The role of heat treatment on microstructure and mechanical properties of Ti–13Zr–13Nb alloy for biomedical load bearing applications

J O U R N A L O F T H E M E C H A N I C A L B E H AV I O R O F B I O M E D I C A L M AT E R I A L S

4 (2011) 1132–1144

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Research paper

The role of heat treatment on microstructure and mechanical properties of Ti–13Zr–13Nb alloy for biomedical load bearing applications P. Majumdar ∗ , S.B. Singh, M. Chakraborty 1 Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur, India

A R T I C L E

I N F O

A B S T R A C T

Article history:

The suitability of heat treated Ti–13Zr–13Nb (TZN) alloy for biomedical load bearing

Received 28 September 2010

applications has been investigated. Depending upon the heat treatment conditions, the

Received in revised form

microstructure of TZN alloy mainly consists of α, β or α′′ martensite phases. In general, for

7 March 2011

all the deformation and solution treatment temperatures the variation of the hardness and

Accepted 18 March 2011

tensile strength with cooling rate is similar. The elastic modulus of TZN alloy decreases

Published online 25 March 2011

with an increase in cooling rate from the solution treatment temperature. Relatively fine α + β microstructure increases the hardness and tensile strength. The presence of

Keywords:

martensite and/or retained β in the microstructure decreases the hardness and elastic

Titanium alloy

modulus and increases the ductility substantially whereas higher amount of α phase in

Microstructure

the matrix increases the elastic modulus. Decomposition of martensite and retained β

Hardness

into α phase during aging increases the hardness, elastic modulus and tensile strength

Elastic modulus

and decreases the ductility. Among the samples studied, the aged TZN sample, which

Tensile properties

was deformed and solution treated at 800 ◦ C followed by water quenching, is a promising candidate for the application as implant material. c 2011 Elsevier Ltd. All rights reserved. ⃝

1.

Introduction

The ideal biomaterial for orthopedic implant applications, especially for load bearing joint replacements, is expected to exhibit excellent biocompatibility with no adverse cytotoxicity, excellent corrosion resistance, and a good combination of mechanical properties such as high strength and good fatigue resistance, low elastic modulus, sufficient ductility, and good wear resistance (Banerjee et al., 2005; Long and Rack, 1998). Titanium and its alloys have become one of the most attractive classes of biomedical implant ∗ Corresponding author. Tel.: +91 3222 283290; fax: +91 3222 282280. E-mail address: [email protected] (P. Majumdar). 1 Indian Institute of Technology, Bhubaneswar, India. c 2011 Elsevier Ltd. All rights reserved. 1751-6161/$ - see front matter ⃝ doi:10.1016/j.jmbbm.2011.03.023

materials. They are generally preferred to stainless steels and Co–Cr alloys because of their high specific strength to weight ratio, superior biocompatibility and corrosion resistance, good mechanical properties and low elastic modulus (Hao et al., 2002; Kobayashi et al., 1998; Majumdar et al., 2008a; Niinomi et al., 1999; Velten et al., 2003; Williams, 2001). Conventionally used titanium based biomaterials such as Ti–6Al–4V, Ti–6Al–7Nb and Ti–5Al–2.5Fe (in wt%) are now found to be unsuitable for biomedical applications because of the toxic effect of both Al and V and may cause long-term health problems (Nag et al., 2005). On the other hand, high

