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Acta Materialia 55 (2007) 5718–5727

In situ observation of texture evolution during a! band b! a phase transformations in titanium alloys investigated

by neutron diffraction

I. Lonardelli a,b, N. Gey c, H.-R. Wenk a,*, M. Humbert c, S.C. Vogel d, L. Lutterotti b

a Department of Earth and Planetary Science, University of California, Berkeley, CA 94720, USAb Dipartimento Ingegneria dei Materiali e Tecnologie Industriali, Universita di Trento, 38050 Trento, Italy

c LETAM CNRS UMR 7078, University of Metz, LETAM-ISGMP, F-57045 Metz Cedex 01, Franced Lujan Center, Los Alamos National Laboratory, Los Alamos, NM 87545, USA

Received 3 November 2006; received in revised form 27 May 2007; accepted 18 June 2007Available online 4 September 2007

Abstract

Texture changes during recrystallization and the a–b–a phase transformation in two titanium alloys were investigated in situ by time-of-flight neutron diffraction by heating in a vacuum furnace to 950 �C. In commercially pure titanium, a strong texture memory effect isobserved. This effect is a direct consequence of an orientation-selective a! b transformation, favoring new orientations produced duringnucleation and grain growth. The b–a transformation favors b orientations with minimal misorientations, resulting in a strong final atexture that emphasizes the grain growth component. In Ti–6Al–4V, the texture memory effect is less pronounced. The high-temperatureb texture is obtained by growth of pre-existing b nuclei. In a similar way, during cooling, the growth of a domains is controlled by high-temperature a orientations inherited from the b grains with Burgers orientation relation.� 2007 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Titanium; Texture; Phase transformation; Variant selection; Neutron diffraction

1. Introduction

Titanium and its alloys are widely used in aerospace,electronics, biomedical and energy applications due to theirlow density–strength ratio, high biocompatibility and highcorrosion resistance. In all these fields, the crystallographictexture developed during thermomechanical treatmentsplays a crucial role on the physical properties of the mate-rial [1–3]. For this reason, texture changes in titaniumalloys during cold work, primary recrystallization, graingrowth and the a/b phase transformation are of long-standing interest [4–6]. Concerning the phase transforma-tion, a texture memory effect is often reported [7–11] when

1359-6454/$30.00 � 2007 Acta Materialia Inc. Published by Elsevier Ltd. All

doi:10.1016/j.actamat.2007.06.017

* Corresponding author. Tel.: +1 510 642 7431; fax: +1 510 643 9980.E-mail address: wenk@berkeley.edu (H.-R. Wenk).

the hexagonal a metal is heated, transforms into the cubic bphase and returns to the hexagonal a phase after cooling.

To study this texture memory effect, knowledge aboutthe high-temperature b texture is required. However, notmuch information is generally available because it is diffi-cult to perform in situ texture measurements above800 �C. Recently new experiments have been proposedfor such investigations, using high-temperature scanningelectron microscopy with electron backscatter diffractionmeasurements [12,13], synchrotron X-ray diffraction [14]or time-of-flight neutron diffraction [8,9] at high tempera-ture. Furthermore, indirect methods have been developedto calculate the high-temperature b grain orientations fromthe a inherited orientations [15–19]. In all these studies, theorientation relation f0 001gh11�20i (hexagonal closepacked (hcp)-a)! f110gh1�11i (body-centered cubic(bcc)-b) observed by Burgers for single crystals [20] was

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I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727 5719

confirmed. During the a! b phase transition and accord-ing to this orientation relation, six potential crystallo-graphic b variants can occur from a single a orientation.Conversely, during the b! a phase transition, there are12 a potential variants. A variant selection can occur dur-ing the heating and/or the cooling phase transformations.The proportion of variants which can be modified duringnucleation and growth (variant selection) determines thefinal texture. Because this final texture influences the mate-rial properties at room temperature, it seemed to us inter-esting to study the texture evolutions of different samplesduring their thermal treatments.

In this contribution, we present the texture evolutionsof two titanium alloys during heating to 950 �C withemphasis on the a�b–a phase transformation. The tex-tures have been measured in situ by time-of-flight (TOF)neutron diffraction. The two titanium alloys had differentchemical compositions and microstructures. A Rietveldanalysis of diffraction spectra was applied to quantifythe changes of texture components during the transforma-tion. The a–b–a texture memory effect was analyzed byconsidering the a! b and b! a texture changes. Themain microstructural features that seemed to control thetransformation textures upon heating and cooling arediscussed.

2. Experimental

2.1. Material

Two titanium alloys were investigated. The first one wasa commercially pure (CP) a titanium sheet 5 mm thick. Itcontained the following impurities in weight ppm: C 52,Fe 237, H 3, N 41, O 1062 and Si <100 (ASTM grade 2).It was first hot rolled in the b and a/b field, quenchedand then cold rolled. Optical micrographs revealed a heav-ily deformed a structure.

The second alloy was a Ti–6Al–4V a/b titanium sheet. Itwas hot rolled in the a/b phase field and annealed. The con-centration of alloying elements and impurities in weightpercent was Al 6.03, V 3.97, Fe 0.15, O 0.18, C 0.03, N0.01 and H <0.01 (ASTM grade 5). It contained less than5% of residual b phase. The a matrix mainly comprised

Fig. 1. Optical micrographs of (a) CP–Ti (crystallographic contrast) and (b) T

equiaxed 20 lm grains; some coarse lamellae were alsopresent.

