Effect of Heat Treatment within Alpha/Beta Dual-Phase Field on the Structureand Tensile Properties of Binary Ti­Mo Alloys

Yu-Po Peng+1, Chien-Ping Ju and Jiin-Huey Chern Lin+2

Department of Materials Science and Engineering, National Cheng-Kung University, Tainan 70101, Taiwan, ROC

The present study investigated the effect of heat treatment within the alpha (¡)/beta (¢) dual-phase field on the structure and tensileproperties of Ti­(1.5­9.5)mass% Mo alloys. The alloys were prepared using an arc-melting vacuum-pressure type casting system. The castalloys 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 (¢) phase intensities increased while ¡/alpha prime (¡A) intensities decreased with increased heat treatment temperature(HTT) and Mo concentration. The ¢ phase was observed to dominate the 800°C-treated Ti­9.5Mo alloy, while the highest alpha double prime(¡AA) phase content was observed in the 800°C-treated Ti­7.5Mo alloy. Both optical and scanning electron microscopy indicated that a relativelycoarse ¡ platelet was always observed in Ti­1.5Mo. A fine, uniformly-distributed acicular microstructure was observed in Ti­7.5Mo, while anequi-axed ¢ granular microstructure was clearly seen in Ti­9.5Mo. 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 withMo concentration. The highest YS and UTS were found in Ti­7.5Mo and Ti­9.5Mo. When heat-treated at 750°C, the strength of Ti­5.5Mo wasimproved without reducing elongation. With Mo concentration increased to 7.5% or higher, the elongation further increased while the strengthmaintained a similar level. When treated at 800°C, the YS of Ti­3.5Mo, Ti­5.5Mo and Ti­7.5Mo maintained a lower level than Ti­1.5Mo andTi­9.5Mo. A fully satisfactory interpretation for the tensile properties and their relationships to the complicated microstructures might not be asimple task due to several different factors simultaneously involved, yet practically it is interesting to note that selected alloys heat-treated withinthe dual-phase field demonstrated quite promising overall mechanical properties. [doi:10.2320/matertrans.M2017371]

(Received December 4, 2017; Accepted February 19, 2018; Published April 6, 2018)

Keywords: titanium alloy, titanium-molybdenum alloy, ¡/¢ dual-phase

1. Introduction

Due to their light weight, high corrosion resistance andspecific strength, pure titanium and titanium alloys have beenwidely used for dental and orthopedic applications such ascrown and bridge, removable partial denture, dental implants,hip prosthesis, artificial knee joint and trauma-fixationdevices,1­6) but not without drawbacks. For example, thepopularly-used c.p. Ti has a relatively low mechanicalstrength.5,7­9) Ti­6Al­4V ELI, despite its excellent mechani-cal properties,10­12) has the potential problem of releasingaluminum and particularly vanadium ions from the alloyimplant which might cause long term health problems such asAlzheimer’s disease and cytotoxicity.13­16) Another potentialproblem of c.p. Ti and Ti­6Al­4V ELI is their much higherelastic modulus values (typically between 100 and 120GPa3)than that of natural bones (about 10­20GPa17)). These muchhigher modulus values could activate the stress-shieldingeffect potentially leading to bone atrophy or even failure ofthe implant.9,18,19)

The search for Al and V-free, biocompatible Ti alloys fororthopedic implant applications was initiated in the mid-80s,and among them ¢ and near-¢ alloys have caught the mostattention.8,20­23) The ¢-type Ti alloys containing largeamounts of such heavy alloy elements as Mo, Nb, Taand/or W demonstrated a better biocompatibility, betterformability, and lower elastic modulus level than ¡ and ¡/¢-type Ti alloys. Nevertheless, their relatively high meltingtemperatures, high densities and high costs are some majorconcerns for the application of these ¢-type alloys.

Using a different approach, the present authors’ team has

developed an Al and V-free, low modulus ¡AA-type Ti­7.5Moalloy with strength/modulus ratios significantly higher thanthose of popularly-used 316L, Co­Cr­Mo and Ti­6Al­4Valloys.24) The results of the study indicated that ¡AA phase hada lower modulus than all other phases in the binary Ti­Moalloy system. To obtain the ¡AA phase, the alloy underwent asolution treatment (ST) in the ¢ phase field followed by awater quench (WQ) process.24,25) In another early study fromthe authors’ lab,26) the Ti­7.5Mo alloy with a tensile modulusof 78GPa and a Ti­6Al­4V alloy with a tensile modulus of110GPa were implanted into rabbit femur. It was interestingto find that, after 26 weeks, the amount of new bone attachedonto the Ti­7.5Mo implant was >5 times larger than thatonto the Ti­6Al­4V implant. This large difference wasthought to derive from a combined effect of chemistry (thepresence of harmful Al and V in Ti­6Al­4V) and elasticmodulus (the stress-shielding effect).

