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Influence of phase transformations on dynamical elastic modulus and anelasticity of beta Ti-nb-Fe alloys for biomedical applications J.M. Chaves, O. Florêncio, P.S. Silva Jr., P.W. B. Marques, C.R.M. Afonso

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S1751-6161(15)00073-9 http://dx.doi.org/10.1016/j.jmbbm.2015.02.030 JMBBM1407

To appear in: Journal of the Mechanical Behavior of Biomedical Materials

Received date:5 December 2014 Revised date: 21 February 2015 Accepted date: 26 February 2015 Cite this article as: J.M. Chaves, O. Florêncio, P.S. Silva Jr., P.W.B. Marques, C.R.M. Afonso, Influence of phase transformations on dynamical elastic modulus and anelasticity of beta Ti-nb-Fe alloys for biomedical applications, Journal of the Mechanical Behavior of Biomedical Materials, http://dx.doi.org/ 10.1016/j.jmbbm.2015.02.030 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Influence of phase transformations on dynamical elastic modulus and anelasticity of beta Ti-Nb-Fe alloys for biomedical applications J. M. Chavesa*, O. Florêncioa, P. S. Silva Jr.a, P. W. B. Marquesa, C. R. M. Afonsob a

Department of Physics, UFSCar, C.P. 676, CEP 13565-905, São Carlos-SP, Brazil.

b

Department of Materials Engineering, DEMa-UFSCar, CEP 13565-905, São Carlos-SP, Brazil.

Abstract Recent studies in materials for biomedical applications have focused on β-titanium alloys that are highly biocompatible, free of toxic elements and with an elastic modulus close to that of human bone (10-40 GPa). Beta Ti-xNb-3Fe (x = 10, 15, 20 and 25 wt.%) alloys were obtained by rapid solidification and characterized by anelastic relaxation measurements at temperatures between 140 K and 770 K, using a free-decay elastometer, as well as analysis by Differential Scanning Calorimetry (DSC), X-ray Diffraction (XRD) and Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM). The observed stabilization of the β-phase with rising Nb content was linked to the strength of the relaxation peak around 570 K. The phase transformations detected in the anelastic relaxation spectra agreed with those observed in the DSC curves. However, the results from anelastic relaxation spectra provide more detailed information about the kinetics of phase transformations. At temperatures between 140 K and 300 K, there was an indication of a reversible transformation in the alloys studied. The elastic modulus measurements showed a hardening of the material, between 400 K and 620 K, related to the ω-phase precipitation. However, the starting temperature of ω-phase precipitation was clearly influenced by the Nb content, showing a shift to high temperature with increasing percentage of Nb. At temperatures above 620 K, a fall was observed in the dynamical elastic modulus, accompanied by a relaxation peak centered at 660 K, which was attributed to the growing αphase arising from the ω-phase, which acts as a nucleation sites or from the decomposition of the metastable β-phase. XRD patterns confirmed the formation of β, α and ω phases after mechanical relaxation measurements. A predominant β phase with dendritic morphology was observed, which became more stable with 25 wt.% Nb. The lowest elastic modulus was of 65 GPa obtained in the Ti-25Nb-3Fe alloy, representing a good low value for a β-Ti alloy with a relatively low addition of β stabilizing elements (Nb and Fe). Keywords: β titanium alloys, anelastic relaxation, elastic modulus, phase transformation. * Corresponding author: e-mail address: [email protected]

1. Introduction Research on titanium alloys designed for biomedical applications has intensified a great deal in recent decades, on account of their interesting combination of properties, such as excellent chemical biocompatibility, low elastic modulus and corrosion resistance. These properties can be tuned by the addition of α- and β-stabilizing alloying elements, adjusting the processing parameters or heating and thermo-mechanical treatments (Niinomi, 2002). This tuning depends on the close relationship between the microstructure and phases formed, their distribution, size and morphology(Elmay et al., 2013). Thus, β-Ti alloys free of toxic elements have been designed in recent research, with low values of elastic modulus, which are closer to that of human bone (10-40 GPa) that other metal alloys used as biomaterials: Co-Cr (220 GPa), stainless steels (190 GPa), CP-Ti (100 GPa) and α+β type Ti-6Al-4V (110 GPa) (Abdel-Hady Gepreel and Niinomi, 2013; Long and Rack, 1998; Niinomi, 2008). The similarity in the elastic modulus between human bone and the implant is very important in alloys used as orthopedic implants, because a great difference can lead to the stress-shielding phenomenon. This phenomenon, which relieves the bone of external stress, reduces the reposition of the bone surrounding the metallic implant and consequently leads to loss of the implant or new fracture of the bone after the implant is removed (Niinomi and Nakai, 2011). However, control of the crystalline phases is of great importance in the design of new biomaterials alloys, which should favor the presence of metastable phases such as α’ and α” martensites, the latter responsible for the shape memory effect and superelasticity (Baker C, 1971; Kim et al., 2006); on the other hand, the ω-phase is deleterious to mechanical properties of the β-Ti alloys, increasing their hardness and elastic modulus and can occur during rapid cooling (quenching), aging heat treatments or upon deformation (Lee et al., 2002b; Li et al., 2013; Mantani and Tajima, 2006). In this context, the mechanical spectroscopy technique can be very useful for the characterization of dynamical processes, including phase transformations in β-Ti alloys. Thus, internal friction is a physical property very sensitive to processes that involve absorption of mechanical energy by atomic rearrangement, phase transformations, matrix-solute interactions (interstitial and/or substitutional), diffusion process and so on (Nowick and Berry, 1972; Schaller et al., 2001). Ti-Nb alloys have been widely studied, on account of their low elastic modulus, shape memory and superelasticity (Kent et al., 2013; Lopes et al., 2011; Tobe et al., 2013). It has been observed that the mechanical properties, structure and morphology are sensitive to Nb content. In this system, alloys containing around 15 wt.% or less of Nb exhibit α’ martensite (hexagonal); at 17

