Effect of Heat Treatment within Alpha/Beta Dual-Phase Field on the Structure and Tensile Properties of Binary Ti–Mo Alloys
2018 Volume 59 Issue 5 Pages 734-740
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2018 Volume 59 Issue 5 Pages 734-740
The present study investigated the effect of heat treatment within the alpha (α)/beta (β) dual-phase field on the structure and tensile properties of Ti–(1.5–9.5) mass% Mo alloys. The alloys were prepared using an arc-melting vacuum-pressure type casting system. The cast alloys were heat-treated at 700, 750 and 800°C in vacuum for 30 minutes followed by quenching in ice water. The X-ray diffraction (XRD) results indicated that beta (β) phase intensities increased while α/alpha prime (α′) intensities decreased with increased heat treatment temperature (HTT) and Mo concentration. The β phase was observed to dominate the 800°C-treated Ti–9.5Mo alloy, while the highest alpha double prime (α′′) phase content was observed in the 800°C-treated Ti–7.5Mo alloy. Both optical and scanning electron microscopy indicated that a relatively coarse α platelet was always observed in Ti–1.5Mo. A fine, uniformly-distributed acicular microstructure was observed in Ti–7.5Mo, while an equi-axed β granular microstructure was clearly seen in Ti–9.5Mo. The tensile properties were found sensitive to the HTT and Mo concentration. When heat-treated at 700°C, the yield strength (YS) and ultimate tensile strength (UTS) increased while the elongation generally decreased with Mo concentration. The highest YS and UTS were found in Ti–7.5Mo and Ti–9.5Mo. When heat-treated at 750°C, the strength of Ti–5.5Mo was improved without reducing elongation. With Mo concentration increased to 7.5% or higher, the elongation further increased while the strength maintained a similar level. When treated at 800°C, the YS of Ti–3.5Mo, Ti–5.5Mo and Ti–7.5Mo maintained a lower level than Ti–1.5Mo and Ti–9.5Mo. A fully satisfactory interpretation for the tensile properties and their relationships to the complicated microstructures might not be a simple task due to several different factors simultaneously involved, yet practically it is interesting to note that selected alloys heat-treated within the dual-phase field demonstrated quite promising overall mechanical properties.
Due to their light weight, high corrosion resistance and specific strength, pure titanium and titanium alloys have been widely used for dental and orthopedic applications such as crown and bridge, removable partial denture, dental implants, hip prosthesis, artificial knee joint and trauma-fixation devices,1–6) but not without drawbacks. For example, the popularly-used c.p. Ti has a relatively low mechanical strength.5,7–9) Ti–6Al–4V ELI, despite its excellent mechanical properties,10–12) has the potential problem of releasing aluminum and particularly vanadium ions from the alloy implant which might cause long term health problems such as Alzheimer’s disease and cytotoxicity.13–16) Another potential problem of c.p. Ti and Ti–6Al–4V ELI is their much higher elastic modulus values (typically between 100 and 120 GPa3) than that of natural bones (about 10–20 GPa17)). These much higher modulus values could activate the stress-shielding effect potentially leading to bone atrophy or even failure of the implant.9,18,19)
The search for Al and V-free, biocompatible Ti alloys for orthopedic implant applications was initiated in the mid-80s, and among them β and near-β alloys have caught the most attention.8,20–23) The β-type Ti alloys containing large amounts of such heavy alloy elements as Mo, Nb, Ta and/or W demonstrated a better biocompatibility, better formability, and lower elastic modulus level than α and α/β-type Ti alloys. Nevertheless, their relatively high melting temperatures, high densities and high costs are some major concerns for the application of these β-type alloys.
Using a different approach, the present authors’ team has developed an Al and V-free, low modulus α′′-type Ti–7.5Mo alloy with strength/modulus ratios significantly higher than those of popularly-used 316L, Co–Cr–Mo and Ti–6Al–4V alloys.24) The results of the study indicated that α′′ phase had a lower modulus than all other phases in the binary Ti–Mo alloy system. To obtain the α′′ phase, the alloy underwent a solution treatment (ST) in the β phase field followed by a water quench (WQ) process.24,25) In another early study from the authors’ lab,26) the Ti–7.5Mo alloy with a tensile modulus of 78 GPa and a Ti–6Al–4V alloy with a tensile modulus of 110 GPa were implanted into rabbit femur. It was interesting to find that, after 26 weeks, the amount of new bone attached onto the Ti–7.5Mo implant was >5 times larger than that onto the Ti–6Al–4V implant. This large difference was thought to derive from a combined effect of chemistry (the presence of harmful Al and V in Ti–6Al–4V) and elastic modulus (the stress-shielding effect).