Conventionally the ¡AA phase in a binary Ti­Mo alloy isobtained by direct casting from the molten state24) or heatingthe alloy into the ¢ phase field followed by fast cooling, oftenwith a water or ice water quench.1,6) Most studies conductedthe solution treatment of Ti­Mo-based alloys in the ¢ phasefield, but few studies were devoted to investigating the heattreatment (HT) effect within the ¡/¢ dual-phase field. Thestudies of Cardoso et al.,1) Jiao et al.27) and Lu et al.28)

involved HTwithin the ¡/¢ zone but were not focused on theHT parameter-structure-mechanical property relationships ofthe Ti­Mo system in a systematic way. The primary purposeof the present study was therefore to investigate the effect ofHT within the ¡/¢ dual-phase field on the structure andtensile properties of a series of binary Ti­Mo alloys. It alsoattempted to clarify whether the ¢ phase formed in the ¡/¢regime at much lower temperatures could transform into theinherently low modulus ¡AA phase.

+1Graduate Student, National Cheng-Kung University+2Corresponding author, E-mail: chernlin@mail.ncku.edu.tw

Materials Transactions, Vol. 59, No. 5 (2018) pp. 734 to 740©2018 The Japan Institute of Metals and Materials

2. Materials and Methods

The Ti­Mo alloys of different compositions (1.5, 3.5, 5.5,7.5 and 9.5mass%Mo) used for this study were prepared fromgrade 2 commercially pure (99.8mass% pure) titanium (ChinaSteel Co., Taiwan) and 99.95mass% 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 evacuatedand purged with argon gas. An argon pressure of 0.147MPawas maintained during melting. To prepare each alloy,appropriate amounts of Ti and Mo metals were melted in aU-shaped copper hearth with a tungsten electrode. The ingotwas re-melted three times to improve chemical homogeneityof the alloy.

Prior to casting, the alloy ingot was re-melted in anopen-based copper hearth in argon gas under a pressure of0.147MPa. The difference in pressure between the twochambers allowed the molten alloy to quickly drop into amold at room temperature. To investigate the effect ofheat treatment in the ¡/¢ dual-phase field (Fig. 1), the castsamples were heat-treated at three different temperatures(700, 750 and 800°C) in vacuum for 30 minutes followed byquenching in ice water.

X-ray diffraction (XRD) for phase analysis was conductedusing a Bruker D2 Phaser diffractometer operated at 30 kVand 10mA with scanning speeds of 2°/min and 0.1°/min.A Ni-filtered CuK¡ radiation was used for the study. Asilicon standard was used for the calibration of diffractionangles. The various phases were identified by matching eachcharacteristic peak in the diffraction patterns with JCPDSfiles (Joint Committee on Powder Diffraction Standards, nowcalled International Centre for Diffraction Data, ICDD).

Microstructural examination of the series of samples wasperformed using an optical microscope (Leica TMX 100,Germany). The surfaces of the materials for light microscopywere mechanically polished via a standard metallographicprocedure to a final level of 0.05 µm alumina powder,followed by chemical etching in a mixture of water, nitricacid, 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 alsoused for microstructural examination in more details. Thesamples for SEM examination were prepared under the sameprocedure 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 usingan electrical discharge machining system (V50, ExcetekTechnologies Co., Taiwan). The dogbone-shaped, reduced-sized specimens for testing were 55mm long, 10mm wideand 1.0mm thick with a gage length of 6mm and gage widthof 3mm. The testing was performed at room temperaturewith a constant crosshead speed of 8.33 © 10¹6ms¹1. Theaverage ultimate tensile strength (UTS), yield strength (YS)at 0.2% offset, Young’s modulus of elasticity, and elongationto failure were taken from six tests under each condition.The measurement of Young’s modulus of elasticity wasconducted following the method set forth in ASTM E111-17,wherein the value of Young’s modulus was obtained bydetermining the slope of the straight-line portion of the stress-strain profile in the tensile test.