wt.% - 25 wt.% Nb, the α” martensite phase (orthorhombic) predominates and, in alloys with more than 30 wt.% Nb, the β phase is dominant with equiaxal grain structure (Lee et al., 2002b). Therefore, the study of the new or modified alloys by this technique can help to improve the performance of the alloys used in biomedical applications. The present study is focuses on the Ti-Nb-Fe system, in which the fractions of β-stabilizing elements Nb and Fe influence phase stability. Fe is added on the basis its strongly effect stabilizing on the β-phase, mechanical strengthening potential, low cost and its lowering effect on the melting point of the alloy, which can facilitate the production of alloys, with the aim of achieving a high performance-to-cost ratio. On the other hand, in compositions close to the eutectic one, this alloy can provide a greater thermal supercooling from the liquid, favoring the formation of refined microstructures, dendritic growth or metastable phases (Reed-Hill, 1973). Thus, the anelastic behavior of the Ti-xNb-3Fe alloys (x=10, 15, 20 and 25 wt.%) was characterized by mechanical spectroscopy and its correlation with the sequence of phase transformations during cyclic heat treatment will be demonstrated, in order to map the optimization of heat treatments for these β-Ti alloys, so as to guarantee the desired phases and microstructure, with the aim of biomedical application. 2. Materials and methods Ti-xNb-3Fe alloys (x=10, 15, 20 and 25 wt.%) were processed from pure elements such as Ti (sponge: 99.5%), Nb (99.8%) and Fe (99.97%). Ingots were fabricated in an arc furnace (AM Arc Melter, Edmund Bühler) in an inert atmosphere of argon in a water-cooled copper hearth crucible and then remelted to ensure compositional homogeneity. From these ingots, smaller samples were rapidly solidified in a commercial arc furnace (EDG Discovery Plasma®) by suction casting-technique, the suction arising from a vacuum between the copper crucible (upper chamber) and the copper mold (lower chamber), with shallow grooves for the vacuum suction. The cooling rates imposed during the copper-mold casting are estimated at 101 K/s to 103 K/s. The copper mold used plate shape, with varying thickness, had dimensions 20x50 mm2 and thickness 2.0 mm, 1.0 mm and 0.5 mm. Plates were cut in dimensions of around 20x5x0.5 mm3 as specimens for mechanical spectroscopy tests. Anelastic spectra (internal friction and elastic modulus relative variation to the room temperature value (∆E/E) plotted against temperature) were obtained with a free decay elastometer equipment (Vibran Technologies AE-102) from the logarithmic decay of free oscillations of the first tone in flexural vibration mode. Dynamical elastic modulus (E) was calculated from the

resonance frequency of the fundamental mode of flexural vibration (f1), in the clamp-free configuration, by the relationship (Nowick and Berry, 1972): f1 = 0.1615

h l2

E

ρ

(1)

where h is the thickness, l the length and ρ the density of the sample. Analysis in the free decay elastometer consisted of a series of cyclic measurements, during which the sample was cooled, by immersing the chamber containing it in liquid nitrogen, from room temperature (RT) down to ~140 K, then heated from 140 K up to 770 K, at a cooling and heating rate of 1 K/min. In this equipment, the sample was held in a vacuum chamber with an internal resistive furnace, at a pressure below 4x10‐5 Torr. The data for the anelastic spectra were collected in measurements made at each 0.5 K. Complementary data on the crystalline structure in the as-quenched condition and after heat treatments carried out during the anelastic relaxation measurements, were provided by X-ray Diffraction (XRD) in a Siemens D5005 X-ray Diffractometer. The microstructure was characterizated by Scanning Electron Microscopy (SEM) whit a JEOL JSM-5800LV and Transmission electron microscopy (TEM) using a FEI Tecnai G2 F20 equipment with the 200 KV beam for field emission (FEG) coupled with energy dispersive spectroscopy (EDS) EDAX. The samples were prepared metallographically by gridding, polishing and etching with Kroll reagent (6 mL HNO3, 3 mL HF and 91 mL of H2O). Measurements of interstitial elements content were carried out using a LECO TC-436 analyzer. Additionally, Vickers hardness measurements were carried out with a Stiefelmayer KL2 microhardenss tester, taking the mean value for six indentations with load of 200 gf during 15 s. To compare the kinetics of phase transformations, analysis by Differential Scanning Calorimetry (DSC) was carried out in Netzch STA 404 calorimeter during two thermal cycles between 140 K and 770 K at a heating rate of 1 K/min.

3. Results and discussion 3.1. Microstructural characterization and phase identification Fig. 1 displays the micrographs of rapidly solidified Ti-xNb-3Fe (x=10, 15, 20 and 25 wt.%) alloys, exhibiting an isotropic dendritic growth morphology, without preferred orientation, in which the microstructure become more refined with increasing of Nb content. This refinement is due to the effect of the β-stabilizer eutectoid alloying element (Fe) that increases the thermal supercooling that occur during rapid solidification.