Conventionally the α′′ phase in a binary Ti–Mo alloy is obtained by direct casting from the molten state24) or heating the alloy into the β phase field followed by fast cooling, often with a water or ice water quench.1,6) Most studies conducted the solution treatment of Ti–Mo-based alloys in the β phase field, but few studies were devoted to investigating the heat treatment (HT) effect within the α/β dual-phase field. The studies of Cardoso et al.,1) Jiao et al.27) and Lu et al.28) involved HT within the α/β zone but were not focused on the HT parameter-structure-mechanical property relationships of the Ti–Mo system in a systematic way. The primary purpose of the present study was therefore to investigate the effect of HT within the α/β dual-phase field on the structure and tensile properties of a series of binary Ti–Mo alloys. It also attempted to clarify whether the β phase formed in the α/β regime at much lower temperatures could transform into the inherently low modulus α′′ phase.
The Ti–Mo alloys of different compositions (1.5, 3.5, 5.5, 7.5 and 9.5 mass%Mo) used for this study were prepared from grade 2 commercially pure (99.8 mass% pure) titanium (China Steel Co., Taiwan) and 99.95 mass% pure molybdenum wire (Alfa Aesar, USA) using a commercial arc-melting vacuum-pressure type casting system (Castmatic Iwatani Corp., Japan). Prior to melting, the melting chamber was evacuated and purged with argon gas. An argon pressure of 0.147 MPa was maintained during melting. To prepare each alloy, appropriate amounts of Ti and Mo metals were melted in a U-shaped copper hearth with a tungsten electrode. The ingot was re-melted three times to improve chemical homogeneity of the alloy.
Prior to casting, the alloy ingot was re-melted in an open-based copper hearth in argon gas under a pressure of 0.147 MPa. The difference in pressure between the two chambers allowed the molten alloy to quickly drop into a mold at room temperature. To investigate the effect of heat treatment in the α/β dual-phase field (Fig. 1), the cast samples were heat-treated at three different temperatures (700, 750 and 800°C) in vacuum for 30 minutes followed by quenching in ice water.
The highlighted α/β dual-phase regime (700–800°C; 1.5–9.5 mass%)Mo in the Ti–Mo phase diagram29) investigated in this study.
X-ray diffraction (XRD) for phase analysis was conducted using a Bruker D2 Phaser diffractometer operated at 30 kV and 10 mA with scanning speeds of 2°/min and 0.1°/min. A Ni-filtered CuKα radiation was used for the study. A silicon standard was used for the calibration of diffraction angles. The various phases were identified by matching each characteristic peak in the diffraction patterns with JCPDS files (Joint Committee on Powder Diffraction Standards, now called International Centre for Diffraction Data, ICDD).
Microstructural examination of the series of samples was performed using an optical microscope (Leica TMX 100, Germany). The surfaces of the materials for light microscopy were mechanically polished via a standard metallographic procedure to a final level of 0.05 µm alumina powder, followed by chemical etching in a mixture of water, nitric acid, and hydrofluoric acid (100:3:1 by volume). A scanning electron microscope (SEM) (JEOL JSM-6510, Japan) operated at 5 kV under secondary electron mode was also used for microstructural examination in more details. The samples for SEM examination were prepared under the same procedure as for optical microscopy.
A servo-hydraulic type testing machine (EHF-EG, Shimadzu Co., Tokyo, Japan) was used for tensile testing. The specimens for the tensile testing were wire-cut using an electrical discharge machining system (V50, Excetek Technologies Co., Taiwan). The dogbone-shaped, reduced-sized specimens for testing were 55 mm long, 10 mm wide and 1.0 mm thick with a gage length of 6 mm and gage width of 3 mm. The testing was performed at room temperature with a constant crosshead speed of 8.33 × 10−6 ms−1. The average ultimate tensile strength (UTS), yield strength (YS) at 0.2% offset, Young’s modulus of elasticity, and elongation to failure were taken from six tests under each condition. The measurement of Young’s modulus of elasticity was conducted following the method set forth in ASTM E111-17, wherein the value of Young’s modulus was obtained by determining the slope of the straight-line portion of the stress-strain profile in the tensile test.