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modulus of elasticity of the implant materials as compared with that of bone causes “stress shielding effect” (Lin et al., 2005; Mändl and Ruschenbach, 2002; Okazaki et al., 1996). This can potentially cause bone resorption that eventually leads to failure of the implant (Banerjee et al., 2004; Niinomi, 1998). Therefore, presence of non-toxic elements and low modulus of elasticity are the two most important criteria for the development of materials for orthopedic applications. Low modulus Ti alloys can be developed by designing β-Ti alloys containing non-toxic alloying elements like Nb, Zr, Ta etc. (Akahori et al., 2005; Banerjee et al., 2005; Hao et al., 2002; Kuroda et al., 1998, 2005; Li et al., 2004; Niinomi, 2003; Niinomi et al., 2005; Sagaguchi et al., 2005; Takahashi et al., 2000; Zorn et al., 2005) and/or by the formation of porous structures (Arciniegas et al., 2007; Brailovski et al., 2011; Davis et al., 2001; Gibson and Ashby, 1997; Goi et al., 2008; Prymak et al., 2005; Wang et al., 2009; Wen et al., 2001, 2007). The β titanium alloys mainly consist of low modulus single β phase and therefore the modulus of these alloys is low (Boehlert et al., 2005; Boyer et al., 1994). On the other hand, the elastic modulus of the porous material is proportional to the elastic modulus of the cell edge material and the square of the relative density which is the ratio of the density of the porous material to that of the solid material (Wen et al., 2007). Thus, the modulus of the porous material is lower than that of the solid material. Titanium alloys are mainly used as replacement materials for hard tissues. Therefore, the elastic modulus is an important parameter, as a value close to that of the bone material leads to a better transfer of functional loads to the bone, enhancing the stimulation for new bone growth (Mändl and Ruschenbach, 2002; Okazaki et al., 1996). The elastic modulus of Ti and α + β Ti alloys (100–110 GPa) is much smaller than stainless steel (210 GPa) and Co–Cr alloys (204–240 GPa). However, it is significantly higher than that of bone tissue which has elastic modulus in the range of 10–40 GPa (Banerjee et al., 2004; Chang et al., 1998; Nelissen et al., 1995; Niinomi, 1998). In addition to a low modulus of elasticity, titanium alloys for load bearing biomedical applications should have adequate mechanical strength and wear resistance (Akahori et al., 2005). The yield strength (YS) of biomedical titanium alloys varies between 500 and 1000 MPa, tensile strength varies in the range of 900–1200 MPa and the elongation lies between 10% and 20% (Gasser, 2001; Niinomi, 1998). The mechanical properties depend on chemical composition and microstructure of the material (Nag et al., 2007). The mechanical behavior of α + β titanium alloys is influenced by the morphology of α and β phases and their relative amounts which can be controlled by applying proper heat treatment (Banerjee and Krishnan, 1981). Okazaki (2001) have investigated the mechanical properties of the Ti–15Zr–4Nb–4Ta alloy and concluded that uniform distribution of the fine α phase by solution treatment and aging leads to good balance of strength and ductility. Fatigue performance of an alloy is an important mechanical property to confirm its reliability as metallic biomaterials. Fatigue strength is essential for materials for medical implant applications such as bone plates, screws and nails, artificial spines, and artificial femoral and hip joints because they are used under fatigue conditions. These implants have to withstand not only one time peak stress, but also several

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million cycles of fluctuating loads which they usually experience during their lifetime (Akahori et al., 2005; Boehlert et al., 2005). The fatigue property of a material varies according to the microstructures obtained by heat treatment or thermomechanical treatment (Akahori et al., 2005; Long and Rack, 1998). Moreover, the deformation behavior of the alloys under cyclic loading is influenced by their surface condition (Leinenbach and Eifler, 2006). The other most important parameters in determining the suitability of a material for biomedical applications are its biocompatibility and corrosion resistance. The corrosion resistance of Ti and its alloys in chloride environment is found to be excellent and better than stainless steel and Co alloys (Strietze et al., 1998). This is because a protective oxide surface is formed on the Ti-based materials which is highly inert and reforms easily after damage (López et al., 2002; Manivasagam et al., 2003). The surface oxide layer also imparts tissue biocompatibility (Williams, 2001). The work presented here primarily aims at the synthesis of ‘near-β’ titanium alloys and proper design of the processing and heat treatment schedule of the alloys in order to produce an appropriate microstructure to provide suitable mechanical properties for implant applications. An in-depth study on the role of heat treatment on the microstructure of the alloys was considered to obtain the optimum combination of properties. The present work focuses on hardness, elastic modulus and tensile properties of the material. The influence of heat treatment on wear and fatigue behavior of Ti–13Zr–13Nb has been described in detail elsewhere (Majumdar et al., 2008a, 2010; Majumdar, 2009).

2.

Materials and methods

Ti–13Zr–13Nb (TZN) alloy (compositions in wt%) was prepared by arc melting with a non-consumable tungsten electrode in a high purity argon atmosphere in a vacuum arc melting unit supplied by Vacuum Techniques Pvt. Ltd., Bangalore. The differential scanning calorimetry (DSC) on the alloy was conducted using a SETARAM Labsys DSC supplied by SETARAM Instrument, France. The sample was heated from room temperature to 800 ◦ C and then cooled again to room temperature at a controlled heating/cooling rate of 10 ◦ C/min (Fig. 1). Peak of the curve was automatically estimated with the help of the equipment software package. The as-cast samples were given 30%–40% reduction by rolling at two different temperatures (800 and 650 ◦ C) and then air cooled to room temperature. The hot rolling temperatures were selected in such a way that it was above the β transition temperature in one case (800 ◦ C) and below that temperature in another (650 ◦ C). The hot rolled TZN samples were solution treated at 800 ◦ C (above β transus), 700 ◦ C (slightly below β transus) and 650 ◦ C (below β transus) for 1 h in a dynamic argon atmosphere; this was followed by furnace cooling (FC), air cooling (AC) or water quenching (WQ). It has been observed that high strength in ‘near β’ titanium alloys can be achieved if aging is done between 500 and 600 ◦ C for 4–6 h (Geetha et al., 2004). Based on this observation, water quenched samples were aged at 500 ◦ C for 5 h. The heat treatment schedule of TZN alloy is shown schematically in Fig. 2.