The chemical analysis by EDS and the interstitials solute atoms content measured by gas analysis, shown in tables 1 and 2, respectively, indicate that the preparation of the alloy was adequate in regards as nominal composition. Fig. 2 shows the XRD patterns of rapidly solidified Ti-xNb-3Fe (x=10, 15, 20 and 25 wt.%) alloys in the as-quenched condition. In the diffractograms there are typical diffraction peaks of the β-phase (bcc) stabilized at room temperature, besides diffraction peaks of the metastable ωphase (hcp) for the Ti-10Nb-3Fe and Ti-15Nb-3Fe alloys. It can be seen that increasing the Nb percentage from 10 to 25 wt.% stabilized the β-phase formed. Additionally, the presence of 3 wt.% Fe, a strong β-stabilizing eutectoid element, favored the stabilization of the β-phase at just 20 wt.% Nb, instead of a minimum of 35 wt.% Nb to obtain the stable β-phase in the binary TiNb system (Lee et al., 2002a; Lin et al., 2002). Regarding the splitting in β-Ti phase peak, observed in the Fig. 2, is due to segregation in dendritic growth, according to SEM images in BSE (backscattered electrons) mode (Fig. 3), resulting in variation of crystalline cell parameter with Nb-rich β-Ti phase (brighter) in the inner region of the dendrites and Ti and Fe-rich β-Ti phase (darker) in the interdendritic region of the microstructure. That differ from β-Ti phase separation that occurs in stable β-Ti alloys (nanometric scale) with higher contents of β-stabilizing elements, such as TNZT, depending on the processing route and heat treatment conditions (Afonso et al., 2010). Thermal analysis by DSC, collected at a heating rate of 1 K/min, of samples with 10 and 15 wt.% Nb (Fig. 4) showed an exothermic process around 790 K, whereas during the second heating cycle, no process was revealed. For samples with 20 and 25 wt.% Nb, an endothermic process around 715 K was observed, and in the second heating cycle the process still takes place. From literature data (Mantani and Tajima, 2006), precipitation of the α-phase is associated with the exothermic peak and precipitation of β and ω-phases with the endothermic peak. 3.2. Anelasticity and dynamic elastic modulus characterization In Fig. 5, the elastic modulus is plotted against electron concentration (e/a), it can be observed that the addition of Nb leads to an increase in the (e/a) ratio, which is linked to a decrease in the elastic modulus. This fact is related to β-phase stabilization, since it has been suggested in (Ikehata et al., 2004) that, for e/a values less than 4.20, the bcc structure is unstable and, on decreasing the valence electron concentration, the stability of the hcp structure rises, so that metastable phases, such as ω or martensites, can be formed, increasing the elastic modulus. For

this alloy system, higher stabilization of the β-phase was achieved from e/a= 4.23 (20 wt.% Nb) upwards, while at lower values the ω-phase appears, as shown in XRD patterns. This result is consistent with that established by Collings (Boyer R. et al., 1994), who showed that in Ti-TM alloys (TM= transition metal), the ω-phase forms when the content of the alloy lead to an electron concentration (e/a) between 4.12 and 4.21. Comparing these results with those obtained by Fedotov (Collings, 1986) in the Ti-Nb system, and the curve Ti-TM for binaries systems proposed by Collings (Collings, 1984), it can be seen that the points for the Ti-Nb-Fe alloy system are to the left of the curves for binary systems, indicating the higher stability of the β-phase, caused by suppression of the metastable phases α” and ω. Thus, the 3 wt.% Fe leads to a stabilization of the β-phase in samples with low Nb content, a fact linked to the higher e/a ratio in the Ti-Nb-Fe system. However, TEM micrographs of Ti-10Nb-3Fe and Ti-25Nb-3Fe samples in the as-quenched condition, in Figs. 6 and 7 respectively, showed a nanoscale structure. Fig. 6 shows a bright field (BF) image (Fig. 6.a) and dark field (DF) image (Fig. 6.b) of athermal nanometric ω-phase precipitate dispersed in β-Ti (bcc) grain matrix, due to rapid quenching from the melt. The respective selected area electron diffraction (SAED) pattern of Ti-10Nb-3Fe (inset) is shown with orientation relationship of zone axis [2 1 2]β//[001]ω. Fig. 7 shows TEM micrographs of the Ti-25Nb-3Fe sample, with BF (Fig. 7.a) and DF micrographs (Fig. 7.b) showing athermal ωphase precipitates, together with respective SAED pattern (inset Fig.7.a) with orientation relationship of zone axis [1 1 5]β. Although XRD patterns of a Ti-25Nb-3Fe sample in Fig. 2 do not have clear peaks of ω-phase, TEM analysis confirms that there is still some fraction of this phase precipitated on the nanoscale. It can be seen as well that the ω-phase is finer in the Ti10Nb-3Fe alloy and in a greater fraction than that in the Ti-25Nb-3Fe alloy, owing to Nb content. Thus, the elastic modulus values and e/a ratio are consistent with the ω-phase fraction in alloys Ti-10Nb-3Fe and Ti-25Nb-3Fe. 3.2.1. Anelastic relaxation spectra for Ti-10Nb-3Fe alloy Fig. 8 shows the anelastic relaxation spectra (elastic modulus relative variation to the room temperature value (∆E/E) and internal friction, plotted against temperature) and XRD patterns for the Ti-10Nb-3Fe alloy during three measurement cycles between 140 K and 770 K. In these spectra a relaxation structure can be seen that depends on the temperature and is related to phase transformations, since heat treatment of metastable phases promotes changes in the