Cubes 5 · 5 · 5 mm3 were cut from each titanium sheetfor in situ neutron diffraction measurements up to950 �C. The microstructure after the in situ experimentsat 950 �C are given in Fig. 1a and b, respectively for theCP–Ti and Ti–6Al–4V alloy. Large grain growth isobserved in CP–Ti (Fig. 1a), the Ti–6Al–4V alloy is fine-grained, with b-lamellae in some grains (Fig. 1b).

2.2. Texture measurements

The texture measurements were performed on the neu-tron TOF diffractometer HIPPO (high pressure preferredorientation) at the Los Alamos Neutron Science Center(LANSCE) [21,22]. HIPPO is equipped with 1360 3Hedetector tubes arranged on 50 detector panels positionedon rings at five different diffraction angles and each detec-tor probes a different sample direction. For the presentstudy, only the 140� bank (8 detector panels), the 90� bank(10 detector panels) and the 40� bank (12 detector panels)were used. The incident beam size was 10 mm in diameter.The sample was clamped into a vanadium fixture in a vac-uum furnace with vanadium heating elements and heatshields. Vanadium has the advantage that it does not pro-duce coherent scattering, i.e. the diffraction spectra showonly peaks from the sample. For each texture measurementthe sample was rotated around the furnace axis (perpendic-ular to the incident neutron beam) into four positions toincrease pole figure coverage. At each position data werecollected for 18 min for CP–Ti and 10 min for Ti–6Al–4V, resulting in a total time for a temperature interval of72 and 40 min (including time for data readout and rota-tion), respectively.

Figs. 2a and b show the temperature profiles duringboth experiments. The pole figure measurements took placeduring heating at 400, 800, 850, 900, 950 �C and duringcooling at 400 �C. This corresponds to measurements inthe a field before recrystallization, then after recrystalliza-tion just before a! b transformation, in the a/b and bfield and finally after the reversed b! a transformation.Heating and cooling are rather sluggish, particularly atlow temperatures.

i–6Al–4V (topological contrast), after heating experiments up to 950 �C.

Fig. 3. Typical time-of-flight diffraction spectra of (a) CP–Ti and (b) Ti–6Al–4V, collected at 800, 900 and 950 �C with a and b phases coexisting.Experimental data (dotted line) and recalculated fit (solid line) are shown.Hexagonal peaks are marked with 4-digit Miller-Bravais indices, cubicpeaks with 3-digit Miller indices.

Fig. 2. Temperature as function of time during the neutron diffractiontexture measurements at LANSCE: (a) CP–Ti, (b) Ti–6Al–4V.

5720 I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727

Some experimental details are significant. With spalla-tion sources, neutron production is occasionally inter-rupted and thus the actual time for data collection maybe much longer than the counting time. Diffraction mea-surements with CP–Ti were straightforward; with Ti–6Al–4V the analysis was complicated by the fact that Tihas a negative and Al a positive scattering length, whereasV does not contribute to the coherent scattering, resultingin weak scattering.

2.3. Data analysis

The neutron diffraction spectra were analyzed with theRietveld method [23,24] as implemented in the programpackage MAUD [25]. The procedure is analogous to thatdescribed by Wenk et al. [21]. Data are refined taking intoaccount, first, instrument parameters and background, thencrystallographic and microstructural parameters, andfinally all parameters, including the texture and volumefractions of the phases. The input data are 4 · 30 = 120spectra that were refined simultaneously. Examples for bothCP–Ti and Ti–6Al–4V at different temperatures (800, 900and 950 �C) are shown in Fig. 3 with experimental data(dots) and refined profile (solid line). It is possible to observeclearly the a! b phase transformation due to the relativeintensity changes in a (hcp) and b (bcc) allotropic diffractionpeaks. Note that in Ti–6Al–4V, a and b phases coexist at alltemperatures, though with different volume fractions; inCP–Ti both polymorphs are present only at 900 �C.

The texture was refined inside the least-squares Rietveldrefinement procedure but in a separate cycle at each com-pleted iteration, using texture factors extracted by the LeBail algorithm as input and determining the orientationdistribution function (ODF) with the tomography algo-rithm WIMV [26], in a special version that allows data tobe entered at arbitrary pole figure positions (EWIMV)[27,28]. Separate ODFs of the a and b phases at varioustemperatures have been calculated. We do not show theODF but instead selected pole figures recalculated fromthe ODF on which the main textural features are clearlyvisible.

In the case of CP–Ti, the a/b textures are furtherdescribed using ideal texture components and the volumefraction evolution of each component during thea! b! a texture change is quantified. For this purpose,the a/b ODFs are decomposed into ideal spherical Gausscomponents in Euler space [29]. Each component is definedby five parameters: a, b, c (Euler angles in Matthies/Roeconvention), the Gauss width at half maximum and thevolume fraction (in % of multiples of a random distributionor m.r.d.). The MAUD program assigns to each ideal com-ponent a position and Gauss width (in this case fixed at18�). Finally, the volume fraction, position and width ofeach component are refined to obtain the best fit of the ori-ginal experimental spectra [30]. Standard deviations arebased on the least-squares refinement. For the refinementof component volumes, orthorhombic sample symmetrywas imposed. This reduces the number of independentcomponents. In the following, the set of symmetrical com-ponents is described by only one equivalent, assumingorthorhombic sample symmetry.

