- Email: [email protected]
Alpha phase precipitation in Ti-30Nb-1Fe alloys – phase transformations in continuous heating and aging heat treatments
Materials Science & Engineering A 677 (2016) 222–229
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Alpha phase precipitation in Ti-30Nb-1Fe alloys – phase transformations in continuous heating and aging heat treatments Fernando Henrique da Costa, Camilo A.F. Salvador, Mariana G. de Mello, Rubens Caram n University of Campinas (UNICAMP), School of Mechanical Engineering, Campinas, SP 13083-860, Brazil
art ic l e i nf o Article history: Received 11 July 2016 Received in revised form 5 September 2016 Accepted 6 September 2016 Keywords: Titanium alloys Phase transformations Aging heat treatments Differential scanning calorimetry
a b s t r a c t Microstructure and mechanical properties of β metastable Ti-30Nb and Ti-30Nb-1Fe alloys subjected to different aging heat treatments were investigated. Various temperatures and heating rates were employed in order to determine their influence on α phase precipitation. The alloys were characterized by means of differential scanning calorimetry (DSC), X-ray diffraction (XRD), scanning and transmission electron microscopy. Several mechanical properties were evaluated, such as Vickers hardness and elastic modulus, and tensile tests were also performed. The results indicate that the addition of Fe reduces the α phase precipitation and ω dissolution temperatures, although ω phase was still detected in Ti-30Nb-1Fe aged at 400 °C. The addition of Fe caused refinement of α phase precipitates, probably due to the higher diffusion coefficient of Fe. As for the influence of the heating rate, 2 and 30 °C/min did not cause significant microstructural changes. Finally, α phase refinement in Ti-30Nb alloy was favored by aging at 550 °C, at a heating rate of 600 °C/min. & 2016 Elsevier B.V. All rights reserved.
1. Introduction The use of pure titanium and titanium alloys as engineering materials increased significantly during the second half of the 20th century. This growth was due to the development of new titanium alloys with improved mechanical properties [1]. Titanium alloys include β metastable alloys, whose wide range of mechanical properties make them interesting for industrial applications in general, as they enable combinations of α and β phases in different ratios and morphologies, depending on the processing route [2]. A promising form of heat treatment for these alloys involves aging in the β þ α phase field, aiming to produce intragranular αphase precipitates distributed homogeneously in the β phase matrix. However, it is well known that isothermal ω phase and α phase compete to precipitate in Ti-Nb based alloys. Moffat and Larbalestier [3] demonstrated that aging prequenched Ti-20Nb (at%) alloys at 300, 350 and 400 °C for 4 days enabled the growth of ω precipitates to the exclusion of α precipitates, although a marked acceleration of α precipitation using ω precipitates as nucleation sites occurs at 450 and 500 °C in Ti-20Nb (at%), according to Kobayashi et al. [4]. This heterogeneous α phase nucleation mechanism has been discussed in detail by Prima et al. [5], and more recently, Zheng et al. demonstrated that it is not favored by higher heating rates ( 100 °C/min), although it plays n
Corresponding author. E-mail address: [email protected] (R. Caram).
http://dx.doi.org/10.1016/j.msea.2016.09.023 0921-5093/& 2016 Elsevier B.V. All rights reserved.
an important role in Ti-5553 alloy when slowly heated at 20 °C/ min to 600 °C [6]. Moreover, Ivasishin et al. [7] reported results demonstrating that the heating rate influences α precipitation in several Ti alloys, such as VT22, TIMETAL LCB and Ti-15-3-3-3. Although a microstructure of fine well-dispersed α precipitates has a slightly higher elastic modulus than that of the parent β phase obtained by water-quenching, appreciably higher yields and ultimate tensile strengths, as well as remarkable ductility, can be obtained in β þ α microstructures [8–10]. Another way to increase the strength of Ti alloys is to add elements with a remarkable solid solution strengthening effect on the crystal lattice, such as Ta, Nb, Zr [11], Mo, Si, Sn, and particularly Fe [12]. One atomic percent of Fe can increase the yield strength of binary Ti-22Nb (at%) by nearly 200 MPa [13]. Despite previous reports by Lee et al. [14], Hsu et al. [15] and Lopes et al. [13] about the addition of iron to Ti-Nb alloys in particular, none of the aforementioned researchers evaluated the influence of iron on continuous phase transformations in these modified alloys, i.e., during continuous heating at several heating rates or during isothermal aging heat treatments. In addition to the solid solution effect, our experiments suggest that Fe also enables α phase to form earlier during aging, which is interesting from a technological standpoint. In this work, an evaluation was made of Ti-30Nb and Ti-30Nb-1Fe alloys aged at 400, 500 and 550 °C. The aging heat treatments were performed at heating rates of 2, 30 and 600 °C min 1.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
2. Experimental Hundred-gram ingots of Ti-30Nb and Ti-30Nb-1Fe (wt%) alloys, approx. 12 mm thick, were arc melted under an argon atmosphere (99.9%). They were encapsulated in quartz tubes filled with argon for 24 h of homogenization heat treatment at 1000 °C, followed by Table 1 Composition of the experimental alloys. Alloy
Ti
Nb (wt%)
Fe (wt%)
O (wt%)
N (wt%)
Ti-30Nb Ti-30Nb-1Fe
Balance
28.9 29.5
0.2 1.3
0.18 0.14
0.0092 0.0083
223
furnace cooling. The Ti-30Nb and Ti-30Nb-1Fe ingots were then hot rolled at 1000 °C to produce 3 mm thick sheets. Specimens of 1 cm2 were cut from the sheets and subjected to chemical pickling to remove surface oxide layer produced by hot rolling. This procedure reduced the thickness of the specimens by about 0.5 mm. The composition was analyzed by X-ray fluorescence (XRF) in a Rigaku RIX 3100 spectrometer, and interstitial O and N contents were examined using a LECO 400 analyzer (Table 1). All the samples were solution-treated at 1000 °C for 1 h, waterquenched, and then subjected to various aging heat treatments at
Fig. 1. Schematic diagram of tensile tests.
Fig. 4. DSC results of samples aged at 300 °C and 600 °C for 24 h.
Table 2 Phase transformation temperatures determined by DSC measurements. Dominating process
Fig. 2. XRD patterns of samples aged at 300 °C and 600 °C for 24 h.
Suggested phase transformation
Endo
α“-β
Endo
ω-β
Exo
β/ω-α
Endo
α-β
Aging temperature
300 °C 600 °C 300 °C 600 °C 300 °C 600 °C 300 °C 600 °C
Fig. 3. BSE-SEM images of α phase precipitation after aging heat treatment at 600 °C for 24 h: (a) Ti-30Nb, (b) Ti-30Nb-1Fe.
Temperature range Ti-30Nb
Ti-30Nb1Fe
360–402 – 439–566 418–483 550 – – 614–752
– – 372–529 418–491 510 – – 616–749
224
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Fig. 5. TEM dark field (DF: a, b, d, e) and selected area diffraction (SAD: c, f) patterns of Ti-30Nb-1Fe alloy aged at 400 °C for 1 h showing massive formation of small ω precipitates, which appear with an elongated morphology through [315] beta zone axis (f).
