Texture Evolution in a Ti-Ta-Nb Alloy Processed by SeverePlastic Deformation

VASILE-DANUT COJOCARU,1,3 DOINA RADUCANU,1

THIERRY GLORIANT,2 and ION CINCA1

1.—University Politehnica of Bucharest, Spl. Independentei 313, 060042 Bucharest, Romania.2.—UMR CNRS 6226 Sciences Chimiques de Rennes, 35043 Rennes Cedex, France. 3.—e-mail:dan.cojocaru@mdef.pub.ro

Titanium alloys are extensively used in a variety of applications because oftheir good mechanical properties, high biocompatibility, and corrosion resis-tance. Recently, b-type Ti alloys containing Ta and Nb have received muchattention because they feature not only high specific strength but also bio-corrosion resistance, no allergic problems, and biocompatibility. A Ti-25Ta-25Nb b-type titanium alloy was subjected to severe plastic deformation (SPD)processing by accumulative roll bonding and investigated with the aim toobserve the texture developed during SPD processing. Texture data expressedby pole figures, inverse pole figures, and orientation distribution functions forthe (110), (200), and (211) b-Ti peaks were obtained by XRD investigations.The results showed that it is possible to obtain high-intensity share texturemodes ({001}h110i) and well-developed a and c-fibers; the most important fiberis the a-fiber ({001} 1�10

� �to {114} 1�10

� �to {112} 1�10

� �). High-intensity texture

along certain crystallographic directions represents a way to obtain materialswith high anisotropic properties.

INTRODUCTION

Ultrafine-grained (UFG) and nanocrystalline(NC) materials are found to exhibit outstandingmechanical properties, such as high strength,toughness, and superelasticity at ambient temper-atures compared with their coarse grained coun-terparts.1–4 Severe plastic deformation (SPD) is oneof the most effective methods of producing UFG/NCmaterials in bulk dimensions. Various SPD pro-cesses such as equal-channel angular pressing(ECAP),5–7 accumulative roll-bonding (ARB),8 high-pressure torsion (HPT),9,10 repetitive corrugationand straightening (RCS),11 torsion extrusion,12 andsevere torsion straining13 have been developed.

ARB is a promising production method for con-tinuous production of large bulk multilayeredsheets for different industrial applications. TheARB processing consists of rolling of two metalsheets (ARB precursor) with equal dimensionsusing a 50% thickness reduction, in one pass,resulting in one single bonded sheet that has thesame thickness as the originals sheets. The ARBcycle is repeated until the desired multilayered

sheet stacks are obtained. Using ARB processing, itis possible to achieve an extremely high plasticstrain because theoretically, the number of ARBpasses can be repeated without a limit. At each ARBpass, each single layer can support an addeddeformation up to theoretical limit.

In previous studies, the ARB process has suc-ceeded in producing pancake-shaped or elongatedUFG and NC structures on different material types.It was also suggested that the formation mechanismof UFG/NC during ARB can be explained in terms ofgrain subdivision at a submicron scale,14–16 whereinitial coarse grains have been subdivided bydeformation-induced, high-angle grain boundaries.In the case of body centered cubic (bcc) materials,during thermomechanical (TM) processing, thegrains can align toward preferred orientations,resulting different texture modes and texture fibers,such as a-fiber, c-fiber, g-fiber, h-fiber, etc.17–19

Titanium and its alloys are widely used in aero-space, automotive, electronics, biomedical, andenergy applications because of their low density–strength ratio, high biocompatibility, and highcorrosion resistance.20,21 During the last decade,

JOM, Vol. 64, No. 5, 2012

DOI: 10.1007/s11837-012-0312-6� 2012 TMS

572 (Published online April 20, 2012)

special attention has been directed to titanium-based orthopedic implants, where b-type alloysstand out because of their higher biocompatibilityand low elastic modulus.22,23 In this context, addi-tions of Nb and Ta to Ti are usual based on theirbiological passivity and capacity of reducing theelastic modulus.24–26 These elements are widelyused in combination with Zr, Sn, and Mo.27,28

In all these application fields, the crystallographictexture developed during TM processing plays acrucial role on the physical properties of the materi-als.29–31 For this reason, texture changes in titaniumalloys during TM processing are of interest.32–36

In the case of Ti-25Ta-25Nb (wt.%) alloy, atambient temperatures, the microstructure mayconsist of an equiaxed body-centered cubic betaphase (b-Ti), which may also may containing mid-dle-like orthorhombic secondary-alpha phase(a¢¢-Ti). The a¢¢-Ti phase appears and grows becauseof the stress-induced phase transition. The pro-cessing route of Ti-25Ta-25Nb alloy plays animportant role in steering the phase structure(constitutive phases and phase quantities) as wellas the texture formation toward desired properties.