3. Results

The lattice parameters for both samples at different tem-peratures are determined from the diffraction pattern bythe Rietveld method (Table 1). Discrete thermal a- and

Fig. 4. Recalculated pole figures for a CP–Ti measured in situ at (a)400 �C, (b) 800 �C, (c) 900 �C and (d) 400 �C after b–a phase transfor-mation. Equal area projection and linear pole density scale. Rollingdirection (RD) and transverse direction (TD) are indicated.

Table 1Lattice parameter changes during heat treatment

Sample T (�C) Lattice parameters CTE (a · 10�6 K�1)

a a (A) c (A) b a (A) c/a a-axis c-axis

CP–Ti 400 2.9494(2) 4.6813(3) – 1.5872 10.4 14.2800 2.9601(1) 4.7058(2) – 1.5897850 2.9615(2) 4.7112(3) – 1.5908900 2.9637(1) 4.7146(3) 3.2952(1) 1.5908950 – – 3.2974(3)400 2.9489(1) 4.6819(2) – 1.5877

Ti–6Al–4V 400 2.9247(1) 4.6713(2) 3.2065(6) 1.5972 9.5 11.3800 2.9343(2) 4.6879(4) 3.2544(2) 1.5976850 2.9362(2) 4.6915(4) 3.2615(3) 1.5978900 2.9376(1) 4.6951(3) 3.2677(2) 1.5982950 2.9401(3) 4.7002(8) 3.2732(1) 1.5986400 2.9212(1) 4.6671(2) 3.2068(6) 1.5976

CTE, coefficient of thermal expansion. Note: There is some uncertainty in absolute values of lattice parameters determined with TOF. Estimated standarddeviations are in parentheses.

I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727 5721

c-axes strain for a-CP–Ti and a-Ti–6Al–4V were calculatedfrom changes in the lattice parameters as a function of thetemperature using:

eTa ¼ ðaT � a400Þ=a400 and eT

c ¼ ðcT � c400Þ=c400;

where aT (cT) and a400 (c400) are the lattice parameters ofthe a-Ti at temperature T and at 400 �C. Coefficients ofthermal expansion (CTEs) in Table 1 were determinedfrom the slope of a linear fit to the temperature vs. thermalstrain curve [31] between 400 and 900 �C for a-CP–Ti andbetween 400 and 950 �C for a-Ti–6Al–4V.

As reported in the literature [31], it is reasonable thatCTEs calculated using a- and c-lattice parameters arecompatible with the CTEs from (10�10) and (00 02) singlepeak values, respectively. The results clearly indicate thatthe difference in CTE is related to a texture effect of thepolycrystal.

Figs. 4 and 5 show the experimental a and b pole figuresfor CP–Ti, respectively, and Figs. 7 and 8 the pole figuresfor Ti–6Al–4V. All pole figures display more or less ortho-rhombic symmetry, consistent with the prior rolling geom-etry of the sheets. Table 2 summarizes statisticalinformation about the texture (ODF maxima/minima, polefigure maxima/minima and texture index). Table 3 lists theideal spherical components used to describe the CP–Ti tex-tures. Table 4 gives the refined volume fraction of eachcomponent, to obtain the best fit of the experimental tex-tures. Fig. 6 illustrates the volume fraction evolution ofeach CP–Ti texture component during the heat treatment.

3.1. CP–Ti sheet

For the CP–Ti alloy the a! b! a phase transforma-tion produces a strong orientation density for one texturecomponent that is weak in the cold rolling texture.

At 400 �C, the a texture is relatively sharp (ODFmax = 12.9 m.r.d.) and typical of a cold-rolled titaniumsheets (Fig. 4a) [32]. It is characterized by a main compo-nent with c-axes in a symmetrical double maximum

between normal and transverse direction (ND–TD) andwith f10�10g poles mainly in the rolling direction (RD).This symmetrical texture component can be described bythe ideal orientation ga

1 {90�, 35�, 0�} and covers over 40%of the orientation space (Table 4, Fig. 6). A second compo-nent, described by ga

2 {90�, 35�, 30�}, is much lower. It ischaracterized by similar c-axis positions but with f11�20gpoles in RD. Two additional components ga

3 and ga4 have

been introduced to represent more accurately the startinga texture. These additional components mainly contributeto the background.

At 800 �C the material is still completely in a phase (thetransition temperature being around 880 �C). However, a

Table 3Spherical components during the a–b–a phase transformation use to describe

Phase during thermal cycling Texture c

a before phase transformation ga1

ga2

ga3

ga4

b after phase transformation from ga2 using Burgers variants gb

1

gbII

gbIII

a after phase transformation from gb1 ; g

b2 using Burgers variants ga

A1

gaB1

gaC1

gaD1

gaE1

gaF 1

gaG1

gaH1

gaI1

gaA2

gaB2

gaC2

Euler angles (a, b, c) are expressed according to the Matthies/Roe convention

Table 2Experimental data and quantitative texture information

Sample T (�C) a phase (%) Texture index ODF max/min (m

CP–Ti 400 100 2.7 12.9/0.1800 100 4.6 21.4/0.0900 47 3.4 13.1/0.1

(1.5)bcc (6.4/0.3)bcc

950 0 3.8 16.3/0.2400 100 7.2 41.5/0.0

Ti–6Al–4V 400 95 1.3 5.2/0.3800 80 1.4 6.5/0.3

(1.3)bcc (3.8/0.5)bcc

900 57 1.2 4.4/0.4(1.4)bcc (7.8/0.3)bcc

950 23 1.2 4.6/0.4(2.2)bcc (9.8/0.2)bcc

400 89 2.1 8.1/0.3(1.1)bcc (2.4/0.7)bcc

Fig. 5. Recalculated pole figures for b CP–Ti measured in situ (a) at900 �C, where the sample is not completely transformed, and (b) at 950 �C.Equal area projection and linear scale. Rolling direction (RD) andtransverse direction (TD) are indicated.