3. Results and Discussion

The XRD patterns of Ti­(1.5­9.5)Mo alloys heat-treatedat 700, 750 and 800°C with a scan speed of 2°/min from 30to 90° (2ª) are shown in Figs. 2a, 2b and 2c, respectively.Figure 2 indicates that, at any given heat treatmenttemperature (HTT), the XRD intensities of ¢ phase increased,while the ¡/¡A peak intensities decreased with increased Moconcentration, as expected from the Ti­Mo phase diagram.29)

It should be reminded that ¡ and ¡A phases have the samecrystal structure (hcp) and are indistinguishable from theirXRD patterns. These two phases are often distinguishedfrom each other by their different microstructures. ¡A phasehas a fine, martensitic-type acicular microstructure usuallyobtained from a fast cooling process, while ¡ phase usuallyexhibits a coarser plate-shaped microstructure.30) When thepresent alloys were heat-treated within the dual-phase fieldfollowed by WQ, it was possible for ¡ and ¡A phases to bothform in the alloys. For this reason, the term “¡/¡A peaks”was used in this study to represent the hcp peaks possiblyattributed to both ¡ and ¡A phases in all the XRD patterns.According to the binary Ti­Mo phase diagram, the relativeamount of ¡ phase should increase with decreased Moconcentration when the alloys were heat-treated in the dual-phase field followed by WQ.

Figure 2 also indicates that, when HTT increased, theXRD intensities of ¢ phase increased while the ¡/¡A peakintensities decreased, which was also expected from Ti­Mophase diagram. The most-dominated ¢ phase was observedin the alloy containing the highest Mo concentration(9.5mass%) heat-treated at the highest HTT (800°C). In theearly study of Ho et al.,24) a Mo concentration of 10mass%or higher was capable of retaining substantially all the ¢

phase during the fast-cooling casting process. When thealloys of the present study were heat-treated at lower HTT,their retained ¢ phase amounts were lower due to the loweramounts of ¢ phase formed in the dual-phase field.

Fig. 1 The highlighted ¡/¢ dual-phase regime (700­800°C; 1.5­9.5mass%)Mo in the Ti­Mo phase diagram29) investigated in this study.

Effect of Heat Treatment within Alpha/Beta Dual-Phase Field on the Structure and Tensile Properties of Binary Ti­Mo Alloys 735

Although distinguishing ¡A from ¡ from XRD patterns wasdifficult, the identification of the orthorhombic ¡AA phase wasrather straightforward. ¡AA phase could be identified in thesplitting of single ¡/¡A peaks into double peaks of ¡AA phase.According to the early studies of Baker31) and Brownet al.,32) this fast cooling-induced athermal orthorhombicstructure was derived from a distorted hexagonal cell with thec-axis of the orthorhombic cell corresponding to the c-axis ofthe hexagonal cell and a/b corresponding to the orthogonalaxis of the hexagonal cell. In the early study of Ho et al.24) ofa series of binary Ti­Mo alloys, the amount of ¡AA phase wasfound very sensitive to the Mo concentration, and thefast-cooled binary Ti­7.5mass% Mo was dominated by theorthorhombic ¡AA phase.

Due to the often low and diffuse intensity distribution ofthe split characteristic peaks, the XRD conducted at low scanspeeds was found helpful in the identification of ¡AA phase.The XRD patterns scanning from 52 to 55° (2ª) of the alloysheat-treated at 800°C at a low scan speed (0.1°/min) areshown in Fig. 3. The splitting of the individual ¡/¡A (102)peak into its corresponding double ¡AA (112/022) peaks isclearly shown in these low-scan-speed XRD patterns. It wasalso observed that, with increased Mo concentration, the

spacing between the ¡AA double peaks became larger and the¡AA (022) peak shifted toward the high angle side, consistentwith the early finding of Bagariatskii et al.33)

The highest ¡AA phase content accompanied with a low ¢

content was observed in the 800°C-treated Ti­7.5Mo alloy,which is reasonable since this temperature is quite closeto the ¢-transus of the alloy. As mentioned earlier, a Moconcentration of 7.5mass% produced the highest content of¡AA phase in the Ti­Mo binary system during WQ. The presentXRD results indicated that, depending on HTT and Moconcentration, the ¢ phase formed in the dual-phase regimecould be partially retained, transformed into ¡A phase ortransformed into ¡AA phase during WQ. These different phasescould sensitively affect the mechanical properties of the Ti­Mo alloys, as will be discussed later.