The focus of this investigation was to assess theTi-25Ta-25Nb alloy’s capacity to form UFG/NCstructures through SPD and investigate developedtexture during ARB processing.

METHODS

The investigated Ti-25Ta-25Nb (wt.%) alloy wasproduced from commercially pure elements usinga levitation induction melting furnace FIVECELES—MP25 (Fives Group, Lautenbach, France)with nominal power 25 kW and melting capacity30 cm3, under an argon protective atmosphere. Thechanges in chemical composition caused by thesynthesis by induction melting in levitation are lessthan 0.1 wt.% and therefore are judged to be negli-gible. Because of the large difference in meltingtemperatures (Ti: 1.660�C; Ta: 2.996�C; and Nb:2.468�C), the ingots were remelted two times toobtain a high degree of chemical homogeneity.

The Ti-25Ta-25Nb alloy was processed by a com-plex TM processing route (Fig. 1) to obtain a thinprecursor strip used to obtain multilayered stackedstrips, by ARB, consisting in 4- and 16-layer stacks.

The TM processing consists of first cold rollingwith a thickness reduction of approximately 84.32%(equivalent strain e = 1.85) in 20 rolling passes,followed by a recrystallization treatment in argonprotective atmosphere, at 850�C for 30 min, using aGERO SR 100 9 500 heat treatment oven (GEROHochtemperaturofen GmbH & Co. KG, Neuhausen,Germany). The recrystallization treatment wasperformed to remove the strain-hardening effectsthat resulted during first cold rolling. After recrys-tallization, a second cold rolling was performed,with a thickness reduction of approximately 71.05%

(equivalent strain e = 1.24) in 16 rolling passes, toachieve a final strip thickness of 141 lm. This finalstrip consists as ARB precursor layer.

The ARB processing consists of four ARB passes toobtain 4-layer (after second ARB pass) and 16-layer(after fourth ARB pass) stacks.

All cold-rolling and ARB passes were performedusing a LQR120AS rolling mill (Mario di MaioS.p.A., Saronno, Italy), at 3 m/min rolling speed,without any lubrication.

The as-cast, ARB precursor and ARB-processedspecimens (4-layer and 16-layer stacks) were x-raydiffraction (XRD) characterized, using a PanalyticalX’Pert PRO MRD (PANalytical, Almelo, the Nether-lands) diffractometer with a wavelength of Cu K-a(k = 1.5418 A), to determine the phase structure andphase characteristics. The recorded XRD spectra werefitted using PeakFit v4.11 software package (Systat

Fig. 1. Thermomechanical and ARB processing route of Ti-25Ta-25Nb alloy.

Texture Evolution in a Ti-Ta-Nb Alloy Processed by Severe Plastic Deformation 573

Software, Inc., San Jose, CA) to determine for eachdiffraction peak the position, intensity, and peakbroadening at fullwidthathalfmaximum parameters.

To analyze the resulted texture for ARB precursorand ARB-processed specimens (4-layer and 16-layerstacks), the (110), (200), and (211) b-Ti pole figures(PFs) was measured, using a Philips PW 3710 dif-fractometer (Philips, Amsterdam, the Netherlands),with Cu K-a (k = 1.5418 A) wavelength. The PFraw data were fitted and analyzed using MTEXv3.2.2 (Dr. Ralf Hielscher, TU Chemnitz, Chemnitz,Germany and Florian Bachmann TU BergakademieFreiberg, Freiberg, Germany) open source softwarepackage37,38 to calculate inverse pole figures (IPFs)and orientation distribution functions (ODF’s).