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significant texture change is observed. It does not apprecia-bly affect the c-axes distribution, but reorients the a-axes,producing a f11�20g maximum in RD (Fig. 4b). Geometri-cally this corresponds to a 30� rotation of crystals aroundtheir c-axis. At the same time, the texture becomes sharper(ODF max = 21.4 m.r.d.). Component ga

1 decreases andcomponent ga

2 becomes stronger (>40%), with very smallchange in components ga

3 and ga4. Such a texture evolution

is typically observed during grain growth following pri-mary recrystallization (the texture after primary recrystalli-zation being not very different from the deformationtexture) [32,33]. A microstructure analysis was performedon a sample after a similar heat treatment up to 800 �C.A fully recrystallized microstructure was observed with

the texture evolution for sample CP–Ti

omponents Euler angles (�)

a b c

90 35 090 35 3090 30 1590 50 15

47.1 �62.9 �202.0 Close to gb1 ¼ ð0; 35; 135Þ

62.9 121.9 �240.1 Close to gb1 ¼ ð0; 35; 135Þ

6.0 94.2 �259.8 Close to gb2 ¼ ð0; 0; 0Þ

90 30 31.5 Close to ga2

31.5 73 57.60 55 52.7

90 90 87.731.5 73 12.231.5 107 12.231.5 107 57.690 30 16.990 90 17.790 45 22.7 Close to ga

2

45 90 52.70 45 22.7

. Orthorhombic sample symmetry is assumed.

.r.d.) (0001) ð10�10Þ ð11�20Þ (100) (110) (111)

(Max/min values of the pole figures in m.r.d.)

6.2/0.2 4.1/0.3 2.6/0.37.4/0.0 2.9/0.2 4.2/0.27.0/0.1 3.7/0.3 3.0/0.3 1.7/0.6 1.9/0.5 2.3/0.4

2.7/0.3 2.8/0.3 3.0/0.29.2/0.1 3.5/0.3 6.5/0.12.3/0.4 2.0/0.6 1.7/0.62.5/0.3 2.3/0.6 1.9/0.6 1.4/0.6 1.9/0.6 1.7/0.6

2.0/0.5 1.7/0.7 1.6/0.7 1.4/0.5 2.0/0.5 1.9/0.4

2.2/0.4 1.6/0.7 1.5/0.7 1.9/0.4 3.3/0.3 2.8/0.3

2.9/0.3 2.9/0.5 2.5/0.5 1.2/0.8 1.3/0.8 1.2/0.8

Fig. 7. Recalculated pole figures for a Ti–6Al–4V measured in situ at (a)400 �C, (b) 800 �C, (c) 900 �C and (d) 400 �C after b–a phase transfor-mation. Equal area projection and linear scale. Rolling direction (RD) andtransverse direction (TD) are indicated.

Fig. 6. Volume fractions for each a and b texture component duringa! b! a phase transformation in CP–Ti. For explanation of symbols,see Table 3.

Table 4Refined volume fractions (%) for the spherical components of CP–Ti described in Table 3 during heat treatment

T (�C) ga1 ga

2 ga3 ga

4 gb1 gb

2 gaA1 ga

A2 gaB2 ga

B1 gaC1 ga

C2 gaD1

400 46(4) 12(1) 30(3) 12(1)800 13(1) 42(4) 33(3) 12(1)900 15(2) 9(1) 23(3) 0 44(5) 9(1)950 66(7) 34(3)400 33(3) 22(3) 14(2) 10(1) 9(1) 5(.6) 4(.5)

Only components >4% are shown; estimated standard deviation in parentheses.

I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727 5723

an average grain size of 20 lm. It confirms a recrystalliza-tion and grain growth for the CP–Ti sheet, during heatingto 800 �C.

At 900 �C, 53% of the a volume is transformed into theb phase. Fig. 4c shows the residual a texture observed at900 �C. The f11�20g pole density in RD is significantlylower than before the initiation of the transformation. Thiscorresponds to a large decrease of a grains oriented close toga

2 (in fact, components ga3 and ga

1 are dominating in theresidual a phase, see Fig. 6).

The b texture at 900 �C is relatively sharp. It is charac-terized by the symmetrical component gb

1 {0�, 35�, 135�}with a {111} maximum in RD (Fig. 5a). A small volumefraction of b grains is oriented around a minor cube texturecomponent gb

2 {0�, 0�, 0�}.Upon heating to 950 �C, the material is fully trans-

formed to b. One observes the same texture componentsas at 900 �C, but with a strong increase of orientation den-sity (Fig. 5b). Moreover component gb

2 becomes aspronounced as the symmetrical component gb

1. Similarhigh-temperature b textures were observed for low-alloyedtitanium sheets [18]. In these studies, the b textures werenot measured in situ, but calculated from the inherited atextures. The comparison of a and b textures before andafter transformation (Figs. 4b and 5a and b) makes it clearthat the {110}b maxima are close to the (0001)a maxima,in agreement with the Burgers orientation relation.

