texture evolution during grain growth in recrystallized commercially pure titanium

Materials Science and Engineering A 397 (2005) 346–355

Texture evolution during grain growth in recrystallizedcommercially pure titanium

N. Bozzolo∗, N. Dewobroto, T. Grosdidier, F. WagnerLaboratoire d’Etude des Textures et Application aux Mat´eriaux (LETAM-UMR CNRS 7078),

Universite de Metz, Ile du Saulcy, F-57045 Metz-Cedex, France

Received 19 July 2004; received in revised form 17 February 2005; accepted 22 February 2005

Abstract

The evolutions of microstructure and crystallographic texture in low-alloyed titanium sheets are investigated by electron backscatteringdiffraction at different grain growth stages. Recrystallization of 80% cold-rolled sheets and subsequent grain growth lead to equiaxed mi-crostructures. The texture obtained at the end of primary recrystallization is very close to that of the cold-rolled state, with the maximumvalue of the orientation distribution function at{0◦, 35◦, 0◦}. The orientations developing during grain growth correspond to a broad peakcentered around{0◦, 35◦, 30◦}which is a minor component in the initial texture. The disappearing orientations are widely scattered throughouto he growingt t the end oft oportion ofm©

K

1

cnpictnatt

pIa

tal-lly

d toori-yhisnta-

uchtionf the

atephic

ed asme-

edate-, the

ud-

0d

rientation space and present two major disadvantages in the growth competition: (i) they are highly misoriented with respect to texture component and (ii) the grains having these orientations belong to the smallest size range in the microstructure obtained ahe primary recrystallization. The grain boundaries remaining after extended grain growth are characterized by an increasing prisorientations below 30◦ and random rotation axes.2005 Elsevier B.V. All rights reserved.

eywords:Grain growth; Texture; Orientation distribution function; Electron backscattering diffraction; Titanium (�-phase)

. Introduction

A complete understanding of the physical and metallurgi-al phenomena occurring in a metal all over its processing isecessary to control its final properties. These properties de-end both on the microstructure and on the texture developed

n the material under processing. The mechanisms of mi-rostructure and texture evolution during thermomechanicalreatment (including cold and hot deformation, static and dy-amic recrystallization, phase transformations, etc.) of BCCnd FCC metals and alloys have been intensively studied in

he recent years[1,2] but much less attention has been paido HCP ones.

One of the main industrial applications of commerciallyure (cp-)titanium is the manufacture of heat exchangers.

n this field, the ductility is of the uppermost importancend known to be determined by the annealing of the cold-

∗ Corresponding author. Tel.: +33 387 31 57 19; fax: +33 387 31 53 77.E-mail address:[email protected] (N. Bozzolo).

rolled strips[3]. The industrial parameters for the recryslization annealing of a cold-rolled cp-Ti strip are typica2 h at 700◦C. The resulting texture, sometimes referreas “recrystallization texture”, is characterized by a mainentation identified as{1 0 1 3} 〈1 21 0〉 or characterized b{ϕ1 = 0◦, Φ = 35◦, ϕ2 = 30◦} angles in the Euler space. Torientation can be roughly deduced from the main orietion of the cold-rolling texture ({2 1 1 5} 〈01 1 0〉) by a 30◦rotation about the�c-axis. It has been shown that under sannealing conditions, where both primary recrystallizaand grain coarsening actually occur, the modification otexture is essentially controlled by grain growth[4,5]. Thus,the terminology of “recrystallization texture” is inappropriand should not be further used. Although the crystallogratextures of this type of Ti alloy have already been describa function of different thermomechanical processing paraters[4,6–12], limited information is available on the detailtexture evolution during annealing of cold-deformed mrial and on the related mechanisms. In the present worktexture evolution during grain growth of cp-titanium is st

921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2005.02.049

N. Bozzolo et al. / Materials Science and Engineering A 397 (2005) 346–355 347

ied. The main investigation technique employed was electronbackscattering diffraction (EBSD). It allows to correlate thetexture and microstructure evolutions and to discuss the rea-sons for texture evolution in connection with microstructuralparameters like the grain size or the nature of the grain bound-aries.

2. Experimental procedures

The material investigated is a cp-titanium (ASTM grade2) having the chemical composition shown inTable 1. Thesheets (initially 3 mm thick) were cold-rolled by reductionsteps of 3%, down to a final thickness reduction ratio of 80%.They were subsequently heat-treated at 540◦C for 24 min tocomplete the primary recrystallization process. This recrys-tallized material is the initial state for the present grain growthstudy. Its microstructure consists of equiaxed grains havinga mean grain size (MGS) of 3.3�m. Grain growth heat treat-ments were performed under argon atmosphere at 600, 700and 800◦C (respectively, 0.45, 0.50 and 0.55T/TM ratios,whereTM is the melting temperature), for 5, 10, 15, 30, 60and 100 min. The heat treatments were terminated by waterquenching.