Fig. 6. BSE-SEM images of α phase precipitation after aging heat treatment at 500 °C for 1 h: (a–c) Ti-30Nb, (b–d) Ti-30Nb-1Fe.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
400, 500 and 550 °C for 1 h, applying heating rates of 2 and 30 °C/ min, followed by water quenching (WQ). Metallographic specimens were prepared by sanding with 1500 grit sandpaper and polishing with 0.02 mm colloidal silica. Scanning electron microscopy (SEM) was performed in a FEI Inspect F50 and a Hitachi TM1000 microscopes, while transmission electron microscopy (TEM) analyses were performed using a JEOL 3010 microscope operating at 300 kV. TEM thin foils were prepared via conventional dimpling and argon ion milling procedures. For a detailed microstructural analysis, X-ray diffraction (XRD) experiments were performed under copper radiation (λ ¼0.15406 nm) in a PANalytical X-Pert Pro diffractometer operating at 40 kV and 30 mA. The diffractometer was equipped with a PIXcel fast detector and a spinner sample holder. Vickers hardness measurements were taken with a Buehler Vickers 2100 hardness tester, applying an indentation load of 500 g.f. for 15 s. Two additional 24 h heat treatments were performed at 300 °C and 600 °C in order to produce equilibrium microstructures composed of β þ ω (300 °C) and β þ α (600 °C) for evaluation by differential scanning calorimetry (Netzsch STA409). DSC characterization was carried out under an argon flux at a heating rate of 30 °C/min. Lastly, the mechanical properties of microstructures subjected to aging were evaluated by tensile testing in an MTS 810 universal testing machine, applying a strain rate of 1 mm/min. Specimens with a gage section of 3 mm 6 mm were carefully prepared by waterjet cutting and machining, as illustrated in Fig. 1. The extensometer was not used due to the small size of the specimens.
225
about 40 °C the temperature at which α precipitation begins in Ti30Nb and Ti-30Nb-1Fe alloys. This is a key point to evaluate the results obtained by aging. 3.2. Aging heat treatments and
α phase precipitation
Considering the transformation temperatures of the two alloys described in Section 3.1, if the aging temperature were lower than 430 °C, the amount of ω phase would be expected to increase with aging time owing to the formation of isothermal ω (ωiso). At temperatures above 500 °C, α precipitation should prevail. On the other hand, α phase begins to transform into β phase close to 650 °C (near beta transus), so the optimal aging heat treatment temperature may be lower than that. Therefore, aging heat treatment temperatures of 500 and 550 °C were chosen, since copious α precipitation is expected to occur within this temperature interval. These were the typical aging temperatures employed in several studies on Ti-Nb based alloys [1,19,20]. The TEM images in Fig. 5 show that ω phase was still detectable in the Ti-30Nb-1Fe alloy aged at 400 °C for 1 h, which is consistent with the DSC results. Moreover, at this temperature, heterogeneous α phase precipitation appeared to begin at the β/ω interfaces, given the ω phase elongated morphology (Fig. 5d) – even
3. Results 3.1. Phase transformation temperatures Barriobero et al. [16] successfully combined in situ high energy X-ray diffraction (HEXRD) and DSC data for Ti-10V-2Fe-3Al alloy. To this end, they heated specimens continuously at the same heating rates (5, 20 and 50 °C min 1) in both experiments. They then made a detailed description of the exothermic and the endothermic events (DSC) associated with each phase transformation (XRD), as a guide for future evaluations of phase transformation temperatures in other alloy systems. These results can certainly be extrapolated and compared to the Ti-Nb system, as shown by Bönisch et al. [17]. In this work, to make a specific phase transformation more visible for detection by DSC, we performed two aging heat treatments at 300 °C and at 600 °C for 24 h, so as to force the detection of the ω-solvus temperature and the β transus, respectively. The same procedure was applied to both Ti-30Nb and Ti-30Nb-1Fe alloys, and the results (XRD, SEM and DSC) are depicted in Figs. 2–4, respectively. A summary is presented in Table 2. The Ti-30Nb alloy sample aged at 300 °C showed α“-β transformation (endothermic) followed by another endothermic peak associated with ω-β transformation, beginning at 439 °C. An adjacent exothermic peak corresponding to the copious nucleation of α phase that occurs in ωiso precipitates was visible at close to 550 °C. In the analysis of the sample aged at 600 °C, a sharp peak corresponding to the beta transus temperature (α-β) was detected, starting at 614 °C and ending at 752 °C. Both temperatures are in accordance with the phase diagram of Ti-Nb based alloys simulated by Zhang et al. [18]. In the Ti-30Nb-1Fe alloy, no significant amount of α“ was detected in the initial condition, so the first endothermic peak pertains to β-ω transformation, which starts at 372 °C, while α precipitation starts close to 510 °C. The two alloys appear to fall within in a similar range of beta transus temperatures. An interesting point noted in the DSC analysis (samples aged at 300 °C) is that 1 wt% of Fe reduces by
Fig. 7. XRD patterns of Ti-30Nb and Ti-30Nb-1Fe samples aged at for 1 h at 500 °C and 550 °C, applying heating rates of 2 and 30 °C/min.
226
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
taking into account that α phase diffraction was not clearly detected in the SAD diffraction patterns shown in Fig. 5c and f. At 400 °C, α phase precipitation appeared to occur slowly, non-optimally, owing to the small size of the precipitates and their heterogeneous distribution. Compared to previous works, it can be stated that these ωiso precipitates are slightly larger than the ωath precipitates produced by WQ in this alloy [13]. The SEM images in Fig. 6 correspond to the samples aged at 500 °C at two heating rates, 2 and 30 °C/min. Aging Ti-30Nb alloy at a heating rate of 30 °C/min resulted in the formation of ω phase, detectable by X-ray diffraction (Fig. 7). This phase was probably formed upon water-quenching (performed after aging) from a compositional pocket or solute-lean β matrix. According to the DSC data (Fig. 4), it may also be residual isothermal ω that did not
decompose at 500 °C due to the short heating ramp, which competes with α phase precipitation, making α precipitates sparse [3]. In the Ti-30Nb-1Fe sample, since 30 °C/min is the same heating rate as that applied in the DSC experiments, considerable α phase formation would be expected at 500 °C, which is confirmed in Fig. 6b. On the other hand, no previous DSC information was available for the lower heating rate (2 °C/min). However, based on SEM and XRD data, the Ti-30Nb-1Fe sample appeared to present fewer α phase precipitates compared to those produced at the higher heating rate. Moreover, copious α phase was formed in the Ti-30Nb samples owing to the longer heating ramp, which enables the diffusion of Nb that results in these long α laths (Fig. 6c). With regard to the aging heat treatments performed at 550 °C, Ti-30Nb alloy should only undergo massive α formation close to
Fig. 8. BSE-SEM images of α phase precipitation after aging heat treatment at 550 °C for 1 h: (a–c) Ti-30Nb, (b–d) Ti-30Nb-1Fe.
