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Materials Science and Engineering A 527 (2010) 4262–4269

Contents lists available at ScienceDirect

Materials Science and Engineering A

journa l homepage: www.e lsev ier .com/ locate /msea

igh-strength nanostructured Ti–12Mo alloy from ductile metastable beta staterecursor

. Suna, F. Primab, T. Glorianta,∗

INSA Rennes, UMR CNRS 6226 SCR/Chimie-Métallurgie, F-35043 Rennes cedex, FranceENSCP, UMR CNRS 7045 LPCS/Métallurgie Structurale, F-75231 Paris cedex 05, France

r t i c l e i n f o

rticle history:eceived 3 November 2009eceived in revised form 10 March 2010ccepted 11 March 2010

a b s t r a c t

In this study, the metastable beta Ti–12Mo (wt.%) alloy was synthesized by the cold crucible levitationmelting (CCLM) method and then quenched in water from the beta domain. The tensile test carried outon the as-quenched alloy showed an important ductility reaching 45% of elongation before fracture. Fromthe as-quenched metastable beta state, non-isothermal electrical resistivity and dilatometry measure-

eywords:itanium alloyetastable phases

ransmission electron microscopy

ments were carried out to detect the phase transition in order to control the nanophase precipitationsequence. After heating, XRD analysis and TEM observations revealed a very fine nano-scale omega andalpha precipitation in the beta matrix. Tensile test results indicated a very high strengthening effectafter appropriate thermo-mechanical treatments, which was observed particularly huge after a two-step annealing process where a tensile strength as high as around 1600 MPa was obtained. This highly

h waspha n

ensile test enhanced tensile strengtconsisting of two-scale al

. Introduction

Over the last few decades, the interest for titanium alloys hasontinuously increased due to their properties of high strength,xcellent hardenability, low density and good corrosion resis-ance making these alloys very interesting for many industrialpplications. In titanium-based alloys, it has been reported a pos-ible nanostructuration leading to a general improvement of theechanical properties [1,2]. The optimization of these improvedechanical properties in Ti-based alloys depends on the volume

raction, the size, the morphology and the distribution of theifferent phases or nanophases that composed the microstruc-ure. This microstructure can be optimized by controlling thehermo-mechanical parameters such as the temperature and timef annealing, the heating and cooling rate, the ageing sequencend/or the deformation rate. It has already been reported thathe nanostructuration of finely distributed � phases in � matrixan be developed with the help of intragranular nucleation sitesor � phase. This nanostructuration is due to the presence ofhe nanoscale � phase (20–30 nm in size), which is observed in

etastable � Ti-based alloys (quenching from the � phase field)ontaining appropriate �-stabilizer elements such as Fe, Mo, Ta, Nb,r and/or V [3–7]. In high misfit systems, such as binary Ti–V alloyswhere the isothermal � nanophases form in cuboidal precipitates),

∗ Corresponding author. Tel.: +33 2 2323 8241/8240; fax: +33 2 2323 8241/8240.E-mail address: Thierry.Gloriant@insa-rennes.fr (T. Gloriant).

921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2010.03.044

attributed to the complex intragranular nanostructure observed by TEManoprecipitates inside sub-micrometer beta grains.

© 2010 Elsevier B.V. All rights reserved.

� phases were observed to nucleate at ledges of dislocations at the�/� interfaces [6,8]. In low misfit systems, such as binary Ti–Moalloys [9], low-cost � (LCB) alloy, Ti-5553 alloy (where isothermal� nanophases form in ellipsoidal precipitates), there are differentresults regarding the mechanism of the �-assisted nucleation of �phase. The role of � nanophases in �/� transformation in LCB alloyhas been well discussed by Prima et al. [10] with micrographic evi-dences of � nanophase heterogeneous nucleation from the core of �particles by high-resolution electron microscopy. On the contrary,Azimzadeh and Rack [11] indicate that the � precipitates in LCBtitanium alloy nucleate near, but at a certain distance away from,the �/� interfaces due to the enrichment of Al (the element of �-destabilizer) in the vicinity around the � precipitates. Anyway, themost recent investigation on �-assisted nucleation and growth of �precipitates in � titanium Ti-5553 alloy proposed by Nag et al. [12]clearly confirms that � precipitates are definitely associated with �precipitates, which in most cases act as heterogeneous nucleationsites within the � matrix. Furthermore, these precipitates possessa relatively finer scale (usually nanometer to submicrometer) andoften exhibit morphologies which are different from � precipitatesnucleated at � grain boundaries.