After the reverse b! a phase transformation inducedby cooling, the microstructure observed at room tempera-ture indicates a colonies with a large spread in size. Theyare inherited from millimeter-sized b grains (Fig. 1a). Noa grain boundary precipitation is visible at the earlier bboundaries.

The a texture is surprisingly sharp with an ODF maxi-mum of 41.5 m.r.d (Table 2). The material is like two singlecrystals oriented close to ga

2, the a grain growth component(Fig. 4d). As seen in Fig. 6, the volume fraction of compo-nent ga

2 that increased during grain growth increases fur-ther during the transformation.

As a consequence, the b heat treatment strongly sharp-ens the symmetrical components ga

2. These componentsare favored during grain growth prior to the transforma-tion and during the transformation itself.

3.2. Ti–6Al–4V sheet

The evolution of the a and the high-temperature b tex-ture during the heat treatment is given in Figs. 7 and 8,

Fig. 8. Recalculated pole figures for b Ti–6Al–4V measured in situ at(a) 800 �C, (b) 900 �C, (c) 950 �C, and (d) 400 �C after thermal cycling.Equal area projection and linear scale. Rolling direction (RD) andtransverse direction (TD) are indicated.

5724 I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727

respectively. The textures are much weaker than for theCP–Ti sample (Table 2). No recrystallization of the amatrix is expected before transformation. In this sample,a and b phases are always present at all conditions, butwith different volume fractions (Table 2).

At 400 �C the c-axes describe a small circle girdle nearTD, with a weak distribution extending over ND.f10�10g poles concentrate in RD (Fig. 7a). The amountof b phase at this temperature (<5%) is too low for a reli-able estimate of the texture.

At 800 �C the texture of the hexagonal phase does notsignificantly change (Fig. 7b). The volume fraction of theb phase was estimated at 20%. The high-temperature bgrains exhibit specific orientations from the Æ110æ//RDand the Æ111æ//ND fibers (Fig. 8a).

At 900 �C the volume fraction of the b phase increasesto 43%. The b texture sharpens (Table 2), but without sig-nificant changes in the distribution pattern (Fig. 8b). Theresidual a phase has a similar pattern as observed at800 �C but with weaker densities (Fig. 7c).

At the end of the heating (950 �C), the amount of the bphase reaches 77%. Fig. 8c shows a strong b texture with aÆ110æ maximum in RD and a Æ111æ maximum in ND. Thistexture can be described by two main components: compo-nent (90�,55�,45�) with a Æ111æ maximum in ND and com-ponent (0�,0�,45�) with a Æ110æ maximum in RD. Theresidual a phase that reaches 23% at 950 �C exhibits a sim-ilar texture as at 900 �C (Fig. 7c).

After the b! a phase transformation, the volume frac-tion of the b phase at 400 �C decreased to 11%. The micro-structure after cooling is bimodal with equal volume

fractions of equiaxed a grains and grains with a lamellarstructures, composed of a and thinner b layers (Fig. 1b).

The residual b texture is characterized by the same tex-ture components as at 950 �C. It is, however, much weakerthan at 950 �C, with a significant decrease in densitiesaround Æ110æ//RD and Æ111æ//ND fibers.

The a texture after cooling (Fig. 7d) contains the com-ponents present before the b! a transformation, withsome increase in orientation densities. Moreover, new aorientations appear. They correspond to an a phase withc-axes in RD and f11�20g poles in ND. Again, after thea! b! a phase transformation, the main component ofthe initial a texture has been reinforced (see Table 2).

4. Discussion

The a! b! a transformation cycle implies a largemultiplication of a orientations if no variant selectionoccurs. Indeed, one starting a orientation can theoreticallygive rise to six b orientations, each of which can transforminto 12 a orientations. Such a transformation schemewould multiply the a orientation components 72-fold andthus reduce the a texture sharpness. However, this is notobserved in the present investigation. For both titaniumalloys, the heat treatment has significantly strengthenedthe main texture components of the starting a texture(observed just before transformation). This suggests astrong variant selection.

4.1. Analysis of the a! b texture changes

4.1.1. CP–Ti alloy: a! b texture changes from a

recrystallized a matrix

For the CP–Ti alloy, the a! b phase transition beginsabove 880 �C from a fully recrystallized a microstructure.Above 500 �C, recrystallisation and grain growth occur.During this process, ga

2 orientations grow at the expenseof ga

1. It is interesting to remark that component ga2 corre-

sponds to the orientation of the most highly twinned grainsand component ga

1 to the orientation of the less twinnedgrains in originally cold-rolled hexagonal metals [8,14,34].This suggests growth of nuclei in highly deformed hetero-geneous domains. The part of material, where twinningproduces local deformation and fragmentation, has a highdriving force in the recrystallization process [34].

Between 880 �C and 900 �C, the a! b phase transfor-mation reduces significantly the volume fraction of a grainsoriented close to ga

2 (Fig. 4b and c) and produces b grainsmainly oriented around components gb

1 and gb2 (Fig. 5a).

In order to explain these b orientations, all six b variantsthat can form from each of the two symmetrical orienta-tions ga

2 have been calculated, according to a strict Burgersrelation [15]. The study of 12 b variants can be reduced tothe study of three b variants according the orthorhombicsample symmetry. These three b variants (gb

I;II;IIIÞ are givenin Table 3 and the corresponding {100} and {110} polefigures in Fig. 9a (Roman numbers indicate that orienta-

Fig. 9. (a) {100}b and {110}b pole figures of b orientations calculatedfrom component ga

2 using Burgers relationships (see Table 3). (b, c) {100}b

and {110}b pole figures of the b texture simulated from component ga2 at

900�C and 950�C, respectively. The volume fraction of each b variant arerefined with the Rietveld method to fit best the experimental b texture ofCP–Ti (Table 4). Equal area projection and linear scale.

