α layers/feather-like structure along the β grain

boundary, mainly because of the fact that the hcpα phase is incoherent to the parent bcc β

phase. During continuous cooling, theα phase will grow into the β grains as parallel plates

that form an α colony [4][29].

1.5. Stability of the β phase and its relation to the MS temper-

ature

One of the most important characteristics of metastable β titanium alloys is the β phase

stability. It is the dominating factor in determining if the martensitic transformation will

8 1. Introduction

Figure 1.4: A schematic of the solute partitioning effect during the course of the precipitation ofα phase from

parent β phase.

take place as well as the amount of martensite formed in response to either a thermal driv-

ing force (e.g. quenching) or a mechanical driving force (e.g. deformation). For instance,

deformation-induced martensite is very sensitive to the composition of β phase that also

relates closely to MS temperature [30][31].

1.5.1. Alloying elements and β phase stability

β phase stability, in essence, is a function of the in-situ composition of β phase at a given

heat treatment state in response to the solute partitioning effect of the precipitation of α

phase that takes away α-stabilizers but leaves β-stabilizers accumulating in the parent β

phase, as illustrated in Figure 1.4. The solute partitioning phenomenon has been reported

and verified in several recent works [5][32]. Taxonomically, β-stabilizers – such as Mo, V,

Nb, Fe, Cr, Mn and H – stabilize the β phase, whereas α-stabilizers – represented by Al, N,

C, and O – destabilize β phase. In other words, the more the β-stabilizers accumulate in

the β phase, the more easily the β phase can be secured after quenching and the lower the

MS temperature.

Therefore, β phase stability and thus the MS temperature can be tuned via designing

appropriate heat treatment schemes which control the precipitation of α phase.

1.6. Martensitic transformation in metastable β titanium alloys 9

1.5.2. Molybdenum equivalence

Molybdenum equivalence (Moeq ) is by far the most widely used approach to evaluate and

also quantify the β phase stability in titanium alloys, which is expressed as a linear sum of

weight averages of alloying elements in weight percent (w t%). Molybdenum equivalence

equation is given in Equation 1.1 [4]:

Moeq = 1.00Mo +0.28N b +0.22Ta +0.67V +1.6Cr +2.9Fe −1.00Al (1.1)

1.5.3. MS temperature

Neelakantan et al. [5][33] proposed a a linear dependence between the in-situ composition

of the β phase, expressed in weight percent (w t%) and atomic fraction (N ), and the MS

temperature:

MS [w t%] = 1156−150Fe −107Mn −96Cr −67Ni −49Mo

−41Cu −37V −17N b −7Z r +15Al (1.2)

MS [N ] = 1156−17480Fe −12186Mn −11299Cr −8096Ni −9463Mo

−5250Cu −4354V −3160N b −1200Z r +865Al (1.3)

Hence, with Equation 1.2 and Equation 1.1, the stability ofβ phase and MS temperature

can be calculated from in-situ β phase composition. Further, it can be easily see that the

more stable the β phase, the lower the MS temperature.

1.6. Martensitic transformation in metastable β titanium al-

loysMartensitic transformation is observed in many metals. According to Nishiyama [34], the

martensitic transformation occurs by the cooperative movement of atoms, featuring dif-

fusionless character, surface relief and the presence of many surface imperfections; these

characteristics can also be employed to validate the existence of martensite.

10 1. Introduction

1.6.1. Martensite formed during quenching

Martensite, first observed in steel by A. Martens, is a metastable non-equilibrium phase

that is not present in an equilibrium phase diagram. It can exist indefinitely at ambient

temperatures, but will transform back into equilibrium phases upon annealing at an el-

evated temperature. Martensite forms very rapidly if the required cooling rate (thermal

driving force) is applied. Each atom only displaces a small (sub-atomic) distance during

the transformation, along with a sudden reorientation at the same time.

Martensitic transformation in titanium alloys comprises the cooperative movement of

atoms by shearing, resulting in a microscopically homogeneous transformation from body-

centered cubic β phase into the hexagonal α′ phase [13]. With increasing solute element

content, the hexagonal α′ martensite becomes distorted and, from the prospective of crys-

tallography,α′ martensite loses its hexagonal symmetry and then changes into orthorhom-

bic crystal structure which is designated as α′′ martensite [4][14].

The extent of martensitic transformation under a fixed thermal driving force in the

metastable β alloys can be engineered by tuning the stability of β phase and thus be em-

bodied by tailored MS temperature.

Martensites in titanium (both α′ and α′′) are much softer than their counterpart in

steels, because interstitial oxygen atoms can only induce a small elastic distortion of the

titanium martensite with hexagonal lattice in bcc β matrix [4]; at the same time, substi-

tutional alloying elements also contribute to martensitic transformation in titanium al-

loys [34]. Therefore, the hcp α′ martensite does not possess high hardness and strength,

and the mechanical properties of orthorhombic α′′ martensite are close to that of stable

α and β solid solution with similar chemical composition [8]. Although at the first sight

the strengthening effect of martensite in titanium alloys is not that obvious compared with

steels, its complexity and heat-treatability still leaves physical metallurgists a lot of space

for strength enhancement.

1.6.2. Athermal and isothermal martensitic transformation

Athermal and isothermal martensitic transformation may take place in titanium alloys.

Generally speaking, as for isothermal martensitic transformation, the fraction of marten-

1.7. Objectives of this work 11

site increases with time at fixed transition temperature, whereas the fraction remains con-

stant at fixed temperature in the case of athermal martensitic transformation [35].

Athermal martensitic transformation is clearly recognized as a transformation with no

thermal activation needed. The isothermal martensitic transformation is mentioned more

frequently in fields related to shape memory effect such as in Ti-Ni alloys [36]. However, a

satisfactory explanation of the nature of isothermal and athermal martensitic transforma-

tion does not yet exist [36][37][38].

1.6.3. Low-temperature martensitic transformation

In theory, by appropriate heat treatment design, retained β phase in metastable β titanium

alloys can be obtained at room temperature, which means the alloy still possesses the ca-

pability for martensitic transformation. A martensitic transformation can be initiated by

further quenching – i.e. providing an additional thermal driving force.

In the case of steels, the kinetics of martensitic transformation at temperature below

room temperature points at an isothermal process [39][40].

1.6.4. Stress-induced martensitic transformation

Deformation mechanisms, in general, are classified into four basic categories: dislocation

glide, twinning, plasticity derived from phase transformation, and thermally activated glide

by diffusion (for instance, creep and grain boundary sliding). The deformation mecha-

nisms of metastable β alloys mainly include slip, Stress-Induced Martensitic transforma-

tion (SIM), deformation twinning or any combination of them [41]. As a means to provide

energy for martensitic transformation, mechanical driving force can also initiate the trans-

formation process in retained β phase. The SIM effect in metastable β alloys features the

β→α′′ martensitic transformation [5][18][19].

1.7. Objectives of this work

As an important category of aerospace engineering materials, metastable β titanium al-

loys have been increasingly used on, for instance, aircraft in which the structural mate-

rial may experience low-temperature work environment. Therefore, investigating the low-

12 1. Introduction

temperature martensitic transformation phenomenon of metastable β alloys is of practical

significance to evaluate the alloy’s performance for aerospace applications.

However, to the best of the author’s knowledge from the available published work, this

phenomenon with regard to metastable β titanium alloys has not been effectively inves-

tigated. Therefore, this exploratory study is devoted to investigate the phenomenon in

metastable β titanium alloys as a function of the effectiveness of thermal driving force in

activating further martensitic transformation in this category of titanium alloys.

A variety of experimental methods was employed in this work. Starting from computer-

aided thermodynamic equilibrium simulations, Thermo-Calc® was utilized for finding out

the appropriate heat treatment scheme in order to better tune the β phase stability. Experi-

mental work was carried out subsequently by virtue of a combination of Thermo-Calc® sim-

ulation and previous experimental experience. Then, the phenomenon of low-temperature

martensitic transformation was systematically investigated both in in-situ and ex-situ man-

ners by performing DSC testing and further-quenching experiments in selected cooling

baths, respectively. XRD testing was performed to examine the experiment results. EPMA

testing was conducted to study the formation of Martensite-Free Zone. Metallographic

specimen preparation and metallographic image processing were conducted throughout

this work.

2Materials

Three novel metastable β titanium alloys designed by the NovAM group at TU Delft were

employed in this work – they are, Ti-10V-1Fe-3Al (T i −1013), Ti-10V-2Cr-3Al (T i −102Cr 3)

and Ti-10V-2Fe-3Al-B (T i −1023B), of which T i −1013 and T i −102Cr 3 were fabricated by

the Institute of Metal Research, Chinese Academy of Science, Shenyang, China, while the

manufacturer of T i −1023B was GfE Gesellschaft für Elektrometallurgie mbH, Nuremberg,

Germany.

2.1. Composition of as-received materials

Presented in Table 2.1 are the nominal and the achieved compositions (in w t%) of the

three alloys, and calculated molybdenum equivalence Moeq and MS temperature (◦C) from

Equation 1.1 and Equation 1.2, respectively, at fully β phase solution-treated condition; β-

transus temperature Tβ (◦C) of the three metastable β alloys derived from the calculation

results of Thermo-Calc® TCW™ v.5 with TTTI3 Thermotech Ti-based Alloys Database.

In addition, as far as T i −1023B was concerned, a minor amount of TiB intermetallic

compound exists at room temperature irrespective of the solution treatment [5], to control

prior β grain size by Zener pinning, which is properly predicted by Thermo-Calc®.

13

14 2. Materials

Tab

le2.

1:T

he

ach

ieve

dco

mp

osi

tio

ns

(in

wt%

)o

fas-

rece

ived

Ti−

102C

r3,T

i−10

13,a

nd

Ti−

1023

Bal

loys

[5][

19];

calc

ula

ted

mo

lyb

den

um

equ

ival

ence

Mo e

qan

d

MS

tem

per

atu

reu

nd

erth

eco

nd

itio

no

ffu

llβ

ph

ase

mic

rost

ruct

ure

;cal

uca

ledβ

-tra

nsu

ste

mp

erat

ure

Tβ

oft

he

thre

em

etas

tab

leβ

allo

ys.b

al.s

tan

ds

for

bal

ance

.

VC

rFe

Al

BC

uSi

OH

NC

Ti

Mo e

q

MS

Tβ

[◦C

][◦

C]

Ti−1

02C

r3(T

i-10

V-2

Cr-

3Al)

10.6

2.18

–3.