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Fig. 1 – Differential scanning calorimetric analysis of TZN alloy.

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Leco Corporation, USA, at an indentation load of 10 kg and a dwell time of 10 s. Hardness measurements were done on the polished surface of the specimen. At least eight indentations were taken for each specimen and the average values were reported. The modulus of the heat treated samples was measured by ultrasonic method using an ultrasonic velocity gauge, 35 DL, Panametric, USA. A normal incident probe of 5 MHz (model: M110) and a shear probe of the same frequency (model: V221) were used for the measurement of normal and shear velocities of the wave, respectively. The density of the samples was measured using Archimedes principle. The principle of measurement of elastic properties is briefly described below. Young’s modulus (E) is a measure of the stiffness of a material. Within the limits of elasticity, the ratio of the linear stress to the linear strain is termed as the modulus of elasticity or Young’s modulus and is expressed as (Majumdar et al., 2008b; Raj et al., 2003) E = ρV2S (3VL2 −4V2S )/(VL2 −V2S )

The average composition of the alloys was determined by scanning electron microscopy-energy dispersive spectroscopy (SEM-EDS) on a Scanning Electron Microscope (Model: JSM-5800, JEOL, Japan), attached with Energy Dispersive X-ray system (Model: ISIS 300, Oxford Instruments Limited, UK). The operating voltage of the SEM was 20 kV. For chemical analysis of the alloy, a number of regions were examined for each alloy in order to have an average composition. In addition, interstitial gas analyses were carried out using Laboratory Equipment Corporation (LECO) analyzers in Wah Chang lab, USA in order to measure the oxygen, nitrogen, hydrogen and carbon content of the alloy. Microstructure analysis of the heat treated samples was carried out on Leica-DMLM optical microscope and scanning electron microscope (SEM), FEI-Nova Nanolab 200 with FEGSEM column operating at 20 kV. For this purpose the metallographically polished samples were etched with Kroll’s reagent (10 vol% HF and 5 vol% HNO3 in water). Room temperature X-ray diffraction analysis was carried out on an X-ray Diffractometer, Philips, Holland, PW 1710 with Cu Kα radiation at 40 kV and 20 mA. The scanning rate was kept at 3◦ − 2θ/60 s. The Vickers hardness measurements were carried out on a Diamond Vickers Hardness tester, LV 700, supplied by

where VL and VS are the ultrasonic longitudinal and shear wave velocities, respectively and ρ is the density of the material. Tensile testing of the specimens was carried out on an Instron Testing Machine (Static, Model 4204) using appropriate static extensometer for elongation measurements. The procedure specified in ASTM standard E8 was followed. The cross head speed was kept at 8.33 × 10−6 m s−1 . The tensile specimen dimensions are given in Fig. 3. After tensile testing, the fractured surfaces were examined on a scanning electron microscope (Model: JSM-5800, JEOL, Japan) at 20 kV operating voltage.

3.

Results

The composition, including interstitial content, of the investigated TZN alloy, obtained by SEM-EDS studies and LECO analysis, is given in Table 1. The average composition (in wt%) of the TZN alloys can be expressed as Ti–13Zr–13Nb. From the DSC curve, the β transition temperature, which represents the lowest temperature during heating above which the material is completely β, was determined to be 740 ◦ C (Fig. 1). Based on this transition temperature, the heat treatment was performed on TZN alloy.

Fig. 2 – Schematic diagram illustrating the thermo-mechanical treatment of TZN alloy.

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Fig. 3 – Dimension of sub-size round tensile specimen proportional to the ASTM E-8 standard.

Fig. 4 – Microstructures of the TZN alloy deformed and solution treated for 1 h at (I) 800 ◦ C and (II) 650 ◦ C followed by (a) furnace cooling, (b) air cooling and (c) water quenching.

Table 1 – The chemical composition (wt%) of the TZN alloy. Ti 73.42

3.1.

Nb

Zr

O

13.41

13.17

0.12

C 0.0122

N 0.0047

H 0.012

Microstructure

In the case of samples hot rolled and solution treated at 800 ◦ C, i.e. above the β transus, the β phase transforms to the α phase on furnace cooling or air cooling through the transus. The microstructure of furnace cooled sample consisted of a basket-weave structure (Widmanstätten lath like α) formed from prior β grains (Fig. 4(I)(a)). Multiple variants of α lath

were observed within prior β grains. In air cooled sample, fine Widmanstätten α laths were observed within pre-existing β grains (Figs. 4(I)(b) & 5(a)). Water quenching from 800 ◦ C resulted in the formation of martensite and retained β (Figs. 4(I)(c) & 5(b)). When the hot rolling was performed above β transus (800 ◦ C) and subsequently solution treatment was done well below β transus (650 ◦ C), the presence of primary α and transformed β within pre-existing β phase was observed for furnace cooled sample whereas air cooled sample showed similar microstructure on a relatively fine scale. Martensite was not observed in the water quenched sample, only primary α and retained β were found in the microstructure. In the case of TZN samples subjected to deformation at 650 ◦ C (well below β transus) and solution treatment at 700 ◦ C

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Fig. 5 – Scanning electron micrographs of the TZN alloy deformed and solution treated at 800 ◦ C for 1 h followed by (a) air cooling (AC) and (b) water quenching (WQ).