Fig. 9. BSE-SEM images of α phase precipitation after aging heat treatment at 550 °C for 1 h, at a heating rate of 600 °C/min: (a) Ti-30Nb, (b) Ti-30Nb-1Fe.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
this temperature, according to Table 2, which in fact occurred, as shown in Fig. 8a and c; additionally, a small fraction of α” formed in response to WQ was also detected. Bönisch et al., who investigated Ti-29Nb (wt%) aged at 545 °C followed by water quenching, also obtained α”, α and β phases [17]. Banumathy et al. studied the effects of 24 h aging heat treatments at 450 and 550 °C on Ti-27Nb (wt%) alloy, followed by WQ. They reported the formation of martensite at 450 °C, but not at 550 °C [21]. These results suggest that α” formation occurs only in the initial stages of α phase precipitation [22]. Ti-30Nb-1Fe samples likewise presented β þ α microstructures when aged at 550 °C. After a high density of α phase precipitates at 550 °C was reached, regardless of the heating rate, we found that, notwithstanding the findings of some authors about the strong influence of the heating rate on the size of α precipitates [6], in our preliminary study with 2 and 30 °C/min, the experimental alloys only appeared to be sensitive to the heating rate when aged at lower temperatures, i.e., below or close to 500 °C. For a more detailed analysis of the influence of the heating rate on the microstructure, we designed an additional experiment at a preset temperature in a tubular furnace, quickly inserting the samples in the furnace once it had reached this temperature. These samples were subjected to a much higher heating rate of about 600 °C/min. Fig. 9 shows SEM images of samples subjected to a heating rate of 600 °C/min to 550 °C, isothermally aged for 1 h, followed by WQ. The Ti-30Nb-1Fe alloy (Fig. 9b) appears to contain fewer precipitates compared to those produced at the other heating rates (Figs. 6 and 8), as was expected based on the literature [7]. In this case, ω-assisted α precipitation was somewhat inhibited by the high heating rates (430 °C/min) [23,24]. According to Tong Li et al., oxygen rich sites associated with ωiso may play a major role in this fine-scaled α precipitation [25], but it takes longer times at lower temperatures for these zones to be formed. Unlike Ti-30Nb-1Fe, the Ti-30Nb sample exhibited an unusual behavior. The higher the heating rate applied to the Ti-30Nb alloy the smaller the size of the precipitates, as can be seen by comparing Fig. 9a with Fig. 8a and c. In this case, these fine and well dispersed α precipitates were probably generated by another mechanism, since the initial microstructure of Ti-30Nb is composed of β þ α”, and we believe that heating occurs so fast that there is likely no ωiso phase to support ω-assisted α formation (which would also lead to the same behavior as that of 1Fe alloy). Considering this fine scale α phase precipitation, it would be possible that this microstructure was produced by a non-classical mechanism [26]. However, based on thermodynamic MatCalc [27] simulations, we concluded that Ti-Nb based alloys with Nb content higher than 15 wt% could not undergo pseudospinodal decomposition at 550 °C, because that would require a very significant compositional fluctuation, due to a wide gap between the Gibbs free energy curves at this temperature (Fig. 10). To date, the mechanism discovered by Nag et al. and Boyne et al. has been only detected in Ti–Mo based alloys, such as Timetal-5553 [26,28]. Therefore, we believe that massive and fine α precipitation, in this particular case, must occur via another mechanism, such as heterogeneous nucleation at dislocations, twin boundaries or other lattice defects resulting from the rapid decomposition of α” phase. In essence, α precipitation in Ti-30Nb-1Fe alloy is influenced very little by the heating rate. In addition to altering some transformation temperatures (Table 2), we believe that α formation occurs in the heat treatment route in Ti-30Nb-1Fe due to the faster diffusion of Fe in the Ti-Nb matrix. Based on the diffusion coefficients found in Smithells Metals Reference Book [29], we estimated that Nb diffusion in Ti at 550 °C (823 K) occurs with D ¼7.6 10 19 m2/s, while Fe diffusion in Ti is almost 90 times faster than that, reaching D¼6.5 10 17 m2/s. Thus, Fe alters the
227
kinetics of α precipitation, making it viable at any heating rate. The rejection of Fe from α phase also leads to enrichment of the neighbor β phase in Fe, restricting the growth of α laths and resulting in consistently small α precipitates. On the other hand, in Ti-30Nb alloy, an exceptionally refined α precipitation was observed only at the highest heating rate (600 °C/min).
Fig. 10. Gibbs free energy curves of β and α phases containing Nb content at 550 °C. The optimal composition for the pseudospinodal is indicated as 14.9 wt%.
Table 3 Vickers hardness, elastic modulus and phases detected by XRD in the Ti-30Nb and Ti-30Nb-1Fe samples aged at 500 and 550 °C for 1 h, at heating rates of 2 and 30 °C/ min.
Ti-30Nb
Heating rate (°C/ min)
Temperature (°C) – Heating rate (°C/ min)
Vickers hardness (HV)
Elastic modulus (GPa)
2
400 500 550 400 500 550
428 7 9 2747 11 219 73 4217 13 269 7 6 210 73
1167 1 797 1 727 1 1137 2 1157 2 697 1
400 500 550 400 500 550
423 7 5 2617 8 2317 4 399 7 16 251 720 226 7 8
1157 2 927 1 807 2 1067 1 987 3 827 1
30
Ti-30Nb-1Fe 2
30
Fig. 11. Stress-strain curves of the Ti-30Nb and Ti-30Nb-1Fe samples aged at 550 °C for 1 h, at a heating rate of 2 °C/min. WQ samples studied by Lopes et al. [13].
228
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Fig. 12. Fracture surface of the samples aged at 550 °C for 1 h, at a heating rate of 30 °C/min: (a) Ti-30Nb; (b) Ti-30Nb-1Fe.
3.3. Mechanical properties The results of Vickers hardness and elastic modulus of the aged alloys are listed in Table 3. Note that these results reflect the presence of different phases in the microstructure. The addition of Fe limits the elastic modulus to a minimum of 80 GPa, which is undesirable for biomedical applications. Ti-30Nb heated to 550 °C at 30 °C/min showed the lowest elastic modulus due to the formation of α” phase after WQ. Tensile tests were performed to evaluate the mechanical properties of aged Ti-30Nb and Ti-30Nb-1Fe samples (Fig. 11). The mechanical properties of solution heat-treated and water quenched Ti-30Nb and Ti-30Nb-1Fe alloys were investigated by Lopes et al. and their curves are plotted in Fig. 11 for comparison [13]. The aging heat treatments for tensile testing were carried out at a temperature of 550 °C, heating rate of 2 °C/min, and isothermal holding time of 1 h. As expected, α phase precipitation during the aging heat treatment increased the ultimate tensile strength of both alloys. Although the heat treatment decreased elongation (strain), the values obtained are still acceptable, resulting in mixed-mode fracture with prevalence of dimples, as can be seen in Fig. 12, with dimpled fracture surfaces. The behavior of the tensile test curve of Ti-30Nb differed from that of the alloy containing 1% of Fe. Fig. 11 shows this curve with the double yield point, which means that this alloy presents shape memory effect and superelasticity [30].
4. Conclusions In this work, Ti-Nb and Ti-Nb-Fe alloys were subjected to various aging heat treatments. The results indicate a correlation between the mechanical properties and the precipitated phases. The main results were as follows: – DSC analyses revealed peaks corresponding to the decomposition of α”, ω and α phases. The results suggest that the addition of 1 wt% Fe allows α phase nucleation at lower temperatures. – The two alloys reach maximum hardness upon precipitation of isothermal ω phase. The highest hardness values were found in samples aged at 400 °C. – The addition of 1 wt% Fe refines the precipitation of α phase in comparison to that of the base alloy, probably due to the higher diffusion coefficient of Fe. – Reducing the heating rate from 30 °C/min to 2 °C/min did not lead to the refinement of α phase precipitates in either of the alloys. – For Ti-30Nb alloy, the heating rate of 600 °C/min resulted in the
finest α phase precipitates, which can be explained by the different mechanism activated during martensite decomposition. In the case of Ti-30Nb-1Fe alloy, the heating rate did not interfere in the refinement of α phase precipitates.
Acknowledgments The authors gratefully acknowledge the financial support of the Brazilian research funding agencies CAPES (Federal Agency for the Support and Improvement of Higher Education), CNPq (National Council for Scientific and Technological Development) and FAPESP (São Paulo Research Foundation) (Grants #2013/13867-2 and #2014/06099-1).