Typical optical and scanning electron micrographs of theTi­(1.5­9.5)Mo alloys heat-treated at 700, 750 and 800°Care given in Figs. 4­6 and 7­9, respectively. As shown inthe optical micrographs, at any HTT, the relatively coarseplatelet-shaped ¡ phase microstructure appeared less inquantity with increased Mo concentration. In Ti­1.5Moalloy, the ¡ platelet microstructure was observed at alltemperatures due to its high ¡ phase content in the dual-phasefield (Figs. 4­6). The higher-magnification SEM micrographsrevealed that there existed numerous fine needles (arrow inFig. 7a) within the primary coarse platelets retained from thehigh temperature which could not be detected from opticalmicroscopy. Whether these fine needles were formed duringcooling as excess Mo atoms were expelled from the low-

Fig. 2 XRD patterns of Ti­(1.5­9.5mass%)Mo alloys heat-treated at(a) 700°C, (b) 750°C and (c) 800°C with a scan speed of 2°/min from30 to 90° (2ª).

Fig. 3 XRD patterns of Ti­(1.5­9.5mass%)Mo alloys heat-treated at800°C with a scan speed of 0.1°/min from 52 to 55° (2ª).

Y.-P. Peng, C.-P. Ju and J.-H. C. Lin736

temperature, low-Mo concentration ¡ phase (as shown inTi­Mo phase diagram) is not certain at this moment. It seemshighly unlikely that these fine needles were ¢-transformed ¡Aphase due to the small amount of ¢ phase in the dual-phaseregime, especially at lower temperatures. With increased

HTT, these needles appeared thicker and the retained ¢ phasegrain boundaries became more visible (arrows in Fig. 9a).

As the Mo concentration increased to 3.5 or 5.5mass%,the optical micrographs showed that the acicular-shapedfine needle microstructure increased in quantity due to the

Fig. 5 Optical micrographs of Ti­(1.5­9.5mass%)Mo alloys heat-treatedat 750°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and(e) Ti­9.5Mo.

Fig. 4 Optical micrographs of Ti­(1.5­9.5mass%)Mo alloys heat-treatedat 700°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and(e) Ti­9.5Mo.

Fig. 6 Optical micrographs of Ti­(1.5­9.5mass%)Mo alloys heat-treatedat 800°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and(e) Ti­9.5Mo.

Fig. 7 Scanning electron micrographs of Ti­(1.5­9.5mass%)Mo alloysheat-treated at 700°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and (e) Ti­9.5Mo.

Effect of Heat Treatment within Alpha/Beta Dual-Phase Field on the Structure and Tensile Properties of Binary Ti­Mo Alloys 737

increased ¢ phase content in the dual-phase field whichpartially transformed into fine ¡A/¡AA needles during WQ(Figs. 4­6). According to the aforementioned XRD patterns,these needles were most likely ¡A/¡AA needles. Again, thisneedle type microstructure revealed in the higher magnifica-

tion SEM micrographs indicated that the needles becamethicker and the retained ¢ phase became more recognizabledue to the higher ¢ phase content with increased HTT. Theformation of the even finer structures observed within certainzones, e.g., along the retained ¢ boundaries (arrows in Fig. 9c)is not fully understood. Identification of this fine structure isdifficult with the present SEM. One hypothesis might be dueto the fluctuations in Mo concentration in the 800°C-treated ¢regime. During WC, the higher-Mo zones tended to retain as ¢phase while the lower-Mo zones tended to transform into ¡AAphase. The size of this fine structure seems to be limited by thedimension of these lower-Mo zones.

When the Mo concentration increased to 7.5mass%, theoptical micrographs showed a fine needle ¡AA (identified byXRD) microstructure uniformly distributed throughout thealloy. The aforementioned XRD patterns indicated that thehighest orthorhombic ¡AA phase content was seen at this Moconcentration. Although ¡A and ¡AA had a similar martensitic-type fine needle microstructure, as mentioned earlier,distinguishing the orthorhombic ¡AA phase from the hcp ¡Aphase was rather easy by XRD. When the Mo concentrationfurther increased to 9.5mass%, the equi-axed ¢ granular typemicrostructure was clearly revealed in the optical as well asSEM micrographs. Consistent with XRD, although ¢ phasedominated the XRD pattern of the 800°C-treated Ti­9.5Moalloy, both optical and SEM micrographs clearly showed thepresence of other phase, most likely ¡AA (identified by XRD),indicating that the ¢ phase in this study could neither beentirely retained, nor entirely transformed into ¡A or ¡AA phaseduring WQ.