It is well known that a strengthening effect from nanos-tructuration in �-metastable alloys can be achieved by proper

thermo-mechanical processes including solution treatment, coldrolling, recovery/recrystallization and subsequent ageing. Inorder to control the finer � precipitates, low-temperature-high-temperature two-step ageing was often applied on the quenchedmetastable �-type Ti alloys [13–15]. At low ageing temperature,

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F. Sun et al. / Materials Science an

etastable phases (�′ precipitate formed by the phase separa-ion of � and isothermal � precipitate [16]) was observed toct as nucleation sites for � phase formation at higher ageingemperature. Furthermore, heavy cold rolling and well-controlledecovery/recrystallization prior to the final two-step ageing areuggested to be advantageous in the � grain size refinement andlso in the nanostructure formation as observed with the �–Ti-15-alloy [17]. The ageing after recovery/recrystallization treatment

an provide a uniform distribution of � precipitates with multi-ariants because the heterogeneous structure induced by coldolling is replaced by the formation of fine � subgrains beforehe precipitation of intragranular � phase. Therefore, the recov-ry/recrystallization treatment prior to ageing is very effective tobtain more uniform distribution of � precipitates.

In this study, a metastable � Ti–12Mo (wt.%) alloy was synthe-ized by cold crucible levitation melting. The phase transformationequence of the quenched alloy was investigated by employinglectrical thermo-resistivity and thermo-dilatometry measure-ents. These methods are well known to be very sensitive

o minor constitutional changes produced by quenching, age-ng and precipitation treatments [18–21] and then particularlyseful to investigate the �iso and �nano phase transformationequence on heating. In addition, X-ray diffraction (XRD), opti-al microscopy (OM) and transmission electron microscopy (TEM)ere used for the structural characterization. With these com-

ined approaches, critical temperatures of transition and phaseransformation sequences were determined in order to control theormation of micro- and nano-scaled � phases from the metastable

state. Besides, thermo-mechanical processes including solutionreatment, heavy cold rolling, recrystallization and two-step ageingere applied to optimize the formation of the desired nanostruc-

ure. Tensile tests were employed to investigate the tensile strengthnd the ductility of the treated Ti–12Mo alloy.

. Experimental

The binary Ti–12Mo (wt.%) alloy was synthesized by the coldrucible levitation melting (CCLM) technique by using an induc-ion furnace (CELES high frequency generator). The melting was

arried out under a pure Ar atmosphere after several cycles ofigh vacuum pumping. Notable features of CCLM are that it canelt metals with a high melting point, create an alloy of uniform

omposition and prevent the crucible contamination. The obtainedngots with a weight about 20 g were remelted in a laboratory

Fig. 1. Optical microscopy images showing the microstructure of the Ti–12

neering A 527 (2010) 4262–4269 4263

scale arc furnace (7400 TUBINGEN, Edmund Buhler), with a tung-sten alloy electrode on a water-cooled copper module in a purityargon atmosphere. Near cylindrical ingots with a diameter of about10 mm and a length of about 25 mm were obtained. Once synthe-sized, thermo-mechanical treatments were carried out. The ingotswere firstly solution treated in the �-phase domain at 1223 K underhigh vacuum (10−6 mbar) for 2–3 h in a tubular furnace and thenquenched in water at room temperature to obtain the metastable �microstructure (as-quenched state). Then the ingots were cut into3 mm thick plates and cold rolled with a laboratory scale laminator(Caltex Ursa Heavy) into 0.5 mm in thickness, which correspond toa reduction rate of about 80%. The thin plates were reheated upto 1143 K for 30 min for the recrystallization and then quenchedinto water (as-quenched recrystallized state). Two-step annealingprocesses were applied after complete recrystallization in order toachieve high-strength properties. The recrystallization and anneal-ing treatments were carried out in tubular furnaces under highvacuum (10−7 mbar) to prevent oxidation. All the specimens werecleaned in an acidic bath (10% HF in nitric acid HNO3) before andafter every step.