19–

––

0.1

0.00

10.

013

0.05

bal.

7.4

329

810

Ti−1

013

(Ti-

10V

-1Fe

-3A

l)9.

97–

0.97

3.02

––

–0.

110.

001

0.01

40.

051

6.5

414

837

Ti−1

023B

(Ti-

10V

-2Fe

-3A

l-B

)10

.69

–1.

993.

140.

089

0.00

70.

057

0.21

00.

014

0.00

80.

034

9.8

236

835

2.2. Initial heat treatment state and resulting microstructure 15

Table 2.2: The applied solution treatment scheme of as-recieved T i−102Cr 3, T i−1013 and T i−1023B alloys.

Solution treatment schemes

Cooling modeTemperature domain

Temperature Soaking time

[◦C] [mi n]

T i −102Cr 3

β phase field 900 60 Water quenchT i −1013

T i −1023B

2.2. Initial heat treatment state and resulting microstructureThe three as-received metastable β alloys have gone through a solution heat treatment.

Their solution treatment schemes are given in Table 2.2, and the initial microstructures of

the three as-received alloys are shown in Figure 2.1.

For T i − 1013 and T i − 102Cr 3, as they were directly quenched from β single-phase

region, the lath-like phase present throughout the parent β grains is due to martensite

formed during quenching. In comparison, no martensite was seen in T i−1023B . It showed

the β phase as background and TiB intermetallics (black spots in Figure 2.1(c)).

2.3. Sample preparationCylindrical samples with 4mm diameter and 7mm length were machined by means of Elec-

trical Discharge Machining (EDM). After the machining, a thin oxide layer of unknown

thickness was formed, which was carefully removed by mechanical polishing before fur-

ther heat treatment.

16 2. Materials

(a) T i −1013 (b) T i −102Cr 3

(c) T i −1023B

Figure 2.1: Initial microstructure of T i −1013, T i −102Cr 3 and T i −1023B .

3Thermo-Calc® Simulation

Thermo-Calc® is the most powerful and widely used computer-aided simulation software

for thermodynamic calculations. It is capable of calculating the amounts of phases and

their compositions at different temperatures under the equilibrium condition, liquidus and

solidus, and calculating phase diagrams (be it binary, ternary and multi-component), as

long as the composition of the alloy and conditions, such as temperature, pressure and

system size, are given.

Using the composition of the three metastable β alloys (given in Table 2.1), Thermo-

Calc® simulations were carried out to simulate which β phase compositions can be ob-

tained when annealing at theα+β two-phase region. In this work, two versions of Thermo-

Calc® database were used – i.e. Thermo-Calc® TCW™ v.5 and Thermo-Calc® v.3.1.

3.1. Objectives of the simulation workThe main goal of performing computer-aided simulation is to find out the optimal anneal-

ing temperature of the three metastable β titanium alloys in the α+β two-phase region in

order to better tune the β phase composition and hence stability and to secure enough β

phase at room temperature after quenching for the purpose of realizing further martensitic

transformation.

17

18 3. Thermo-Calc® Simulation

As stated in Section 1.5.1, β phase will become more stable because of the accumu-

lation of β-stabilizers (e.g. Mo, V, Nb, Fe, Cr and Mn) in the β phase, because of the so-

lute partitioning effect owing to precipitation of α phase. The idea of conducting Thermo-

Calc® simulation is to design the heat treatment scheme that could effectively control the

α precipitation in practice.

In addition, it is worthwhile to note that the simulation is only effective provided the

equilibrium state is reached for a given heat treatment.

3.2. Simulation results

The most important data expected from the simulation results are the in-situ composition

of the β phase and the volume fraction of α precipitate. Furthermore, β phase stability and

MS temperature can be calculated from Equation 1.1 and Equation 1.2.

Thermo-Calc® TCW™ v.5

From the outset, Thermo-Calc® TCW™ v.5 was employed to perform the simulation with

TTTI3 Thermotech Ti-based Alloys Database. The simulation is mainly performed within

the promising temperature range between β-transus and 200◦C lower than it. A summary

of the simulation results is given in Table 3.1.

Thermo-Calc® v.3.1

Afterward, the simulation work was conducted on Thermo-Calc® v.3.1 again for further

sifting out appropriate intercritical annealing temperature, based on the experience and

results from Thermo-Calc® TCW™ v.5. The results are summarized in Table 3.2.

3.3. Analysis and discussion of simulation results

From Table 3.2 and Table 3.1, four major alloying elements – vanadium, chromium , iron

and aluminum – can be identified that have the major influence upon the stability of β

phase in response to heat treatment, of which vanadium, iron and chromium are the β-

stabilizers, whereas aluminum is α-stabilizer.

3.3. Analysis and discussion of simulation results 19

Table 3.1: Simulation results from Thermo-Calc® TCW™ v.5 with TTTI3 Thermotech Ti-based Alloys

Database of β phase compositions (in w t%) and α phase fraction in equilibrium condition for T i −102Cr 3,

T i −1013 and T i −1023B alloys at selected annealing temperatures (Tanne.) within α+β two-phase region

to obtain different β phase stability (Moeq ) and MS temperature by tuning α phase precipitation; alloying

elements presented are those directly related to the calculations of Moeq and MS given in Equation 1.1 and

Equation 1.2, respectively.

Tanne.V Cr Fe Al Cu Moeq

MS α phase fraction

[◦C] [◦C] [%]

T i −102Cr 3

750 11.8 2.4

0

5.3

0

6.5 291 20.5

700 14.6 3.1 5.1 9.6 123 37.5

680 16.0 3.4 5.1 11.0 45 43.7

670 16.8 3.5 5.0 12.0 -5 46.8

660 17.7 3.7 5.0 12.8 -53 49.5

650 18.6 3.9 4.9 13.9 -108 52.3

630 20.5 4.4 4.8 16.0 -228 57.2

T i −1013

750 12.1

0

1.1 4.9

0

6.4 341 28.1

700 15.1 1.5 4.7 9.6 177 44.6

680 16.6 1.6 4.7 11.2 93 50.3

670 17.4 1.7 4.6 12.0 50 52.6

650 19.0 1.9 4.5 13.7 -35 57.2

630 20.8 2.1 4.5 15.6 -137 61.2

T i −1023B

750 11.6

0

2.0 5.3

≈ 0

8.3 231 18.2

700 13.7 2.5 5.0 11.3 79 31.7

680 14.7 2.7 5.0 12.5 17 36.8

670 15.2 2.8 4.9 13.3 -22 39.3

660 15.8 2.9 4.8 14.1 -61 41.8

630 17.7 3.3 4.7 16.6 -189 48.3


Table 3.2: Simulation results from Thermo-Calc® v.3.1 with TTTI3 Thermotech Ti-based Alloys Database of β

phase compositions (in w t%) and α phase fraction in equilibrium condition for T i −102Cr 3, T i −1013 and

T i−1023B alloys at selected annealing temperatures (Tanne.) withinα+β two-phase region to obtain different

β phase stability (Moeq ) and MS temperature by tuning α phase precipitation; alloying elements presented

are those directly related to the calculations of Moeq and MS given in Equation 1.1 and Equation 1.2, respec-

tively.

Tanne.V Cr Fe Al Cu Moeq

MS α phase fraction

[◦C] [◦C] [%]

T i −102Cr 3

780 11.7 2.4

0

3.1

0

8.6 266 11.5

750 12.6 2.7 3.1 9.7 205 20.4

700 15.7 3.4 3.0 12.9 25 37.6

680 17.2 3.7 2.9 14.5 -61 43.8

T i −1013

780 11.7

0

1.2 2.9

0

8.3 319 17.9

750 13.1 1.4 2.8 9.9 236 28.5

700 16.3 1.8 2.7 13.4 54 44.8

680 17.8 1.9 2.6 14.8 -23 50.3

T i −1023B

780 11.9

0

2.3 3.1

≈ 0

11.4 151 10.9

750 12.8 2.4 3.1 12.6 89 18.4

700 14.9 2.9 2.9 15.6 -65 31.8

680 16.0 3.1 2.9 16.9 -135 37.2

3.3. Analysis and discussion of simulation results 21

3.3.1. Comparison of the simulation results

Comparing the simulation results from Thermo-Calc® TCW™ v.5 with the results from

Thermo-Calc® v.3.1, it is obvious that the two versions have different interpretation of the

accumulation and also the dissipation of β- and α-stabilizers.

Specifically, as for the in-situ composition of the β phase, in Thermo-Calc® v.3.1, ap-

proximately 1w t% more vanadium can be found in β phase for the three metastable β

alloys at the same annealing temperature in comparison with Thermo-Calc® TCW™ v.5.

Around 0.3w t% less chromium is shown in v.3.1 than TCW™ v.5 in the case of T i −102Cr 3,

whereas roughly 0.25w t% and 0.45w t% more iron are presented in v.3.1 than TCW™ v.5 in

the cases of T i −1013 and T i −1023B , respectively. On the other hand, for the α-stabilizer,

an approximately 2w t% more aluminum can be found in v.3.1 than TCW™ v.5.

3.3.2. Effects of the redistribution of solute elements

Furthermore, it can be seen from both Table 3.2 and Table 3.1 that the lower the anneal-

ing temperature in α+β two-phase region, the more the vanadium, iron and chromium

– β-stabilizers – accumulated in the β phase, and the more the aluminum – α-stabilizer –

dissipated from the β phase as a result of the solute partitioning effect. This corresponds

well with the EPMA results of Neelakantan et al. [5] and Li et al. [32].

The impact ofα precipitation on MS temperature (closely related to β phase stability) is

further illustrated in Figure 3.1, in which the more the α precipitation, the more intensive

the solute partitioning effect that brings about more stable retained β phase and having a

lower MS temperature.

3.3.3. Selection of the most relevant result

An apparent difference in interpreting the behavior of the solute partition of aluminum

(α-stabilizer) exists between v.3.1 and TCW™ v.5.

In theory, along with α precipitation, the aluminum content should only decrease in β

phase and should not go higher than the total aluminum content of the alloy.

In TCW™ v.5, for all of the three alloys, the aluminum content in β phase during inter-

critical annealing is even higher than the composition of the alloys, even though a decreas-


Figure 3.1: Comparison between α volume fraction and calculated MS temperature from the in-situ β phase

composition, on the basis of the simulation results from Thermo-Calc® v.3.1.

ing trend of aluminum content can be seen along with the lowering intercritical annealing

temperature.