Fig. 6 – Microstructures of the TZN alloy aged at 500 ◦ C for 5 h after hot rolling at and water quenching from (a) 800 ◦ C and (b) 650 ◦ C.

(below but near β transus), equiaxed α in the β matrix was seen in the furnace cooled samples. Fine globular/elongated α and martensite were found in water quenched samples. The cooling rate of the air cooled sample is intermediate between furnace cooled and water quenched samples and hence, mixed morphology was observed in the microstructure. For the sample deformed and solution treated at 650 ◦ C which was well below β transus, substantial coarsening of primary equiaxed α was noticed for the furnace cooled sample (Fig. 4(II)(a)). The microstructure of the air cooled sample consisted of primary α and transformed β on a very fine scale (Fig. 4(II)(b)). After water quenching, fine globular and elongated α in the β matrix was seen in the microstructure (Fig. 4(II)(c)). Aging at 500 ◦ C for 5 h after water quenching from a solution treatment temperature of 800 ◦ C transformed the martensite of the water quenched sample into α and β phases and produced fine distribution of globular α along the preexisting martensite plates (Fig. 6(a)). A similar aging of the water quenched sample that was deformed at 800 ◦ C and solution treated at 650 ◦ C resulted in β to α transformation and at the same time it enhanced the growth of primary α. The aging of the water quenched samples subjected to deformation at 650 ◦ C and solution treatment at 700 or 650 ◦ C showed the growth of α while retaining the overall morphology of the water quenched samples (Fig. 6(b)). This growth was more pronounced in the case of sample solution treated at 650 ◦ C than 700 ◦ C.

To better understand the microstructures of the heat treated TZN alloy, phase analysis was carried out using Xray diffraction. The phase constituents of the heat treated TZN alloy samples were identified from the X-ray diffraction patterns shown in Fig. 7. Peaks of α and β phases were observed in furnace cooled, air cooled and aged TZN samples whereas those of martensite was identified in samples water quenched from 800 (Fig. 7(a)) or 700 ◦ C. The microstructural features of all the heat treated TZN samples are summarized in Table 2.

3.2.

Mechanical properties

3.2.1.

Hardness

The effect of heat treatment on macrohardness of the TZN samples is presented in Fig. 8. In general, for all the deformation and solution treatment temperatures the variation of the hardness with respect to the cooling rate was similar. Air cooled samples showed higher hardness than the furnace cooled and water quenched samples. Among the furnace cooled samples, the sample deformed and solution treated at 650 ◦ C showed the lowest hardness (Fig. 8(b)). A substantial decrease in hardness was found for all the water quenched samples. This decrease was more pronounced for the samples that were solution treated above or near β transus (800 or 700 ◦ C). In general, aging treatment of water quenched samples increased the hardness and this increase was substantially higher for the samples which were solution

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Table 2 – Microstructural features of the heat treated TZN alloy. Heat treatment Hot working temperature (◦ C)

Major microstructural features Solution treatment temperature (◦ C)

Cooling condition FC

Basket-weave structure (Widmanstätten lath like α) formed from prior β grains Fine Widmanstätten α laths within pre-existing β grains Martensite with retained β phase Fine distribution of globular α along the pre-existing martensite plates

800 AC

800

WQ WQ & Aged

Primary α and transformed β within β phase. Growth of primary α because of slow cooling Primary α and transformed β on a relatively fine scale Primary α and retained β. Partitioning effect of alloying elements prevents any martensite formation Growth of α

FC 650 AC WQ

WQ & Aged

Equiaxed primary α due to recrystallization of elongated α and transformed β Equiaxed/elongated primary α on a fine scale and transformed β Globular/elongated α and martensite Growth of primary α while retaining the overall morphology of the water quenched samples

FC 700 AC

650

WQ WQ & Aged

Equiaxed primary α due to recrystallization of any elongated α and transformed β. Growth of primary α Equiaxed primary α on a fine scale and transformed β No martensite due to partitioning effect. Elongated α due to incomplete recrystallization and fine globular α. Substantial growth of α

FC 650

AC WQ

WQ & Aged

treated at higher temperatures (800 or 700 ◦ C). However, in the case of the water quenched sample that was deformed and solution treated at 650 ◦ C, aging treatment had little effect on hardness. In the present study, among all the heat treated samples, the sample which was deformed at 800 ◦ C and solution treated at the same temperature followed by water quenching showed the lowest hardness (HV10 = 220±5) and aging of this sample offered the highest hardness (HV10 = 329 ± 2) (Fig. 8(a)).