The resistivity measurements were performed by using thefour-probe method with a serial electrical circuit set-up composedof: the sample (in contact with the thermocouple for measur-ing the temperature), a reference resistor Rref (2 �), a d.c. sourceand a computer-controlled data acquisition system (self-designedapparatus). All the experiments were carried out under high vac-uum (10−6 mbar). Samples for electrical resistivity measurementswere cut into a thin 20 mm × 2 mm × 0.5 mm lamella. For thedilatometry measurements (self-designed apparatus), the mea-surement of dilatometry variation, �L, of the sample (in contactwith the thermocouple for measuring the temperature) was carriedout under high vacuum (10−6–10−7 mbar). Samples for dilatom-etry measurements were shaped into a square rod (dimension:18 mm × 4.2 mm × 4.2 mm). A constant heating rate of 5 K/min wasused for both thermal analysis methods.

X-ray diffraction analysis (XRD) was used to examine themicrostructure (generator Philips PW3710 using CuK� radi-ation operating at 30 kV-20 mA). Samples were observed byoptical microscopy (MO, PME OLYMPUS), by scanning electron

microscopy (SEM, JSM6400 with EDS OXFORD Link Isis) and bytransmission electron microscopy (TEM, JEOL 2000FX TEM oper-ating at 200 kV). Thin foils were prepared by twin-jet electropolishing technique using a solution of 4% perchloric acid inmethanol.

Mo alloy: (a) as-quenched state, (b) as-quenched recrystallized state.

4264 F. Sun et al. / Materials Science and Engineering A 527 (2010) 4262–4269

Fig. 2. Electrical resistivity and dilatometry curves from the as-quenched Ti–12Moalloy (heating rate of 5 K/min).

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ig. 3. XRD profiles of the Ti–12Mo alloy: (a) as-quenched state, (b) annealed at51 K for 40 h, (c) annealed at 773 K for 8 h, (d) self-cooled from 1223 K.

Tensile tests were performed using a INSTRON 3369

est machine with an extensometer (10 mm long) to obtaintrain–stress curves and to calculate Young’s modulus. The tensileest specimens possess a normalized shape: 4 mm width, 0.5 mmn thickness and a gage length of 24 mm.

Fig. 5. TEM observation of the Ti–12Mo alloy annealed at 473 K for 1

Fig. 4. TEM electron diffraction pattern of the as-quenched Ti–12Mo alloy.

3. Results and discussion

3.1. Microstructure and phase transformation sequence from theas-quenched state

Optical micrographs showing the microstructure of the as-quenched state (after solution treatment) and the as-quenchedrecrystallized state (after cold rolling and recrystallization treat-ment) are presented in the Fig. 1(a and b), respectively. Typicalequiaxe �-grain microstructure with a grain dimension of few hun-dreds �m in size is observed in the as-quenched alloy (Fig. 1(a)).After cold-rolling and recrystallization treatment (Fig. 1(b)), the �grain size was observed to be much smaller with a dimension offew tens �m in size. It has to be mentioned the excellent workhardening behavior of this Ti–12Mo alloy composition because thecold rolling at room temperature was done very easily without any

problem of cracking.

The non-isothermal characterization of the transformation pro-cess from the as-quenched metastable � Ti–12Mo alloy wasperformed by electrical resistivity and dilatometry measurements

h: (a) dark-field image, (b) corresponding diffraction pattern.