References [1] D. Banerjee, J.C. Williams, Perspectives on titanium science and technology, Acta Mater. 61 (2013) 844–879, http://dx.doi.org/10.1016/j. actamat.2012.10.043. [2] M. Niinomi, M. Nakai, J. Hieda, Development of new metallic alloys for biomedical applications, Acta Biomater. 8 (2012) 3888–3903, http://dx.doi.org/ 10.1016/j.actbio.2012.06.037. [3] D.L. Moffat, D.C. Larbalestier, The competition between the alpha and omega phases in aged Ti-Nb alloys, Metall. Trans. A 19 (1988) 1687–1694. [4] S. Kobayashi, T. Takeichi, K. Nakai, T. Sakamoto, Acceleration or suppression of α-phase precipitation using isothermal ω phase in Ti-20 at.pct Nb Alloy, Metall. Mater. Trans. A 45 (2013) 1217–1229, http://dx.doi.org/10.1007/ s11661-013-2092-1. [5] F. Prima, P. Vermaut, G. Texier, D. Ansel, T. Gloriant, Evidence of α-nanophase heterogeneous nucleation from ω particles in a β-metastable Ti-based alloy by high-resolution electron microscopy, Scr. Mater. 54 (2006) 645–648, http://dx. doi.org/10.1016/j.scriptamat.2005.10.024. [6] Y. Zheng, R.E.A. Williams, D. Wang, R. Shi, S. Nag, P. Kami, et al., Role of ω phase in the formation of extremely refined intragranular α precipitates in metastable β-titanium alloys, Acta Mater. 103 (2016) 850–858, http://dx.doi.org/ 10.1016/j.actamat.2015.11.020. [7] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Aging response of coarse- and fine-grained β titanium alloys, Mater. Sci. Eng. A 405 (2005) 296–305, http://dx.doi.org/10.1016/j.msea.2005.06.027. [8] P. Majumdar, S.B. Singh, M. Chakraborty, The role of heat treatment on microstructure and mechanical properties of Ti–13Zr–13Nb alloy for biomedical load bearing applications, J. Mech. Behav. Biomed. Mater. 4 (2011) 1132–1144, http://dx.doi.org/10.1016/j.jmbbm.2011.03.023. [9] E.S.N. Lopes, R.J. Contieri, S.T. Button, R. Caram, Femoral hip stem prosthesis made of graded elastic modulus metastable β Ti alloy, Mater. Des. 69 (2015) 30–36, http://dx.doi.org/10.1016/j.matdes.2014.11.040. [10] S. Nag, R. Banerjee, H.L. Fraser, Intra-granular alpha precipitation in Ti–Nb–Zr– Ta biomedical alloys, J. Mater. Sci. 44 (2008) 808–815, http://dx.doi.org/ 10.1007/s10853-008-3148-2. [11] A. Biesiekierski, J. Wang, M.A.-H. Gepreel, C. Wen, A new look at biomedical Tibased shape memory alloys, Acta Biomater. 8 (2012) 1661–1669, http://dx.doi. org/10.1016/j.actbio.2012.01.018. [12] D.C. Zhang, Y.F. Mao, Y.L. Li, J.J. Li, M. Yuan, J.G. Lin, Effect of ternary alloying elements on microstructure and superelastictity of Ti–Nb alloys, Mater. Sci.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Eng. A 559 (2013) 706–710, http://dx.doi.org/10.1016/j.msea.2012.09.012. [13] É.S.N. Lopes, C.A.F. Salvador, D.R. Andrade, A. Cremasco, K.N. Campo, R. Caram, Microstructure, mechanical properties, and electrochemical behavior of Ti-NbFe alloys applied as biomaterials, Metall. Mater. Trans. A 47 (2016) 3213–3226, http://dx.doi.org/10.1007/s11661-016-3411-0. [14] C.M. Lee, W.F. Ho, C.P. Ju, J.H. Chern Lin, Structure and properties of Titanium25 Niobium-x iron alloys, J. Mater. Sci. Mater. Med. 13 (2002) 695–700 http:// www.ncbi.nlm.nih.gov/pubmed/15348579. [15] H.-C. Hsu, S.-K. Hsu, S.-C. Wu, C.-J. Lee, W.-F. Ho, Structure and mechanical properties of as-cast Ti–5Nb–xFe alloys, Mater. Charact. 61 (2010) 851–858, http://dx.doi.org/10.1016/j.matchar.2010.05.003. [16] P. Barriobero-Vila, G. Requena, F. Warchomicka, et al., Phase transformation kinetics during continuous heating of a β-quenched Ti–10V–2Fe–3Al alloy, J. Mater. Sci. 50 (2015) 1412, http://dx.doi.org/10.1007/s10853-014-8701-6. [17] M. Bönisch, M. Calin, T. Waitz, A. Panigrahi, M. Zehetbauer, A. Gebert, et al., Thermal stability and phase transformations of martensitic Ti–Nb alloys, Sci. Technol. Adv. Mater. 14 (2013) 055004, http://dx.doi.org/10.1088/1468-6996/ 14/5/055004. [18] Y. Zhang, H. Liu, Z. Jin, Thermodynamic assessment of the Nb-Ti system, Calphad 25 (2001) 305–317, http://dx.doi.org/10.1016/S0364-5916(01)00051-7. [19] M. Geetha, A.K. Singh, A.K. Gogia, R. Asokamani, Effect of thermomechanical processing on evolution of various phases in Ti–Nb–Zr alloys, J. Alloy. Compd. 384 (2004) 131–144, http://dx.doi.org/10.1016/j.jallcom.2004.04.113. [20] L. Wang, Z. Lin, X. Wang, Q. Shi, W. Yin, D. Zhang, et al., Effect of aging treatment on microstructure and mechanical properties of Ti27Nb2Ta3Zr Titanium alloy for implant applications, Mater. Trans. 55 (2014) 141–146, http: //dx.doi.org/10.2320/matertrans.M2013187. [21] S. Banumathy, K.S. Prasad, R.K. Mandal, A.K. Singh, Effect of thermomechanical processing on evolution of various phases in Ti-Nb alloys, Bull. Mater. Sci. 34 (2011) 1421–1434, http://dx.doi.org/10.1007/s12034-011-0338-3. [22] E. Aeby-Gautier, A. Settefrati, F. Bruneseaux, B. Appolaire, B. Denand, M. Dehmas, et al., Isothermal α″ formation in β metastable titanium alloys, J.
[23]
[24]
[25]
[26]
[27]
[28]
[29]
[30]
229
Alloy. Compd. 577 (2013) S439–S443, http://dx.doi.org/10.1016/j. jallcom.2012.02.046. Y. Zheng, R.E.A. Williams, J.M. Sosa, T. Alam, Y. Wang, R. Banerjee, et al., The indirect influence of the ω phase on the degree of refinement of distributions of the α phase in metastable β-Titanium alloys, Acta Mater. 103 (2016) 165–173, http://dx.doi.org/10.1016/j.actamat.2015.09.053. Y. Zheng, R.E.A. Williams, J.M. Sosa, Y. Wang, R. Banerjee, H.L. Fraser, The role of the ω phase on the non-classical precipitation of the α phase in metastable β-titanium alloys, Scr. Mater. 111 (2016) 81–84, http://dx.doi.org/10.1016/j. scriptamat.2015.08.019. T. Li, D. Kent, G. Sha, J.M. Cairney, M.S. Dargusch, The role of ω in the precipitation of α in near-β Ti alloys, Scr. Mater. 117 (2016) 92–95, http://dx.doi. org/10.1016/j.scriptamat.2016.02.026. S. Nag, Y. Zheng, R.E. a. Williams, A. Devaraj, A. Boyne, Y. Wang, et al., Nonclassical homogeneous precipitation mediated by compositional fluctuations in titanium alloys, Acta Mater. 60 (2012) 6247–6256, http://dx.doi.org/10.1016/ j.actamat.2012.07.033. E. Kozeschnik, MatCalc––a simulation tool for multicomponent thermodynamics, diffusion and phase transformation kinetics, in: E. Kozeschnik, B. Buchmayr (Eds.), Mathematical Modelling of Weld Phenomena 5, 738, 2001, pp. 349–361. A. Boyne, D. Wang, R.P. Shi, Y. Zheng, A. Behera, S. Nag, et al., Pseudospinodal mechanism for fine α/β microstructures in β-Ti alloys, Acta Mater. 64 (2014) 188–197, http://dx.doi.org/10.1016/j.actamat.2013.10.026. Y. Sohn, Diffusion in metals, In: C.J. Smithells, (Ed), Smithells Metals Reference Book, 2004, Oxford: Elsevier, 1–120, http://dx.doi.org/10.1016/B978-075067509-3/ 50016-6. A.M.G. Tavares, W.S. Ramos, J.C.G. de Blas, E.S.N. Lopes, R. Caram, W.W. Batista, et al., Influence of Si addition on the microstructure and mechanical properties of Ti–35Nb alloy for applications in orthopedic implants, J. Mech. Behav. Biomed. Mater. 51 (2015) 74–87, http://dx.doi.org/10.1016/j. jmbbm.2015.06.035.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Alpha phase precipitation in Ti-30Nb-1Fe alloys – phase transformations in continuous heating and aging heat treatments Fernando Henrique da Costa, Camilo A.F. Salvador, Mariana G. de Mello, Rubens Caram n University of Campinas (UNICAMP), School of Mechanical Engineering, Campinas, SP 13083-860, Brazil
art ic l e i nf o Article history: Received 11 July 2016 Received in revised form 5 September 2016 Accepted 6 September 2016 Keywords: Titanium alloys Phase transformations Aging heat treatments Differential scanning calorimetry
a b s t r a c t Microstructure and mechanical properties of β metastable Ti-30Nb and Ti-30Nb-1Fe alloys subjected to different aging heat treatments were investigated. Various temperatures and heating rates were employed in order to determine their influence on α phase precipitation. The alloys were characterized by means of differential scanning calorimetry (DSC), X-ray diffraction (XRD), scanning and transmission electron microscopy. Several mechanical properties were evaluated, such as Vickers hardness and elastic modulus, and tensile tests were also performed. The results indicate that the addition of Fe reduces the α phase precipitation and ω dissolution temperatures, although ω phase was still detected in Ti-30Nb-1Fe aged at 400 °C. The addition of Fe caused refinement of α phase precipitates, probably due to the higher diffusion coefficient of Fe. As for the influence of the heating rate, 2 and 30 °C/min did not cause significant microstructural changes. Finally, α phase refinement in Ti-30Nb alloy was favored by aging at 550 °C, at a heating rate of 600 °C/min. & 2016 Elsevier B.V. All rights reserved.