Microstructures and textures were investigated at half-w Them ea oft aps.E dis-t

cor-r ucha tiona idt r thec F).D ition,d nd inR RDt romE y theϕ them scc[

TCT

CFHNOS

3. Results

3.1. Microstructure evolution

No noticeable abnormal grain growth occurred duringthe grain growth process. The microstructures were alwaysequiaxed with single-modal grain size distributions. It shouldnevertheless be noted that some of the microstructures ob-tained after recrystallization and in the early stages of graingrowth showed locally an inhomogeneous spatial distributionof the grain sizes.Fig. 1a shows an example of such a situa-tion (top part of the image) where an area contains grains thatare larger than the grains in the rest of the material. The largegrains did not correspond to any particular orientation. Theseheterogeneities are produced by the primary recrystallizationprocess and are probably due to different recrystallizationmechanisms and kinetics according to local heterogeneitiesof the deformed material. They disappeared with grain coars-ening, probably because the driving force for grain growthand, consequently, the grain growth rate is higher in the smallgrained areas. These local heterogeneities do not affect thesingle-modal nature of the global grain size distribution his-togram (Fig. 1e).

In Fig. 1a, the grains are colored in light or dark grey ac-cording to the location of theirc-axis (seeFig. 1b definingthis color code) in order to show the spatial distribution oft theo aref canc ents.Ta -e igh-b rs ofp ribu-t ion.B y ofl iont istri-b ingr

al-i raing per-a , ands driv-i arys

edf

D

Tt e

idth of the sheets in order to avoid surface artifacts.ean grain size was determined from the surface ar

he grains obtained on optical micrographs or EBSD mquivalent circle diameters were used to build grain size

ribution histograms.EBSD was used for texture measurements in order to

elate texture evolution with microstructural features ss grain size and morphology, and also to get informabout grain boundaries[13,14]. Special attention was pa

o the statistical relevance of the EBSD data used foalculation of the orientation distribution functions (ODetails about the applied procedures for data acquisata post-processing and ODF calculation can be fouef. [15]. Global textures were also measured by X

o check the statistical validity of the ones deduced fBSD data. Concerning the texture representation, onl1 = 0◦ ODF section is given here, since it containsajor texture components. The{ϕ1, Φ, ϕ2} Euler angle

orrespond to the definition given by Bunge[16] and therystal coordinate system is{X = [1 0.0], Y = [1 2.0], Z =0 0.1]}.

able 1hemical composition of the studied cp-Ti alloy (weight (ppm)− balancei)

52e 237

341

1062i <100

he two components in symmetric position with regards torthorhombic sample symmetry. The two types of grains

airly well mixed in the microstructure, even if some areasontain more grains belonging to one of the two componhe low and high angle misorientation boundaries (Fig. 1c)re also well distributed spatially.Fig. 1d shows two misorintation angle distributions: the one computed from neoring data point pairs and the one obtained from paiixels randomly chosen in the map. The latter is the dist

ion resulting from the texture without any spatial correlatoth distributions are very similar except that the densit

ow misorientations is slightly higher in the real distributhan in the one resulting from the texture. The spatial dution of the orientations is therefore not so far from beandom.

Fig. 2 shows the evolution of the grain size with anneng time for the different temperatures. As expected, growth is faster at higher temperatures. At a given temture, the grain size increases rapidly at the beginningubsequently more slowly because of the weakening inng force associated with the reduction in grain boundurface.

The data ofFig. 2 were analyzed using the generalizorm of the Burke and Turnbull relation(1) [17]:

n − Dn0 =

[A0 exp

(−Q

RT

)]t (1)

he activation energy for grain boundary migration (Q) ob-ained from this formula was 204 kJ mol−1. This value is clos

348 N. Bozzolo et al. / Materials Science and Engineering A 397 (2005) 346–355

Fig. 1. (a) EBSD map showing the microstructure of a sample heated for 5 min at 600◦C (MGS = 5.1�m). The black lines correspond to misorientations higherthan 3◦. The two grey levels allow to distinguish the grains belonging to both sides of the{0 0.2} pole figure (b). (c) Spatial distribution of the low and highangle misorientations: 5–30◦ in white and >30◦ in black. (d) Misorientation angle distribution histograms built either with the misorientation data calculatedbetween neighboring pixels (black) or with the ones calculated between pixels randomly chosen in the map. (e) Grain size distribution histogram of this sample.

to the activation enthalpy for grain boundary self-diffusion inhigh-purity�-Ti [18]. This seems to indicate that the impuri-ties (which may be present in the form of solute atoms or inthe form of few small precipitates) do not have a significanteffect on grain boundary motion kinetics in cp-Ti. This is con-sistent with the absence of abnormal grain growth, a process,which is often caused by some amount of grain boundarypinning.