F. Sun et al. / Materials Science and Engineering A 527 (2010) 4262–4269 4265

Table 1Phase transformation sequence from the metastable � Ti–12Mo alloy.

Phase transformation sequence

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Alloy Transition temperatures (K) Zones (phases after annealing)

T�iso T�nano T�eq T� A B C DTi–12Mo 451 573 866 1093 � + �ath � + �iso � + �nano � + �

nder continuous heating condition. Fig. 2 shows the resistivitynd the derivative dilatometry variation curves from 300 K to about200 K (heating rate: 5 K/min).

The transformation sequence from the metastable � state car-ied out on different Ti-based alloys has been recently establishedy our group and the curves obtained in the present study are inccordance with previous observations [18–22]. The phase trans-ormation sequence can be separated into 4 zones: A, B, C and Dnd typical X-ray diffraction patterns (XRD) obtained after appro-riate thermal treatment related to each zone are presented inhe Fig. 3. Thus, it has already been reported that the negativeemperature dependence observed on the resistivity curve at lowemperature (T < 450 K zone A in Fig. 2) is due to the presence of thethermal omega phase (�ath), which is formed during quenching3,18,23–25]. This athermal omega phase is usually very difficult toe detected by X-ray diffraction in the as-quenched state (Fig. 3(a))ut by electron diffraction technique in a transmission electronicroscope (TEM), the typical weak diffraction spots correspond-

ng to two �ath variants in the � matrix can be observed as shownn the TEM micrograph presented in Fig. 4.

At higher temperature, the hexagonal isothermal omegaanophase (�iso) was observed to precipitate from �ath implyingiffusional phenomenon. The formation of this �iso phase on heat-

ng is shown to occur at 451 K (identified by T�iso in Fig. 2), wheren increase of the resistivity and a decrease of the dilatometry arebserved simultaneously (zone B in Fig. 2). XRD profile clearly indi-ates the presence of �iso phases after long time annealing (40 h)t T�iso (Fig. 3(b)).

The presence of �iso was confirmed by TEM observations ashown in the Fig. 5. These micrographs are related to the Ti–12Molloy annealed at 473 K for 1 h (zone B) and a very fine �iso precipi-ates of few nanometers in size are observed in the dark-field imageFig. 5(a)) and the corresponding diffraction pattern (Fig. 5(b)).

rom the XRD profile (Fig. 3(b)), some additional peaks indicatehe formation of the metastable �′ phase, which appears by phaseeparation mechanism from � phase. A previous study showed thatsothermal � and �′ phases could both act as a nucleation site for �

Fig. 6. Tensile curve from the as-quenched recrystallized Ti–12Mo alloy.

Fig. 7. Optical microscope image of the Ti–12Mo alloy after tensile test.

phase [26]. However, the annealing time required for the develop-ment of � phase separation is relatively long prior to � precipitation[14,27]. Consequently, the � phase separation will be avoided in thepresent study concerning the two-step annealing treatments thatwe will apply on the Ti–12Mo alloy.

From 573 K (zone C), the resistivity drop and the dilatome-try rate change observed (indicated by T�nano in Fig. 2) are dueto the progressive vanishing of �iso phase through a �/� phasetransformation. At about 790 K a continuous decrease in resistivity(followed by a sudden increase in derivative of dilatometry) after anear flat stage (rectangle area indicated in Fig. 2), is clearly detected,which probably indicates the end of phase transformation from �isophase to �nano phase and the subsequent growth of intragranular �phase. XRD analysis confirms the formation of the � phase in zoneC as shown in Fig. 3(c).

By heating the alloy at a higher temperature, the precipitationof a conventional intragranular and intergranular � phase occursat 866 K, where a sudden change in both resistivity and dilatome-try curves is observed (T�eq in Fig. 2). This state correspond of thereturn to the equilibrium � + � state (zone D) and example of typ-ical XRD profile from Ti–12Mo after being heated up to 1223 K ispresented in Fig. 3(d).