1. Introduction The use of pure titanium and titanium alloys as engineering materials increased significantly during the second half of the 20th century. This growth was due to the development of new titanium alloys with improved mechanical properties [1]. Titanium alloys include β metastable alloys, whose wide range of mechanical properties make them interesting for industrial applications in general, as they enable combinations of α and β phases in different ratios and morphologies, depending on the processing route [2]. A promising form of heat treatment for these alloys involves aging in the β þ α phase field, aiming to produce intragranular αphase precipitates distributed homogeneously in the β phase matrix. However, it is well known that isothermal ω phase and α phase compete to precipitate in Ti-Nb based alloys. Moffat and Larbalestier [3] demonstrated that aging prequenched Ti-20Nb (at%) alloys at 300, 350 and 400 °C for 4 days enabled the growth of ω precipitates to the exclusion of α precipitates, although a marked acceleration of α precipitation using ω precipitates as nucleation sites occurs at 450 and 500 °C in Ti-20Nb (at%), according to Kobayashi et al. [4]. This heterogeneous α phase nucleation mechanism has been discussed in detail by Prima et al. [5], and more recently, Zheng et al. demonstrated that it is not favored by higher heating rates ( 100 °C/min), although it plays n
Corresponding author. E-mail address: [email protected] (R. Caram).
http://dx.doi.org/10.1016/j.msea.2016.09.023 0921-5093/& 2016 Elsevier B.V. All rights reserved.
an important role in Ti-5553 alloy when slowly heated at 20 °C/ min to 600 °C [6]. Moreover, Ivasishin et al. [7] reported results demonstrating that the heating rate influences α precipitation in several Ti alloys, such as VT22, TIMETAL LCB and Ti-15-3-3-3. Although a microstructure of fine well-dispersed α precipitates has a slightly higher elastic modulus than that of the parent β phase obtained by water-quenching, appreciably higher yields and ultimate tensile strengths, as well as remarkable ductility, can be obtained in β þ α microstructures [8–10]. Another way to increase the strength of Ti alloys is to add elements with a remarkable solid solution strengthening effect on the crystal lattice, such as Ta, Nb, Zr [11], Mo, Si, Sn, and particularly Fe [12]. One atomic percent of Fe can increase the yield strength of binary Ti-22Nb (at%) by nearly 200 MPa [13]. Despite previous reports by Lee et al. [14], Hsu et al. [15] and Lopes et al. [13] about the addition of iron to Ti-Nb alloys in particular, none of the aforementioned researchers evaluated the influence of iron on continuous phase transformations in these modified alloys, i.e., during continuous heating at several heating rates or during isothermal aging heat treatments. In addition to the solid solution effect, our experiments suggest that Fe also enables α phase to form earlier during aging, which is interesting from a technological standpoint. In this work, an evaluation was made of Ti-30Nb and Ti-30Nb-1Fe alloys aged at 400, 500 and 550 °C. The aging heat treatments were performed at heating rates of 2, 30 and 600 °C min 1.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
2. Experimental Hundred-gram ingots of Ti-30Nb and Ti-30Nb-1Fe (wt%) alloys, approx. 12 mm thick, were arc melted under an argon atmosphere (99.9%). They were encapsulated in quartz tubes filled with argon for 24 h of homogenization heat treatment at 1000 °C, followed by Table 1 Composition of the experimental alloys. Alloy
Ti
Nb (wt%)
Fe (wt%)
O (wt%)
N (wt%)
Ti-30Nb Ti-30Nb-1Fe
Balance
28.9 29.5
0.2 1.3
0.18 0.14
0.0092 0.0083
223
furnace cooling. The Ti-30Nb and Ti-30Nb-1Fe ingots were then hot rolled at 1000 °C to produce 3 mm thick sheets. Specimens of 1 cm2 were cut from the sheets and subjected to chemical pickling to remove surface oxide layer produced by hot rolling. This procedure reduced the thickness of the specimens by about 0.5 mm. The composition was analyzed by X-ray fluorescence (XRF) in a Rigaku RIX 3100 spectrometer, and interstitial O and N contents were examined using a LECO 400 analyzer (Table 1). All the samples were solution-treated at 1000 °C for 1 h, waterquenched, and then subjected to various aging heat treatments at
Fig. 1. Schematic diagram of tensile tests.
Fig. 4. DSC results of samples aged at 300 °C and 600 °C for 24 h.
Table 2 Phase transformation temperatures determined by DSC measurements. Dominating process
Fig. 2. XRD patterns of samples aged at 300 °C and 600 °C for 24 h.
Suggested phase transformation
Endo
α“-β
Endo
ω-β
Exo
β/ω-α
Endo
α-β
Aging temperature
300 °C 600 °C 300 °C 600 °C 300 °C 600 °C 300 °C 600 °C
Fig. 3. BSE-SEM images of α phase precipitation after aging heat treatment at 600 °C for 24 h: (a) Ti-30Nb, (b) Ti-30Nb-1Fe.
Temperature range Ti-30Nb
Ti-30Nb1Fe
360–402 – 439–566 418–483 550 – – 614–752
– – 372–529 418–491 510 – – 616–749
224
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Fig. 5. TEM dark field (DF: a, b, d, e) and selected area diffraction (SAD: c, f) patterns of Ti-30Nb-1Fe alloy aged at 400 °C for 1 h showing massive formation of small ω precipitates, which appear with an elongated morphology through [315] beta zone axis (f).
Fig. 6. BSE-SEM images of α phase precipitation after aging heat treatment at 500 °C for 1 h: (a–c) Ti-30Nb, (b–d) Ti-30Nb-1Fe.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
400, 500 and 550 °C for 1 h, applying heating rates of 2 and 30 °C/ min, followed by water quenching (WQ). Metallographic specimens were prepared by sanding with 1500 grit sandpaper and polishing with 0.02 mm colloidal silica. Scanning electron microscopy (SEM) was performed in a FEI Inspect F50 and a Hitachi TM1000 microscopes, while transmission electron microscopy (TEM) analyses were performed using a JEOL 3010 microscope operating at 300 kV. TEM thin foils were prepared via conventional dimpling and argon ion milling procedures. For a detailed microstructural analysis, X-ray diffraction (XRD) experiments were performed under copper radiation (λ ¼0.15406 nm) in a PANalytical X-Pert Pro diffractometer operating at 40 kV and 30 mA. The diffractometer was equipped with a PIXcel fast detector and a spinner sample holder. Vickers hardness measurements were taken with a Buehler Vickers 2100 hardness tester, applying an indentation load of 500 g.f. for 15 s. Two additional 24 h heat treatments were performed at 300 °C and 600 °C in order to produce equilibrium microstructures composed of β þ ω (300 °C) and β þ α (600 °C) for evaluation by differential scanning calorimetry (Netzsch STA409). DSC characterization was carried out under an argon flux at a heating rate of 30 °C/min. Lastly, the mechanical properties of microstructures subjected to aging were evaluated by tensile testing in an MTS 810 universal testing machine, applying a strain rate of 1 mm/min. Specimens with a gage section of 3 mm 6 mm were carefully prepared by waterjet cutting and machining, as illustrated in Fig. 1. The extensometer was not used due to the small size of the specimens.