3.2. Global texture evolution

3.2.1. General description of the texture evolutionFig. 3a–e shows the evolution of the crystallographic tex-

ture during the grain growth process. As already mentioned,

the starting state of the grain growth stage obtained after pri-mary recrystallization completion (Fig. 3a) is, in terms of tex-ture, very close to the deformed state[5]. The major volumefraction of the material is indeed distributed around{ϕ1 = 0◦,Φ = 35◦, ϕ2 = 0◦}. During grain growth, the texture evolvesslowly to reach, after extensive annealing, the state shownin Fig. 3e. This final texture is characterized by orientationslocated around{ϕ1 = 0◦, Φ = 35◦, ϕ2 = 30◦} with still a largespread around this ideal component. The intermediate stagesof the grain growth process (e.g.Fig. 3b–d) are characterizedby a mixture of different amounts of the{0◦, 35◦, 0◦} and the{0◦, 35◦, 30◦} components with a wide spread around them.

The evolution of the texture indexJ during grain growthis shown inFig. 4. J is defined by Eq.(2):


Fig. 2. Evolution of the grain size as a function of annealing time for differenttemperatures (the initial grain size obtained after primary recrystallizationwas 3.3�m).

J =∮ [

f (g)]2 dg (2)

wheref(g) is the orientation density function (ODF) andgthe orientation.

It increases significantly at the beginning of the graingrowth process. Starting with an initial value of 3.5, it reachesrapidly the value of 6.0 for a mean grain size (MGS) of 30�m.During further grain growth, the texture index keeps increas-

Fig. 4. Evolution of the texture index as a function of the MGS.

ing but much more slowly and tends to stabilize at about6.5. It can thus be concluded that even after extensive an-nealing, the annealing textures in titanium are never verysharp.

3.2.2. Developing and disappearing orientationsIn order to get more information about texture changes

(identification of the developing and the disappearing com-ponents; quantification in terms of volume fractions), texture

Fe

ig. 3. ϕ1 = 0◦ ODF sections of samples with increasing mean grain size (a–ach ODF cut) and relatedϕ1 = 0◦ sections of ODF differences (f–h).

e; the values of the lowest and the highest density contour lines are indicated on


differences were computed. The ODF difference, previouslyused in Ref.[5], is a function defined as

�f (g) = f2(g) − f1(g) (3)

wheref1(g) is the ODF of the reference state andf2(g) theODF after subsequent grain growth.

Positive values correspond to orientations for which thevolume fraction has increased while negative ones are asso-ciated to decreasing volume fractions.

Fig. 3f is the difference of the ODF corresponding to aMGS of 8.3�m and the initial one (MGS = 3.3�m). It clearlyevidences that the texture change corresponds mainly to theincrease of the{0◦, 35◦, 30◦} component, while the vol-ume fraction oriented around the initially major{0◦, 35◦, 0◦}component does not change significantly. The disappearingorientations are located atϕ2 close to 0◦ (≡ 60◦) andΦ be-low 20◦ or above 40◦. ODF cuts at otherϕ1 values showedglobally the same trends forϕ1 values up to 35◦ while ODFcuts atϕ1 higher than 40◦ exhibited only disappearing orien-tations.

Fig. 3g compares the ODFs ofFig. 3d to the one ofFig. 3c,which, respectively, correspond to MGS of 32.8 and 8.3�m.It indicates that the{0◦, 35◦, 30◦} component continues togrow. Interestingly, it must be noticed that the volume fractionoriented around the initially major{0◦, 35◦, 0◦} componenti f theod tagetc 3ar

GSit ch isc aluesft twog -w ar-iϕ

eDFs

a art oft ph.

3.2.3. Volume fractions associated to the orientationchanges

During grain growth, an orientationg0 can grow at the ex-pense of another oneg′

0 while, somewhere else in the sample,a reciprocal mechanism (g′

0 consumingg0) can take place.Such a type of mechanism cannot be revealed by the ODFdifferences. Indeed, the volume fractions ofg0 andg′

0 wouldnot change if this mechanism was to take place. The bal-ance between such a “reciprocal” mechanism and situationswhere the disappearance of an orientation is not counterbal-anced (and therefore visible in the ODF difference) can beestimated using the ODF difference functions as follows.