Finally, the � transus temperature was also detected by resis-tivity and dilatometry on heating. As indicated in the Fig. 2, T� wasfound at about 1093 K. All the transition temperatures detected byresistivity and dilatometry measurements and the different phasesobserved by XRD and TEM in each zone are summarized in theTable 1.

3.2. Tensile test of the as-quenched recrystallized state

Fig. 6 presents the stress-elongation curve of the recrystallizedTi–12Mo alloy obtained by tensile test. The tensile curve indicatesa yield stress at about 550 MPa, followed by an excellent ductilitywith an elongation percentage as high as 45%. The Young’s modulus

4266 F. Sun et al. / Materials Science and Engineering A 527 (2010) 4262–4269

age of the �′′ martensites (b) dark-field image of a single �′′ martensite.

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Fig. 8. TEM images of Ti–12Mo after tensile test: (a) bright-field im

as measured by using an extensometer and a value close to 90 GPaas found.

Fig. 7 shows the optical micrograph of the deformed microstruc-ure after fracture in tension. It can be clearly observed on this

icrograph the presence of dark zones in the � grain microstruc-ure related to the presence of the �′′ martensite as reported by theiterature. It has been well known that the deformation mechanismf metastable � Ti alloys is strongly affected by the concentrationf the � stabilizer [7,28]. In less stable � titanium alloys, stress ortrain induced �′′ martensite has been reported as in the presentase [29]. Besides, the formation of {3 3 2} 〈 1 1 3〉 twins can prob-bly happen in the present case of having such a relatively largeetastability of � phase [30,31]. Furthermore, it has been reported

hat the �ath phase formed here by quenching is one of the mostmportant conditions necessary for the {3 3 2} 〈1 1 3〉 twin forma-ion [31]. During the tensile test, such twinning requires only halfhe shear strain needed for {1 1 2} 〈1 1 1〉 twin, and its formations also interrelated with the �′′ martensite transformation [29].herefore, the presence of both strain induced �′′ martensite and3 3 2} 〈1 1 3〉 twinning and the interaction between them are ableo contribute a ductile behavior in tensile test.

The twinned �′′ martensite in the deformed as-quenchedi–12Mo alloy was observed by TEM and typical micrographs areresented in Fig. 8. From the bright-field image (Fig. 8(a)), it can belearly seen the �′′ martensite (in dark) formed parallel along theame direction and, from the dark-field image at higher magnifica-ion (Fig. 8(b)), the nano-scaled twinning inside (in dark). Therefore,he Ti–12Mo alloy exhibits excellent ductility in spite of the twin

ode and stress or strain induced �′′ phases.

.3. Microstructures and tensile tests after two-step annealingrocesses

The as-quenched recrystallized Ti–12Mo alloy was furtherhermally treated in two-step mode and its tensile behavior

able 2ensile properties of the Ti–12Mo alloy.

Thermal treatments and tensile properties

Alloy Heat treatment Phase YS (MPa) U

Ti–12Mo As-quenched � + �ath 550 6573 K(8 h) + 773 K(8 h) � + �nano 950 10573 K(8 h) + 793 K(8 h) � + �nano 1600*

* Fracture before yield.

Fig. 9. Stress–strain curves of the Ti–12Mo alloy after two-step annealing treat-ments.

was investigated. It is well-known that the two-step annealing(low/high or high/low) is an efficient procedure to control the phasetransformation and final microstructure. In this work, tensile spec-imen was first heated with a low 5 K/min rate to obtain a finedispersion of �iso nanoprecipitates and then annealed at 573 K (atT�nano) for 8 h to ensure the transformation �iso/�nano. The sec-ond annealing step was carried out at 773 K or 793 K for 8 h inorder to control the concentration and distribution of intragran-ular � phases. In the Fig. 9 are presented the stress–strain curves of

the thermally treated alloy obtained by tensile tests carried out atroom temperature.