The tensile properties of the investigated Ti­Mo alloysheat-treated at different temperatures are demonstrated inFig. 10. Figures 10a, 10c and 10e respectively showed theYS, UTS and elongation of the alloys treated at 700, 750 and800°C, while 10b, 10d and 10f respectively showed theYoung’s modulus values of the alloys treated at the same threetemperatures. When the alloy was heat-treated at 700°C, bothYS and UTS increased with increasing of Mo concentration,while the elongation decreased in general. However, theproperties of Ti­7.5Mo and Ti­9.5Mo were almost the same.Since the ¢ phase content increased with increasing Moconcentration while the Mo concentration in the ¢ phaseremained the same, the strong ¢-strengthening effect seenat 700°C was probably due to the high Mo concentration ofthe ¢ phase in the dual-phase field by a solute-strengtheningmechanism. According to the Ti­Mo phase diagram, the Moconcentration of the ¢ phase at 700°C could be estimatedas 21mass%Mo, which could effectively strengthen the¢-phase. However, the Young’s modulus of Ti­7.5Mo wasconsiderably higher than that of the ¡AA-dominated binary Ti­7.5Mo.24) In the absence of ¡AA phase, all the modulus valueswere observed between 90­100GPa (Fig. 10b) probably dueto the similar modulus level of ¡, ¡A and ¢ at 700°C.

When heat-treated at 750°C, the YS, UTS and elongationof Ti­1.5Mo and Ti­3.5Mo alloys were almost identical tothe same alloys treated at 700°C. When the Mo concentrationincreased from 3.5 to 5.5mass%, the strength increasedwithout reducing elongation. When the Mo concentrationfurther increased to 7.5% or higher, the strength maintained asimilar level, while the elongation further increased to 31%.

Fig. 8 Scanning electron micrographs of Ti­(1.5­9.5mass%)Mo alloysheat-treated at 750°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and (e) Ti­9.5Mo.

Fig. 9 Scanning electron micrographs of Ti­(1.5­9.5mass%)Mo alloysheat-treated at 800°C. (a) Ti­1.5Mo, (b) Ti­3.5Mo, (c) Ti­5.5Mo, (d) Ti­7.5Mo and (e) Ti­9.5Mo.

Y.-P. Peng, C.-P. Ju and J.-H. C. Lin738

The lower strength and higher elongation values (comparedto 700°C-treated alloys) may be explained by the lower Moconcentration of the ¢ phase at 750°C, although the ¢ fractionwas higher than that of 700°C. The modulus values of the750°C-treated alloys also maintained a level of 90­100GPa(Fig. 10d).

The YS, UTS, elongation and Young’s modulus values ofthe 800°C-treated Ti­1.5Mo alloy were almost the same asthose of the 700°C and 750°C-treated alloys, indicating thedominant role of the ¡ phase in this low Mo alloy. Whenthe Mo concentration increased, however, the YS and UTSbehaved quite differently in 700°C or 750°C-treated alloys.While the UTS continued to increase to the maximum value(877.6MPa) found in Ti­7.5Mo alloy, the YS decreased andmaintained a much lower level (556.0MPa) in Ti­3.5Mo,Ti­5.5Mo and Ti­7.5Mo alloys than in Ti­1.5Mo (658.7MPa)and Ti­9.5Mo (667.6MPa). It is interesting to note that thesethree alloys with constant low YS demonstrated consistentlylower modulus values (about 91GPa) than the other two alloyswhen treated at 800°C (Fig. 10f ). While a large differencebetween YS and UTS in an alloy accompanied with a lowmodulus is commonly observed in a fast-cooled ¡AA-dominated