RESULTS AND DISCUSSION

Microstructural Characterization

To investigate the changes in phase structure,phase quantities, lattice parameters, and coherentcrystalline domain size during SPD, XRD experi-ments were performed on as-cast state, ARB pre-cursor, ARB 4-layer state, and ARB 16-layer states.In Fig. 2, one can observe the Ti-25Ta-25Nb XRDspectra for investigated structural states; in allcases, the presence of b-Ti and a¢¢-Ti phases(Figs. 2a–d) can be observed.

Analyzing the XRD spectra for as-cast state(Fig. 2a), it was found that the b-Ti phase (indexedin body-centered cubic system—Im-3m) has the

lattice parameter a = 3.288 A whereas a¢¢-Ti phase(indexed in orthorhombic system—Cmcm) has thelattice parameters as follows: a = 3.181 A, b =4.813 A, and c = 4.631 A. It was calculated that theb-Ti and a¢¢-Ti phase quantities were 99.52% ±0.23% and 0.48% ± 0.11%, respectively. The aver-age coherent crystalline domain size for both b-Tiand a¢¢-Ti phases were calculated using simpleScherrer equation,39 taking into account b-Ti (110)and a¢¢-Ti (002) diffraction lines, respectively (fittedposition, intensity, and peak broadening). Theresulting average values for coherent crystallinedomain size in the case of b-Ti and a¢¢-Ti phaseswere as follows: 82 nm and 14 nm, respectively.

Analyzing the XRD spectra for ARB precursorstate (Fig. 2b), it was found that the lattice param-eter of the b-Ti phase changes to a = 3.284 Awhereas the lattice parameters of a¢¢-Ti phasechanges also to a = 3.211 A, b = 4.731 A, and c =4.632 A to accommodate the stress-induced a¢¢-Tiphase. The b-Ti and a¢¢-Ti phase quantities resultedfor ARB precursor state were as follows:93.79% ± 1.17% and 6.21% ± 0.68%, respectively.The a¢¢-Ti phase quantity grows because of thestress-induced transformation.40 The averagecoherent crystalline domain size, in the case of ARBprecursor state, for both b-Ti and a¢¢-Ti phases tak-ing into account b-Ti (200) and a¢¢-Ti (200) diffrac-tion lines, respectively, were as follows: 28 nm and31 nm, respectively. The same observations werereported in the case of Ti-22Nb-6Ta alloy41; theparent b-Ti showed a lattice parameter a = 3.289 Aand lattice parameters of a¢¢-Ti phase: a = 3.221 A;b = 4.766 A, and c = 4.631 A.

In the case of the ARB 4-layer state (Fig. 2c), itwas found that for both b-Ti and a¢¢-Ti phases,almost no lattice parameters change took place, onlycoherent crystalline domain size and phase quanti-ties changes. It was calculated that coherent crys-talline domain size reaches 24 nm in b-Ti case and21 nm in a¢¢-Ti case, respectively. Concerning phasequantities, it was calculated that the b-Ti phasereaches 61.19% ± 2.84% and a¢¢-Ti phase reaches38.81 ± 3.09, respectively.

Analyzing the XRD spectra for ARB 16-layer state(Fig. 2d), no changes in lattice parameters werefound for both b-Ti and a¢¢-Ti phases. The calculatedcoherent crystalline domain size reaches 22 nm inb-Ti phase case and 16 nm in a¢¢-Ti phase case.Concerning phase quantities, it was calculated thatthe b-Ti phase reaches 57.38% ± 4.31% and a¢¢-Tiphase reaches 42.62% ± 3.11%.

As a general trend, one can say that during ARBprocessing, the b-Ti phase coherent crystallinedomain size shows a minimal decrease of approxi-mately 21%, from 28 nm (for ARB precursor state)to 22 nm (for ARB 16-layer state), whereas the a¢¢-Tiphase shows a moderate decrease of approximately48%, from 31 nm (for ARB precursor state) to 16 nm(for ARB 16-layer state). Also, one can say thatduring ARB processing, a¢¢-Ti phase quantity

Fig. 2. XRD spectra of Ti-25Ta-25Nb alloy in as-cast state (a), ARBprecursor state (b), ARB 4-layer processed state (c), and ARB16-layer processed state (d).