After the two-step annealing treatment carried out at 573 K(8 h) and then at 773 K (8 h), an enhanced yield stress at around

TS (MPa) Elongation (�l/l0) Young’s modulus (0.05–0.4%) (GPa)

60 45% 9020 9% 120

– <1% 120

F. Sun et al. / Materials Science and Engineering A 527 (2010) 4262–4269 4267

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ig. 10. Optical microscope images from the Ti–12Mo alloy after two-step annealin73 K for 8 h, (b) and (d) 573 K for 8 h and 793 K for 8 h.

50 MPa and an acceptable ductility with an elongation of about% are obtained. In a second tentative, the second annealing tem-erature was brought to 793 K (8 h) (instead of 773 K). The results

ndicate that the ductility was completely lost but a very impressive

ig. 11. Scanning electron microscope observations of the Ti–12Mo alloy after two-step annb) and (c) SEM images at lower magnification with EDS profiles across grain boundaries.

tments (observation at two different magnifications): (a) and (c) 573 K for 8 h and

strengthening as high as about 1600 MPa was obtained before frac-ture. In order to verify such a high tensile strength, two more tensiletests were realized. Although the alloy become very brittle after thetwo-step annealing treatment, the very high strengthening effect

ealing at 573 K for 8 h and 793 K for 8 h: (a) SEM image of the general microstructure,

4268 F. Sun et al. / Materials Science and Engineering A 527 (2010) 4262–4269

F (573 Kt field

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ig. 12. TEM images from the Ti–12Mo alloy after a two-step annealing treatmenthe � subgrain microstructure, (c) bright-field image of � nanoprecipitates, (d) dark

as confirmed and values reaching 1520 and 1570 MPa at ruptureere obtained, respectively. The tensile properties obtained in theresent work are summarized in the Table 2. It can be noticed an

ncreasing of the Young’s modulus after the two-step annealingrocesses from 90 to 120 GPa.

Fig. 10 shows the optical micrographs obtained from thei–12Mo alloy after the two two-step annealing treatments.ccording to the temperatures of annealing used, � precipitatesust be present in the microstructure. With the first two-step

nnealing treatment, where a lower secondary annealing tem-erature at 773 K was applied, the alloy presents a beta grainicrostructure and the � phase, probably nanometer in scale, that

emains invisible under optical microscope (Fig. 10(a) and (c)). Onhe contrary, when the second annealing temperature is higher793 K), small dark intragranular precipitates can be observed inhe beta grains by optical microscopy (Fig. 10(b) and (d)).

On these optical micrographs, the � grain boundaries seemery thick and intergranular � precipitation might occurs duringhe thermal treatment, although the intergranular � precipitationould appears only at T�eq = 866 K (see Table 1.). Consequently, the

i–12Mo microstructure after the two-step annealing treatmentas also observed by scanning electron microscopy (SEM) in order

o clarify this point. SEM images of the microstructure, observedith 3 different magnifications, are presented in the Fig. 11. The

eta grain microstructure is well visible on these micrographs and

/8 h and 793 K/8 h): (a) bright-field image of the � phase, (b) bright-field image ofimage of one � grain showing the low misorientation of the subgrains.

the chemical composition across the grain boundaries was evalu-ated by energy dispersive spectroscopy (EDS analysis in Fig. 11(b)and (c)). The EDS results indicates that no important variation ofMo concentration exists across the � grain boundary, which meansthat the � formation at the boundaries did not happen yet after two-step annealing treatments and the refined � phase was uniformlydistributed inside the � grains.

In order to observe the nanostructured intragranular � precipi-tation, transmission electron microscope observations (TEM) wereperformed. TEM micrographs obtained from the Ti–12Mo alloyafter a two-step annealing treatment at 573 K for 8 h and 793 K for8 h are presented in the Fig. 12.