I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727 5725

tions are calculated and not observed.) The calculated bvariants gb

I;II are close to the observed component gb1 and

the calculated b variant gbIII is close to the minor component

gb2. This demonstrates that the observed b orientations can

effectively be produced from the transformation of only agrains oriented close to ga

2 (the a grain growth component).A similar mechanism was observed in zirconium [8,19].

The b texture observed at 900 �C has been simulatedfrom only the a grain growth component and with variantselection. Indeed, the volume fractions assigned to the bvariants are: gb

I 41%, gbII 41%, gb

III 18%. The simulated tex-ture (Fig. 9b) matches well with the experimental one(Fig. 5a).

This calculation suggests that a grains which preferen-tially grow at 800�C are those that preferentially transform.This transformation may be described as ‘‘orientation-selective a transformation’’. Moreover these a grains maytransform with variant selection. It can be deduced thatwhen the a! b transformation occurs from a fully recrys-tallized matrix, the high-temperature b texture is mainlycontrolled by the preferential transformation of the a graingrowth component.

From 900 to 950 �C, the b grains remain orientedaround components gb

1 and gb2. However, the relative vol-

ume fraction of each component evolves with the tempera-ture (Fig. 5a and b, Table 4). Indeed, component gb

2, whichis weak at 900�C, reaches the volume fraction of the sym-metrical component gb

1 at 950 �C. The b texture at 950 �Chas also been simulated from the a grain growth compo-nent. The best simulation is obtained with identical contri-bution of each b variant (gb

I , gbII, gb

IIIÞ in terms of volumefractions assigned (Fig. 9c).

The changes in b texture from 900 to 950 �C may berelated to the competitive growth between specific b grainorientations. Comparing the recrystallized a and the finalb grain size (20 lm for a; >1 mm for b), it is clear thatthe a! b phase transformation competes with b graingrowth. The b nuclei form preferentially, according to theBurgers relation, from a grains belonging to the graingrowth component. Then, these b nuclei mainly grow atthe expense of surrounding a grains. This is consistent within situ SEM observations that documented that the b phasenucleates and grows along the a grain boundaries as allot-riomorphs and within the a grain as platelets [13]. The highvolume fraction of b variants close to the gb

2 componentwas interpreted as the result of preferred nucleation andgrowth of b allotriomorphs.

The texture changes in CP–Ti during heating to 950 �Care summarized in Fig. 6. They document complex mecha-nisms during recrystallization and grain growth and thenduring phase transformation with a! b preferential trans-formation of the a grain growth component, formation ofb nuclei with variant selection, and finally orientation-selective growth of the b nuclei.

4.1.2. Ti–6Al–4V alloy: a! b texture changes in presence of

residual b phase

For the Ti–6Al–4V alloy, the a! b phase transforma-tion occurs in presence of a residual b phase. The b textureat high temperature is again relatively strong. The compar-ison between {000 1}a/{1 10}b and f11�20ga=f11 1gb polefigures at 900 �C (Figs. 7c and 8b) shows that the highestpole densities do not fit and thus the Burgers relationshipdoes not apply.

In fact, in the presence of residual b phase, thenucleation of new b grains from the a matrix accordingto Burgers orientation relations is thermodynamicallyunfavorable. It can reasonably be assumed that the high-temperature b phase is mainly obtained by the growth ofpre-existing b nuclei. This was confirmed for Ti–6Al–4Valloys in Refs. [9,16]. Consequently, the high-temperatureb texture observed at 800 �C probably resembles the b tex-ture in the starting material which could unfortunately notbe determined (the volume fraction of b phase being toolow). The b texture evolution from 800 to 950 �C tendsto favor specific b grain orientations belonging to the initialfiber-like texture. This can be linked to a preferentialgrowth of b grains according to their orientations. Thus,in the presence of residual b phase (even in low volumefraction), the high-temperature b texture is controlled bythe orientations of pre-existing b nuclei and by preferentialgrowth of specific b grain orientations.

4.2. Analysis of the b! a texture changes

4.2.1. CP–Ti alloy: b! a texture changes from a textured bmatrix

At high temperature, the b orientations of CP–Ti sampleare centered around two main texture components: gb

1 and

5726 I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727

gb2 (Fig. 5). These texture components correspond to the

Burgers variants inherited from initial a grains orientedclose to ga

2 (Fig. 9a). After cooling, the inherited a textureof the CP–Ti alloy becomes extremely sharp, with strongreinforcement of density around the grain growth compo-nents ga

2 (Fig. 4d).The comparison between {0001}a/{1 10}b and

f11�20ga=f1 11gb pole figures (Figs. 4d and 9a) shows thatthe b! a texture change is consistent with the Burgers ori-entation relation. However, without b! a variant selec-tion, the {0001}a pole density distribution should besimilar to {110}b. This is not the case. {0001}a pole den-sities around ga

2 are much higher than expected from thecorresponding {110}b pole distribution. It reveals ab! a variant selection mechanism during cooling, favor-ing a variants oriented close to the initial grain growthcomponents. This is also confirmed by the sample micro-structure showing significant differences in size, for coloniesinherited from the same parent (Fig. 1a).