Cojocaru, Raducanu, Gloriant, and Cinca574

continuously increase, from 6.21% ± 0.68% (forARB precursor state) to 42.62% ± 3.11% (for ARB16-layer state). This finding shows that it is possibleto steer the phase structure (phase quantities andcoherent crystalline domain size) by ARB processingto obtain the desired final material nanostructureswith outstanding properties. If we are interested inobtaining high material strength, then we need tosteer the structure toward small coherent crystal-line domains, but if we intend to obtain a desiredelastic modulus in a certain interval range, then weneed to steer the structure toward a certain phasequantity b-Ti/a¢¢-Ti ratio (b-Ti phase exhibit a lowelastic modulus whereas the a¢¢-Ti phase exhibits ahigher elastic modulus).

Texture Characterization

Using (110) raw PF data, the IPFs and ODFs werecalculated and plotted. The following assumptionswere made: The material crystalline symmetry wasindexed in cubic Im-3m system and the specimensymmetry in orthorhombic mmm system.

Figure 3 shows collected raw data for (110), (200),and (211) b-Ti PFs in the case of ARB precursorstate (see Fig. 3a), ARB 4-layer state (see Fig. 3b),and ARB 16-layer state (see Fig. 3c). All PFs wereplotted with respect to the following directions(sample reference frame): [100]—rolling direction(RD); [010]—transversal direction (TD), and[001]—normal direction (ND).

Fig. 3. Images of (110), (200), and (211) raw data pole figure of Ti-25Ta-25Nb alloy for ARB precursor state (a), ARB 4-layer processed state (b),and ARB 16-layer processed state (c).

Texture Evolution in a Ti-Ta-Nb Alloy Processed by Severe Plastic Deformation 575

A well-developed rolling-texture is confirmed inARB precursor specimens (Fig. 3a), the (110) PFshows four high-intensity peaks symmetricallylocated around 45� from the center and 45� fromboth RD and TD. This means that the RD is parallelto [110] crystal directions. The (200) PF shows onlythe presence of one high-intensity central peak,whereas the (211) PF shows four symmetric high-intensity peaks located around 28� form the centerand 45� from both RD and TD, and eight symmet-rically peaks located around 60� form the center and15� from RD and TD. In all observed (110), (200),and (211) PFs, the identified peaks show distribu-tion characteristic for bcc crystalline structures.

In the case of (110) and (200) PFs for both ARB4-layer state (Fig. 3b) and ARB 16-layer state(Fig. 3c), one can observe a similar peak distribu-tion like in the case of (110) and (200) PFs of ARBprecursor state. Differences can be observed only in

the case of (211) PF. These differences in peaksobserved in (211) PF may be explained by the stress-induced a¢¢-Ti phase presence (smaller coherentcrystalline size for both parent b-Ti phase andtransformed a¢¢-Ti phase and similar phase quan-tity) and b-Ti to a¢¢-Ti lattice accommodation.

Although the PF shows how the specified crys-tallographic directions are distributed in the samplereference frame, the inverse PF shows how the se-lected direction in the sample reference frame isdistributed in the reference frame of the crystal.Figure 4 shows the (110) inverse PF for RD, TD, andND directions in the case of ARB precursor state(Fig. 4a), ARB 4-layer state (Fig. 4b), and ARB 16-layer state (Fig. 4c).

If we consider the ARB precursor state (Fig. 4a),in the case of RD direction, the inverse PF shows uswhich crystallographic directions in the polycrys-talline material are most likely parallel to the

Fig. 4. (110) Inverse pole figure for RD, TD, and ND directions in the case of ARB precursor state (a), ARB 4-layers processed state (b), andARB 16-layer processed state (c).