Needle-like � precipitates of about 45 nm in width and of250–450 nm in length are observed in the beta matrix as shownin the bright-field image presented in Fig. 12(a). These sub-micrometric intragranular � precipitates grew in two directions(indicated by black arrows in Fig. 12(a)) with an angle of about 70◦,which is in accordance with the literature [14]. Besides, other �precipitates in much smaller dimension (about 20 nm in size) werealso observed in � in some areas separated from the needle-like �

precipitation zone (Fig. 12(c)). Furthermore, TEM observations atlower magnification reveal that these two kinds of � nanoprecipi-tates are distributed through a beta subgrain microstructure. Thisbeta subgrain microstructure is shown on the bright-field imagepresented in Fig. 12(b). The presence of � subgrains of few hun-

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red nanometers in size inside micrometer � grains (as observedy optical microscopy) in the Ti–12Mo alloy is probably inducedy the heavy cold-rolling and recrystallization treatment [17]. Theark-field image presented in Fig. 12(d) shows one � grain com-osed by several � subgrains of about 500 nm in size, and the lowlack/white contrast across the subboundaries indicates a low mis-rientation between each other. The bright-field image in Fig. 12(b)llustrates the general microstructure inside a � grain, where small

nanoprecipitates (around 20 nm) are finely distributed in the �ubgrains (labeled A and E) and needle-like � precipitates (about0 nm in width and 200–400 nm in length) both inside the sub-rains (labeled B, C and D) and at the subgrain boundaries (indicatedy the white arrows). The combination of � nanoprecipitates inwo scales and subgrains in � grain constitutes the intragranu-ar microstructure observed under optical microscopy (dark zonesn Fig. 10(d)). It is clear that the raising density of this kind ofntragranular microstructure at higher secondary annealing tem-erature contributes to the huge increasing of tensile strengthhen the Ti–12Mo alloy was annealed at 573 K for 8 h and then

t 793 K for 8 h. Nevertheless, this strengthening effect is accompa-ied by a strong reduction of the ductility. In addition, it is observedhat a higher secondary annealing temperature within the temper-ture range 760–810 K (the flat stage indicated by dash square inig. 2) was able to generate also a larger fraction of � precipitationt the � subgrain boundaries leading to a totally lost of ductility.onsequently, the best compromise between high strength and suf-cient ductility is the key challenge. Works are in progress with thisbjective.

. Conclusion

In this work, the transformation sequence from the metastableTi–12Mo alloy was investigated by electrical resistivity and

ilatometry measurements. Under heating, the Ti–12Mo alloyhowed a nanophase transformation sequence into � and � pre-ipitates from the metastable � as-quenched state.

Tensile test carried out on the recrystallized as-quenched stateas demonstrated an important ductility, where 45% of elongationas reached before fracture. The reason can be explained by the fact

hat the �ath nanoparticles contributed to the formation of peculiar3 3 2} 〈1 1 3〉 twins and stress or strain induced �′′ martensites,

hich lead to the excellent ductility observed.

Thermal treatments were applied on the Ti–12Mo alloy in ordero favor the formation of well-dispersed nanoprecipitates in theeta matrix. Tensile test results indicated a very high strengthen-

ng effect after appropriate thermo-mechanical treatments, which

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neering A 527 (2010) 4262–4269 4269

is particularly huge with the Ti–12Mo alloy where a tensile strengthas high as around 1600 MPa was obtained after a two-step anneal-ing treatment. This highly enhanced tensile strength was attributedto the complex intragranular nanostructure observed by TEM con-sisting of two-scale alpha nanoprecipitates inside sub-micrometerbeta grains.

Thus, the present Ti–12Mo alloy composition that combines animportant ductility at room temperature and a very high strength-ening effect by controlling the thermal treatment seems verypromising in the view to develop a new generation of high-strengthnanostructured Ti-based alloys.