system toward an energetically more favorable equilibrium. These spectra show more detailed evidence of the kinetic of phase transformations than the DSC curves. In the first thermal cycle at low temperature (300 K to 140 K), there is a relaxation structure that represent high absorption of elastic energy. In this step, the internal friction (Q-1) reaches values around 4x10-3, and that energy is recovered in the subsequent heating up to room temperature. Similar behavior is observed in the elastic modulus, as its variation does not exceed 1%. This indicates that the process is almost completely reversible, so that it must be related to the reverse martensite transformation of the β→α” phase, in agreement with the literature (Bertrand et al., 2013); the reversible movement of the lattice by austenite/martensite transformation dissipates mechanical energy constituting a source of structural damping. These observations indicate that this alloy probably display a shape memory effect in this range temperature, but more specific study is necessary and this will be discussed in a forthcoming publication. At higher temperatures, a minimum in ∆E/E is reached around 400 K; this behavior may be associated with some characteristic temperature of this phase transformation. In the temperature range between 400 K and 700 K, during the first thermal cycle, ∆E/E reaches a maximum value of around 10% at 615 K that could be related to the precipitation or growth of the metastable ω-phase that provokes the hardening of the alloy, since, as reported frequently in the literature (Li et al., 2013; Tane et al., 2013), this phase has the highest elastic modulus among the phases formed in β-Ti alloys. At higher temperatures ∆E/E decreases, while Q-1 continues to rise; this behavior is associated with the growth of α-phase. In the subsequent cooling, a hardening of the alloy is observed, achieving 20% of variation on ∆E/E, as can be observed from the beginning of the second thermal cycle at room temperature. The phase transformations are confirmed by XRD patterns after first thermal cycle, as observed in Fig 8d. In the following heat treatment, it is observed that the presence of ω and α phases affect considerably the reversible behavior at low temperature, since Q-1 decreases to 1x10-3 and 0.5x10-3 during the second and third thermal cycle, respectively. The final ∆E/E value, after third thermal cycle, was around 12%, corresponding to 84 GPa. The intermediate value for ∆E/E could be associated with competition among the series of phases transformations, in sequence: β → ω → α, leading from the initial microstructure β+ω to a final microstructure composed of β+α. This transformation sequence has already been reported in the literature (Chaves et al., 2014; Lopes et al., 2011; Ohmori et al., 2001), since that metastable ω-phase (isothermal) can precipitate under aging heat treatments or during cooling at

lower cooling rates from the β-phase field. Besides that, growth of α-phase can occur from ωphase that acts as a nucleation site, or from the decomposition of the β-phase metastable.

3.2.2. Anelastic relaxation spectra for Ti-15Nb-3Fe alloy Anelastic relaxation spectra ∆E/E and Q-1 plotted against temperature) and XRD patterns for the alloy Ti-15Nb-3Fe, during three measurement cycles between 140 K and 770 K, are shown in Fig. 9. The alloy Ti-15Nb-3Fe behaves similarly at low temperature, to Ti-10Nb-3Fe. In the first thermal cycle, a reversible process occurs and it is affected by the appearance of new phases, as well as changes in the existing phases in each new cycle. Also the weak minimum in ∆E/E is shifted to around 425 K during first heating, which may be due to the higher proportion of βphase stabilized. In this alloy, at temperatures above 425 K, the hardening due to the growth of nanometric ωphase is not clearly observed, as ∆E/E reaches an apparent maximum of 5% around 580 K and then decreases at higher temperatures. This confirms that Ti-15Nb-3Fe has a β-phase and a lower proportion of ω-phase, leading to a rapid transformation ω→α and thus no apparent rise in of the

ω-phase fraction during heating cyclic. Thus, α-phase is formed by the decomposition of metastable β-phase. In this way, the initial β+ω microstructure is transformed to the equilibrium

β+α microstructure during the thermal cycles in the anelastic relaxation measurements. The addition of Fe as a strong eutectoid stabilizer of β-Ti acted as a strong suppressor of precipitation of the ω-phase, even at low content (3 wt.%.), preventing the massive ω-phase formation. The phases transformations were confirmed by XRD patterns observed in Fig. 9.d. Thus, the elastic modulus remained almost constant, reaching a final value of 67 GPa. 3.2.3. Anelastic relaxation spectra for Ti-20Nb-3Fe alloy. For Ti-20Nb-3Fe alloy, in the anelastic relaxation spectra in Fig. 10, the maximal relaxation strength at the minimum temperature was less intense, since the Q-1 maximum value at 140 K was 3x10-3 and little affected by the subsequent thermal cycles. At temperatures between 400 K and 550 K, a slight hardening of material is observed in ∆E/E and, above this temperature, there is a sharp increase in ∆E/E that reaches a maximum of 18% at 630 K. As stated above, this behavior arises from the β→ω transformation, which can be favored by the low heating rate (1 K/min) during the thermal cycle. In the subsequent thermal cycles, a hardening of up to 25% was observed, which stabilized at 20% in the final cycle. This behavior may be associated firstly with