225
about 40 °C the temperature at which α precipitation begins in Ti30Nb and Ti-30Nb-1Fe alloys. This is a key point to evaluate the results obtained by aging. 3.2. Aging heat treatments and
α phase precipitation
Considering the transformation temperatures of the two alloys described in Section 3.1, if the aging temperature were lower than 430 °C, the amount of ω phase would be expected to increase with aging time owing to the formation of isothermal ω (ωiso). At temperatures above 500 °C, α precipitation should prevail. On the other hand, α phase begins to transform into β phase close to 650 °C (near beta transus), so the optimal aging heat treatment temperature may be lower than that. Therefore, aging heat treatment temperatures of 500 and 550 °C were chosen, since copious α precipitation is expected to occur within this temperature interval. These were the typical aging temperatures employed in several studies on Ti-Nb based alloys [1,19,20]. The TEM images in Fig. 5 show that ω phase was still detectable in the Ti-30Nb-1Fe alloy aged at 400 °C for 1 h, which is consistent with the DSC results. Moreover, at this temperature, heterogeneous α phase precipitation appeared to begin at the β/ω interfaces, given the ω phase elongated morphology (Fig. 5d) – even
3. Results 3.1. Phase transformation temperatures Barriobero et al. [16] successfully combined in situ high energy X-ray diffraction (HEXRD) and DSC data for Ti-10V-2Fe-3Al alloy. To this end, they heated specimens continuously at the same heating rates (5, 20 and 50 °C min 1) in both experiments. They then made a detailed description of the exothermic and the endothermic events (DSC) associated with each phase transformation (XRD), as a guide for future evaluations of phase transformation temperatures in other alloy systems. These results can certainly be extrapolated and compared to the Ti-Nb system, as shown by Bönisch et al. [17]. In this work, to make a specific phase transformation more visible for detection by DSC, we performed two aging heat treatments at 300 °C and at 600 °C for 24 h, so as to force the detection of the ω-solvus temperature and the β transus, respectively. The same procedure was applied to both Ti-30Nb and Ti-30Nb-1Fe alloys, and the results (XRD, SEM and DSC) are depicted in Figs. 2–4, respectively. A summary is presented in Table 2. The Ti-30Nb alloy sample aged at 300 °C showed α“-β transformation (endothermic) followed by another endothermic peak associated with ω-β transformation, beginning at 439 °C. An adjacent exothermic peak corresponding to the copious nucleation of α phase that occurs in ωiso precipitates was visible at close to 550 °C. In the analysis of the sample aged at 600 °C, a sharp peak corresponding to the beta transus temperature (α-β) was detected, starting at 614 °C and ending at 752 °C. Both temperatures are in accordance with the phase diagram of Ti-Nb based alloys simulated by Zhang et al. [18]. In the Ti-30Nb-1Fe alloy, no significant amount of α“ was detected in the initial condition, so the first endothermic peak pertains to β-ω transformation, which starts at 372 °C, while α precipitation starts close to 510 °C. The two alloys appear to fall within in a similar range of beta transus temperatures. An interesting point noted in the DSC analysis (samples aged at 300 °C) is that 1 wt% of Fe reduces by
Fig. 7. XRD patterns of Ti-30Nb and Ti-30Nb-1Fe samples aged at for 1 h at 500 °C and 550 °C, applying heating rates of 2 and 30 °C/min.
226
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
taking into account that α phase diffraction was not clearly detected in the SAD diffraction patterns shown in Fig. 5c and f. At 400 °C, α phase precipitation appeared to occur slowly, non-optimally, owing to the small size of the precipitates and their heterogeneous distribution. Compared to previous works, it can be stated that these ωiso precipitates are slightly larger than the ωath precipitates produced by WQ in this alloy [13]. The SEM images in Fig. 6 correspond to the samples aged at 500 °C at two heating rates, 2 and 30 °C/min. Aging Ti-30Nb alloy at a heating rate of 30 °C/min resulted in the formation of ω phase, detectable by X-ray diffraction (Fig. 7). This phase was probably formed upon water-quenching (performed after aging) from a compositional pocket or solute-lean β matrix. According to the DSC data (Fig. 4), it may also be residual isothermal ω that did not
decompose at 500 °C due to the short heating ramp, which competes with α phase precipitation, making α precipitates sparse [3]. In the Ti-30Nb-1Fe sample, since 30 °C/min is the same heating rate as that applied in the DSC experiments, considerable α phase formation would be expected at 500 °C, which is confirmed in Fig. 6b. On the other hand, no previous DSC information was available for the lower heating rate (2 °C/min). However, based on SEM and XRD data, the Ti-30Nb-1Fe sample appeared to present fewer α phase precipitates compared to those produced at the higher heating rate. Moreover, copious α phase was formed in the Ti-30Nb samples owing to the longer heating ramp, which enables the diffusion of Nb that results in these long α laths (Fig. 6c). With regard to the aging heat treatments performed at 550 °C, Ti-30Nb alloy should only undergo massive α formation close to
Fig. 8. BSE-SEM images of α phase precipitation after aging heat treatment at 550 °C for 1 h: (a–c) Ti-30Nb, (b–d) Ti-30Nb-1Fe.