To analyze this variant selection with components, wehave calculated the b! a Burgers variants obtained fromeach of the main high-temperature b orientations gb

1 andgb

1. For each b orientation, there are 12 a variants; how-ever, superpositions again occur because of orthorhombicsample symmetry. For the symmetrical component gb

1 thereare nine distinct variants and for the cube component gb

2

there are only three distinct variants (Table 3). Fig. 10shows the {00 01}a,f10�10ga and f11�20ga pole figurescomputed with all variants equally weighted (Fig. 10a)and after the Rietveld component refinement (Fig. 10b).For the computed texture shown in Fig. 10b, the volumefraction of each a variant was refined with a fixed Gausswidth of 18�, so that the texture fits the best the experimen-tal one (Fig. 4d). The volume fractions are reported inTable 4 and Fig. 6 for those components that exceed 4%.Because of strong variant selection during cooling, morethan 50% of the volume is oriented around ga

2 (close toga

A1 and gaA2Þ. Additional EBSD analysis has been per-

formed on the sample to study this variant selection on a

Fig. 10. Simulated CP–Ti {0001}a, f10�10ga and f11�20ga pole figuresafter heat treatment obtained using ideal spherical components andorthorhombic symmetry. (a) Using all a texture components inheritedfrom gb

1 and gb2 with the same volume fraction. (b) After Rietveld analysis.

The volume fractions for each texture component are shown in Table 4.Equal area projection and linear scale.

local scale. This analysis is not shown here but is similarto that performed in Ref. [18]. The acquired orientationmaps confirmed that adjacent a variants belonging to dif-ferent parent b grains often exhibit low misorientations.The same a variant can even grow over two b grains whenit is strictly common to both grains. In this case, the priorb/b boundaries are no longer visible, as often observed onthe microstructure in Fig. 1a. This confirms that the variantselection during cooling favors the a variants being com-mon to adjacent b grains.

The texture memory effect observed in CP–Ti after aa! b! a transformation cycle is a direct consequenceof orientation-selective a transformation during the a! btransformation followed by a strong b! a variant selec-tion favoring common variants between adjacent b grains.

4.2.2. Ti–6Al–4V alloy: b! a texture changes in presence of

primary ap phase

At 950 �C, the Ti–6Al–4V alloy contains 23% of a phase(volume fraction). The a phase present before cooling is clas-sically called primary a phase and written ap. The ap orien-tations are relatively similar to those observed before heattreatment (Fig. 7a and c). In particular, the c-axes describea small circle girdle near TD with a weak distributionextending over ND. f10�10g poles concentrate in RD. Aftercooling, the main a orientations are again very similar tothose observed in the starting material (Fig. 7a and d). Apronounced increase in density around the main texturecomponents is observed. However, new components alsoappear, leading to an increased (0001) density in RD andreinforcing the f11�20g density in ND. In the microstructureat room temperature (Fig. 1b), the primary ap phase corre-sponds to the equiaxed grains, whereas the lamellar struc-ture (called secondary as phase) results from the classicalb! a transformation according to the Burgers relationship.

The as texture can be estimated by subtracting the ap tex-ture measured at 900 �C (Fig. 7c) from the a texture mea-sured at 400 �C after cooling (Fig. 7d), taking into accountthe volume fraction of each phase (Fig. 11). The differencetexture shows pronounced concentration of c-axes close toRD and f11�20g poles in ND. This component correspondsto the new a components created during the heat treatment.No significant b! as variant selection is detected.

In the micrograph Fig. 1b, the volume fraction of ap

phase is significantly higher than the volume fraction

Fig. 11. as texture in Ti–6Al–4 V after cooling, estimated by subtractingthe a texture at 900 �C (Fig. 7c) from the a texture after cooling to 400 �C(Fig. 7d).

I. Lonardelli et al. / Acta Materialia 55 (2007) 5718–5727 5727

observed at 950 �C (23%). As a consequence the ap grainshave grown during cooling. Coarsening of primary ap

phase was studied under static and dynamic conditions[35,36]. For Ti–6Al–4V, the ap coarsening kinetics are con-trolled by the diffusion of vanadium solutes through the bmatrix. Thus, in contrast to pure Ti alloy, the alloying ele-ments in Ti–6Al–4V control the bulk diffusion and influ-ence the transformation mechanism. In the present study,the ap coarsening explains the similarities between the atexture before and after heat treatment in the a/b field.

For the Ti–6Al–4V alloy, the texture memory effect afterthe heat treatment in the a/b field is mainly related to thegrowth of the ap phase present at high temperature.

5. Conclusions

The texture changes during the a! b! a phase trans-formation are investigated in CP–Ti cold-rolled andTi–6Al–4V hot-rolled sheets with in situ high-temperaturetexture measurements by neutron diffraction.

In single-phase a Ti, the stored energy due to cold work-ing promotes recrystallization and grain growth before thetransformation. This process increases the fraction of agrains belonging to the a grain growth component, withc-axes in symmetrical maxima between normal and trans-verse directions and f11�2 0g poles in the rolling direction.At the transformation temperature, the a grains belongingto this a grain growth component are those that preferen-tially transform, controlling at the same time the high-temperature b texture. The difference in the b texture at900 and at 950 �C could be related to some local a! bvariant selections and competitive growth between bgrains. During the b! a transformation, a strong selectionof a variants occurs, favoring common variants betweenadjacent b grains, probably to minimize the grain bound-ary energy. The combination of all these mechanisms isat the origin of the strong texture memory effect in cold-rolled a Ti during phase transformations.