3.2.2.

Elastic modulus

The modulus of elasticity of TZN alloy subjected to different heat treatment conditions is presented in Fig. 9. In all the cases, for any particular solution treatment temperature, the elastic modulus decreased with an increase in cooling rate from the solution treatment temperature. The furnace cooled samples showed the highest modulus value whereas the water quenched samples showed the lowest value (Fig. 9(a) & (b)). It was observed that the solution treatment temperature had a relatively large effect on the modulus of the water quenched samples. The high temperature solution treatment

(800 or 700 ◦ C) led to lower elastic modulus than the low temperature solution treatment (650 ◦ C). The aging treatment of the water quenched samples increased the modulus and the increase was significant for solution treatment at 800 or 700 ◦ C. The elastic modulus of the heat treated TZN samples varied from 66 to 92 GPa. Among all the samples tested, the water quenched sample that was hot rolled and solution treated at 800 ◦ C exhibited the lowest elastic modulus (66 ± 1.2 GPa). Aging treatment of this sample increased its elastic modulus (74 ± 1.7 GPa). However, this value was lower than most of the other heat treated TZN samples.

3.2.3.

Tensile properties

The load vs. displacement plots obtained from the tension tests of the TZN samples were converted into stress–strain curves and the yield strength (YS) as represented by 0.2% proof stress, ultimate tensile strength (UTS) and total elongation to failure are presented in Figs. 10 and 11. From the tensile data obtained for the TZN samples deformed at 800 ◦ C and solution treated at the same temperature followed by furnace cooling, air cooling or water

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Fig. 7 – XRD pattern of the TZN alloy deformed and solution treated for 1 h at (a) 800 ◦ C and (b) 650 ◦ C followed by furnace cooling (FC), air cooling (AC) and water quenching (WQ). The WQ samples were aged at 500 ◦ C for 5 h. quenching, it was seen that the YS, UTS and elongation of these samples varied from 486 to 562 MPa, 732 to 775 MPa and 12%–20%, respectively (Fig. 10(a)). The tensile strength of air cooled sample was higher than furnace cooled or water quenched sample. However, the water quenched sample showed higher ductility than furnace cooled or air cooled sample. Compared with the above heat treatment conditions, aging treatment of the water quenched sample led to significantly higher strength (YS and UTS) and substantially lower ductility. The YS, UTS and elongation of the aged sample were found to be 924 MPa, 1047 MPa and 4%, respectively (Fig. 10(a)). In the case of samples deformed at 800 ◦ C and solution treated at 650 ◦ C, the furnace cooled, air cooled and water quenched samples showed higher tensile strength and lower ductility than the corresponding samples deformed and solution treated at 800 ◦ C (Fig. 10(b)). In this heat treatment condition, the air cooled sample exhibited higher strength than furnace cooled or water quenched sample whereas the water quenched sample showed lower YS and better

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Fig. 8 – Vickers hardness of TZN alloy deformed at (a) 800 ◦ C and (b) 650 ◦ C and solution treated at different temperatures for 1 h followed by furnace cooling (FC), air cooling (AC) and water quenching (WQ). The WQ samples were aged at 500 ◦ C for 5 h.

ductility than furnace cooled or air cooled samples. The YS, UTS and elongation of the furnace cooled, air cooled and water quenched samples under this heat treatment condition varied in the range of 432–626 MPa, 792–871 MPa and 11%–17%, respectively. After aging of the water quenched sample, the YS, UTS and elongation were found to be 775 MPa, 881 MPa and 8%, respectively (Fig. 10(b)). Like the previous case, here also aging treatment substantially increased the YS and reduced the elongation of the alloy. The tensile strength of aged sample was comparable with that of air cooled sample. Deformation at 650 ◦ C and solution treatment at 700 ◦ C followed by different rates of cooling or aging led to a good combination of strength and ductility. Solution treatment at 700 ◦ C and subsequent cooling at different rates offered YS, UTS and total elongation in the range of 468–690 MPa, 803–878 MPa and 16%–20% respectively (Fig. 11(a)). The YS, UTS and elongation of the aged sample were 906 MPa, 1000 MPa and 12% respectively (Fig. 11(a)). Air cooled sample showed higher tensile strength and lower ductility than furnace cooled or water quenched sample. Here also aging

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Fig. 9 – Elastic modulus of the TZN alloy deformed at (a) 800 ◦ C and (b) 650 ◦ C and heat treated under different conditions.