During a grain growth step from a MGSD1 (correspond-ing to a mean grain volumeV1) towards a larger MGSD2(corresponding toV2), the total volume fraction�V throughwhich the grain boundaries have moved can be estimated by

�V = V2 − V1

V2(4)

The mean grain volumesV1 andV2 were calculated from thegrain volume histogram corresponding to each consideredstate. These histograms were built from the 3D grain size dis-tributions, deduced from the 2D histograms by the Saltykovmethod[19]. The volume fraction�V does not distinguishbetween the two types of mechanisms, counterbalanced orn

n-c e cal-c raloe

V

T n-tV

t( o6 att if-fT lwaysp

TE ounter wth

G tion cochange(non-c)Vdif (%

FFF

s also increasing, but to a lower extent. As the regions orientation space atΦ below 20◦ or above 65◦, which wereecreasing before, have been totally cleared out at this s

he disappearing orientations are now located atϕ1 andϕ2lose to 0◦ andΦ between 40◦ and 65◦. Between MGS = 8.nd 32.8�m, ODF cuts atϕ1 higher than 40◦ still exhibitedesidual disappearing orientations.

Fig. 3h shows the ODF difference associated to a Mncrease from 32.8�m (Fig. 3d) to 64.0�m (Fig. 3e). Theexture changes are much weaker than previously, whionsistent with the fact that the grain size reaches vor which the texture index stabilizes (Fig. 4). It is here in-eresting to note that the trend is inverted between therowing components, with{0◦, 35◦, 0◦} increasing somehat more than{0◦, 35◦, 30◦}. At this stage, the disappe

ng orientations were located around{ϕ1 ≈ 30◦, Φ ≈ 40◦ and2 ≈ 0◦}while the ODF cuts atϕ1 larger than 45◦ were nearlympty.

These orientation changes which are visible in the Ond better revealed by ODF differences, concern only a p

he material. This will be detailed in the following paragra

able 2stimation of the importance of the counterbalanced and of the non-c

rain growth step Total volume fraction throughwhich the grain boundaries havemoved�V (%)

Volume fracorientationdifferencesmechanism

rom 3.3 to 8.3�m 94 38rom 8.3 to 32.8�m 98 35rom 32.8 to 64�m 93 21

;

ot.The volume fractionVdif corresponding to the no

ounterbalanced mechanism (with ODF change) can bulated from the ODF difference function. It is the integf �f(g), restricted to the areas where�f(g) is positive, andxpresses as follows:

dif =∮

�f (g)>0

�f (g) dg = 1

2

∮|�f (g)| dg (5)

he volume fractionVC which is concerned with the couerbalanced mechanism is the difference between�V anddif .

Calculations have been carried out for�V, Vdif andVC athe three different stages of grain growth considered inFig. 3i.e. from 3.3 to 8.3�m, from 8.3 to 32.8�m and from 32.8 t4.0�m). The values are given inTable 2. It can be seen th

he volume fraction�V remains in the same range for the derent stages. Comparatively,VC is always higher thanVdif .his means that the counterbalanced mechanism is aredominant. The difference betweenVdif andVC is never-

balanced orientation change mechanisms occurring during grain gro

rresponding to thes visible in ODFounterbalanced)

Volume fraction submitted to thecounterbalanced mechanism (g0 consumingg′

0 while g′0 consumesg0 somewhere else)

VC (%)

566372


theless increasing drastically. Indeed,VC clearly increases asthe grain growth process progresses while the volume fractionVdif concerned by the orientation changes decreases. It canthus be concluded that the counterbalanced mechanism playsa more and more important role as grain growth progresses.

3.3. Grain size-based partial textures

In order to investigate the correlation between size andorientation of the grains, ODFs were computed from thesmallest and from the largest grains. Each subset of grainscorresponded to 15% in surface fraction. Before going into adetailed analysis of these results, three points should be notedconcerning uncertainties related to this procedure. Firstly, agrain of small size in the metallographic section may alsocorrespond to a large grain cut far away from its center.Thus the partial ODF of the small grains will always containsome amount of the texture components of the large grains.It should therefore not be interpreted alone but in compar-ison to the one of the large grains. Secondly, care must betaken also concerning data statistics for ODF calculation. In-deed, for identical surface fractions, the large grain subsetsinclude much less grains than the small grain ones. It is there-fore necessary to start with large area EBSD maps to get arepresentative population of the large grains for partial ODFcalculation. We used data sets including about 5000 grains int 150l kepti lateda dif-f at ag plesi l bes artialO stingi

gestg gf ft rgest( r-e lest (seeF sert lt withϕ tbo g int

eb rainsi ilarto s toe

Fig. 5. (a–c)ϕ1 = 0◦ ODF sections of the initial state (MGS = 3.3�m).(d–f) ϕ1 = 0◦ ODF sections of a sample annealed for 100 min at 600◦C(MGS = 10.4�m). Sections a and d were obtained by computing the ODFusing all the measured grains, b and e by selecting only the largest grainsand c and f only the smallest grains, corresponding, in both cases, to a 15%surface fraction of the original EBSD map. The values of the lowest and thehighest density contour lines are indicated on each ODF cut.