References

[1] S. Ankem, C.A. Green, Mater. Sci. Eng. A263 (1999) 127.[2] H. Nishizawa, E. Sudekai, W. Liu, H. Hashimoto, Mater. Trans. JIM 36 (1995) 924.[3] J.C. Williams, M.J. Blackburn, Met. Soc. AIME-Trans. 245 (1969) 2352.[4] T. Yamane, A. Miyakubi, Met. Soc. AIME 2 (1980) 1309.[5] U.M. Naik, R. Krishman, Soc. AIME 2 (1980) 1335.[6] Y. Takemoto, M. Hida, A. Sakakibara, J. Jpn. Inst. Met. 57 (1993) 261.[7] T.W. Duerig, G.T. Terlinde, J.C. Williams, Metall. Trans. A 11 (1980) 1987.[8] M.J. Blackburn, J.C. Williams, Trans. Met. Soc. AIME 242 (1968) 2461.[9] G.M. Pennock, H.M. Flower, D.R.F. West, Titanium’80′ Sci. Technol. (1980) 1343.10] F. Prima, P. Vermaut, G. Texier, D. Ansel, T. Gloriant, Scripta Mater. 54 (2006)

645.11] S. Azimzadeh, H.J. Rack, Metall. Mater. Trans. A 29 (1998) 2455.12] S. Nag, R. Banerjee, R. Srinivasan, J.Y. Hwang, M. Harper, H.L. Fraser, Acta Mater.

57 (2009) 2136.13] O.P. Karasevskaya, O.M. Ivasishin, S.L. Semiatin, Yu.V. Matviychuk, Mater. Sci.

Eng. A 354 (2003) 121.14] O.M. Ivasishin, P.E. Markovsky, S.L. Semiatin, C.H. Ward, Mater. Sci. Eng. A 405

(2005) 296.15] O.M. Ivasishin, P.E. Markovsky, Yu.V. Matviychuk, S.L. Semiatin, C.H. Ward, S.

Fox, J. Alloys Compd. 457 (2008) 296.16] G.T. Terlinde, T.W. Duerig, J.C. Williams, Metall. Trans. A 14 (1983) 2101.17] T. Furuhara, T. Maki, T. Makino, J. Mater. Process. Technol. 117 (2001) 318.18] T. Gloriant, G. Texier, F. Prima, D. laillé, D.M. Gordin, I. Thibon, D. Ansel, Adv.

Eng. Mater. 8 (2006) 961.19] M.A. Hill, D.H. Polonis, J. Mater. Sci. 22 (1987) 2181.20] F. Prima, P. Vermaut, D. Ansel, J. Debuigne, Mater. Trans. JIM 41 (2000) 1092.21] T. Gloriant, G. Texier, F. Sun, I. Thibon, F. Prima, J.L. Soubeyroux, Scripta Mater.

58 (2008) 271.22] D.M. Gordin, E. Delvat, R. Chelariu, G. Ungureanu, M. Besse, D. Laille, T. Gloriant,

Adv. Eng. Mater. 10 (2008) 714.23] M. Ikeda, S.Y. Komatsu, T. Sugimoto, K. Kamei, Proceedings of Sixth World

Conference on Titanium, les Editions de Physique, Paris, 1988, p. 313.24] D. de Fontaine, Metall. Trans. A 19A (1988) 169.25] E. Sukedai, H. Hashimoto, M. Tomita, Phil. Mag. A 64 (1991) 1201.26] J.C. William, B.S. Hickman, D.H. Leslie, Metall. Trans. 2 (1971) 477.27] T. Makino, T. Furuhara, T. Maki, in: S. Fujishiro, D. Eylon, T. Kishi (Eds.), Metal-

lurgy and Technology of Practical Titanium Alloys, TMS, Warrendale, PA, 1994,p. 63.

28] T. Grosdidier, C. Roubaud, M. Philippe, Y. Combres, Scripta Mater. 36 (1997) 21.29] Y. Takemoto, M. Hida, A. Sakakibara, J. Jpn. Inst. Met. 60 (1996) 1072.30] Y. Takemoto, I. Shimizu, A. Sakakibara, Metall. Trans. 45 (2004) 1571.31] S. Hanada, O. Izumi, Metall. Mater. Trans. A 17 (1986) 1409.