Cojocaru, Raducanu, Gloriant, and Cinca576

sample RD; in this case, one can observe that the[011]//RD pair reaches an intensity close to 2.6. Inthe case of TD direction, it can be observed that alow intensity pair can be formed, [011]//TD, showingan intensity close to 1.5. High intensities areobtained in the case of [001]//ND and [111]//NDpairs; in this case, the intensities reach values closeto 2.5 for both pairs. As a general remark, based onthe inverse PF, most likely the developed texture forARB processed state will show [001]//ND, [111]//ND,and [011]//RD components. In the case of the ARB4-layer state (Fig. 4b), one can observe that thefollowing high-intensity pairs can be formed: [011]//RD, which shows an intensity close to 2.9 and [111]//ND, which shows a higher intensity of about 3.7. Alow-intensity pair can be observed for [011]//TD,showing an intensity close to 2. The same observa-tions can be made also in the case of an ARB16-layer state (Fig. 4c), but in this case, theobserved intensities are smaller; the [011]//RD pairshows an intensity close to 2, [011]//TD pair showsan intensity close to 1.6, and [111]/ND pair showsan intensity close to 2.4.

Analyzing the observed intensities for followingpossible pairs: [011]//RD, [011]//TD, [001]//ND, and[111]//ND, one can say that the texture index will bemaximum for an ARB 4-layer state and almostidentical for the ARB precursor and ARB 16-layerstates. Another remark is that the coherent crys-talline volume textured in a specific modal orienta-tion will be minimum in the case of ARB precursor(four average intensity texture pairs can coexist)and maximum in the case of ARB 4-layer state (twohigh-intensity texture pairs).

Because the properties of many important engi-neering materials are strongly direction dependent,the inverse PF is useful in predicting and calculat-ing the average properties of polycrystalline mate-rial along a chosen direction.42,43

The crystallographic textures are typically pre-sented in the reduced Euler space using the Bungesystem (u1–U–u2) (u1 = 0�–90�; U = 0�–90�; u2 = 0�–90�). The most relevant texture fibers developed inbcc metals are as follows17–19:

– a-fiber (crystallographic fiber axis h110i parallel tothe RD, including major components: {001}h110i,{112}h110i, and{111}h110i)

– c-fiber (crystallographic fiber axis h111i parallel tothe ND, including major components: {110}h110iand {111}h112i)

– g-fiber (crystallographic fiber axis h001i parallel tothe RD, including major components: {001}h100iand {011}h100i)

– f-fiber (crystallographic fiber axis h011i parallel tothe ND, including major components: {011}h100i,{011}h211i, {011}h111i, and {011}h011i)

– e-fiber (crystallographic fiber axis h011i parallel tothe TD, including major components: {001}h011i,{112}h111i, {111}h112i, and {011}h100i)

– h-fiber (crystallographic fiber axis h001i parallelto the ND, including major components: {001}h100i and {001} h110i).

In the case of bcc metals, the most important ODFsections correspond to rotation angle u2 equal to 0�and 45�.

In Fig. 5, one can observe the calculated ODFsections for ARB precursor state. The plotted ODFsection corresponding to u2 = 0� shows the presenceof the following texture modes:

– High-intensity texture modes {001} 1�10� �

and{010}h101i, showing a maximum orientation den-sity close to 5.9. Both texture modes belong toh-fiber ({001}h110i family) and are formed becauseof the grain shear during SPD processing.

The ODF section corresponding to u2 = 45� showsthe presence of the following texture modes andfibers:

– High-intensity texture modes 001f g 1�10� �

and001f g 1�10

� �: this mode shows a maximum orien-

tation density close to 5.9. Both texture modes001f g 1�10

� �and 001f g �1�10

� �belong to h-fiber

({001}h110i family) and are formed because ofthe grain shear during SPD processing.

– Well-developed c-fiber: the main component ofc-fiber spreads from 111f g 1�10

� �to 111f g �1�12

� �.

The c-fiber is the most important developed fiberduring TM processing, with an orientation den-sity close to 4.

– Well-developed a-fiber: the main component ofa-fiber spreads from 001f g �1�10

� �to 112f g 1�10

� �.

The orientation density of a-fiber shows a value closeto 4.6, which is bigger than in the case of c-fiber.

The ODF section corresponding to u2 = 0� of anARB 4-layer state (Fig. 6) shows the presence offollowing texture modes:

– Low-intensity texture modes 001f g 1�10� �

and{010}h101i: these modes show a maximum orien-tation density close to 3. Both texture modesbelongs to h-fiber ({001}h110i family) and areformed because of grain shear during ARBprocessing.