3.4. Grain boundary characterization

When dealing with grain growth, characterizing grainboundaries is of the outmost importance because the driv-ing force for grain growth is essentially the decrease in thefree enthalpy associated with the grain boundaries. With con-ventional EBSD, only the misorientation across the bound-aries but not the crystallographic position of the planes can bedetermined. However, it is assumed that the misorientationangle is of some significance for the energy and the mobilityof a grain boundary. In order to check if there were any partic-ular grain boundaries, which could have a specific behaviorduring grain growth due to a particularly low energy or highmobility, the EBSD maps were analyzed also to detect CSLboundaries, as defined for Ti and other hexagonal crystals byBonnet et al.[20]. The results were that the occurrence ofCSL boundaries was very low, not higher than statisticallyforeseeable.

Fig. 6a shows the evolution of misorientation angle distri-bution histograms during the grain growth process.Fig. 6band d are inverse pole figures showing the rotation axes corre-sponding to different ranges of misorientation angle.Fig. 6cand e are pole figures indicating the location of these rota-tion axes within the macroscopic sample coordinate system.Misorientation angles lower than 3◦ were cut off in orderto avoid artifacts in local misorientation data, which may bed pix-e ibu-t nd2 e-c first

otal, this led to the selection of 2500 small grains andarge ones for 15% surface fraction. Finally, it should ben mind that the grain size-based partial textures calcut different grain growth stages probably correspond to

erent populations of grains. Indeed, the smallest grainsiven stage will probably have disappeared in the sam

nvestigated after some additional grain growth. As wilhown, if sufficient care is taken, the grain size-based pDFs can nevertheless provide meaningful and intere

nformation.Fig. 5shows the partial texture of the smallest and lar

rains at the initial stage (MGS = 3.3�m) and after annealinor 100 min at 600◦C (MGS = 10.4�m). In the beginning ohe grain growth process, the partial textures of the laFig. 5b) and the smallest (Fig. 5c) grains are very diffent. Theϕ1 = 0◦ ODF section of the largest grains resemb

he global texture obtained after extended grain growthig. 3) while the partial texture of the smallest grains is clo

o the initial global texture (Fig. 3a or Fig. 5a). The partiaexture of the small grains exhibits preferred orientations2 close to 0◦. The minor texture components located aΦ

elow 20◦ and above 60◦ in the global ODF ofFig. 5a arenly occupied by small grains because they are missin

he partial ODF of the large grains.After further grain growth (Fig. 5e and f), the differenc

etween the partial ODFs of the largest and the smallest gs less pronounced. Both grain fractions evolve very simo the evolution of the global texture shown inFig. 3with thenly difference that the texture of the largest grains seemvolve a little faster.

ue to the experimental fluctuations between adjacentls within single grains. The misorientation angle distr

ion histograms (Fig. 6a) consist of two broad peaks, arou0–30◦ and around 60–70◦. These two contributions are boming more visible when grain growth progresses. The


Fig. 6. Grain boundary characterization at the initial stage (b, c) and after grain growth (d, e). (a) Evolution of the misorientation angle distribution histogramas a function of the mean grain size (MGS). (b–e) Rotation axes corresponding to different ranges of misorientation angle, represented in the crystalcoordinatesystem (b, d) and in the sample coordinate system (c, e).


peak shifts to low misorientations while grain growth pro-gresses, which means that the fraction of misorientations be-low 30◦ increases.

Low misorientations (<15◦) around�c axis were observedinside grains at the beginning of the process (Fig. 6b), anddisappeared upon grain growth (Fig. 6d). They probably cor-respond to slightly misoriented subgrains resulting from therecrystallization mechanisms. These subgrains will subse-quently coalescence or be consumed by the growing neigh-boring grains. As a general trend, the rotation axes for mis-orientation angles below 40◦ are randomly distributed in thesample (Fig. 6c and e) while the rotation axes for misorienta-tion angles above 50◦ are close to particular crystallographicdirections:〈1 2.0〉 and/or〈0 1.0〉 (Fig. 6b and d). In addition,they are localized around the rolling direction (Fig. 6c ande), which clearly indicates that the second peak (at high an-gles) in the misorientation angle distribution histograms isrelated to the sample symmetry. It corresponds to the mis-orientations between grains having the�c-axis tilted from NDby +35◦ and−35◦ around RD, which coincides with either〈1 2.0〉 or 〈0 1.0〉, depending on the texture evolution degree.

4. Discussion

In order to shed more light on our understanding of the tex-t s atc dis-a rains ries,a tionw -s eckedt etryi

4

de-c Thisd ariesa ain-i callyu over,f ear:a withw Asg s likeC riesa gy isa them riesi

SLbf ries.

As suggested by the kinetics data, the presence of a verysmall amount of precipitates or the effect of solute atoms donot appear to have a predominant effect in controlling thegrain growth. Thus, the grain boundary misorientation angleappears to be the major parameter controlling grain boundarymobility, and consequently the criterion on the basis of whichgrain growth should be discussed.