In contrast, the aluminum content in β phase from v.3.1 is lower than the composition

of the alloys; and the increasingly intensifying aluminum dissipation fromβ phase with the

lowering intercritical annealing temperature is also embodied.

Therefore, the simulation done by Thermo-Calc® v.3.1 is more convincing, and the sub-

sequent work regarding heat treatment scheme design is constructed upon the results of

Thermo-Calc® v.3.1.

4Heat Treatment and Metallography

In this chapter, heat treatment schemes are designed and carried out to tune the stability

of β phase for the three metastable β alloys; metallographic preparation and examination

are performed to examine the effectiveness of designed heat treatment schemes.

4.1. Heat treatment design processThe main objective of heat treatment is to tune the microstructure of the alloys to such an

extent that enough β phase can be secured at room temperature. Meanwhile, the β phase

stability and also the MS temperature of the alloys can be effectively tailored.

4.1.1. Selection of heat treatment scheme

In this work, β+ (α+β) heat treatment is selected as the initial solution treatment in the

β single-phase region is able to remove the original α precipitates and martensites com-

pletely. Therefore, the influence of original microstructure and previous heat treatment

can be effectively offset or minimized. Further, 15 minutes holding time in the β single

phase is found to be enough to solutionize the alloy according to previous work [19][32].

when annealing in α+β two-phase region, acicular α phase is expected to precipitate,

and acicular α′ and/or α′′ martensite is also expected to form during quenching [19]. Suf-

23

24 4. Heat Treatment and Metallography

ficiently long annealing in the two-phase region should ensure a completed solute parti-

tioning process (i.e. the accumulation of β-stabilizers and the dissipation of α-stabilizers)

and thus a stabilized β phase.

4.1.2. Selection of annealing temperature

Stated in Section 3.3, the simulation results of Thermo-Calc® v.3.1 is chosen to be referred

to. As it is practically unavoidable that martensite will form during quenching to room tem-

perature from the two-phase region [5][42], it would be more advisable to select annealing

temperatures that will lead to not too muchαphase precipitation in order to secure enough

β phase for further transformation.

On the other hand, both α′ and α′′ martensites in β titanium alloys are supersaturated

with solute elements [8]. This can be easily understood that kinetically, there is no time for

martensitic transformation to proceed with solute partitioning process during quenching.

As a consequence, the martensitic transformation will not affect the stability of retained β

phase, but it will consume the total amount of retained β phase. Therefore, higher anneal-

ing temperatures are preferred because of less α precipitation (shown in Table 3.2).

However, from previous work, a discrepancy is found between experiment and the sim-

ulation. Specifically, as for T i −102Cr 3, reported by Li et al. [19], holding at 700◦C for an

hour leads to α phase of 78% volume fraction. In comparison with the simulation results

in Table 3.2, regardless of the big difference between α volume fractions and suppose the

prediction of MS temperature is close to the real case, quenching to room temperature will

only lead to a small amount of retained β phase left, which would not be appropriate for

further transformation. Likewise, for T i−1013, annealing 30 minutes at 750◦C leads to 42%

α phase, let alone such an annealing time falling far short of what is required by reaching

the equilibrium state where more α phase is expected to precipitate. Last but not least,

according to the work done by S. Neelakantan [5], only 29.7% β phase was retained, even

though the alloy was annealed at 700◦C for mere 15 minutes after solution treatment in β

single-phase region.

A higher annealing temperature in the two-phase region implies less stable β phase

because of less α precipitation. As a result, it becomes a crucial task to choose an effective

annealing temperature that is optimal for further martensitic transformation.

4.2. Heat treatment experiments 25

Given the β-transus of the three alloys in Table 2.1 and having considered the trade-

off mentioned above, 780◦C was selected as the annealing temperature in the two-phase

region in order to securing enough β phase at room temperature for further transforma-

tion.

4.2. Heat treatment experiments

Table 4.1 summarizes the heat treatment schemes applied to the three novel metastable β

titanium alloys. Experiments were done in a Nabertherm burn-Off furnace (model LV5).

A special air intake and exhaust system is present that allows the application of inert at-

mosphere or shielding gas. Moreover, the incoming gas is preheated before entering the

interior of the furnace. Thus, a good temperature uniformity is ensured.

Argon gas is applied in this work as inert atmosphere to mitigate the issue of oxidation.

The heating rate from room temperature to 900◦C was approximately 15◦C mi n−1. The

cooling mode between 900◦C and 780◦C is natural furnace cooling. When the annealing

treatment in α+β two-phase region was done, samples were immediately quenched in

water at room temperature with very fast cooling rates (order of magnitude: 102 ◦C s−1

[43]).

Despite the application of inert atmosphere, oxidation cannot be absolutely avoided.

Yet the situation of oxidation is still much better than applying no shielding gas. For in-

stance, on average, a 0.458µm oxide layer is formed for the case of T i − 102Cr 3 going

through 900◦C/15mi n + 780◦C/2.5h followed by water quenching.

4.3. Methods for metallographic specimen preparation and ex-

amination

Being aware of the existence of SIM effect, the methodologies designed specifically in order

to examine the microstructure of the samples was applied.


Table 4.1: The applied heat treatment schemes for T i −102Cr 3, T i −1013 and T i −1023B alloys; WQ stands

for water quenching.

Metastable β alloys Heat treatment schemes

Ti-10V-1Fe-3Al (T i −1013)

900◦C/15mi n + WQ

900◦C/15mi n + 780◦C/15mi n + WQ

900◦C/15mi n + 780◦C/30mi n + WQ

900◦C/15mi n + 780◦C/2h + WQ

780◦C/2h + WQ

Ti-10V-2Cr-3Al (T i −102Cr 3)

900◦C/15mi n + WQ

900◦C/15mi n + 780◦C/15mi n + WQ

900◦C/15mi n + 780◦C/30mi n + WQ

900◦C/15mi n + 780◦C/2.5h + WQ

900◦C/15mi n + 780◦C/8h + WQ

900◦C/15mi n + 780◦C/24h + WQ

780◦C/2.5h + WQ

Ti-10V-2Fe-3Al-B (T i −1023B)

900◦C/15mi n + WQ

900◦C/15mi n + 780◦C/15mi n + WQ

900◦C/15mi n + 780◦C/30mi n + WQ

900◦C/15mi n + 780◦C/2h + WQ

4.3.1. Electropolishing

To avoid retainedβ phase transforming to martensite during normal mechanical polishing,

electropolishing was utilized instead. The electropolishing solution is made up of 36mL

perchloric acid, 390mL methanol, 350mL ethylene glycol and 24mL distilled water. The

instrument used was the Struers electropol setup. Samples were electropolished for 30s at

a voltage of 45V with flow rate of 16 by using the mask of 0.5cm2 at a constant temperature

of 5◦C.

4.3.2. Etching

To better reveal microstructural features, the Standard Kroll’s reagent – 3mL HF, 6mL HNO3

and 100mL distilled water – was used as etchant.

4.4. Results and discussion 27

4.3.3. Optical microscopy

Olympus® BX60M microscope was used to observe the microstructure of heat-treated sam-

ples, which was also used to take metallographic images for further analysis.

4.3.4. Metallographic image processing

Metallographic images were processes by using Adobe Photoshop in combination with Im-

ageJ software for the purpose of calculating the volume fraction of different phases. The

results to be presented were determined by taking the average of processing five metallo-

graphic images (different areas of the same sample) for each sample.

4.4. Results and discussion

4.4.1. Quenching from β single-phase region

Theoretically, the most unstable β phase is obtained when quenching the alloy directly

from the β single-phase region. At the single-phase region, the composition of β phase is

the same as the composition of the alloy, so that, without the β-stabilizers’ accumulation

and the α-stabilizers’ dissipation, the least stable β is obtained.

From Figure 4.1, it is obvious that a huge amount of acicular martensite forms inside

the β grains for both T i −1013 and T i −102Cr 3, and the high density of martensite can be

more clearly seen from Figure 4.1(b).

However, as can be seen from Figure 4.1(d), for T i −1023B , no martensite in any mor-

phology is observed. This implies that even at the most unstable situation, β phase is still

stable enough to refrain from martensitic transformation upon quenching to room tem-

perature. In other words, the β phase of T i − 1023B is stable enough that makes the MS

temperature below room temperature. This also indicates that no matter how this alloy

is annealed in α+β two-phase region, martensite will not form upon quenching to room

temperature, as annealing in two-phase region will only make β phase more stable.


(a) T i −102Cr 3 (b) T i −102Cr 3

(c) T i −1013 (d) T i −1023B

Figure 4.1: The microstructure of 900◦C/15mi n followed by water quenching to room temperature for T i −1013, T i −102Cr 3 and T i −1023B .

4.4.2. Annealing for reaching equilibrium state

Stated in Section 4.1.2, the annealing temperature in α+β two-phase region has been cho-

sen as 780◦C. Several annealing times in the two-phase region of β+ (α+β) heat treatment

were applied as specified in Table 4.1. The results are shown in Figure 4.2, 4.3 and 4.4.

It is clear from the result of T i − 1013 in Figure 4.2 that with increasing holding time

in the α+β two-phase region, less acicular phase was obtained after quenching to room

temperature, and apparently more β phase was retained (shown in Figure 4.2(d)). The ex-

cessive amount of acicular phase is believed to be martensite (α′ and/or α′′), as the max-

imum amount of α phase precipitated while holding 15 minutes at 780◦C must be lower


(a) 900◦C/15mi n + 780◦C/15mi n + WQ (b) 900◦C/15mi n + 780◦C/15mi n + WQ

(c) 900◦C/15mi n + 780◦C/30mi n + WQ (d) 900◦C/15mi n + 780◦C/2h + WQ

Figure 4.2: The microstructure of T i −1013 going through β+ (α+β) heat treatment with different annealing

time in the α+β two-phase region.

than the case of annealing 2 hours at the same temperature, which is formed instantly dur-

ing quenching but the α phase, in comparison, needs a lot more time to precipitate. This

phenomenon, in effect, verifies the importance of realizing the equilibrium state.

Otherwise, with shorter annealing in the two-phase region, the α phase precipitation

process is far from complete. As a result, less α phase is formed, and the expected solute

partitioning process for β-stabilizers and α-stabilizers is halted at a primary stage. There-

fore, a shorter annealing does not stabilize theβphase, and more acicular martensite forms

during quenching from retained β phase. The amount of retained β phase after quenching

is minor, which can indeed be observed in Figure 4.2(a).