In two-phase a/b Ti–6Al–4V, the nucleation of new bgrains from the a matrix according to Burgers orientationrelations is thermodynamically less favorable than thegrain growth of residual b phase orientations in the startingmaterial. Consequently, the high-temperature b phase ismainly obtained by the growth of pre-existing b nuclei.In the same way, if a primary a phase is present at hightemperature, it influences the transformation mechanismupon cooling. In particular, it favors the growth of a grainsin addition to the creation of new b! a orientationsaccording to the Burgers orientation relations (especiallyfor low cooling rates). The texture memory effect afterthe heat treatment in the a/b field is then mainly relatedto the growth of the ap phase present at high temperature.

Acknowledgments

Support of this work through NSF (EAR 0337006),CDAC-DOE, UCRPI and IGPP-LANL is gratefully

acknowledged. This work has benefited from the use ofthe Lujan Neutron Scattering Center at LANSCE, whichis funded by the Department of Energy’s Office of BasicEnergy Sciences. Los Alamos National Laboratory is oper-ated by Los Alamos National Security LLC under DOEcontract DE-AC52-06NA25396.

References

[1] Weaver ML, Garmestani H. Mater Sci Eng A 1998;247:229.[2] Bache MR, Evans WJ. Mater Sci Eng A 2001;319:409.[3] Kalinyuk AN, Trigub NP, Zamkov VN, Ivasishin OM, Markovsky

PE, Teliovich RV, et al. Mater Sci Eng A 2003;346:178.[4] Zaefferer S. Mater Sci Eng A 2003;344:20.[5] Semiatin SL, Fagin PN, Glavicic MG, Sukonnik IM, Ivasishin OM.

Mater Sci Eng A 2001;299:225.[6] Zeng L, Bieler TR. Mater Sci Eng A 2005;392:403.[7] Jourdan C, Gastaldi P, Marzo P, Grange G. J Mater Sci

1991;26:4355.[8] Wenk H-R, Lonardelli I, Williams D. Acta Mater 2004;52:1899.[9] Bhattacharyya D, Viswanathan GB, Vogel SC, Williams DJ, Venk-

atesh V, Fraser HL. Scr Mater 2006;54:231.[10] Ciurchea D, Pop AV, Gheorghiu C, Furtuna I, Todica M, Dinu A,

et al. J Nucl Mater 1996;231:83.[11] Zhu ZS, Gu JL, Liu RY, Chen NP, Yan MG. Mater Sci Eng A

2000;280:199.[12] Seward GGE, Prior DJ, Wheeler J, Celotto S, Halliday DJM, Paden

RS, et al. Scanning 2002;24:232.[13] Seward GGE, Celotto S, Prior DJ, Wheeler J, Pound RC. Acta Mater

2004;52:821.[14] Puig-Molina A, Wenk H-R, Berberich F, Graafsma H. Z Metallk

2003;94:1199.[15] Gey N, Humbert M. J Mater Sci 2003;38:1289.[16] Gey N, Humbert M, Moustahfid H. Scr Mater 2000;42:525.[17] Humbert M, Gey N, Esling C. J Appl Crystallogr 2000;33:206.[18] Gey N, Humbert M. Acta Mater 2002;50:277.[19] Gey N, Gautier E, Humbert M, Cerqueira A, Bechade JL, Archa-

mbault P. J Nucl Mater 2002;302:175.[20] Burgers WG. Physica 1934;1:561.[21] Wenk H-R, Lutterotti L, Vogel SC. Nucl Instr Meth A 2003;515:575.[22] Vogel SC, Lutterotti L, Von Dreele RB, Wenk H-R, Williams DJ.

Powder Diffr 2004;19:65.[23] Rietveld HM. Acta Cryst 1967;22:151.[24] Rietveld HM. J Appl Cryst 1969;2:65.[25] Lutterotti L, Matthies S, Wenk H-R, Schultz AS, Richardson Jr JW.

J Appl Phys 1997;81:594.[26] Matthies S, Vinel GW. Phys Stat Sol (b) 1982;112:K111.[27] Lutterotti L, Chateigner D, Ferrari S, Ricote J. Thin Solid Films

2004;450:34.[28] Matthies S, Pehl J, Wenk H-R, Lutterotti L, Vogel SC. J Appl Cryst

2005;38:462.[29] Matthies S, Vinel GW, Helming K. Standard distributions in texture

analysis, vol. I. Berlin: Springer; 1987.[30] Lutterotti L, Bortolotti M. Acta Cryst A 2005;61:C158.[31] Choo H, Rangaswamy P, Bourke MAM, Larsen JM. Mater Sci Eng

A 2002;325:236.[32] Bozzolo N, Dewobroto N, Grosdidier T, Wagner F. Mater Sci Eng A

2005;397:346.[33] Wagner F, Bozzolo N, Van Landuyt O, Grosdidier T. Acta Mater

2002;50:1245.[34] Bozzolo N, Dewobroto N, Wenk H-R, Wagner F. J Mater Sci

2007;42(7):2405.[35] Semiatin SL, Kirby BC, Salishchev GA. Metall Mater Trans A

2004;35:2809.[36] Semiatin SL, Corbett MW, Fagin PN, Salishchev GA, Lee CS. Metall

Mater Trans A 2006;37:1125.