As shown inFig. 6a, the resisting grain boundaries in cp-Ti are characterized by an increasing proportion of misori-entation angles below 30◦ and the corresponding boundaryenergies should then be low. The disappearing grains are mis-oriented by more than 30◦ with respect to the ideal orientation{0◦, 35◦, 30◦} (Fig. 3), which indicates that the boundarieswith a misorientation angle higher than 30◦ can then be con-sidered as fast moving. Such statement needs of course to beconfirmed, either by fundamental characterization of grainboundaries or by using grain growth models. It is worth notic-ing that all the high misorientation angle boundaries do notdisappear over grain growth: the misorientation angle distri-butions ofFig. 6 also contain a second contribution around70◦ which remains even after extended grain growth. The rea-son for this is that those ones result from the orthorhombicsample symmetry. Therefore, this contribution cannot disap-pear even if the corresponding boundaries have high energyand mobility. During grain growth, some of these boundariesdisappear but others will be formed because they are requiredt

rainbm osem s de-t

4

inc f anAi eso mi-n mlyo e andt one,b ons.T asedo wase ajorc com-p pecialg p toM im-i withms akeo ypeo lose

ure development during grain growth, this discussion aimlarifying why certain orientations develop while othersppear, in connection with (i) the correlation between gize and orientation and (ii) the nature of the grain boundaccording to energy and mobility criteria. Texture evoluas illustrated here in theϕ1 = 0◦ ODF section, but the reults on which the present discussion is based were cho be valid for the whole ODF. In addition, sample symmnduces similar situations atϕ1 close to 180◦.

.1. Grain boundaries

During grain growth, the free enthalpy of the materialreases by reducing the grain boundary total surface.ecrease is more effective if high energy grain boundre eliminated. Thus, the resisting grain boundaries, rem

ng after grain growth, should be the less thermodynaminfavorable ones, i.e. the ones with lowest energy. More

ast moving boundaries have higher probability to disappboundary keeps moving fast until it reaches a grainhich the growing grain forms a slow moving interface.eneral trends (not considering special grain boundarieSLs), it is usually assumed that (i) fast moving boundare associated to high misorientation angles; (ii) enerlso higher for highly misoriented grain boundaries; (iii)isorientation limit between low and high angle bounda

s usually around 15–20◦ [21].Our results do not highlight a significant amount of C

oundaries and do not reveal any special rotation axes (Fig. 6)or the misorientations associated to the grain bounda

o maintain the orthorhombic sample symmetry.In the frame of this work, accepting the idea that the g

oundaries with misorientation angles lower than 30◦ areoving more slowly than others, it is possible to propechanisms responsible for the texture evolution. This i

ailed below.

.2. Mechanisms responsible for texture evolution

The present texture evolution during grain growthp-titanium shows some similarities with the case ol–1 wt.% Mn reported by Rios and Gottstein[22], where the

nitial texture of the Al alloy was constituted of three typf grains belonging to (i) a major cube component, (ii) aor component and (iii) a random distribution. The randoriented grains disappeared due to a smaller grain siz

he minor component grew at the expend of the majoroth under normal and abnormal grain growth conditihe discussion of the mechanisms in this Al alloy was bn the fact that the growth rate of the minor componentnhanced by a CSL orientation relationship with the momponent. In the present case of titanium, the textureonents are much less sharp and the occurrence of srain boundaries is very low. The Ti texture evolution uGS = 32.8�m can nevertheless be explained in a very s

lar manner by considering that the grain boundariesisorientation angles higher than 30◦ move faster.Fig. 7is a

chematic view of the microstructure evolution. For the sf clarity, this simplified microstructure contains three tf grains: the A- and B-types of grains are oriented c


Fig. 7. Schematic grain growth evolution of a simplified microstructure constituted with three types of grains: (i) A grains (light grey) have the mainorientationof the texture obtained after primary recrystallization: around{ϕ1 = 0◦, Φ = 35◦, ϕ2 = 0◦}, (ii) B grains (white) have similar size distribution as A grains, butthey are less numerous at the beginning of grain growth; they are oriented close to{ϕ1 = 0◦, Φ = 35◦, ϕ2 = 30◦}, (iii) C grains (dark grey) are smaller than theA- and B-type, their orientation is either close to{ϕ1 = 0◦, Φ > 65◦, ϕ2 = 0◦} or {ϕ1 = 0◦, Φ < 20◦, ϕ2 = 0◦}, or random. Due to their small size, C-grains haveno chance to grow, while A and B grains can develop when locally favored; the stars indicate the grains which have a local (size and/or morphology) advantageto grow. Thick lines represent high misorientations between grains of different types. Thin lines represent low misorientations between grains of the same type.

to{0◦, 35◦, 0◦} and{0◦, 35◦, 30◦}, respectively. The C-typeof grains are characterized by orientations widely dispersedthroughout the Euler space and by smaller sizes than the A-and B-ones.