Fig. 9. BSE-SEM images of α phase precipitation after aging heat treatment at 550 °C for 1 h, at a heating rate of 600 °C/min: (a) Ti-30Nb, (b) Ti-30Nb-1Fe.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
this temperature, according to Table 2, which in fact occurred, as shown in Fig. 8a and c; additionally, a small fraction of α” formed in response to WQ was also detected. Bönisch et al., who investigated Ti-29Nb (wt%) aged at 545 °C followed by water quenching, also obtained α”, α and β phases [17]. Banumathy et al. studied the effects of 24 h aging heat treatments at 450 and 550 °C on Ti-27Nb (wt%) alloy, followed by WQ. They reported the formation of martensite at 450 °C, but not at 550 °C [21]. These results suggest that α” formation occurs only in the initial stages of α phase precipitation [22]. Ti-30Nb-1Fe samples likewise presented β þ α microstructures when aged at 550 °C. After a high density of α phase precipitates at 550 °C was reached, regardless of the heating rate, we found that, notwithstanding the findings of some authors about the strong influence of the heating rate on the size of α precipitates [6], in our preliminary study with 2 and 30 °C/min, the experimental alloys only appeared to be sensitive to the heating rate when aged at lower temperatures, i.e., below or close to 500 °C. For a more detailed analysis of the influence of the heating rate on the microstructure, we designed an additional experiment at a preset temperature in a tubular furnace, quickly inserting the samples in the furnace once it had reached this temperature. These samples were subjected to a much higher heating rate of about 600 °C/min. Fig. 9 shows SEM images of samples subjected to a heating rate of 600 °C/min to 550 °C, isothermally aged for 1 h, followed by WQ. The Ti-30Nb-1Fe alloy (Fig. 9b) appears to contain fewer precipitates compared to those produced at the other heating rates (Figs. 6 and 8), as was expected based on the literature [7]. In this case, ω-assisted α precipitation was somewhat inhibited by the high heating rates (430 °C/min) [23,24]. According to Tong Li et al., oxygen rich sites associated with ωiso may play a major role in this fine-scaled α precipitation [25], but it takes longer times at lower temperatures for these zones to be formed. Unlike Ti-30Nb-1Fe, the Ti-30Nb sample exhibited an unusual behavior. The higher the heating rate applied to the Ti-30Nb alloy the smaller the size of the precipitates, as can be seen by comparing Fig. 9a with Fig. 8a and c. In this case, these fine and well dispersed α precipitates were probably generated by another mechanism, since the initial microstructure of Ti-30Nb is composed of β þ α”, and we believe that heating occurs so fast that there is likely no ωiso phase to support ω-assisted α formation (which would also lead to the same behavior as that of 1Fe alloy). Considering this fine scale α phase precipitation, it would be possible that this microstructure was produced by a non-classical mechanism [26]. However, based on thermodynamic MatCalc [27] simulations, we concluded that Ti-Nb based alloys with Nb content higher than 15 wt% could not undergo pseudospinodal decomposition at 550 °C, because that would require a very significant compositional fluctuation, due to a wide gap between the Gibbs free energy curves at this temperature (Fig. 10). To date, the mechanism discovered by Nag et al. and Boyne et al. has been only detected in Ti–Mo based alloys, such as Timetal-5553 [26,28]. Therefore, we believe that massive and fine α precipitation, in this particular case, must occur via another mechanism, such as heterogeneous nucleation at dislocations, twin boundaries or other lattice defects resulting from the rapid decomposition of α” phase. In essence, α precipitation in Ti-30Nb-1Fe alloy is influenced very little by the heating rate. In addition to altering some transformation temperatures (Table 2), we believe that α formation occurs in the heat treatment route in Ti-30Nb-1Fe due to the faster diffusion of Fe in the Ti-Nb matrix. Based on the diffusion coefficients found in Smithells Metals Reference Book [29], we estimated that Nb diffusion in Ti at 550 °C (823 K) occurs with D ¼7.6 10 19 m2/s, while Fe diffusion in Ti is almost 90 times faster than that, reaching D¼6.5 10 17 m2/s. Thus, Fe alters the
227
kinetics of α precipitation, making it viable at any heating rate. The rejection of Fe from α phase also leads to enrichment of the neighbor β phase in Fe, restricting the growth of α laths and resulting in consistently small α precipitates. On the other hand, in Ti-30Nb alloy, an exceptionally refined α precipitation was observed only at the highest heating rate (600 °C/min).
Fig. 10. Gibbs free energy curves of β and α phases containing Nb content at 550 °C. The optimal composition for the pseudospinodal is indicated as 14.9 wt%.
Table 3 Vickers hardness, elastic modulus and phases detected by XRD in the Ti-30Nb and Ti-30Nb-1Fe samples aged at 500 and 550 °C for 1 h, at heating rates of 2 and 30 °C/ min.
Ti-30Nb
Heating rate (°C/ min)
Temperature (°C) – Heating rate (°C/ min)
Vickers hardness (HV)
Elastic modulus (GPa)
2
400 500 550 400 500 550
428 7 9 2747 11 219 73 4217 13 269 7 6 210 73
1167 1 797 1 727 1 1137 2 1157 2 697 1
400 500 550 400 500 550
423 7 5 2617 8 2317 4 399 7 16 251 720 226 7 8
1157 2 927 1 807 2 1067 1 987 3 827 1
30
Ti-30Nb-1Fe 2
30
Fig. 11. Stress-strain curves of the Ti-30Nb and Ti-30Nb-1Fe samples aged at 550 °C for 1 h, at a heating rate of 2 °C/min. WQ samples studied by Lopes et al. [13].
228
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Fig. 12. Fracture surface of the samples aged at 550 °C for 1 h, at a heating rate of 30 °C/min: (a) Ti-30Nb; (b) Ti-30Nb-1Fe.
3.3. Mechanical properties The results of Vickers hardness and elastic modulus of the aged alloys are listed in Table 3. Note that these results reflect the presence of different phases in the microstructure. The addition of Fe limits the elastic modulus to a minimum of 80 GPa, which is undesirable for biomedical applications. Ti-30Nb heated to 550 °C at 30 °C/min showed the lowest elastic modulus due to the formation of α” phase after WQ. Tensile tests were performed to evaluate the mechanical properties of aged Ti-30Nb and Ti-30Nb-1Fe samples (Fig. 11). The mechanical properties of solution heat-treated and water quenched Ti-30Nb and Ti-30Nb-1Fe alloys were investigated by Lopes et al. and their curves are plotted in Fig. 11 for comparison [13]. The aging heat treatments for tensile testing were carried out at a temperature of 550 °C, heating rate of 2 °C/min, and isothermal holding time of 1 h. As expected, α phase precipitation during the aging heat treatment increased the ultimate tensile strength of both alloys. Although the heat treatment decreased elongation (strain), the values obtained are still acceptable, resulting in mixed-mode fracture with prevalence of dimples, as can be seen in Fig. 12, with dimpled fracture surfaces. The behavior of the tensile test curve of Ti-30Nb differed from that of the alloy containing 1% of Fe. Fig. 11 shows this curve with the double yield point, which means that this alloy presents shape memory effect and superelasticity [30].
4. Conclusions In this work, Ti-Nb and Ti-Nb-Fe alloys were subjected to various aging heat treatments. The results indicate a correlation between the mechanical properties and the precipitated phases. The main results were as follows: – DSC analyses revealed peaks corresponding to the decomposition of α”, ω and α phases. The results suggest that the addition of 1 wt% Fe allows α phase nucleation at lower temperatures. – The two alloys reach maximum hardness upon precipitation of isothermal ω phase. The highest hardness values were found in samples aged at 400 °C. – The addition of 1 wt% Fe refines the precipitation of α phase in comparison to that of the base alloy, probably due to the higher diffusion coefficient of Fe. – Reducing the heating rate from 30 °C/min to 2 °C/min did not lead to the refinement of α phase precipitates in either of the alloys. – For Ti-30Nb alloy, the heating rate of 600 °C/min resulted in the
finest α phase precipitates, which can be explained by the different mechanism activated during martensite decomposition. In the case of Ti-30Nb-1Fe alloy, the heating rate did not interfere in the refinement of α phase precipitates.
Acknowledgments The authors gratefully acknowledge the financial support of the Brazilian research funding agencies CAPES (Federal Agency for the Support and Improvement of Higher Education), CNPq (National Council for Scientific and Technological Development) and FAPESP (São Paulo Research Foundation) (Grants #2013/13867-2 and #2014/06099-1).