The XRD patterns of Ti–(1.5–9.5)Mo alloys heat-treated at 700, 750 and 800°C with a scan speed of 2°/min from 30 to 90° (2θ) are shown in Figs. 2a, 2b and 2c, respectively. Figure 2 indicates that, at any given heat treatment temperature (HTT), the XRD intensities of β phase increased, while the α/α′ peak intensities decreased with increased Mo concentration, as expected from the Ti–Mo phase diagram.29) It should be reminded that α and α′ phases have the same crystal structure (hcp) and are indistinguishable from their XRD patterns. These two phases are often distinguished from each other by their different microstructures. α′ phase has a fine, martensitic-type acicular microstructure usually obtained from a fast cooling process, while α phase usually exhibits a coarser plate-shaped microstructure.30) When the present alloys were heat-treated within the dual-phase field followed by WQ, it was possible for α and α′ phases to both form in the alloys. For this reason, the term “α/α′ peaks” was used in this study to represent the hcp peaks possibly attributed to both α and α′ phases in all the XRD patterns. According to the binary Ti–Mo phase diagram, the relative amount of α phase should increase with decreased Mo concentration when the alloys were heat-treated in the dual-phase field followed by WQ.
XRD patterns of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at (a) 700°C, (b) 750°C and (c) 800°C with a scan speed of 2°/min from 30 to 90° (2θ).
Figure 2 also indicates that, when HTT increased, the XRD intensities of β phase increased while the α/α′ peak intensities decreased, which was also expected from Ti–Mo phase diagram. The most-dominated β phase was observed in the alloy containing the highest Mo concentration (9.5 mass%) heat-treated at the highest HTT (800°C). In the early study of Ho et al.,24) a Mo concentration of 10 mass% or higher was capable of retaining substantially all the β phase during the fast-cooling casting process. When the alloys of the present study were heat-treated at lower HTT, their retained β phase amounts were lower due to the lower amounts of β phase formed in the dual-phase field.
Although distinguishing α′ from α from XRD patterns was difficult, the identification of the orthorhombic α′′ phase was rather straightforward. α′′ phase could be identified in the splitting of single α/α′ peaks into double peaks of α′′ phase. According to the early studies of Baker31) and Brown et al.,32) this fast cooling-induced athermal orthorhombic structure was derived from a distorted hexagonal cell with the c-axis of the orthorhombic cell corresponding to the c-axis of the hexagonal cell and a/b corresponding to the orthogonal axis of the hexagonal cell. In the early study of Ho et al.24) of a series of binary Ti–Mo alloys, the amount of α′′ phase was found very sensitive to the Mo concentration, and the fast-cooled binary Ti–7.5 mass% Mo was dominated by the orthorhombic α′′ phase.
Due to the often low and diffuse intensity distribution of the split characteristic peaks, the XRD conducted at low scan speeds was found helpful in the identification of α′′ phase. The XRD patterns scanning from 52 to 55° (2θ) of the alloys heat-treated at 800°C at a low scan speed (0.1°/min) are shown in Fig. 3. The splitting of the individual α/α′ (102) peak into its corresponding double α′′ (112/022) peaks is clearly shown in these low-scan-speed XRD patterns. It was also observed that, with increased Mo concentration, the spacing between the α′′ double peaks became larger and the α′′ (022) peak shifted toward the high angle side, consistent with the early finding of Bagariatskii et al.33)
XRD patterns of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 800°C with a scan speed of 0.1°/min from 52 to 55° (2θ).
The highest α′′ phase content accompanied with a low β content was observed in the 800°C-treated Ti–7.5Mo alloy, which is reasonable since this temperature is quite close to the β-transus of the alloy. As mentioned earlier, a Mo concentration of 7.5 mass% produced the highest content of α′′ phase in the Ti–Mo binary system during WQ. The present XRD results indicated that, depending on HTT and Mo concentration, the β phase formed in the dual-phase regime could be partially retained, transformed into α′ phase or transformed into α′′ phase during WQ. These different phases could sensitively affect the mechanical properties of the Ti–Mo alloys, as will be discussed later.