Likewise, in the case of T i −102Cr 3 given in Figure 4.3, obviously more acicular phase



(c) 900◦C/15mi n + 780◦C/2.5h + WQ (d) 900◦C/15mi n + 780◦C/8h + WQ

Figure 4.3: The microstructure of T i−102Cr 3 going throughβ+(α+β) heat treatment with different annealing


formed at shorter annealing time (30 minutes) at 780◦C than at longer times (2.5 and 8

hours). Furthermore, comparing Figure 4.3(c) with Figure 4.3(d), the amount of acicular

phase does not show an apparent increase, which implies that 2.5 hours annealing has

realized or at least closely approached the equilibrium state.

As for the results of T i −1023B in Figure 4.4, given the fact that TiB (black spots in the

metallographic images) forms unavoidably in this alloy [5], at shorter annealing time – 15

and 30 minutes – no acicular precipitate formed. For longer annealing time like 2 hours,

acicular phase formed throughout the parent β grain, which was supposed to be α phase

instead of a combination ofα phase and martensite upon quenching because, as discussed

in Section 4.4.1, no thermal martensite formed in the case of the most unstable β phase, let



(c) 900◦C/15mi n + 780◦C/2h + WQ (d) 900◦C/15mi n + 780◦C/2h + WQ

Figure 4.4: The microstructure of T i−1023B going throughβ+(α+β) heat treatment with different annealing


alone the more stable β phase obtained via intercritical annealing.

Based on the nucleation and diffusional growth process of α phase described in Sec-

tion 1.4, the feather-like structure observed in Figure 4.2 and Figure 4.3 is believed to be α

precipitates [4].

Moreover, based on the experience from previous researchers on the three alloys, be-

cause of the high content of chromium which is a slow diffuser [19], the speed of α precip-

itation in T i −102Cr 3 is much slower than the other two alloys. Therefore, it is logical to

expect that if T i −102Cr 3 reached the equilibrium state at 2.5 hours intercritical anneal-

ing at 780◦C, 2 hours is supposed to be enough for the other two alloys to approach the

equilibrium state.


On the other hand, Van Bohemen et al. [44] reported a delayed solute partition dur-

ing the course of α precipitation exists in β titanium alloys. Taking this phenomenon into

account, there could be a difference in the amount of martensite (represented by the total

amount of acicular phase formed after quenching) between 2 and 8 hours annealing. How-

ever, from Figure 4.3(c) and 4.3(d), no obvious difference was observed. Thus, the delayed

solute partition issue does not apply to the current alloys.

Further, as reported by Li et al. [32], solute partitioning was observed in Ti-10V-2Fe-3Al

alloy both at heat treatment 750◦C/60mi n and 900◦C/15mi n + 700◦C/20mi n by means

of EPMA. Neelakantan observed similar effect on the same alloy under heat treatment

700◦C/15mi n and 900◦C/15mi n + 650◦C/15mi n [5]. Kinetically, the speed of solute parti-

tion while annealing at 780◦C is faster than 750◦C and 650◦C. Therefore, one can conclude

that the designed heat treatment condition leads to effective solute partition.

4.4.3. Comparison between experiment and simulation results

By means of Adobe Photoshop in combination with ImageJ software, metallographic im-

ages of the three alloys after heat treatment were systematically processed. The comparison

between experiment and simulation results are given in Figure 4.5.

After β+ (α+β) heat treatment scheme, both α precipitate and martensite (α′ and α′′)

are acicular [19]. As shown in Figure 4.5 for T i −1013 the experiment data perfectly corre-

sponds with the simulation result (around 18% volume fraction).

In contrast, obvious discrepancy exists between experiment and simulation results for

T i −102Cr 3 and T i −1023B . As for T i −102Cr 3, roughly 12% more acicular phase formed

in comparison with the simulation value (11.5%). Therefore, the excessive acicular phase

is very likely to be martensite formed during quenching as Thermo-Calc® cannot give in-

formation on the non-equilibrium martensitic transformation.

For T i−1023B , 4% less acicular phase formed than the simulation result (10.91%). Con-

sidering the martensitic transformation of the most unstable β phase given in Figure 4.1(d)

that no martensite formed after quenching from β single-phase region to room tempera-

ture, the acicular phase obtained from β+(α+β) is believed to be totally α, as annealing in

the two-phase region makes β phase more stable.

4.5. Conclusion of the heat treatment experiments 33

Figure 4.5: Comparison ofα volume fraction between experiment results and Thermo-Calc® v.3.1 simulation

results; heat treatment state: 900◦C/15mi n + 780◦C/2~2.5h + WQ.

4.5. Conclusion of the heat treatment experimentsThe heat treatment scheme for realizing the equilibrium state is found as 900◦C/15mi n +

780◦C/2h for T i −1013 and T i −1023B , and 900◦C/15mi n + 780◦C/2.5h for T i −1023Cr 3,

in which the alloys are solutionized at 900◦C (β single-phase region), and α precipitates at

780◦C (α+β two-phase region) to tune the stability ofβ phase, via solute partitioning effect.

During quenching from 780◦C to room temperature, no martensite forms in the case of

T i −1023B , a minor amount of martensite forms probably in the case of T i −1013, and the

martensitic transformation is likely to take place in T i −102Cr 3.

The experiment result of T i−1013 at the heat treatment state 900◦C/15mi n + 780◦C/2h

+ WQ corresponds very well with the simulation result done by Thermo-Calc® v.3.1 with

TTTI3 database.

Moreover, if quenching directly from the single-phase region, lath-like martensite with

high density formed in the cases of T i −1013 and T i −102Cr 3, whereas no such a phase

formed in the case of T i −1023B because the β phase is very stable.

5Low-Temperature Martensitic

Transformation

The phenomenon of low-temperature martensitic transformation for the three alloys – T i−1013, T i −102Cr 3 and T i −1023B – is systematically studied by using linear cooling (in a

DSC) and isothermal cooling.

5.1. In-situ method via Differential Scanning CalorimetryDifferential Scanning Calorimetry (DSC) was employed in this study to try to follow and

monitor the low-temperature martensitic transformation in the three alloys during contin-

uous cooling condition.

5.1.1. Reason of choosing DSC

To detect and thus investigate the behavior of low-temperature martensitic transforma-

tion of the three alloys in an in-situ way – i.e. where the transformation of the alloys going

through designed heat-treatment starts taking place – DSC and dilatometry are the two

potential candidates. Dilatometry is to measure the volume change during phase transfor-

mation, while DSC is to detect the change of the amount of heat required to increase the

temperature of the alloy and the reference.

35

36 5. Low-Temperature Martensitic Transformation

Previous research showed that dilatometry could not catch the signal of martensitic

transformation during quenching [5][19], whereas DSC has been explored to induce and

investigate the martensitic transformation in an in-situ manner [45][46].

5.1.2. Sample preparation for DSC

Required by the DSC sample pan and the instrument itself, the heat-treated samples of

the three alloys were machined into thin disk with 4mm diameter and 0.3mm thickness

through Electrical Discharge Machining (EDM). The EDM was applied after the heat treat-

ment processes was done and the oxide layer was carefully removed mechanically.

5.1.3. DSC testing and results

PerkinElmer® DSC 8000 Advanced Double-Furnace Differential Scanning Calorimeter was

utilized. The results is summarized in Figure 5.1. The programming of DSC cooling is:

firstly heating the sample to 50◦C and then cooling it to -155◦C with a cooling rate of 20◦C

mi n−1. This cooling rate is chosen because the instrument may not be able to catch up

with the phase transformation signal at very fast cooling rate, and the disturbance of noise

– e.g. moisture – may be intensive at very slow cooling rate that would cover up the possible

peak of low-temperature martensitic transformation.

The heat treatment states of DSC samples of the three metastableβ alloys is the equilib-

rium state of β+ (α+β) heat treatment – i.e. 900◦C/15mi n + 780◦C/2h~2.5h + WQ. More-

over, to examine the low-temperature behavior of the sample at non-equilibrium state, the

T i −1013 sample going through 900◦C/15mi n + 780◦C/15mi n + WQ was also tested.

5.1.4. Discussion of DSC results

From Figure 5.1, the region between 50◦C and -155◦C is rather smooth, except the two un-

expected peaks near inverse peak the starting point for T i −1013 at non-equilibrium heat

treatment state and T i−102Cr 3 at equilibrium heat treatment state which is because of the

programming of the instrument itself. In other words, no peak regarding low-temperature

martensitic transformation was observed.

5.1. In-situ method via Differential Scanning Calorimetry 37

Figure 5.1: DSC testing results of heat-treated T i −1013, T i −102Cr 3 and T i −1023B with cooling rate 20◦C

mi n−1 from 50◦C to -155◦C.

5.1.5. Metallographic examination of DSC samples and discussion

To examine if the low-temperature martensitic transformation took place during DSC test-

ing, metallographic examination was carried out for DSC samples both before and after

DSC testing. The results are given in Figure 5.2 and 5.3.

It can be seen from Figure 5.2 and 5.3 that the microstructure is surprisingly the same –

that is, there is no microstructural change during the cooling process in DSC.

On the other hand, the initial microstructure of DSC samples is very different from what

we have expected them to be if comparing the subfigures in Figure 5.2 correspondingly with

their counterparts in Figure 4.2(d), 4.2(a), 4.3(c) and 4.4(c). Specifically, the area of retained

β phase obtained from designed heat treatment scheme was missing in DSC samples of

T i −1013 and T i −102Cr 3, but a high-density lath-like phase was observed instead. An-

other strong argument comes from the microstructure of T i − 1023B . In Figure 4.4(c), a

number of acicular α precipitates formed throughout the β grains, whereas when made

into DSC disk, the acicular α precipitates were not present.


Therefore, the designed microstructure has changed during the course of EDM while

fabricating the DSC disk. Also, as the initial microstructure of DSC disk is very similar to the

case of quenching directly from β single-phase region (given in Figure 4.1) for all the three

alloys,the EDM process was likely to introduce similar heat treatment for the thin disk – i.e.

heat is generated along with the electrical discharging cutting process, and because of the

thin thickness (0.3mm), cooling rate during EDM can be very fast.

(a) T i −1013, 900◦C/15mi n + 780◦C/2h + WQ (b) T i −1013, 900◦C/15mi n + 780◦C/15mi n + WQ

(c) T i −102Cr 3, 900◦C/15mi n + 780◦C/2.5h + WQ (d) T i −1023B , 900◦C/15mi n + 780◦C/2h + WQ

Figure 5.2: The initial microstructure of DSC samples of T i −1013, T i −102Cr 3 and T i −1023B before DSC

testing.