In the case of ODF section corresponding tou2 = 45�, one can observe the following texturemodes and fibers:

– Low-intensity texture modes 001f g 1�10� �

and001f g �1�10

� �: these modes show a maximum ori-

entation density close to 4. Both texture modes001f g 1�10

� �and 001f g �1�10

� �belong to h-fiber

({001}h110i family) and are formed because ofgrain shear during ARB processing.

– Well-developed c-fiber: the main component ofc-fiber spreads from 111f g 1�10

� �to 111f g �1�12

� �.

The c-fiber is the most important developed fiberduring ARB processing, with an average orienta-tion density close to 7.

Texture Evolution in a Ti-Ta-Nb Alloy Processed by Severe Plastic Deformation 577

Fig. 5. ODF for Ti-25Ta-25Nb alloy in ARB precursor state (ODF sections: u2 = 0�, 15�, 30�, 45�, 60�, and 75�).

Fig. 6. ODF for Ti-25Ta-25Nb alloy in ARB 4-layer processed state (ODF sections: u2 = 0�, 15�, 30�, 45�, 60�, and 75�).

Cojocaru, Raducanu, Gloriant, and Cinca578

– Well-developed a-fiber: the main component ofa-fiber spreads from 001f g 1�10

� �to 114f g 1�10

� �to

112f g 1�10� �

. The maximum orientation density of8.1 is reached for 112f g 1�10

� �texture mode.

In Fig. 7, the calculated ODF sections for ARB16-layer state are presented. Plotted ODF sectioncorresponding to u2 = 0� shows the presence offollowing texture modes:

– Low-intensity h-fiber ({001}h110i family) texturemodes 001f g 1�10

� �and {010}h101i: these modes a

maximum orientation density close to 3.

The ODF section corresponding to u2 = 45� showsthe presence of the following texture modes andfibers:

– Low-intensity h-fiber ({001}h110i family) texturemodes 001f g 1�10

� �and 001f g �1�10

� �: these modes

show a maximum orientation density close to 3.5.– Well-developed c-fiber: the main component of

c-fiber spreads from 111f g 1�10� �

to 111f g �1�12� �

.The c-fiber is the most important developed fiberduring SPD processing, with an average orienta-tion density close to 3.2.

– Well-developed a-fiber: the main component ofa-fiber spreads from 001f g 1�10

� �to 114f g 1�10

� �to

112f g 1�10� �

. The maximum orientation density of5.2 is reached for 112f g 1�10

� �texture mode.

If one knows the ODF’s calculations of texturemaximum modal orientation, the textured volume

with the same modal orientation as well as thetexture index are possible to calculate. To calculatethe textured volume with same maximum modalorientation, we assumed that the misorientationtolerance was 15�. The calculated values are pre-sented in Table I.

After analyzing the ODFs and data presented inTable I, one can say that at the end of TM process-ing, the most important developed texture mode is{001}h110i (as a result of grains share during TMprocessing); 7.76% of the TM processed materialvolume shows this modal orientation. During ARBprocessing, the maximum modal orientation chan-ges from {010}h101i to 112f g 1�10

� �; the a-fiber tex-

ture is the most important texture mode. Thetexture volume with maximum modal orientationincreases to 23.81% in the case of ARB 4-layer state;increasing the number of ARB passes results in adecrease in textured volume with maximum modalorientation. In the case of ARB 16-layer state, thetextured volume decrease to 16.81%. A similarbehavior is observed in the case of texture index.The maximum of texture index, of approximately2.748, is obtained for ARB 4-layer state.

High-intensity texture along certain crystallo-graphic directions represents a way to obtainmaterials with high anisotropic properties. In thisway, it is possible to steer the ARB processing toobtain a desired textured volume showing the samemaximum modal orientation.

Fig. 7. ODF for Ti-25Ta-25Nb alloy in ARB 16-layer processed state (ODF sections: u2 = 0�, 15�, 30�, 45�, 60�, and 75�).