Ti­7.5Mo alloy,6) whether the decreased YS and modulus inthese three 800°C-treated alloys (especially Ti­3.5Mo and Ti­5.5Mo) was attributed to the formation of ¡AA phase was notcertain. The present XRD results indicated that the 800°C-treated Ti­7.5Mo alloy contained much more ¡AA phasethan Ti­3.5Mo and Ti­5.5Mo alloys treated at the sametemperature. If ¡AA phase played amajor role in lowering the YSand modulus, the YS and modulus values of these three alloyscould not be so close. When the Mo concentration furtherincreased to 9.5mass%, the UTS and elongation decreased,while the YS and modulus increased. The decreased UTSmight be attributed to the grain growth effect of the ¢ phase.The relatively low elongation and high modulus seen in800°C-treated Ti­9.5Mo alloy might hypothetically be relatedto the formation of the brittle ½ phase. According to the Ti­Mophase diagram, the equilibrium Mo concentration at 800°Cis very close to 9.5mass%Mo. The early study of Ho et al.24)

indicated that the strongest½ effect in the binary Ti­Mo systemwas found in the alloy containing 10mass%Mo. The study ofZhang et al.34) also indicated that the Ti­Mo alloy containingabout 6 at% (³11mass%) Mo had the strongest ½ effectaccompanied with a high modulus.

Fig. 10 Tensile properties of Ti­(1.5­9.5mass%)Mo alloys heat-treated at different temperatures. (a, c, e) showing YS, UTS andelongation values of the alloys heat-treated at 700, 750 and 800°C, respectively; (b, d, f ) showing Young’s modulus values of the alloysheat-treated at 700, 750 and 800°C, respectively.

Effect of Heat Treatment within Alpha/Beta Dual-Phase Field on the Structure and Tensile Properties of Binary Ti­Mo Alloys 739

From above results and discussion, it can be seen that afully satisfactory interpretation for all the tensile propertiesand their relationships to the complicated microstructures ofthe present Ti­(1.5­9.5)Mo alloys treated within ¡/¢ 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 tensileproperties seem to at least include the contents of differentphases (¡, ¡A, ¡AA, ¢, and possibly ½) and the concentrationsof Mo in the different phases. Yet, however complicatedthe interpretation of the mechanical properties might be,practically it is interesting to note that some of the alloysheat-treated within the dual-phase field demonstrated quitepromising overall mechanical properties. Taking the 700°C-treated Ti­9.5Mo alloy as an example, with a similarelongation (14.5%) to the popularly-used Ti­6Al­4V ELI(10%), the 700°C-treated Ti­9.5Mo alloy demonstrated muchhigher YS and UTS with a lower modulus (952MPa,957MPa and 96GPa, respectively) than Ti­6Al­4V ELI(795MPa, 860MPa and 114GPa, respectively). Anotherexample is the 750°C-treated Ti­7.5Mo which demonstrateda little higher YS (820MPa), lower UTS (837MPa), lowermodulus (101GPa), and a far higher elongation (31%) thanTi­6Al­4V ELI (10%).

4. Conclusions

(1) The XRD patterns indicated that ¢ phase intensities in-creased while ¡/¡A intensities decreased with increasedHTT and Mo concentration. The most-dominated ¢

phase was observed in the 800°C-treated Ti­9.5Moalloy, while the highest ¡AA phase content was observedin the 800°C-treated Ti­7.5Mo alloy. Depending onHTT and Mo concentration, the ¢ phase formed in thedual-phase field could be partially retained, transformedinto ¡A or into ¡AA phase during WQ.

(2) Both optical and scanning electron microscopy in-dicated that, at any HTT, a relatively coarse ¡ plateletmicrostructure was always observed in Ti­1.5Mo. Theplatelet microstructure became finer with increasedMo concentration. Even finer and more uniformlydistributed platelets/needles were observed in Ti­7.5Mo. The equi-axed ¢ granular microstructure wasclearly revealed in Ti­9.5Mo. The ¢ phase in this studycould neither be entirely retained, nor entirely trans-formed into ¡A/¡AA during WQ.

(3) The tensile properties were found sensitive to the HTTand Mo concentration. When heat-treated at 700°C,the YS and UTS increased while the elongationgenerally decreased with Mo concentration. The highestYS and UTS were found in Ti­7.5Mo and Ti­9.5Mo.When heat-treated at 750°C, the strength of Ti­5.5Mowas improved without reducing elongation. Whenthe Mo concentration increased to 7.5% or higher,the elongation further increased while the strengthmaintained a similar level. When treated at 800°C, theYS of Ti­3.5Mo, Ti­5.5Mo and Ti­7.5Mo maintaineda lower level than Ti­1.5Mo and Ti­9.5Mo.

Acknowledgment

The authors would like to acknowledge the support forthis research by the Ministry of Science and Technology,Republic of China under the Research Grant No. MOST104-2221-E-006-142-.

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