(a) T i −1013, 900◦C/15mi n + 780◦C/2h + WQ (b) T i −1013, 900◦C/15mi n + 780◦C/15mi n + WQ

(c) T i −102Cr 3, 900◦C/15mi n + 780◦C/2.5h + WQ (d) T i −1023B , 900◦C/15mi n + 780◦C/2h + WQ

Figure 5.3: The microstructure of DSC samples of T i −1013, T i −102Cr 3 and T i −1023B after DSC testing

with cooling rate 20◦C/min from 50◦C to -155◦C

5.1.6. Ex-situ testing of DSC cooled samples

To further examine if the DSC samples were capable of exhibiting further martensitic trans-

formation, ex-situ experiment was carried out by putting the original DSC samples into a

cooling bath of -160◦C for 60mi n.The results are given in Figure 5.4.

The results of T i −1013 and T i −102Cr 3 (in Figure 5.4(a) and 5.4(b)) are quite the same

as those cooling in DSC testing (in Figure 5.3(a) and 5.3(c)), and the initial microstructure

of DSC samples was maintained (in Figure 5.2(a) and 5.2(c)).

However, a difference is observed in the case of T i − 1023B . Owing to the very sta-

ble β phase in this alloy that almost does not show any martensitic transformation while


quenching to room temperature (see Figure 5.2(d)), acicular martensite apparently formed

throughout the microstructure, and low-temperature martensitic transformation indeed

took place in the -160◦C cooling bath (shown in Figure 5.4(c)), whereas in DSC, there is no

martensitic transformation taking place at terminal temperature of -155◦C (Figure 5.3(d)).

Four possibilities may explain the discrepancy – soaking time, terminal temperature,

cooling rate, and hydrogen attack. Specifically, it might be possible that the soaking time

in DSC is not enough to trigger the transformation. But referring to the ex-situ further

quenching experiment at -160◦C shown in Figure 5.12(a), dramatic martensitic transfor-

mation already took place at merely 1 minute, and the holding time of DSC testing at ter-

minal temperature is more than 1 minute. Thus, the influence of soaking time can be ruled

out. Subsequently, the terminal temperature. From the ex-situ further quenching experi-

ment results in Figure 5.11, low-temperature martensitic transformation of T i −1023B has

already taken place at -52◦C, and thus in theory it is not a problem for it to exhibit marten-

sitic transformation at -155◦C. Therefore, this factor can also be excluded. Lastly, the cool-

ing rate while putting the thin disk sample in cooling bath is much faster than the cooling

process in DSC (20◦C/mi n). More specifically, the higher the cooling rate, the larger the

driving force for the transformation. Lastly, because of the severe electrochemical treat-

ment during EDM, the sample is vulnerable to hydrogen and oxygen attack. Hydrogen

can stabilize the β phase and could make it stable enough not to exhibit low-temperature

martensitic transformation [30], while oxygen (α-stabilizer) does not stabilize β phase.

5.1.7. Final remarks

Combining the results from DSC testing with the ex-situ reproduction cooling scenario of

DSC, it can be found that, apart from the microstructural change introduced during EDM,

the cooling process of DSC testing did not induce the expected transformation.

Furthermore, based on the metallographic examination, it is highly likely that EDM

process induces hydrogen (β-stabilizer) attack to the disk sample that stabilizes β phase

to such an extent that the low-temperature martensitic transformation is not shown when

quenching to -160◦C and/or new round of heat treatment – heating up the sample to β

single-phase region and followed by quenching to room temperature – that destroy the de-

signed microstructure of the three alloys.


(a) T i −1013, 900◦C/15mi n + 780◦C/2h + WQ (b) T i −102Cr 3, 900◦C/15mi n + 780◦C/2.5h + WQ

(c) T i −1023B , 900◦C/15mi n + 780◦C/2h + WQ (d) T i −1023B , 900◦C/15mi n + 780◦C/2h + WQ

Figure 5.4: The microstructure of DSC samples of T i − 1013, T i − 102Cr 3 and T i − 1023B after soaking in

-160◦C cooling bath for 60mi n.


Table 5.1: Cooling bath applied for further-quenching experiment for different soaking time.

Terminal temperature Cooling bath Soaking time

-15◦C Ethylene glycol + dry ice 60mi n

-26◦C Freezing chamber of domestic refrigerator 60mi n

-52◦C Diethyl carbitol + dry ice1mi n

60mi n

-100◦C Diethyl ether + dry ice1mi n

60mi n

-160◦C i-Pentane + liquid nitrogen1mi n

60mi n

-196◦C Liquid nitrogen1mi n

60mi n

5.2. Ex-situ method via further quenching

5.2.1. Further-quenching experiments

Selected cooling baths chosen with reasonable temperature spans are summarized in Ta-

ble 5.1. Thermocouple was utilized throughout the experiment to ascertain that the wanted

temperature was achieved and maintained. Similar cylindrical samples as those in Chapter

4 were employed.

Metallographic specimen preparation and examination were executed, and obtained

metallographic images were further examined via image-processing software with the same

methodologies as mentioned in Section 4.3.

5.2.2. Results and discussion

The only phase transformation that can take place at low-temperature condition is marten-

sitic transformation. The martensite still possesses acicular morphology.

Ranges of starting temperature

At the designed heat treatment state, T i − 1013 started exhibiting the transformation at -

15◦C (Figure 5.7), T i −102Cr 3 did not show the transformation at -15◦C but started from

-26◦C (Figure 5.9), whereas T i −1023B only started showing the transformation from -52◦C

5.2. Ex-situ method via further quenching 43

(Figure 5.11).

It is found that the temperature ranges for the three alloys to show low-temperature

martensitic transformation after the designed heat treatment state are: room temperature

to -15◦C for T i −1013, -15◦C to -26◦C for T i −102Cr 3 and -26◦C to -52◦C for T i −1023B , as

illustrated in Figure 5.5.

According to the Thermo-Calc® simulation results in Table 3.2, T i −1013 has the high-

est MS temperature at designed heat treatment state, followed by T i − 102Cr 3 and T i −1023B in sequence, which corresponds with the starting temperature of low-temperature

martensitic transformation of the three alloys. As the higher the MS temperature, the lower

the β phase stability (Moeq ), the retained β phase of T i −1013 at room temperature is still

the most unstable one among the three alloys, and thus it starts showing low-temperature

martensitic transformation at higher temperature.

Shift of MS temperature in comparison with simulation results

As shown in Figure 4.5, the experiment result of T i −1013 corresponded well with Thermo-

Calc® simulation; the acicular phases in this alloy at room temperature are believed as α

phase, and no martensitic transformation has taken place until this point. Given the start-

ing temperature of low-temperature martensitic transformation of T i −1013 in Figure 5.5,

its MS temperature is slightly lower than room temperature in practice.

As for T i − 1023B , concluded in Section 4.5, there was no martensitic transformation

occurred by quenching to room temperature because of the very stable β phase and thus

the below-room-temperature MS . From Figure 5.5, its MS temperature is therefore found

in the range between -26◦C to -52◦C, which is around 200◦C lower than the MS obtained

from simulation result (Table 3.2).

For T i −102Cr 3, if the MS temperature is higher than room temperature as simulated,

the low-temperature martensitic transformation of this alloy should continuously take place

below room temperature, rather than starting showing from a region between -15◦C and

-26◦C. Thus, its real MS temperature is between -15◦C and -26◦C at the designed heat treat-

ment state.

In sum, there is a shift of MS temperature from what was simulated by Thermo-Calc® ,

and the reason of such a shift is the difference in the compositions of alloying elements.

From Table 3.2, for the three alloys, such an MS temperature can be obtained at the simula-


Figure 5.5: The ranges of low-temperature martensitic transformation starting temperatures of T i − 1013,

T i −102Cr 3 and T i −1023 at the designed heat treatment state.

tion result of around 100◦C lower. This means that, in reality, there is more intensive solute

partitioning and thus more stable retained β phase than what was simulated. Thermo-

Calc®, therefore, predicts less intensive solute partitioning and less stable β phase for the

three alloys at the designed heat treatment state.

Isothermal character

Comparing Figure 5.8(e) with 5.8(f), Figure 5.10(c) with 5.10(d), Figure 5.11(d) with 5.11(e),

and Figure 5.12(b) with 5.12(c), the amount of martensite (acicular morphology) formed is

a function of soaking time. Therefore, the low-temperature martensitic transformation in

the case of metastable β titanium alloy showed isothermal character.

A comparison after metallographic image processing between the amount of acicular

martensite formed for the situations of as-quenched state and low-temperature condition

with short and long soaking time is presented in Figure 5.6. From the figure, in the cases

of T i −1013 and T i −102Cr 3, the amount of martensite formed is a function of time. Yet


Figure 5.6: The amount of martensite formed after further quenching in liquid nitrogen for 1 and 60 minutes

in comparison with the as-quenched state (initial state).

for T i −1023B , the amount of martensite formed is roughly the same at 1 minute and 60

minutes soaking time, which means the speed of low-temperature martensitic transforma-

tion occurred in T i −1023B is faster than the speeds of the other two alloys at -196◦C. But

the isothermal character of T i − 1023B showed at higher terminal temperatures, such as

further quenching to -52◦C (see Figure 5.11(d) and 5.11(e)).

However, in some cases such as T i −1013, when soaking in the -52◦C cooling bath for 1

and 60 minutes (shown in Figure 5.7(d) and 5.7(e), respectively), the amount of martensite

does not increase obviously with longer soaking time under optical examination.

Martensite-Free Zone

The Martensite-Free Zone (MFZ) was observed at grain boundaries in the cases of further

quenching T i −102Cr 3 to -26◦C and -52◦C. No MFZ observed at lower further-quenching

temperatures in the case of T i −102Cr 3 (see Figure 5.10). The microstructure is shown in

Figure 5.9(c), 5.9(d) and 5.9(e), of which blank/white areas can be observed near the grain

boundary. This region is optically similar to the Precipitate-Free Zone (PFZ) found in heat-


treatable aluminum alloys [47].