treatment of the water quenched sample increased the strength and decreased the elongation. In general, the tensile strength as well as ductility of the TZN alloy was higher in the case of solution treatment at 700 than at 800 ◦ C. Among all the samples solution treated at this temperature, the highest strength was obtained for the aged sample whereas the ductility of the furnace cooled sample was comparable with that of water quenched sample. In the above three heat treatment conditions, the yield strength of the furnace cooled and air cooled samples was comparable and this was higher than that of water quenched sample. In the case of sample, deformed and solution treated at 650 ◦ C, furnace cooling, air cooling and water quenching resulted in YS in the range of 445–739 MPa, UTS in the range 686–851 MPa and elongation in the range 16%–20% (Fig. 11(b)) whereas the YS, UTS and elongation of the aged sample were 739 MPa, 853 MPa and 15%, respectively (Fig. 11(b)). Here also the air cooled sample showed higher YS and tensile strength than the furnace cooled or the water quenched sample and aging treatment of the water quenched sample increased its strength (YS and UTS). In this heat treatment condition a

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Fig. 10 – Tensile properties of TZN alloy deformed at 800 ◦ C and solution treated for 1 h at (a) 800 ◦ C and (b) 650 ◦ C. Solution treatment was followed by furnace cooling (FC), air cooling (AC) and water quenching (WQ). Water quenched sample was aged at 500 ◦ C for 5 h.

small increase in ductility was observed after aging treatment of the water quenched sample. Compared with the samples solution treated at 700 ◦ C, samples solution treated at 650 ◦ C showed lower tensile strength and this was more pronounced for aged samples. All the above water quenched samples showed higher tensile strength as compared to the yield strength. Among all the heat treated samples, the sample that was deformed and solution treated at 800 ◦ C followed by water quenching and aging showed the highest strength and the lowest ductility. In the present study, it was found that the trend of the variation of tensile strength with heat treatment was similar to that of hardness (Figs. 8, 10 and 11). In general, the trend can be summarized as follows. Air cooled sample showed higher UTS than furnace cooled and water quenched samples. Water quenching treatment offered high ductility coupled with good strength. Aging of the water quenched sample increased the strength of the alloy substantially but

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sample showed a large number of very small size dimples (Fig. 12(d)). This indicates that very fine α phase forms during aging and voids nucleate at these α particles. As a result, a large number of dimples form which can easily cause a fracture. Hence, the aged sample shows lower ductility.

4.

Fig. 11 – Tensile properties of TZN alloy deformed at 650 ◦ C and solution treated for 1 h at (a) 700 ◦ C and (b) 650 ◦ C. Solution treatment was followed by furnace cooling (FC), air cooling (AC) and water quenching (WQ). Water quenched sample was aged at 500 ◦ C for 5 h.

at the same time lowered the ductility except for the sample deformed and solution treated at 650 ◦ C where a small increase in ductility was noticed. The relationship between ductility and microstructure can further be explained through examination of the fracture surfaces. Some typical SEM fractographs of the heat treated TZN samples are shown in Fig. 12. In this study, the samples that were deformed and solution treated at 800 ◦ C, were selected for SEM analysis. Other heat treated samples showed similar fractographs and hence they are not presented here. The furnace cooled (Fig. 12(a)), air cooled (Fig. 12(b)) and water quenched (Fig. 12(c)) TZN samples showed mainly shallow dimples throughout the fracture surfaces indicating ductile fracture. During tensile testing of the furnace cooled and air cooled samples voids are nucleated at the α phase particles. These voids grow and coalesce to form a fracture. Similar results have been reported by Hon et al. (2003) for Ti-14–30 wt% Nb alloys. Compared with the fractured surface of the other heat treated samples, the fractured surface of the aged

Discussion

In the case of rolling at 800 ◦ C (above β transus) it was expected that a dynamically recrystallized equiaxed structure would form. Solution treatment above β transus temperature dissolves the entire alpha that develops during thermomechanical work and led to the formation of β phase in the TZN alloy. During furnace cooling or air cooling from this temperature, a part of the β phase transforms into Widmanstätten α laths which is characterized by a ‘basketweave’ arrangement of ‘packets’ of α plates (Banerjee and Krishnan, 1981). The size of the laths depends on the cooling rate. It has been observed that cracks propagating in a straight line within a packet of similarly aligned α plates would be deflected or arrested at the boundary between packets (Boyer et al., 1994). On the other hand, the α phase is sandwiched between the β matrix within each packet and gives local plastic constraints which leads to strengthening of the sample. Hexagonal closed packed (hcp) α that remains untransformed during solution treatment in the α + β phase field (650 ◦ C in this case) is called primary α. The morphology of primary α is influenced by the thermo-mechanical history and can be lamellar, equiaxed or mixed (Banerjee and Krishnan, 1981; Boyer et al., 1994). Therefore, in the case of the hot rolling at 800 ◦ C (β transus) followed by solution treatment at 650 ◦ C (in the α + β region), presence of primary equiaxed α and transformed β within pre-existing β phase was observed for furnace cooled sample. Air cooled sample showed similar microstructure on a relatively fine scale because of faster cooling. In the case of deformation at 800 ◦ C, compared with the furnace cooled or air cooled sample solution treated at 800 ◦ C, the higher tensile strength of the corresponding TZN sample solution treated at 650 ◦ C is associated with the presence of primary α in the matrix. Heavy deformation (>30%) at temperatures below the transus is expected to produce worked structure in the material and α phase is nucleated during the plastic deformation. Subsequent annealing in the α + β phase field at 700 or 650 ◦ C with a long holding time of 1 h followed by furnace cooling seems to have resulted in the recrystallization of any elongated α that may have formed during deformation at 650 ◦ C. Thus equiaxed α is seen in the furnace cooled samples with noticeable coarsening of primary α in the sample solution treated at 650 ◦ C (Fig. 4.II(a)). Therefore, solution treatment at 650 ◦ C (in the α + β field) followed by furnace cooling gives higher amount of α phase. This coarsening effect of α decreases the hardness and tensile strength of the samples heat treated in this condition. Solution treatment at 700 ◦ C after deformation at 650 ◦ C results in smaller amount of α because a part of the α transforms into β during solution treatment near β transus temperature. On the other hand, deformation in the α + β field provides a large number of nucleation sites for the