the transformation of metastable β-phase, from which the ɷ-phase nucleated and, after the αphase is formed and nucleated, as it was reported that, in β-Ti alloys, α-phase laths are nucleated at the ɷ/β interfaces and grow into both the matrix and ɷ particles, thus consuming the ɷ-phase particles(Ohmori et al., 2001). The evolution of phases in the sample led the original β-phase to a structure β+α, as shown in XRD patterns in each thermal cycle (Fig. 10d). Internal friction (Q-1) curve for the first thermal cycle, in the range 400 K to 750 K shows two anelastic relaxation peaks of strength 2x10-3, centered at 570 K and 680 K, respectively. In the subsequent cooling, the curves does not replicate this behavior, as only the peak at higher temperature appears. In the second and third thermal cycles the peak is centered around 635 K and the relaxation strength falls to 1.6x10-3 and 1.4x10-3, respectively. In addition, a weak inflection is seen in ∆E/E, at the maximum strength of the relaxation; this behavior is characteristic of Snoek-type relaxation, according to the standard anelastic solid theory (Nowick and Berry, 1972). The nature of the relaxation peak will be discussed in section 3.3. 3.2.4. Anelastic relaxation spectra for Ti-25Nb-3Fe alloy The behavior of the anelastic spectra for Ti-25Nb-3Fe alloy submitted to three thermal cycles is presented in the Fig. 11. The lower temperature behavior seen in the previous alloys of system Ti-xNb-3Fe (x = 10, 15, 20 wt.%) is again similar for this composition, where the relaxation strength in Q-1 is affected by each thermal cycle. In the first cycle, the ∆E/E curve shows a weak fall (4%) that stabilizes between 425 K and 550 K, above which ∆E/E rises sharply, reaching a hardening of 20% at 640 K, followed by a drop up to 750 K. In the subsequent cooling ∆E/E increased slowly, reaching 10%. In the second thermal cycle, contrary to what happened at other compositions, around 650 K a new maximum of 8% is formed in ∆E/E. This is possibly related to residual ω-phase in the specimen. In the third cycle, ∆E/E stabilizes around 23%, corresponding to 79 GPa . In the Q-1, during the first cycle, a distinctive relaxation anelastic peak is displayed around 565 K, with a relaxation strength of 5.5 x10-3, much higher than in Ti-20Nb-3Fe alloy. At higher temperatures, a shoulder appears at 680 K, similar to the second peak for Ti-20Nb-3Fe. In the subsequent thermal cycles, the relaxation peaks overlap to form a single peak with relaxation strength reduced to 2.5x10-3 and centered around 620 K. This behavior will be discussed in the next section.

In table 3 are summarized the elastic modulus and microhardness values, as well as the phases observed in alloys Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) in each thermal cycle. Thus, it is seen that the ω-phase is responsible for the higher values of microhardness and elastic modulus. 3.3. Relation of anelastic relaxation peak with structure phases In Fig. 12, it can be seen that the strength of the relaxation peak around 565 K depends on the Nb content in the system Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%). In section 3.2 (Fig. 5), it was observed that the addition of Nb leads to a decrease in the elastic modulus, since the β-phase is more stable. Therefore, the height of the relaxation peak may be linked to the stability and amount of β-phase, related to the Nb content in the system. Thus, it can be observed that in the alloys containing 10 and 15 wt.% Nb, which show β+ω phases in the structure, the peak height is imperceptible, compared to alloys containing 20 and 25 wt.% Nb, which have a small amount of ω-phase, as shown by the XRD and TEM results (Figs. 2 and 7 respectively). Similar behavior was observed by Zhou and coworkers in Ti-Nb alloys (Zhou et al., 2011). The second peak, around 680 K, may be associated with the growth of α-phase, since this peak appears in the region where ∆E/E drops sharply, so that the α-phase may be emerge from the ωphase, which acts as a nucleation site or by decomposition of the metastable β-phase(Ohmori et al., 2001). Curiously, this drop in ∆E/E for Ti-15Nb-3Fe alloy begins at a lower temperature, which suggests that the α-phase arises without hardening from the ω-phase. In the ∆E/E curve (Fig. 12), it is observed that the onset temperature for hardening shifts to higher temperatures with increasing Nb content. This behavior indicates that, with 3 wt.% Fe in the alloys, the addition of Nb, retards the growth of ω-phase, which develops rapidly when it reaches a characteristic temperature, for example 550 K in the alloys containing 20 and 25 wt.% Nb. In the subsequent cycles, for alloys containing 20 and 25 wt.% Nb, relaxation peak overlap, their intensity decreases and an inflection is seen in ∆E/E, characterizing a Snoek-type relaxation, mentioned above. Other authors (Almeida et al., 2009; Lu et al., 2012; Yin et al., 2006) have reported a Snoek-type relaxation in β-Ti alloys, suggesting that this peak can be caused by stressinduced reorientation of interstitial atoms in the octahedral positions of the bcc lattice: these interactions may include Nb-O, Ti-O, Nb-O-O, Ti-O-O, Nb-N in the β-phase. Since the Nb and Fe are acting as β-stabilizers and the α-phase is rich in Ti and the β-phase is rich in Nb and, with increasing contents of Nb, the β-phase is more stabilized, it may be