Nevertheless, in T i −1013 and T i −1023B , no MFZ was observed. Li et al. [19] has sug-

gested a possible reason of the formation of MFZ only in T i −102Cr 3 could be the limited

diffusion distance of chromium, the β stabilizer. Specifically, β-stabilizer will be expelled

from newly-precipitated α phase to the parent β phase. The diffusion distance of iron in

T i − 1013 and T i − 1023B is longer than chromium in T i − 102Cr 3, so that iron is more

likely to be more evenly distributed in β phase than chromium that is likely to accumulate

more in the region of parent β phase near α precipitates and, as a consequence, stabilize

this region. As a result, the region of β phase near α precipitates that is oversaturated with

chromium will be too stable to exhibit martensitic transformation during further quench-

ing.

To validate this assumption, Electron Probe Micro-Analysis (EPMA) was performed re-

garding the MFZ. Results and discussion is presented in Section 5.4.

Specialty of T i −1023B

Comparing Figure 5.11(a) with 5.11(d), 5.11(e), 5.12(b) and 5.12(c), it can be conspicuously

seen that the capacity of T i −1023B of exhibiting low-temperature martensitic transforma-

tion is tremendous. At as-quenched state, no martensite formed (Section 4.5), the amount

ofα precipitate is minor, and thereby more retain β phase. Under further quenching to low

temperature, a huge amount of acicular martensite formed throughout the grains. There-

fore, T i −1023B is the one with the largest capacity of exhibiting low-temperature marten-

sitic transformation among the three alloys.

SIM capability

The designed microstructure of the three alloys are believed to be capable of exhibiting the

SIM behavior – i.e. further martensitic transformation driven by mechanical driving force

– resulting in an improvement of ductility. From the work by Li et al. [19] and Neelakantan

et al. [5], the three alloys going through β+(α+β) heat treatment schemes can successfully

show SIM effect as long as there is enough retained β phase available for the transforma-

tion. T i −1023B is likely to exhibit the SIM behavior with a greater extent in comparison

with T i −1013 and T i −102Cr 3, because of the huge amount of retained β in this alloy.


(a) 900◦C/15mi n + 780◦C/2h + WQ (initial state)

(b) -15◦C/60mi n (c) -15◦C/60mi n

(d) -52◦C/1mi n (e) -52◦C/60mi n

Figure 5.7: The microstructure of T i −1013 of the further-quenching experiment.


(a) -100◦C/1mi n (b) -100◦C/60mi n

(c) -160◦C/1mi n (d) -160◦C/60mi n

(e) -196◦C/1mi n (f) -196◦C/60mi n

Figure 5.8: The microstructure of T i −1013 of the further-quenching experiment.


(a) 900◦C/15mi n + 780◦C/2.5h + WQ (initial state)

(b) -15◦C/60mi n (c) -26◦C/60mi n

(d) -52◦C/1mi n (e) -52◦C/1mi n

Figure 5.9: The microstructure of T i −102Cr 3 of the further-quenching experiment.


(a) -100◦C/60mi n (b) -160◦C/60mi n

(c) -196◦C/1mi n (d) -196◦C/60mi n

Figure 5.10: The microstructure of T i −102Cr 3 of the further-quenching experiment.


(a) 900◦C/15mi n + 780◦C/2h + WQ (initial state)

(b) -15◦C/1mi n (c) -26◦C/60mi n

(d) -52◦C/1mi n (e) -52◦C/60mi n

Figure 5.11: The microstructure of T i −1023B of the further-quenching experiment.


(a) -160◦C/1mi n

(b) -196◦C/1mi n (c) -196◦C/60mi n

Figure 5.12: The microstructure of T i −1023B of the further-quenching experiment.

5.3. X-ray diffraction 53

Table 5.2: Conditions of T i −1013 samples subjected to XRD testing; the heat treatment state of Sample 1 is

the starting state of the further quenching to low temperature of Sample 2 to Sample 7; WQ stands for water

quenching to room temperature.

Numbering Conditions

Sample 1 900◦C/15mi n + 780◦C/2h + WQ

Sample 2 -15◦C/60mi n






Sample Ref 900◦C/15mi n + WQ

5.3. X-ray diffractionTo further identify the phases formed during low-temperature martensitic transformation,

the XRD testing was carried out on T i −1013, the most tunable alloy on the basis of optical

microscopic examination.

5.3.1. Materials

To obtain better X-ray diffraction pattern, plate samples with 20mm × 20mm × 4mm di-

mension were employed. The big surface area of this geometry ensures more grains will be

exposed to the X-ray beam for better statistics.

Samples of eight conditions were heat-treated, further-quenched (see Table 5.2) and

electropolished in the same methods as used in the rest of this work. Specifically, Sample 1

is the initial state for the further-quenching treatments for Sample 2 to Sample 7. Sample

Ref is the homogenized and water-quenched state.

5.3.2. Experimental details

Bruker D8 Advance diffractometer Bragg-Brentano geometry and Lynxeye position sensi-

tive detector were utilized with Cu Kα radiation, 45kV and 40m A. Sample was spinning

during the testing. Divergence slit is V 20, scatter screen height is 5mm. Because the grain


size of the samples is relatively large, approximately in the order of 200 to 600µm, large

sample and large beam sizes were chosen. The V 20 setting (= 20mm irradiated length)

was chosen to maximize the beam size in order to get as many grains as possible in the

beam. For 2θ higher than 70◦, the optics cannot maintain this 20mm setting, and it grad-

ually reduces till approximately 12mm at 135◦. The range of coupled 2θ scan was between

25◦ to 135◦, step size 0.0348◦; counting time per step is 0.5s. Furthermore, X-ray diffraction

data was analyzed and evaluated by Bruker software Diffrac.EVA v4.1.

5.3.3. XRD results and discussions

The X-ray diffraction pattern of the samples are given in Figure 5.13. In order to make the

weak peaks in the pattern more visible, a portrait XRD pattern is also given in Figure 5.14.

The obtained pattern were compared with ICDD (International Center for Diffraction Data)

pattern to identify phases, and the lattice parameters of phases were tuned (by using the

tuning-cell function) to match the diffraction peaks, as listed in Table 5.3 together with the

ICDD patterns used.

It can be clearly seen from the XRD pattern that three phases formed in the samples: β

phase, α phase and α′′ martensite. There is no extra peak left for α′ martensite. Given that

both α and α′ are of hcp structure, while analyzing the pattern, the hcp phase from PDF

04-004-8449 was regarded as α phase because, on one hand, α phase exists in the samples

due to annealing in α+β two-phase region for sure, and on the other hand the Sample Ref,

which was directly quenched from β single-phase region and thus is not supposed to have

α phase precipitated, does not show obvious peaks on the regions of α peaks of Sample 1

to Sample 7.

However, from the work on the same alloy by Li. et al. [42], the martensite formed if

quenching directly after solution treatment was verified as α′ via X-ray diffraction.

The possible reason of the different kinds of martensite could be the extent of oversatu-

ration of martensite. Specifically, with increasing solute content, the hcp α′ martensite will

lose its hexagonal symmetry and transform into orthorhombic α′′ martensite [15][16][17].

The solute partition owing to intercritical annealing will make β phase oversaturated with

β-stabilizers, which may in turn induce the oversaturation of martensite that causes the

martensite formed when quenching from two-phase region at equilibrium heat treatment


Table 5.3: Used lattice parameters that matches the diffraction pattern of T i −1013 samples.

Lattice parametersPDF2

(Å)

β a 3.246 PDF 00-044-1288

αa 2.937

PDF 04-004-8449c 4.693

α′′a 3.018

PDF 04-017-4093b 4.897

c 4.629

state to be α′′ in stead of α′.

The peaks of α′′ martensite are generally broad in comparison with those of β and α

phases. This phenomenon has also been seen in the case of the α′ and α′′ martensites

from the work on T i − 1023 [5][42]. The typical reasons for this phenomenon are micro

strain (deformed crystal lattice) and/or very small grain size. In this case, micro strain is

the likely reason owing to the displacive nature of martensitic transformation that causes

the distortion of crystal lattice. According to the optical micrographs in this thesis work,

the acicular martensite is very thin, of the order of 1 to 3µm, so this is not small enough to

cause such an effect.

However, quantifying phases from the pattern is practically difficult as the texture in

samples is still rather strong. In theory, if the texture is strong, part of the sample would be

etched perfectly that the microstructure can be effectively unveiled, whereas another part

of the sample cannot be effectively etched, which corresponds well with the observation

of some samples from this work. For instance, in Figure 5.15, where the low-temperature

martensitic transformation took place already, the difference of the density of the acicular

martensite at different regions of the sample is obvious. Some grains with high-density

acicular martensite, yet some other grains showing almost no acicular phases. At this point,

the fact that some grains were not showing obvious acicular martensite does not mean that

they did not realize the transformation, but the reason is the existence of strong texture that

results in grains with different responses to chemical etching.


Fig

ure

5.13

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ayd

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tern

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Figure 5.14: X-ray diffraction pattern of T i − 1013 (Ti-10V-1Fe-3Al) further quenching in selected cooling

baths, with Cu Kα radiation at room temperature; 2θ range: 25◦ to 135◦; wavelength: 1.54060; step size:

0.0348◦ 2θ, counting time per step: 0.5s.


Figure 5.15: The microstructure of T i −1013 showing the effect of texture in response to chemical etching;

heat treatment state: further quenching to -15◦C for 60mi n starting from the designed heat treatment state.

5.4. Electron probe micro-analysis 59

5.4. Electron probe micro-analysisIn order to study the formation of Martensite-Free Zone (MFZ) formed in T i − 102Cr 3,

Electron Probe Micro-Analysis (EPMA) was conducted to investigate the relation between

MFZ and the accumulation of β-stabilizers.

5.4.1. Experimental details

T i −102Cr 3 sample at designed heat treatment state and being further quenched to -26◦C

for 60mi n (Figure 5.9(c)) was selected.

The measurements were performed with a JEOL JXA 8900R microprobe using an elec-

tron beam with energy of 15keV and 20n A beam current by means of Wavelength Dis-

persive Spectrometry (WDS). The composition at each analysis location of the sample was

determined by using the X-ray intensities of the constituent elements after background

correction in relation to the corresponding intensities of reference materials. The inten-

sity ratios of elements were processed with a matrix correction program CITZAF [48]. The

points of analysis were located along line, about 108µm long with increments of 1µm. Ele-

ments under analysis were aluminum, vanadium, chromium and titanium.

5.4.2. EPMA results and discussions

The result of EPMA measurement is given in Figure 5.16. The line scan was performed as

indicated in Figure 5.16(a).