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Fig. 12 – Fractographs of TZN alloy, deformed at 800 ◦ C, solution treated at 800 ◦ C for 1 h followed by (a) furnace cooling, (b) air cooling, (c) water quenching and (d) aging of the water quenched sample.

formation of α phase. During furnace cooling or air cooling a part of the β is transformed into α and forms on the large number of nucleation sites. Hence, compared with the solution treatment at 650 ◦ C, solution treatment at 700 ◦ C leads to the formation of fine dispersed α in the matrix of furnace cooled or air cooled sample which increases the hardness and tensile strength. All the furnace cooled and air cooled TZN samples consisted of α and β phases. However, compared with the furnace cooled samples, the higher hardness and strength of the air cooled samples is associated with their finer microstructure due to relatively fast cooling. It is known that martensite structure can be produced if the cooling rate from the β phase field or from the high temperature region of the α + β field is sufficiently high (Banerjee and Krishnan, 1981; Boyer et al., 1994). Accordingly, fast cooling from the above temperature produced martensite in the microstructure. In the case of binary Ti–Nb alloy, quenching from the β phase field can result in the formation of two metastable martensite structures, either hexagonal α′ in alloys with Nb < 13 wt% or orthorhombic α′′ at higher Nb content (Tang et al., 2000). In the present study, the Nb content in the titanium alloy is 13 wt%. Therefore, orthorhombic α′′ martensite is formed after water quenching from 800 or 700 ◦ C. However, water quenching from solution treatment temperature of 700 ◦ C also showed the presence of fine globular/elongated α in the microstructure because here the solution treatment temperature is slightly below the β transus. The hardness of martensite is significantly lower than that of β phase (Ohmori et al., 2001). Hence, the presence of martensite and lack of α or sufficient amount of α in the microstructure of the samples water quenched from 800 or 700 ◦ C showed a decrease in hardness. On the other hand, water quenching from the solution treatment temperature of

800 or 700 ◦ C produces solid solution hardening effect which gives reasonable amount of tensile strength, at the same time presence of soft β and α′′ martensite in the microstructure after water quenching offers lower yield strength and better ductility. Along with β to α transformation, partitioning of alloying elements also takes place during solution treatment at 650 ◦ C (Geetha et al., 2004; Tang et al., 2000). Thus, β phase becomes enriched with Nb during this treatment. This reduces the Ms temperature of the untransformed β to below room temperature (Tang et al., 2000) and thus no martensite was formed on water quenching from 650 ◦ C. Elongated α and retained β is found in samples deformed at 650 ◦ C and water quenched from the same temperature. This is possibly because of incomplete recrystallization of deformed α due to faster cooling (Fig. 4.II(c)). Therefore, this type of microstructure shows higher hardness and tensile strength than the microstructure that consists of martensite and retained β, because of presence of α phase. During aging of the water quenched sample after deformation and solution treatment at 800 ◦ C, the martensite phase decomposes into the α + β phase through diffusion controlled process. Aging of the water quenched sample produces a fine distribution of small globular α along the pre-existing martensite plates. It has been found that the precipitation of the α phase in the matrix increases the hardness of the β titanium alloys (Ikeda et al., 2002; Kuroda et al., 2005). Substantial increase in hardness and tensile strength after aging can therefore be associated with the decomposition of martensite and retained β and the precipitation of fine α phase in the matrix during aging. However, these fine α particles along the pre-existing martensite plates act as a potent site for the nucleation of the cracks and thus, reduce the ductility of the material.