References [1] D. Banerjee, J.C. Williams, Perspectives on titanium science and technology, Acta Mater. 61 (2013) 844–879, http://dx.doi.org/10.1016/j. actamat.2012.10.043. [2] M. Niinomi, M. Nakai, J. Hieda, Development of new metallic alloys for biomedical applications, Acta Biomater. 8 (2012) 3888–3903, http://dx.doi.org/ 10.1016/j.actbio.2012.06.037. [3] D.L. Moffat, D.C. Larbalestier, The competition between the alpha and omega phases in aged Ti-Nb alloys, Metall. Trans. A 19 (1988) 1687–1694. [4] S. Kobayashi, T. Takeichi, K. Nakai, T. Sakamoto, Acceleration or suppression of α-phase precipitation using isothermal ω phase in Ti-20 at.pct Nb Alloy, Metall. Mater. Trans. A 45 (2013) 1217–1229, http://dx.doi.org/10.1007/ s11661-013-2092-1. [5] F. Prima, P. Vermaut, G. Texier, D. Ansel, T. Gloriant, Evidence of α-nanophase heterogeneous nucleation from ω particles in a β-metastable Ti-based alloy by high-resolution electron microscopy, Scr. Mater. 54 (2006) 645–648, http://dx. doi.org/10.1016/j.scriptamat.2005.10.024. [6] Y. Zheng, R.E.A. Williams, D. Wang, R. Shi, S. Nag, P. Kami, et al., Role of ω phase in the formation of extremely refined intragranular α precipitates in metastable β-titanium alloys, Acta Mater. 103 (2016) 850–858, http://dx.doi.org/ 10.1016/j.actamat.2015.11.020. [7] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Aging response of coarse- and fine-grained β titanium alloys, Mater. Sci. Eng. A 405 (2005) 296–305, http://dx.doi.org/10.1016/j.msea.2005.06.027. [8] P. Majumdar, S.B. Singh, M. Chakraborty, The role of heat treatment on microstructure and mechanical properties of Ti–13Zr–13Nb alloy for biomedical load bearing applications, J. Mech. Behav. Biomed. Mater. 4 (2011) 1132–1144, http://dx.doi.org/10.1016/j.jmbbm.2011.03.023. [9] E.S.N. Lopes, R.J. Contieri, S.T. Button, R. Caram, Femoral hip stem prosthesis made of graded elastic modulus metastable β Ti alloy, Mater. Des. 69 (2015) 30–36, http://dx.doi.org/10.1016/j.matdes.2014.11.040. [10] S. Nag, R. Banerjee, H.L. Fraser, Intra-granular alpha precipitation in Ti–Nb–Zr– Ta biomedical alloys, J. Mater. Sci. 44 (2008) 808–815, http://dx.doi.org/ 10.1007/s10853-008-3148-2. [11] A. Biesiekierski, J. Wang, M.A.-H. Gepreel, C. Wen, A new look at biomedical Tibased shape memory alloys, Acta Biomater. 8 (2012) 1661–1669, http://dx.doi. org/10.1016/j.actbio.2012.01.018. [12] D.C. Zhang, Y.F. Mao, Y.L. Li, J.J. Li, M. Yuan, J.G. Lin, Effect of ternary alloying elements on microstructure and superelastictity of Ti–Nb alloys, Mater. Sci.
F.H. da Costa et al. / Materials Science & Engineering A 677 (2016) 222–229
Eng. A 559 (2013) 706–710, http://dx.doi.org/10.1016/j.msea.2012.09.012. [13] É.S.N. Lopes, C.A.F. Salvador, D.R. Andrade, A. Cremasco, K.N. Campo, R. Caram, Microstructure, mechanical properties, and electrochemical behavior of Ti-NbFe alloys applied as biomaterials, Metall. Mater. Trans. A 47 (2016) 3213–3226, http://dx.doi.org/10.1007/s11661-016-3411-0. [14] C.M. Lee, W.F. Ho, C.P. Ju, J.H. Chern Lin, Structure and properties of Titanium25 Niobium-x iron alloys, J. Mater. Sci. Mater. Med. 13 (2002) 695–700 http:// www.ncbi.nlm.nih.gov/pubmed/15348579. [15] H.-C. Hsu, S.-K. Hsu, S.-C. Wu, C.-J. Lee, W.-F. Ho, Structure and mechanical properties of as-cast Ti–5Nb–xFe alloys, Mater. Charact. 61 (2010) 851–858, http://dx.doi.org/10.1016/j.matchar.2010.05.003. [16] P. Barriobero-Vila, G. Requena, F. Warchomicka, et al., Phase transformation kinetics during continuous heating of a β-quenched Ti–10V–2Fe–3Al alloy, J. Mater. Sci. 50 (2015) 1412, http://dx.doi.org/10.1007/s10853-014-8701-6. [17] M. Bönisch, M. Calin, T. Waitz, A. Panigrahi, M. Zehetbauer, A. Gebert, et al., Thermal stability and phase transformations of martensitic Ti–Nb alloys, Sci. Technol. Adv. Mater. 14 (2013) 055004, http://dx.doi.org/10.1088/1468-6996/ 14/5/055004. [18] Y. Zhang, H. Liu, Z. Jin, Thermodynamic assessment of the Nb-Ti system, Calphad 25 (2001) 305–317, http://dx.doi.org/10.1016/S0364-5916(01)00051-7. [19] M. Geetha, A.K. Singh, A.K. Gogia, R. Asokamani, Effect of thermomechanical processing on evolution of various phases in Ti–Nb–Zr alloys, J. Alloy. Compd. 384 (2004) 131–144, http://dx.doi.org/10.1016/j.jallcom.2004.04.113. [20] L. Wang, Z. Lin, X. Wang, Q. Shi, W. Yin, D. Zhang, et al., Effect of aging treatment on microstructure and mechanical properties of Ti27Nb2Ta3Zr Titanium alloy for implant applications, Mater. Trans. 55 (2014) 141–146, http: //dx.doi.org/10.2320/matertrans.M2013187. [21] S. Banumathy, K.S. Prasad, R.K. Mandal, A.K. Singh, Effect of thermomechanical processing on evolution of various phases in Ti-Nb alloys, Bull. Mater. Sci. 34 (2011) 1421–1434, http://dx.doi.org/10.1007/s12034-011-0338-3. [22] E. Aeby-Gautier, A. Settefrati, F. Bruneseaux, B. Appolaire, B. Denand, M. Dehmas, et al., Isothermal α″ formation in β metastable titanium alloys, J.
[23]
[24]
[25]
[26]
[27]
[28]
[29]
[30]
229
Alloy. Compd. 577 (2013) S439–S443, http://dx.doi.org/10.1016/j. jallcom.2012.02.046. Y. Zheng, R.E.A. Williams, J.M. Sosa, T. Alam, Y. Wang, R. Banerjee, et al., The indirect influence of the ω phase on the degree of refinement of distributions of the α phase in metastable β-Titanium alloys, Acta Mater. 103 (2016) 165–173, http://dx.doi.org/10.1016/j.actamat.2015.09.053. Y. Zheng, R.E.A. Williams, J.M. Sosa, Y. Wang, R. Banerjee, H.L. Fraser, The role of the ω phase on the non-classical precipitation of the α phase in metastable β-titanium alloys, Scr. Mater. 111 (2016) 81–84, http://dx.doi.org/10.1016/j. scriptamat.2015.08.019. T. Li, D. Kent, G. Sha, J.M. Cairney, M.S. Dargusch, The role of ω in the precipitation of α in near-β Ti alloys, Scr. Mater. 117 (2016) 92–95, http://dx.doi. org/10.1016/j.scriptamat.2016.02.026. S. Nag, Y. Zheng, R.E. a. Williams, A. Devaraj, A. Boyne, Y. Wang, et al., Nonclassical homogeneous precipitation mediated by compositional fluctuations in titanium alloys, Acta Mater. 60 (2012) 6247–6256, http://dx.doi.org/10.1016/ j.actamat.2012.07.033. E. Kozeschnik, MatCalc––a simulation tool for multicomponent thermodynamics, diffusion and phase transformation kinetics, in: E. Kozeschnik, B. Buchmayr (Eds.), Mathematical Modelling of Weld Phenomena 5, 738, 2001, pp. 349–361. A. Boyne, D. Wang, R.P. Shi, Y. Zheng, A. Behera, S. Nag, et al., Pseudospinodal mechanism for fine α/β microstructures in β-Ti alloys, Acta Mater. 64 (2014) 188–197, http://dx.doi.org/10.1016/j.actamat.2013.10.026. Y. Sohn, Diffusion in metals, In: C.J. Smithells, (Ed), Smithells Metals Reference Book, 2004, Oxford: Elsevier, 1–120, http://dx.doi.org/10.1016/B978-075067509-3/ 50016-6. A.M.G. Tavares, W.S. Ramos, J.C.G. de Blas, E.S.N. Lopes, R. Caram, W.W. Batista, et al., Influence of Si addition on the microstructure and mechanical properties of Ti–35Nb alloy for applications in orthopedic implants, J. Mech. Behav. Biomed. Mater. 51 (2015) 74–87, http://dx.doi.org/10.1016/j. jmbbm.2015.06.035.