suggested that more Nb-O pairs are formed, since as showed in the gas analysis (table 2) these samples contain a certain amount of oxygen interstitial. Thus, this fact may result in an increase in the peak intensity. In the subsequent thermal cycles, the precipitation of ω-phase and/or growth of α-phase occur in each thermal cycle, as observed from XRD patterns. Thus, it is suggested that firstly the β→ω transformation takes place; then, the α-phase is formed from β and ω-phases, consuming ɷ-phase particles and some portion of the existing β-phase, resulting in a competition in the growth of ω and α-phases in each thermal cycle and leading to a decrease in the proportion of β-phase and therefore a reduction of the relaxation peak. This statement agrees with the report by Moffat (Moffat and Larbalestier, 1988), who showed that between 648 K and 698 K the β+ω → β+α transformation can occur in Ti-Nb alloys during aging. Besides, as the sequence of transformation involves redistribution of elements between regions lean and rich in solutes, this process may affect the interactions between elements Nb, Ti and interstitial atoms and the temperature and intensity of the relaxation peak in Ti-20Nb-3Fe and Ti-25Nb-3Fe alloys.

4. Conclusions The study of anelasticity and the variation of elastic modulus during thermal cycles were correlated with phase transformations in the Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) alloy system. At low temperatures, signs of a reversible transformation were observed in these alloys. On the other hand, at higher temperatures, the growth of the metastable ω-phase was observed followed by the nucleation and growth of the equilibrium α-phase. Thence, in the rapidly solidified metastable condition, the thermal cycles led from the β+ω microstructure (initial state) to a β+α microstructure (final state). For Ti-20Nb-3Fe and Ti-25Nb-3Fe alloys there were observed two relaxation peak: a peak around 565 K was associated with stabilization of β-phase, since the height was proportional to the Nb content; the other peak, around 680 K, was linked to α-phase, which grows from the ωphase and decomposition of the metastable β-phase. In these alloys, a fall in the relaxation strength was associated with competition of ω and α-phases, which consume a portion of the existing β-phase. In addition, a behavior of Snoek-type relaxation was identified after these peaks overlap, to form one peak, which contains components of phase transformations and interactions of interstitial solutes. The addition of 3 wt.% Fe, as well as increasing Nb content, induced a fall in the elastic modulus, since the β-phase was more stable. From the ∆E/E curve, retarded growth of ω-phase

was seen, since the onset temperature for hardening shifts to higher temperatures with increasing Nb content. The microstructural changes from the initial rapidly solidified condition were reflected in the elastic modulus values and hardness. For the alloy with the highest addition of β-stabilizing elements (Ti-25Nb-3Fe), the elastic modulus value was 65 GPa in the rapidly quenched condition (β-phase) and 68 GPa after three thermal cycles (β+α phases). In general, Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) alloys, rapidly solidified showed low values of elastic modulus with a relatively low addition of the isomorphous β-stabilizing element (Nb), combined with addition of the eutectoid β-stabilizing element (Fe), relative to the commercial metallic biomedical alloys (stainless steel, Co-Cr-Mo, CP-Ti and Ti-6Al-4V) employed as biomaterials for orthopedic implants. This makes Ti-Nb-Fe alloys an attractive system for applications such as metallic alloys for biomedical implants, in view of the lower melting point and reduced cost associated with the lower addition of isomorphous β-stabilizing element (The noble metal Nb).

Acknowledgments The authors would like to thank the FAPESP, CNPq and the CAPES for the financial support.

References Abdel-Hady Gepreel, M., Niinomi, M., 2013. Biocompatibility of Ti-alloys for long-term implantation. J. Mech. Behav. Biomed. Mater. 20, 407-415. Afonso, C.R., Ferrandini, P.L., Ramirez, A.J., Caram, R., 2010. High resolution transmission electron microscopy study of the hardening mechanism through phase separation in a beta-Ti-35Nb-7Zr-5Ta alloy for implant applications. Acta Biomater. 6, 1625-1629. Almeida, L.H., Grandini, C.R., Caram, R., 2009. Anelastic spectroscopy in a Ti alloy used as biomaterial. Mater. Sci. Eng. A 521-22, 59-62. Baker C, 1971. The Shape-Memory Effect in a Titanium-35 wt.-% Niobium Alloy. Met. Sci. 5, 92-100. Bertrand, E., Castany, P., Gloriant, T., 2013. Investigation of the martensitic transformation and the damping behavior of a superelastic Ti–Ta–Nb alloy. Acta Mater. 61, 511-518. Boyer R., Collings E.W., Welsch G., 1994. Materials Properties Handbook: Titanium Alloys. ASM International. Chaves, J.M., Florêncio, O., Silva, P.S., Marques, P.W.B., Schneider, S.G., 2014. Anelastic relaxation associated to phase transformations and interstitial atoms in the Ti–35Nb–7Zr alloy. J. Alloy Comp. 616, 420-425.