Typical optical and scanning electron micrographs of the Ti–(1.5–9.5)Mo alloys heat-treated at 700, 750 and 800°C are given in Figs. 4–6 and 7–9, respectively. As shown in the optical micrographs, at any HTT, the relatively coarse platelet-shaped α phase microstructure appeared less in quantity with increased Mo concentration. In Ti–1.5Mo alloy, the α platelet microstructure was observed at all temperatures due to its high α phase content in the dual-phase field (Figs. 4–6). The higher-magnification SEM micrographs revealed that there existed numerous fine needles (arrow in Fig. 7a) within the primary coarse platelets retained from the high temperature which could not be detected from optical microscopy. Whether these fine needles were formed during cooling as excess Mo atoms were expelled from the low-temperature, low-Mo concentration α phase (as shown in Ti–Mo phase diagram) is not certain at this moment. It seems highly unlikely that these fine needles were β-transformed α′ phase due to the small amount of β phase in the dual-phase regime, especially at lower temperatures. With increased HTT, these needles appeared thicker and the retained β phase grain boundaries became more visible (arrows in Fig. 9a).
Optical micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 700°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
Optical micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 750°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
Optical micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 800°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
Scanning electron micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 700°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
Scanning electron micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 750°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
Scanning electron micrographs of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at 800°C. (a) Ti–1.5Mo, (b) Ti–3.5Mo, (c) Ti–5.5Mo, (d) Ti–7.5Mo and (e) Ti–9.5Mo.
As the Mo concentration increased to 3.5 or 5.5 mass%, the optical micrographs showed that the acicular-shaped fine needle microstructure increased in quantity due to the increased β phase content in the dual-phase field which partially transformed into fine α′/α′′ needles during WQ (Figs. 4–6). According to the aforementioned XRD patterns, these needles were most likely α′/α′′ needles. Again, this needle type microstructure revealed in the higher magnification SEM micrographs indicated that the needles became thicker and the retained β phase became more recognizable due to the higher β phase content with increased HTT. The formation of the even finer structures observed within certain zones, e.g., along the retained β boundaries (arrows in Fig. 9c) is not fully understood. Identification of this fine structure is difficult with the present SEM. One hypothesis might be due to the fluctuations in Mo concentration in the 800°C-treated β regime. During WC, the higher-Mo zones tended to retain as β phase while the lower-Mo zones tended to transform into α′′ phase. The size of this fine structure seems to be limited by the dimension of these lower-Mo zones.
When the Mo concentration increased to 7.5 mass%, the optical micrographs showed a fine needle α′′ (identified by XRD) microstructure uniformly distributed throughout the alloy. The aforementioned XRD patterns indicated that the highest orthorhombic α′′ phase content was seen at this Mo concentration. Although α′ and α′′ had a similar martensitic-type fine needle microstructure, as mentioned earlier, distinguishing the orthorhombic α′′ phase from the hcp α′ phase was rather easy by XRD. When the Mo concentration further increased to 9.5 mass%, the equi-axed β granular type microstructure was clearly revealed in the optical as well as SEM micrographs. Consistent with XRD, although β phase dominated the XRD pattern of the 800°C-treated Ti–9.5Mo alloy, both optical and SEM micrographs clearly showed the presence of other phase, most likely α′′ (identified by XRD), indicating that the β phase in this study could neither be entirely retained, nor entirely transformed into α′ or α′′ phase during WQ.
The tensile properties of the investigated Ti–Mo alloys heat-treated at different temperatures are demonstrated in Fig. 10. Figures 10a, 10c and 10e respectively showed the YS, UTS and elongation of the alloys treated at 700, 750 and 800°C, while 10b, 10d and 10f respectively showed the Young’s modulus values of the alloys treated at the same three temperatures. When the alloy was heat-treated at 700°C, both YS and UTS increased with increasing of Mo concentration, while the elongation decreased in general. However, the properties of Ti–7.5Mo and Ti–9.5Mo were almost the same. Since the β phase content increased with increasing Mo concentration while the Mo concentration in the β phase remained the same, the strong β-strengthening effect seen at 700°C was probably due to the high Mo concentration of the β phase in the dual-phase field by a solute-strengthening mechanism. According to the Ti–Mo phase diagram, the Mo concentration of the β phase at 700°C could be estimated as 21 mass%Mo, which could effectively strengthen the β-phase. However, the Young’s modulus of Ti–7.5Mo was considerably higher than that of the α′′-dominated binary Ti–7.5Mo.24) In the absence of α′′ phase, all the modulus values were observed between 90–100 GPa (Fig. 10b) probably due to the similar modulus level of α, α′ and β at 700°C.
Tensile properties of Ti–(1.5–9.5 mass%)Mo alloys heat-treated at different temperatures. (a, c, e) showing YS, UTS and elongation values of the alloys heat-treated at 700, 750 and 800°C, respectively; (b, d, f) showing Young’s modulus values of the alloys heat-treated at 700, 750 and 800°C, respectively.