From the aluminum, vanadium, and chromium line scan profile in Figure 5.16, no obvi-

ous gradient of alloying elements can be observed across the MFZ region, except a sudden

variation of the concentrations of elements at the grain boundary. Thus, the formation of

MFZ is not due to the limited diffusion distance of chromium that stabilizes the region of

retained β phase near α precipitates as suggested in previous work [19].


(a) The position of the line scan (blue line and ending at the arrow head)

(b) The composition of the elements Al, V, Cr and Ti along the line scan (in w t%), an arbitrary MFZ region in

accordance with Figure 5.16(a) indicated

Figure 5.16: EPMA line scan of MFZ in T i −102Cr 3 at heat treatment state 900◦C/15mi n + 780◦C/2.5h + WQ

+ -26◦C/60mi n.

5.5. Conclusion of low-temperature experiments 61

5.5. Conclusion of low-temperature experimentsThe designed heat treatment scheme for tuning the β phase stability of T i − 1013, T i −102Cr 3 and T i−1023B was tuned out to be effective in terms of exhibiting low-temperature

martensitic transformation.

The ranges of low-temperature martensitic transformation starting temperature for the

three given metastable β titanium alloys at the designed heat treatment state are: room

temperature to -15◦C for T i −1013, -15◦C to -26◦C for T i −102Cr 3, and -26◦C to -52◦C for

T i −1023B . Further, the transformation is isothermal martensitic transformation.

Meanwhile, Martensite-Free Zone (MFZ) was found in certain further quenching con-

ditions just below the starting temperature in the case of T i −102Cr 3. As experimentally

verified via EPMA, the limited diffusion distance of chromium is not the reason of the for-

mation of MFZ.

On the basis of previous work, the three alloys with designed microstructure are ex-

pected to exhibit the SIM behavior at room temperature under mechanical driving force.

The capacity of T i −1023B to exhibit low-temperature martensitic transformation is high,

in comparison with the other two alloys, because conspicuously more retainedβphase was

obtained from designed heat treatment scheme. In other words, the stability of retained β

phase of T i −1023B is higher than the other two alloys.

In the case of T i − 1013, three major phases – β, α and α′′ – were identified by X-ray

diffraction. Because of the effect of strong texture in the samples, the expected trend ac-

cording with the extent of low-temperature martensitic transformation was not visible in

the XRD pattern. The martensite formed in response to quenching from two-phase re-

gion is verified as α′′ martensite. This may mainly be because of the high solute content in

martensite.

In comparison with the behaviors simulated by Thermo-Calc®, at designed heat treat-

ment state for the three alloys, in practice, the extent of solute partitioning during α phase

precipitation is more intensive, and the retained β phase is more stable.

EDM process for machining the heat-treated samples into disk required by DSC instru-

ment was found to be counterproductive as the process seems to alter the designed mi-

crostructure. As a consequence, the effectiveness of DSC in terms of following and moni-

toring the transformation process is yet to be verified.

6Conclusion

This work effectively designed the microstructure of three metastable β titanium alloys –

T i −1013, T i −102Cr 3 and T i −1023B – via β+ (α+β) heat treatment (i.e. 900◦C/15mi n +

780◦C/2h~2.5h + WQ) and had the stability of β phase tuned for further martensitic trans-

formation under thermal driving force (low-temperature martensitic transformation).

As for β+ (α+β) heat treatment, 2 to 2.5 hours annealing at α+β two-phase region

(780◦C) is concluded to be effective in realizing the equilirbium state, stabilizing β phase,

and securing enough retained β phase at room temperature.

The ranges of starting temperatures of low-temperature martensitic transformation of

the three metastable β alloys at the designed heat treatment state were found – i.e. room

temperature to -15◦C for T i − 1013, -15◦C to -26◦C for T i − 102Cr 3, and -26◦C to -52◦C

for T i − 1023B , and the transformation of the three alloys was found to be of isothermal

character.

The three alloys with designed microstructures are expected to exhibit SIM behavior

that can improve the alloys’ ductility. Among the three alloys, T i −1023B showed the most

dramatic capacity to exhibit the low-temperature martensitic transformation owing to the

higher stability of the β phase and more retained β phase at room temperature in compar-

ison with the other two alloys.

Through X-ray diffraction, the martensite of T i−1013 formed in response to quenching

63

64 6. Conclusion

from two-phase region at equilibrium heat treatment state is the orthorhombicα′′ marten-

site. In theory, the occurrence of α′′ martensite is closely related to oversaturated marten-

site with the high solute content. Meanwhile, the effect of heavy alloying is likely to also

contribute to large MS −M f interval.

Martensite-Free Zone (MFZ) was found in the case of T i − 102Cr 3 if the alloy is fur-

ther quenched to a temperature just below the starting temperature of low-temperature

martensitic transformation. EPMA testing proved that this phenomenon does not relate to

the diffusion behavior of alloying elements.

Thermo-Calc® predicts less intensive solute partitioning during α phase precipitation

and therefore less stable retained β phase for the three alloys.

Electrical Discharge Machining (EDM) was found detrimental to the three heat-treated

metastable titanium β alloys as the designed microstructure was altered during the ma-

chining. EDM is likely to induce such a heat treatment process that heats up the sample to

β single-phase region and then quenches to room temperature.

7Recommendations for Future Work

As for the future work, the effectiveness of DSC in terms of detecting and following the low-

temperature martensitic transformation in metastable β titanium alloys is worthwhile to

be further investigated. The main obstacle is the preparation of DSC sample (small disk).

An effective technique/machining that will not change the designed microstructure is rec-

ommended.

Having investigated the effect of thermal driving force by means of low-temperature

martensitic transformation, the mechanical properties of further-quenched samples could

be examined in comparison with their initial as-quenched state at room temperature. It

would be valuable to see if the low-temperature martensitic transformation could bring

about effective strengthening effect. Furthermore, the formation of MFZ in T i −102Cr 3 at

low-temperature condition is worth to be further investigated.

The idea of making titanium metallic composite by combining the SIM effect with Hall-

Petch relation, as a new route to strengthen titanium alloy, is proposed in Appendix A,

which requires further effort to bring the idea into reality.

65

AFabricating Novel Titanium Metallic

Composites by SPS

In this proposal, general characteristics and merits of Additive Manufacturing (AM) by Se-

lective Laser Melting (SLM) and Spark Plasma Sintering (SPS) to fabricate titanium prod-

ucts from titanium powder are discussed in Section A.1.

Subsequently in Section A.2, the project idea is proposed, theoretical background of

why metastable titanium alloy is used is given, influence of using metal powder (smaller

grain size) instead of traditional casting metal is analyzed, and an in-situ heat treatment

scheme during the course of SPS are suggested.

The outcomes of this research are anticipated in Section A.3.

A.1. IntroductionFeatured by relatively low density, good corrosion resistance, high yield strength, satisfac-

tory high-temperature strength and good creep behavior, titanium and its alloys are impor-

tant structure materials widespread utilized in aerospace, medical implantation, chemical,

marine, offshore and other industries.

Nevertheless, the overhead and direct costs of smelting and processing titanium and its

alloys, due to their inherent chemical properties and crystallographic structures, still hover

67

68 A. Fabricating Novel Titanium Metallic Composites by SPS

at high level and thus considerably impede the development and widespread application

of titanium alloys, a far cry from their steel or aluminum counterparts.

The awkward situation of titanium industry, on the other hand, makes exploring new

processing route to manufacture titanium alloys a worthy endeavor, i.e. techniques that

could lower feed stock materials cost and achieve near net shape [49].

Given the considerations above, powder-based Additive Manufacturing (AM) by Selec-

tive Laser Melting (SLM) has recently been regarded as a promising near-net-shape fabrica-

tion technology. Based on Computer-Aided Design (CAD) models, metal powders on plat-

form are directly melted by high-energy laser beam, ending up with the geometry buildup

in a layer-by-layer manner [50][51].

Previous experience has shown that, as far as the most widely used T i −6Al −4V was

concerned, the products built by SLM could be able to achieve a fully dense state. Mean-

while, they exhibit comparable strength and ductility in comparison with their wrought

counterparts [52][53].

Unlike conventional way of metal making, several parameters play significant role in

AM techniques in determining the final properties of products – specifically, powder mor-

phology, particle size distribution, purity, flowability, powder bed condition, laser spot size,

scanning speed, preheat temperature, the formation and development of melt pool, and

homogeneity [52][54][55][56][57][58].

Nevertheless, more importantly, most published scientific works to date are focusing on

additively manufacturing titanium alloys by selective laser melting single type of titanium

alloy powder, whereas the possibility and effectiveness of doing this by mixing different

titanium alloy powders have not been extensively investigated and studied yet.

Therefore, inspired by the work done by using SLM, a novel idea of fabricating metallic

composite from two different alloy categories of titanium powder by Spark Plasma Sinter-

ing (SPS) is going to be proposed in this work. The idea of utilizing SPS technique is to try

out the idea of mixing two different titanium alloy powders for better mechanical proper-

ties on SPS first, and then further realize it on AM techniques represented by SLM.

A.2. Research work 69

A.2. Research work

A.2.1. Project idea

The main goal of employing SPS to fabricate two different titanium powders – Ti-10V-2Fe-

3Al (transformable grade) and the most widely used Ti-6Al-4V (matrix) – together is to re-

alize the novel idea of composite microscopically in traditional metal alloys. Specifically,

the Stress-Induced Martensitic transformation effect (abbreviated as SIM effect) of Ti-10V-

2Fe-3Al is going to be employed as the reinforcement of this "metallic composite". And

the matrix is Ti-6Al-4V that does not exhibit SIM effect but provide homogeneity for the

metallic composite.

Meanwhile, granted the inherent features of powder metallurgy, finer grain size (deriv-

ing from the minimal size of metal powders) is expected to be realized to such an extent

that, in the light of Hall–Petch relation [59], the resulting products are expected to end up

with high strength.

Furthermore, one of the biggest advantages of SPS is that the density of the product will

be near the theoretical density at lower sintering temperature compared to conventional

sintering techniques, as the heat generation is internal, different from conventional hot

pressing, because of Joule heating.

Given that SPS and also SLM can be performed at elevated temperature, heat treatment

can be carried out simultaneously during the course of production. More imperatively,

heat treatment is an important means to tune and control SIM effect for the powder of

transformable grade, such that by preheating powder bed, practitioners are spared the core

of doing heat treatment separately. In the meantime the coarsening effect is anticipated to

be not very serious and obvious.