Collings, E.W., 1984. The physical metallurgy of titanium alloys. Metals Park, OH : American Society for Metals, Ohio. Collings, E.W., 1986. Applied Superconductivity Metallurgy and Physics of Titanium Alloys. Plenum Press, New York and London. Elmay, W., Prima, F., Gloriant, T., Bolle, B., Zhong, Y., Patoor, E., Laheurte, P., 2013. Effects of thermomechanical process on the microstructure and mechanical properties of a fully martensitic titanium-based biomedical alloy. J. Mech. Behav. Biomed. Mater. 18, 47-56. Ikehata, H., Nagasako, N., Furuta, T., Fukumoto, A., Miwa, K., Saito, T., 2004. Firstprinciples calculations for development of low elastic modulus Ti alloys. Phys. Rev. B 70, 174113. Kent, D., Wang, G., Dargusch, M., 2013. Effects of phase stability and processing on the mechanical properties of Ti-Nb based beta Ti alloys. J. Mech. Behav. Biomed. Mater. 28, 15-25. Kim, H.Y., Ikehara, Y., Kim, J.I., Hosoda, H., Miyazaki, S., 2006. Martensitic transformation, shape memory effect and superelasticity of Ti–Nb binary alloys. Acta Mater. 54, 2419-2429. Lee, C.M., Ho, W.F., Ju, C.P., Chern Lin, J.H., 2002a. Structure and properties of Titanium25 Niobium-x iron alloys. J. Mater. Sci. - Mater. Med. 13, 695-700. Lee, C.M., Ju, C.P., Chern Lin, J.H., 2002b. Structure-property relationship of cast Ti-Nb alloys. J. Oral Rehabil. 29, 314-322. Li, C., Lee, D.-G., Mi, X., Ye, W., Hui, S., Lee, Y., 2013. Phase transformation and age hardening behavior of new Ti–9.2Mo–2Fe alloy. J. Alloy Comp. 549, 152-157. Lin, D.J., Lin, J.H., Ju, C.P., 2002. Structure and properties of Ti-7.5Mo-xFe alloys. Biomaterials 23, 1723-1730. Long, M., Rack, H.J., 1998. Titanium alloys in total joint replacement--a materials science perspective. Biomaterials 19, 1621-1639. Lopes, E.S.N., Cremasco, A., Afonso, C.R.M., Caram, R., 2011. Effects of double aging heat treatment on the microstructure, Vickers hardness and elastic modulus of Ti–Nb alloys. Mater. Charact. 62, 673-680. Lu, H., Li, C.X., Yin, F.X., Fang, Q.F., Umezawa, O., 2012. Effects of alloying elements on the Snoek-type relaxation in Ti–Nb–X–O alloys (X=Al, Sn, Cr, and Mn). Mater. Sci. Eng. A 541, 28-32. Mantani, Y., Tajima, M., 2006. Phase transformation of quenched α″ martensite by aging in Ti–Nb alloys. Mater. Sci. Eng. A 438-440, 315-319. Moffat, D.L., Larbalestier, D.C., 1988. The competition between the alpha and omega phase in aged Ti-Nb alloys. Metall. Trans. A 19, 1687-1694.

Niinomi, M., 2002. Recent metallic materials for biomedical applications. Metall. Mater. Trans. A 33, 477-486. Niinomi, M., 2008. Mechanical biocompatibilities of titanium alloys for biomedical applications. J. Mech. Behav. Biomed. Mater. 1, 30-42. Niinomi, M., Nakai, M., 2011. Titanium-Based Biomaterials for Preventing Stress Shielding between Implant Devices and Bone. Int. J. Biomater. 2011, 836587. Nowick, A.S., Berry, B.S., 1972. Anelastic Relaxation in Crystalline Solids. Academic Press, Now York and London. Ohmori, Y., Ogo, T., Nakai, K., Kobayashi, S., 2001. Effects of ω-phase precipitation on

β→α, α′′ transformations in a metastable β titanium alloy. Mater. Sci. Eng. A 312, 182-188. Reed-Hill, R.E., 1973. Physical Metallurgy Principles. PWS-Kent Publishing, Van Nost. Reinhold, U. S. Schaller, R., Fantozzi, G., Gremaud, G., 2001. Mechanical Spectroscopy Q-1 2001. Trans Tech Publications LTD, Laubisrutistr, CH. Tane, M., Nakano, T., Kuramoto, S., Niinomi, M., Takesue, N., Nakajima, H., 2013. ω Transformation in cold-worked Ti–Nb–Ta–Zr–O alloys with low body-centered cubic phase stability and its correlation with their elastic properties. Acta Mater. 61, 139-150. Tobe, H., Kim, H.Y., Inamura, T., Hosoda, H., Nam, T.H., Miyazaki, S., 2013. Effect of Nb content on deformation behavior and shape memory properties of Ti–Nb alloys. J. Alloy Comp. 577, S435-S438. Yin, F., Iwasaki, S., Ping, D., Nagai, K., 2006. Snoek-Type High-Damping Alloys Realized in β-Ti Alloys with High Oxygen Solid Solution. Adv. Mater. 18, 1541-1544. Zhou, Z.C., Xiong, J.Y., Gu, S.Y., Yang, D.K., Yan, Y.J., Du, J., 2011. Anelastic relaxation caused by interstitial atoms in β-type sintered Ti–Nb alloys. J. Alloy. Comp. 509, 7356-7360.