When heat-treated at 750°C, the YS, UTS and elongation of Ti–1.5Mo and Ti–3.5Mo alloys were almost identical to the same alloys treated at 700°C. When the Mo concentration increased from 3.5 to 5.5 mass%, the strength increased without reducing elongation. When the Mo concentration further increased to 7.5% or higher, the strength maintained a similar level, while the elongation further increased to 31%. The lower strength and higher elongation values (compared to 700°C-treated alloys) may be explained by the lower Mo concentration of the β phase at 750°C, although the β fraction was higher than that of 700°C. The modulus values of the 750°C-treated alloys also maintained a level of 90–100 GPa (Fig. 10d).
The YS, UTS, elongation and Young’s modulus values of the 800°C-treated Ti–1.5Mo alloy were almost the same as those of the 700°C and 750°C-treated alloys, indicating the dominant role of the α phase in this low Mo alloy. When the Mo concentration increased, however, the YS and UTS behaved quite differently in 700°C or 750°C-treated alloys. While the UTS continued to increase to the maximum value (877.6 MPa) found in Ti–7.5Mo alloy, the YS decreased and maintained a much lower level (556.0 MPa) in Ti–3.5Mo, Ti–5.5Mo and Ti–7.5Mo alloys than in Ti–1.5Mo (658.7 MPa) and Ti–9.5Mo (667.6 MPa). It is interesting to note that these three alloys with constant low YS demonstrated consistently lower modulus values (about 91 GPa) than the other two alloys when treated at 800°C (Fig. 10f). While a large difference between YS and UTS in an alloy accompanied with a low modulus is commonly observed in a fast-cooled α′′-dominated Ti–7.5Mo alloy,6) whether the decreased YS and modulus in these three 800°C-treated alloys (especially Ti–3.5Mo and Ti–5.5Mo) was attributed to the formation of α′′ phase was not certain. The present XRD results indicated that the 800°C-treated Ti–7.5Mo alloy contained much more α′′ phase than Ti–3.5Mo and Ti–5.5Mo alloys treated at the same temperature. If α′′ phase played a major role in lowering the YS and modulus, the YS and modulus values of these three alloys could not be so close. When the Mo concentration further increased to 9.5 mass%, the UTS and elongation decreased, while the YS and modulus increased. The decreased UTS might be attributed to the grain growth effect of the β phase. The relatively low elongation and high modulus seen in 800°C-treated Ti–9.5Mo alloy might hypothetically be related to the formation of the brittle ω phase. According to the Ti–Mo phase diagram, the equilibrium Mo concentration at 800°C is very close to 9.5 mass% Mo. The early study of Ho et al.24) indicated that the strongest ω effect in the binary Ti–Mo system was found in the alloy containing 10 mass% Mo. The study of Zhang et al.34) also indicated that the Ti–Mo alloy containing about 6 at% (∼11 mass%) Mo had the strongest ω effect accompanied with a high modulus.
From above results and discussion, it can be seen that a fully satisfactory interpretation for all the tensile properties and their relationships to the complicated microstructures of the present Ti–(1.5–9.5)Mo alloys treated within α/β dual-phase field might not be a simple task. Further research, especially in microanalysis of the various fine structures, is needed. The factors that could possibly affect the tensile properties seem to at least include the contents of different phases (α, α′, α′′, β, and possibly ω) and the concentrations of Mo in the different phases. Yet, however complicated the interpretation of the mechanical properties might be, practically it is interesting to note that some of the alloys heat-treated within the dual-phase field demonstrated quite promising overall mechanical properties. Taking the 700°C-treated Ti–9.5Mo alloy as an example, with a similar elongation (14.5%) to the popularly-used Ti–6Al–4V ELI (10%), the 700°C-treated Ti–9.5Mo alloy demonstrated much higher YS and UTS with a lower modulus (952 MPa, 957 MPa and 96 GPa, respectively) than Ti–6Al–4V ELI (795 MPa, 860 MPa and 114 GPa, respectively). Another example is the 750°C-treated Ti–7.5Mo which demonstrated a little higher YS (820 MPa), lower UTS (837 MPa), lower modulus (101 GPa), and a far higher elongation (31%) than Ti–6Al–4V ELI (10%).
The authors would like to acknowledge the support for this research by the Ministry of Science and Technology, Republic of China under the Research Grant No. MOST 104-2221-E-006-142-.