On the other hand, the cooling rate in SPS is estimated to be up to 103K mi n−1, much

lower than the critical cooling rate for martensitic transformation in Ti-6Al-4V (410K s−1)

[60][61]. However, the cooling rate of SLM is 103 to 105K s−1, which is orders of magnitude

higher than 410K s−1. Therefore, as a consequence, the product built by SPS is expected to

have equilibriumα+βmicrostructure, whereas acicular martensite (α′ and/orα′′) together

with α phase for that built by SLM.


A.2.2. Theoretical basis of applying SIM effect as reinforcement

β phase, comparable to austenite in steels, is the most predominant phase while looking at

martensitic transformation in titanium alloys. Inspired by previous works upon T i −1023

[5][18][21][32][62][63][64], metastable β titanium alloys have the potential to exhibit both

normal martensitic transformation and SIM (or TRIP) transformation (β→ α′/α′′) under

different conditions by careful alloy design and treatments.

Theoretically, martensitic transformation requires energy input to overcome the activa-

tion barrier and to be triggered by means of either quenching or mechanical deformation.

In other words, the more stable the β phase, the lower the MS temperature, and the more

the energy (either thermal or mechanical) is required to initiate and trigger martensitic

transformation. Therefore, the phase transformation and mechanical behavior of metastable

β alloy can be tuned by tailoring the stability of β phase.

Influencing factors for SIM effects

From previous works done by varied researchers, seven main factors affecting SIM effects

have been recognized as following:

• Stability of β phase (molybdenum equivalence Moeq )

• MS temperature

• Volume fraction of α phase

• Grain size of β phase

• Morphology of α phase

• Deformation temperature

• Strain and strain rate

Of these factors, the primary criteria supposed to be taken into account while designing

are β phase stability (Moeq ) and MS temperature. The rest of influencing factors, such as

volume fraction ofα phase, have direct but complicated relation with β phase stability and

MS temperature owing to solute partitioning effect.


Impacts of small grain size deriving from using metal powders

Grain size, as one of the most important parameters in metals science, is well-known by

its potent impact on mechanical properties of metals based on Hall-Petch relation. It also

has significant influence on the triggering stress to initiate TRIP effect. In titanium alloys,

linear relations between grain size and triggering stress have been reported in many works

– i.e. the larger the grain size, the higher the triggering stress [32][65]. Liu and Yang [66]

proposed a thermodynamic expression for the transformation from metastable β phase to

orthorhombic martensite:

∆G =∆H −T∆S +∆Eel +δEi r r −εtr − 1

2[ fσ2

EM− (1− f )

σ2

Eβ] (A.1)

In Equation A.1, enthalpy ∆H and entropy ∆S are not dependent upon grain size [67].

εtr has to do with texture and is the force displacement work because of lattice distortion

during martensitic transformation; the complex term 12 [ f σ2

EM− (1− f )σ

2

Eβ] is elastic energy

due to deformation, which is related to the difference in Young’s modulus between α′′ and

β during elastic deformation [68]. They are independent on grain size as well.

However, stored internal elastic energy ∆Eel and irreversible energy for the work done

in undergoing internal friction δEi r r depend on β grain size [32][68]. Since β grain size

directly relates to mean length and width of martensite plates, and martensite plate length

has a predominant influence on the magnitude of internal stress field in β matrix during

SIM transformation, it can be proposed that the smaller the grain size, the lower the stored

internal elastic energy ∆Eel [69].

Friction resistance δEi r r is in direct proportion to the total volume the interface sweeps.

Therefore, provided given volume of martensite, the smaller the plate size, the larger the

volume swept by interfaces and the higher the δEi r r [68][70].

A positive linear tendency was found between β grain size and SIM effect triggering

stress in Ti-1023 [32][65]. Implied is that as far as the overall effect was concerned, with

decreasing β grain size, the SIM triggering stress decreases, and ∆G in Equation A.1 (driv-

ing force) therefore decreases as well. Thus, the contribution of term (∆Eel +δEi r r ) shows

positive response with regard to changing β grain size.

Therefore, if other variables remain the same, a simplified relation is proposed in this

literature review between β phase grain size, molybdenum equivalence value of β phase


Moβeq (or α phase volume fraction), free energy change ∆G , SIM triggering stress σSI M and

MS temperature (in which " ⇑ " means changing positively and " ⇓ " means changing nega-

tively):

• β grain size ⇑ ∆G ⇑ σSI M ⇑ MS ⇓

• Moβeq ⇑ ∆G ⇑ σSI M ⇑ MS ⇓

As a result, the smaller the β grain size, the smaller the ∆G , the lower the σSI M , and

the higher the MS – ending up with less stable β phase. And thus smaller β phase grain

size resulting from additive manufacturing will render martensitic transformation easier to

occur.

Tailoring MS temperature

β phase stability is quantified by Molybdenum equivalence Moeq ; and Moeq is directly re-

lated with MS (in Kelvin) by β phase composition according to Equation A.2 [4] and Equa-

tion A.3 [5][33]:

Moeq = 1.00Mo +0.28N b +0.22Ta +0.67V +1.6Cr +2.9Fe −1.00Al (A.2)

MS [w t%] = 1156−150Fe −107Mn −96Cr −67Ni −49Mo

−41Cu −37V −17N b −7Z r +15Al (A.3)

Heat treatments are supposed to be designed in such a way that β phase is tuned to a

stability (indicated by the value of Moeq ) that the martensitic transformation will take place

in predicted ways, and enough retained β phase must be secured at room temperature for

SIM transformation.

The intercritical annealing temperature is recommended to be set at the level where the

MS of metastable titanium alloy can be lowered to around room temperature. Alternatively,

in SPS, at least should enough retained β phase be secured for the SIM effect.


A.2.3. Selection of powder bed preheating temperature

The selection of preheating temperature mainly zeros in on tuning the metastable titanium

alloy powders for SIM effect, since it is practically and theoretically difficult to simultane-

ously tailor two different categories of titanium alloys.

The preheating treatment scheme (for both SLM and SPS) suggested here strives to

strike a balance of impacts of heat between matrix and reinforcement titanium alloy pow-

ders.

Influence of preheating on matrix Ti-6Al-4V

The MS of Ti-6Al-4V in full β condition is calculated as 825◦C. Since the preheating tem-

perature will be selected between 700 and 800◦C for tuning the stability of β phase in

transformable grade Ti-10V-2Fe-3Al, martensitic transformation is unavoidable in matrix,

whereas according to the data given in [60], SLM-manufactured Ti-6Al-4V is still able to

retain a yield strength of 900 to 1000MPa with elongation 8.5 to 9%.

Influence of preheating on reinforcement Ti-10V-2Fe-3Al

In comparison, the metastable titanium alloy Ti-10V-1Fe-3Al is unlikely to exhibit marten-

sitic transformation immediately at a temperature above 700◦C; and by means of annealing

and thus α participation owing to powder bed preheating, we would be able to tailor the β

phase composition and thus β stability and MS through controlling the precipitation of α

phase.

A.2.4. Controlled parameters

During real experimentation, several technical parameters are worthwhile to be consid-

ered.

Volume fraction of the transformable grade

The volume fraction of the Ti-10V-2Fe-3Al powder (transformable grade) would be well-

advised to be less than 50%. The optimal range of the volume fraction for transformable

grade powder is recommended to be between 10% to 40%, as enough volume fraction

should be given to the matrix powder that ensures the structural stability and homogene-

ity for the designed metallic composite. The anticipitation, of course, needs to be further


examined by laboratory work.

Powder particle size

Secondly, the original particle size is of primary importance in determining the final me-

chanical properties of the metallic composite. The powder size of transformable grade is

recommended to be at least two times bigger than the matrix powder, because SIM effect is

expected to take place in the transformable grade and this would not be obviously effective

if the powder size, and thus the room available for SIM effect, of transformable grade is too

small.

As for the powder size of matrix alloy, to obtain high product strength on the basis of

Hall-Petch relation [59], small particle size is preferred. Considering the commercial avail-

ability in the market, a mesh size range between 100 and 325 (opening: 44 to 150µm) is

recommended.

Furthermore, in real experimentation, the size of particle utilized is supposed to be

stringently controlled, because the over-irregular powder particle size will have severe ad-

verse effect on the formation of voids, surface finish, the achieved mechanical property of

the product, etc. [49]. Thus, the applied particle size range should be controlled to be as

narrow as possible.

Effects of changing composition during sintering

Lastly, it can be anticipated that, during the course of consolidating the two alloys together

at high temperature, the local composition of powders may change. However, we cannot

predict the extent at current stage, and therefore, we cannot take this effect into account.

A.3. Expected research outcomesEventually, the Ti-10V-2Fe-3Al-plus-Ti-6Al-4V metallic composite made by SPS, as a trial

examination of the idea of mixing two different titanium alloy powders, is expected to have

continuous Ti-6Al-4V matrix in combination with Ti-10V-2Fe-3Al reinforcement exhibiting

SIM effect during deformation.

The final product’s ductility is expected to be better than additive-manufactured Ti-6Al-

4V product owing to the SIM behavior of the transformable grade Ti-10V-2Fe-3Al, while at

the same time high strength of the product would be ensured. The SIM martensite is ex-

A.4. Supplier information of titanium powders 75

pected to occur in the product in a dispersed and isotropic manner with acicular morphol-

ogy, as the transformable grade powder is thoroughly mixed with matrix powder.

Furthermore, some new interesting phenomenon are expected, for instance the inter-

action (e.g. stress field) between martensites from two Ti alloys.

A.4. Supplier information of titanium powders

Matrix Ti-6Al-4V

The commercial Ti-6Al-4V powder is globally available on the market. Taking the supplier

TIMET in the U.S. for instance, their quote is US$3,700.00 for 201b (approx. 9kg ); mesh

range can be chosen between -100 to +325 (minimum particle size achievable: 44µm).

Transformable grade Ti-10V-2Fe-3Al

Ti-10V-2Fe-3Al powder is not available on the market. The most viable solution so far is to

buy the alloy bar from suppliers, such as TIMET, and request company with toll atomization

technique to make the bar into powder.

The author has found TLS Technik GmbH & Co. Spezialpulver KG in Germany (contact

person: Norbert Ludwig; Email: [email protected]) who is willing to offer toll

atomization if the alloy as bar feedstock is provided. The average particle size from their

atomization is 80µm, 90% in the range between 0 to 150µm. The pricing for a quantity up

to 10kg is d1,550.00.

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