The C-type of grains are the ones disappearing duringgrain growth. The disappearing orientations are widely scat-tered in the orientation space but it is clear that the do-main where disappearing orientations are located graduallychanges during the grain growth process. During the firststages of grain growth (MGS from 3.3 to 8.3�m), the orien-tations atϕ1 andϕ2 close to 0◦ andΦ below 20◦ or above65◦ disappear. Over larger grain size ranges (MGS up to32.8�m), the orientations atϕ1 andϕ2 close to 0◦ andΦ be-tween 40 and 65◦ keep disappearing. Globally, all the grainsbeing misoriented by more than 30◦ with the growing ori-entations (around{0◦, 35◦, 30◦}) tend to shrink and disap-pear (Fig. 3). These grains are statistically characterized bysmaller sizes than the ones of the main texture components (Aand B) because the densities of the disappearing orientationsare higher in the partial texture of the small grains (Fig. 5)compared to the one of the large grains. Their disappearanceis due to the combined effect of their smaller size and of theirhigh misorientation angle boundaries.

The growing orientations define a broad peak centered at{0◦, 35◦, 30◦}. The initial major texture component centeredaround{0◦, 35◦ 0◦} (corresponding to A-type grains) is in-c s dog s be-i s)w eset re int ed us,

grains A and B are both growing to the detriment of the small-sized and highly misoriented C-type grains. However, the tex-ture componentgA raises less than the componentgB. Tworeasons can be called upon to explain this, as illustrated inthe schematic view ofFig. 7. Firstly, A grains, because theyare more numerous at the beginning, have a higher proba-bility to have neighbors of the same type, which means lowmisorientation angle and, thus, low mobility grain bound-aries (thin lines on the scheme). Secondly, according to thelocal configuration (especially the size and the morphologyof a grain compared to the ones of its neighbors), either anA grain can grow to the detriment of B grains, or inversely aB grain can consume A ones. OnFig. 7, the grains markedwith a star are favored to grow. Let us consider first the caseof the growth of an A grain. When the first neighbors of B-(and C-) type will be fully consumed, the growing A grainwill meet a second order neighbor. The probability for thelatter to be of an A-type character is higher than being of theB-type. This leads to a slowing down of the growth speedof the considered A grain by orientation pinning. Compar-atively, a growing B grain, after having consumed its firstorder A (and/or C) neighbors, has a low probability to meeta second order neighbor belonging to the B-type. The newgrain boundary will then continue to move and the B grainwill keep growing because, since the grain has grown, it issize-advantaged and can continue to grow until it meets ag lowm

entn eso -i henM is

luded in the domain of the orientation space where grainrow. However, these grains grow much less than the one

ng closer to{0◦, 35◦, 30◦} (corresponding to B-type grainhich represent initially a lower volume fraction. Since th

wo types of orientations do not show any specific featuhe grain size-based partial textures inFig. 5, the grain sizistributions of the A- and B-type grains are similar. Th

rain with which a low misorientation angle (and thusobility) grain boundary is created.The grains belonging to the initially minor compon

eargB ={0◦, 35◦, 30◦} keep growing faster than the onriented neargA ={0◦, 35◦, 0◦} until they become predom

nant. After extensive grain growth has taken place (wGS increases from 32.8 to 64.0�m), the inverse trend


observed. This is because, within the last stages, the A grainsare becoming less numerous than the B grains. Such an in-version in the growth rates of two texture components withinitially different volume fractions has been modeled by us-ing a Monte Carlo model and by considering anisotropicgrain boundary mobility[23]. This occurs in cp-Ti for largegrain sizes (Fig. 3h) but without significant effect on theglobal texture because, at this stage, the reciprocal mecha-nism (with A grains growing to the detriment of B grainswhile B grains grow to the detriment of A grains some-where else in the material) becomes predominant, as shownin Table 2.

5. Conclusions

Recrystallization and grain growth in commercially pureTi sheets lead to equiaxed microstructures without abnormalgrowth. The texture at the end of primary recrystallization isvery close to the one of the deformed state with the majorvolume fraction of the material around{0◦, 35◦, 0◦} in Eulerspace. During grain growth, the texture evolves slowly bydeveloping, after extensive annealing, orientations locatedaround{0◦, 35◦, 30◦}. The major conclusions revealed byour detailed analysis of the grain growth are as follows:

e be-thenta-por-lvesap-

ns ofers,e in

fer-ing

n atjori-

d inrien-ng to

the smallest size range at the end of primary recrystal-lization and have fast moving boundaries (with misori-entation higher than 30◦).