Texture Evolution in a Ti-Ta-Nb Alloy Processed by Severe Plastic Deformation 579

During the plastic deformation of the alloy, grainorientation changes take place as a consequence ofshear on specific, favorable oriented crystal planesand directions. In the case of bcc materials, slip canoccur on {hkl}h111i slip systems, where {hkl} couldbe {110}, {112}, or {123}. Imposed strain and con-straint between neighboring grains affect the choiceand number of slip systems. Slip activation andtheir variation within and between grains deter-mine the deformation microstructure and thechange of grain orientation.

During cold rolling, dislocation glide will force theslip planes to rotate toward the tensile axis becauseof the constraint of the axis, and the rotation axis is[hkl] 9 [uvw], where [hkl] is the normal of the slipplane and the [uvw] is the slip direction. It is pos-sible to activate different texture modes {hkl}huvwithrough a mechanism of formation and rotation ofmicrobands.44

The development of {001}h110i texture is charac-teristic of bcc materials in the presence of share.45,46

The c-fiber ({110}h110i and {111}h112i mods) is well-developed in bcc materials under high loads becauseof the plain strain deformation state during defor-mation.45,46 The {001}h110i texture is a major roll-ing texture component observed in bcc metals suchas a-Fe, Ta, Mo, b-Ti, and W.47–49

It was reported that a correlation exists betweenthe deformation microstructure and the crystalorientation. It was suggested that during cold-rolleddeformation, the a-fiber grains can use as many asseven independent systems, which allow homoge-neous deformation in them. It was also indicatedthat under such circumstances, dislocation struc-tures in grain belonging to a-fiber must constantlyevolve during cold deformation, but how they do soremains unclear.44

In the case of an ARB-processed Ti-25Ta-25Nballoy, one can say that the developed texture mainlyshows high-intensity share texture modes ({001}h110iand well-developed a- and c-fibers, the most importantfiber being the a-fiber ( 001f g 1�10

� �to 114f g 1�10

� �to

112f g 1�10� �

).

CONCLUSIONS

A complex TM and SPD processing route was suc-cessfully used to fabricate NC/UFG Ti-25Ta-25Nballoy multilayered stacked strips. The microstructure

after TM processing showed formation of a NCmicrostructure consisting of b-Ti and a¢¢-Ti grainswith an average NC domain size close to 28 nm in thecase of b-Ti grains and 31 nm in the case of a¢¢-Tigrains. Subsequent grain refinement was obtained byARB processing; in the case of ARB 16-layer stacks,an average NC domain size close to 22 nm wasobtained in the case of b-Ti grains and 16 nm in thecase of a¢¢-Ti grains.

The b-Ti/a¢¢-Ti phase ratio shows that after TMprocessing, a small a¢¢-Ti/b-Ti ratio was obtained(the a¢¢-Ti phase quantity close to 6.21% ± 0.68%),whereas after ARB processing, the final a¢¢-Ti/b-Tiratio was much higher (the a¢¢-Ti phase quantityclose to 42.62% ± 3.11%).

Developed texture during ARB processing showsa strong {001}h110i rolling texture, with main ori-entation component ranges from f001gh1�10i tof001gh�1�10i. High-intensity texture fibers (a-fiberand c-fiber) were obtained also. In the case of c-fiber,the main orientation component ranges fromf111gh1�10i to f111gh�1�12i; whereas in the case ofa-fiber, the main orientation component rangesfromf001gh1�10i to f112gh1�10i. Further optimizationof the TM and ARB processing route to achievehigher texture intensities is desirable. The researchinto this topic is currently being undertaken by thecurrent authors.

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Table I. Texture characteristics developed in Ti-25Ta-25Nb alloy processed by TM-ARB

Maximum Modal Orientation (MO) ARB Precursor State ARB 4-Layer State ARB 16-Layer State

u1 (�) 44.82 6.72 7.13U (�) 89.93 39.75 39.81u2 (�) 0.06 44.22 43.77Texture component {hkl}huvwi {010}h101i 112f g 1�10

� �112f g 1�10

� �

Texture volume with same MO (%) 7.76 23.81 16.81Texture index 1.647 2.748 1.588

Cojocaru, Raducanu, Gloriant, and Cinca580

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Texture Evolution in a Ti-Ta-Nb Alloy Processed by Severe Plastic Deformation 581