Figure captions Fig. 1 - SEM micrographs of rapidly solidified alloys, (a) Ti-10Nb-3Fe (b) Ti-15Nb-3Fe (c) Ti20Nb-3Fe and (d) Ti-25Nb-3Fe, in the as-quenched condition showing a microstructure becoming more refined with rising Nb content. Fig. 2 - XRD patterns of (a) Ti-10Nb-3Fe, (b) Ti-15Nb-3Fe, (c) Ti-20Nb-3Fe and (d) Ti-25Nb3Fe rapidly solidified alloys in the as-quenched condition. The inset depict vertically zoomed portions in a limited 2θ range, showing the presence of ω-phase for samples with lower Nb content. Fig 3 - SEM images in BSE (backscattered electrons) mode show the solute distribution in dendritic microstructure for (a) Ti-10Nb-3Fe and (b) Ti-25Nb-3Fe alloys, with EDS analysis in the inner region of the dendrites (brighter)) and in the interdendritic region (darker). Fig. 4 - DSC thermograms of rapidly solidified Ti-xNb-3Fe (x=10, 15, 20 and 25 wt.%) alloys in the as-quenched condition, showing the phase transformation during two heating cycles at a heating rate of 1 K/min. Fig. 5 - Elastic modulus at room temperature plotted against electron concentration (e/a) for the Ti-xNb-3Fe (
), Ti-xNb () and Ti-TM (----) systems. Fig. 6 - TEM micrographs of as-quenched Ti-10Nb-3Fe sample showing a) bright field (BF) and b) dark field (DF) images of nanometric ω-phase precipitate dispersed in β-Ti (bcc) grain matrix, together with respective selected area electron diffraction (SAED) pattern with orientation relationship of zone axis [2 1 2]β//[001]ω (inset). Fig. 7 - TEM micrographs of as-quenched Ti-25Nb-3Fe sample of β-Ti (bcc) grain boundary showing a) bright field (BF) and b) dark field (DF) images of nanometric ω-phase precipitate dispersed in β-Ti matrix, together with respective selected area electron diffraction (SAED) pattern with orientation relationship of zone axis [1 1 5]β (inset). Fig. 8 - Anelastic relaxation spectra for the alloy Ti-10Nb-3Fe, during (a) first, (b) second and (c) third heating cycles and (d) XRD patterns showing the variation of phases in each cycle.

Fig. 9 - Anelastic relaxation spectra for the alloy Ti-15Nb-3Fe, Ti 3Fe, during (a) first, (b) second and (c) third heating cycles and (d) XRD patterns showing the variation of phases in each cycle. Fig. 10 - Anelastic relaxation spectra for the alloy Ti Ti-20Nb-3Fe, 3Fe, during (a) first, (b) second and (c) third heating cycles and (d) XRD patterns showing the variation of phases in each cycle. Fig. 11 - Anelastic relaxation spectra for the alloy Ti Ti-25Nb-3Fe, 3Fe, during (a) first, (b) second and (c) third heating cycles and (d) XRD patterns showing showing the variation of phases in each cycle. Fig. 12 - Dependence of strength of anelastic relaxation peak on Nb content in Ti Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) alloys.

Fig. 1

Fig. 2

Fig. 3

Fig. 4

Fig. 5

[2 1 2]β

a

b

(-1 -2 1)β (-2 -1 1)β ω2

(-1 1 0)β ω1

50 nm

50 nm

Fig. 6

a

b

[1 1 5]β

50 nm

50 nm

Fig. 7

Fig. 8

Fig. 9

Fig. 10

Fig. 11

Fig. 12

Table captions Table 1. Chemical analyses for Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) alloys. Table 2. Concentration interstitial solute atoms in the Ti-xNb-3Fe (x= 10, 15, 20, 25 wt.%) alloys. Table 3. Elastic modulus (E), Vickers hardness (HV) and phases presents in the alloys Ti-xNb3Fe (x= 10, 15, 20, 25 wt.%) in each thermal cycle.

Table 1. Nominal Composition

EDS [wt.%] Nb

Ti

Ti-10Nb-3Fe

86.9 ± 0.2

Ti-15Nb-3Fe

80.5 ± 0.2

Ti-20Nb-3Fe

75.1 ± 0.2

Ti-25Nb-3Fe

71.3 ± 0.2

Fe

10.2 ± 0.1

2.9 ± 0.1

16.8 ± 0.2

2.7 ± 0.1

21.9 ± 0.2

3.0 ± 0.1

26.0 ± 0.2

2.7 ± 0.1

Table 2. Alloy

Oxygen [wt.%]

Nitrogen [wt.%]

Ti-10Nb-3Fe

0.1520±0.0030

0.0032±0.0001

Ti-15Nb-3Fe

0.1960±0.0039

0.0038±0.0001

Ti-20Nb-3Fe

0.1600±0.0032

0.0027±0.0001

Ti-25Nb-3Fe

0.1820±0.0036

0.0037±0.0001

Table 3. Ti-10Nb-3Fe

Ti-15Nb-3Fe

Ti-20Nb-3Fe

Ti-25Nb-3Fe

Cycle number

E GPa

HV

Phases

E GPa

HV

Phases

E GPa

HV

Phases

E GPa

HV

Phases

As cast

77± 6

479± 4

β+ω

71± 4

411± 5

β+ω

67± 5

326± 3

β

65± 6

289± 3

β

1

95± 8

332± 3

β+ω

67± 4

314± 3

β+ω

84± 6

341± 4

β+α

69± 6

304± 3

β+α

2

84± 7

331± 3

β+α

68± 4

336± 3

β+α

82± 6

348± 4

β+α

76± 7

310± 3

β+α

3

84± 7

331± 3

β+α

67± 4

328± 3

β+α

82± 6

354± 4

β+α

79± 7

294± 3

β+α