(IV) CSL grain boundaries do not appear to have an impor-tant effect during grain growth. The grain boundaries re-maining after extensive grain growth are characterizedby an increasing proportion of misorientation anglesbelow 30◦. A large amount of high misorientation an-gle (around 70◦) grain boundaries nevertheless remainseven after extensive grain growth due to the orthorhom-bic sample symmetry.

References

[1] F.J. Humphreys, M. Hatherly, Recrystallization and Related Anneal-ing Phenomena, Pergamon Press, Oxford, 1996.

[2] R.D. Doherty, D.A. Hughes, F.J. Humphreys, J.J. Jonas, D. JuulJensen, M.E. Kassner, W.E. King, T.R. McNelley, H.J. McQueen,A.D. Rollett, Mater. Sci. Eng. A 238 (1997) 219.

[3] K. Hulse, K.H. Kramer, in: G. Lutjering, U. Zwicker, W. Bunk(Eds.), Titanium Science and Technology, Deutsche Gesellschaft furMetallkunde EV, Munich, 1985, p. 1065.

[4] A.K. Singh, R.A. Schwarzer, Z. Metallkd. 91 (2000) 702.[5] F. Wagner, N. Bozzolo, O. Van Landuyt, T. Grosdidier, Acta Mater.

50 (2002) 1245.[6] H. Hu, R.S. Cline, Trans. Met. Soc. AIME 242 (1968) 1013.[7] K. L ucke, R. Rixen, Metall. Trans. 1 (1970) 259.

ence,

mbe,

[ eed-rials88,

[ 9.[[[[ rum

[ Lon-

[[ A

[ dat,

[ 37

[ et-Ra-

[[

(I) The changes in texture are more pronounced at thginning of the grain growth process. At this stage,size criterion and the correlation between the orietion and the size of the grains seem to be very imtant. Over long-term heat treatments, the texture evomuch more slowly. The predominant mechanismpears then to be a counterbalanced one with graisome orientations growing to the detriment of othwhile a reciprocal situation occurs somewhere elsthe material.

(II) The computation of ODF differences between difent grain growth stages clearly shows that the groworientations are included in a broad peak (30◦ misori-entation wide) centered on{0◦, 35◦, 30◦}. The cor-responding grains belong to the larger-size fractiothe end of primary recrystallization. The initially macomponent{0◦, 35◦, 0◦} is included in this broad orentation domain but the orientations close to{0◦, 35◦,0◦} develop less than the ones closer to{0◦, 35◦, 30◦},due to orientation pinning.

(III) The disappearing orientations are widely scatterethe orientation space and characterized by a misotation angle larger than 30◦ with the ideal orientatio{0◦, 35◦, 30◦}. They disappear because they belon

[8] J. Grewen, Proceedings of the Third European Texture ConferSociete Franc¸aise de Metallurgie, Paris, 1975, p. 195.

[9] S. Naka, J.P. Mardon, P. Dervin, R. Pernot, R. Penelle, P. LacoJ. Less-Common Met. 56 (1977) 277.

10] H. Inoue, N. Inakazu, in: J.S. Kallend, G. Gottstein (Eds.), Procings of the Eigth International Conference on Textures of Mate(ICOTOM 8), The Metallurgical Society Inc., Warrendale, PA, 19p. 997.

11] M.L. Weaver, H. Garmestani, Mater. Sci. Eng. A 247 (1998) 2212] S. Zaefferer, Mater. Sci. Eng. A 344 (2003) 20.13] H.J. Bunge, R.A. Schwarzer, Adv. Eng. Mater. 3 (2000) 25.14] F.J. Humphreys, J. Mater. Sci. 36 (2001) 3833.15] F. Wagner, N. Bozzolo, N. Dewobroto, N. Gey, Mater. Sci. Fo

408–412 (2002) 143.16] H.J. Bunge, Texture Analysis in Material Science, Butterworths,

don, 1982.17] J.E. Burke, D. Turnbull, Prog. Met. Phys. 3 (1952) 220.18] C. Herzig, Y. Mishina, S. Divinski, Metall. Mater. Trans. A 33

(2002) 1.19] S.A. Saltykov, Stereometric Metallography, 2nd ed., Metallurgiz

Moscow, 1958.20] R. Bonnet, E. Cousineau, D.H. Warrington, Acta Crystallogr. A

(1981) 184.21] G. Gottstein, L.S. Shvindlerman, Grain Boundary Migration in M

als: Thermodynamics, Kinetics, Applications, CRC Press, Bocaton, 1999.

22] P.R. Rios, G. Gottstein, Acta Mater. 49 (2001) 2511.23] K. Mehnert, P. Klimanek, Comput. Mater. Sci. 7 (1996) 103.

texture evolution during grain growth in recrystallized commercially pure titanium

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