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The increase in hardness and strength and decrease in ductility after aging of the water quenched sample which is deformed at 800 ◦ C and solution treated at 650 ◦ C are associated with the transformation of retained β into α phase. It is suggested that during transformation major part of the α phase forms on the grain boundary of β grains and hence the elongation decreases. In the case of samples deformed at 650 ◦ C and solution treated at 700 ◦ C, the water quenched samples consist of mainly α, β and relatively small amount of martensite. Aging of the water quenched samples led to the growth of α and at the same time transform a part of retained β into α. In addition, martensite dissociates and forms a small amount of fine α on the pre-existing martensite plates during aging and therefore this small volume fraction of fine α particles has no substantial influence on the ductility. Hence, the ductility of the water quenched sample is not significantly decreased by aging. On the other hand, the sample subjected to deformation and solution treatment at 650 ◦ C followed by water quenching consists of elongated/globular α and retained β. Aging of this water quenched sample changes the morphology of the elongated α into equiaxed α, transforms some amount of β into α and at the same time enhances the growth of α. Hence, during aging the increase in hardness due to the decomposition of small amount of retained β into α may be compensated by the coarsening of the α and therefore no considerable improvement in hardness after aging of the water quenched sample is found. The change of morphology of α from elongated to equiaxed structure also leads to an increase in ductility of the water quenched sample after aging. The elastic modulus is more sensitive to phase/crystal structure than to other factors (Ho et al., 1999). In a multiphase alloy the modulus is determined by the specific modulus of the phases and by their volume fractions. The composition of the constituent phases and their volume fractions depend on the prior thermo-mechanical treatment (like hot working) and heat treatment (like solution treatment and aging treatment) of the alloy (Lee and Welsch, 1990). It has been reported that the β and α′′ phase mixture offers lower modulus than α phase and the modulus of different phases in titanium alloys increase in the following sequence: Eβ < Eα′′ < Eα (Hao et al., 2002; Ho et al., 1999). Hence, the variation of elastic modulus observed in the present study can be explained with respect to the microstructure. The microstructure of the furnace cooled and air cooled samples consisted of α and β phases. However, compared with air cooled sample, the furnace cooled sample has higher volume fraction of α phase as a result of slower cooling during which α phase has sufficient time to grow. Hence, the high elastic modulus of the furnace cooled sample is due to presence of high amount of α phase in the microstructure. Water quenching leads to the formation of martensite and retained β or α and retained β depending upon the solution treatment temperature. In the case of sample water quenched from 800 ◦ C, the microstructure consists of martensite and β phase. Both phases offer lower elastic modulus as compared with α phase. Thus, among all the investigated samples, this sample exhibits the lowest elastic modulus value (66 ± 1.2 GPa). The water quenched sample, which was solution treated at 700 ◦ C, consisted of β, α and martensite and showed low elastic modulus. However, this

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value is still higher than that of previous sample because α is present in the microstructure. In the case of solution treatment at 650 ◦ C, a small decrease in elastic modulus of the water quenched sample when compared with furnace cooled and air cooled samples can be attributed to the presence of smaller amount of retained β and absence of any low modulus α′′ martensite in the microstructure. Aging of the water quenched sample transforms the low modulus martensite and retained β phases into high modulus α phase and hence the modulus of the samples is increased. In the case of sample solution treated at 650 ◦ C the effect of aging on modulus is less pronounced. Aging at 500 ◦ C for 5 h after water quenching from the α + β phase field resulted in the decomposition of only retained β as there is no martensite in this water quenched sample. This decomposition leads to the formation of α and β phases. Small amount of α phase formed on aging does not increase the modulus substantially.

5.

Conclusions

The effect of heat treatment on microstructure and mechanical properties of near-β Ti–13Zr–13Nb alloy for biomedical load bearing applications has been studied. Depending on the heat treatment conditions, the microstructure of the heat treated TZN alloy consists mainly of α, β or α′′ martensite phase with different morphologies. In general, for all the deformation and solution treatment temperatures the variation of the hardness and tensile strength with cooling rate is similar. Air cooled sample shows higher hardness and tensile strength than furnace cooled or water quenched samples. Aging of water quenched TZN samples increases the hardness and tensile strength but decreases the ductility. The elastic modulus of the alloy decreases with an increase in cooling rate from the solution treatment temperature. The presence of high amount of α phase in the microstructure increases the elastic modulus whereas the presence of α′′ martensite and retained β phases lowers the modulus of the samples. The elastic modulus of TZN alloy is in the range of 66–92 GPa, which is significantly lower than that of conventionally used stainless steel (210 GPa), Co–Cr alloys (204–240 GPa) and α + β titanium alloys (100–120 GPa). Among the alloys studied in the present work, the aged TZN sample, which was deformed and solution treated at 800 ◦ C followed by water quenching exhibits good combination of mechanical properties (elastic modulus 74 GPa, yield strength 924 MPa and tensile strength 1047 MPa). Earlier work has already shown that this sample has very good biocompatibility and wear resistance in body fluid (Gupta et al., 2006; Majumdar et al., 2008a). Therefore, TZN alloy in this heat treatment condition is a promising candidate for biomedical load bearing applications.

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