Texture Evolution in Severe Plastic Deformation Processes

Satyam Suwas and Soumita Mondal

Department of Materials Engineering, Indian Institute of Science, Bangalore 560012, India

Severe plastic deformation processes involve large grain rotations due to the action of different modes of plastic deformation and othermicrostructural changes which lead to characteristic texture formation. The present review deals with the evolution of texture during the mostimportant severe plastic deformation processes, namely Equal Channel Angular Pressing (ECAP), High Pressure Torsion (HPT), Friction StirProcessing (FSP), Accumulative Roll Bonding (ARB) and Multi-Axial Forging (MAF). First three of the processes are shear based, while thelatter two are plane-strain based. The textures formed during ECAP are visually different from simple shear textures due to (i) the inclination ofthe shear plane, (ii) additional contribution of non-shear based deformation. The relative intensities of texture components are function ofdeformation micro-mechanisms, amount of straining and configuration of the strain path. The texture evolved during HPT is very similar tosimple shear texture, with additional consequences of microstructural changes that occur due to very large deformations. The textures formed inFSP process also resemble shear textures. On the other hand, texture evolution during ARB and MAF can be described using plane straindeformation. The present review deals with texture evolution during severe plastic deformation as a function of nature of processes and type ofmaterials. [doi:10.2320/matertrans.MF201933]

(Received February 25, 2019; Accepted March 18, 2019; Published June 21, 2019)

Keywords: texture, severe plastic deformation

1. Introduction

The relevance for SPD processing is the production of bulkultrafine grained materials without changing the dimensions.Severe plastic deformation is generally carried out for grainsize reduction, which in turn leads to the enhancement instrength and hardness of a material. SPD techniquesintroduce a large amount of strain in the material, and avariety of strain path changes. As a consequence, substantialmodification/evolution of texture also occurs, that plays anequally important role in the mechanical response of thedeformed material, together with grain size. Texture too playsa major role in the grain refinement process during SPD.

Texture is defined as the orientation distribution ofcrystallites in a polycrystalline aggregate. Texture changesare produced by the rotation of grains. During deformation,small regions within a grain start rotating by the activity ofdifferent slip systems, which lead to the formation ofsubstructures within the grain. The orientation within thesesubstructures will have a slight deviation from that of itsparent grain. Gradually these substructures rotate to preferredorientations which are characteristic to the applied deforma-tion. The same mechanism is applicable when deformationis mediated through twinning, except for the fact that thetwinned region will exhibit a high misorientation angle withrespect to its parent grain. The characteristic orientations arealso dependent on the rate of rotation of the crystallites.

2. Texture Evolution during Severe Plastic Deformation

Many severe plastic deformation techniques have beenproposed. The most important among them are equal channelangular pressing (ECAP), high pressure torsion (HPT),friction stir processing (FSP), accumulative roll bonding(ARB) and multi-axial forging (MAF). Out of theseprocesses, first three are shear based processes, and the lattertwo involve plane strain deformation. Accordingly, thetexture evolved during severe plastic deformation can also

be categorised in two types: (i) shear based, and (ii) planestrain based.

2.1 Ideal orientations after shear based SPD processesAs mentioned earlier, textures evolved during SPD are

basically deformation textures. The parameters that controlthe deformation texture are the imposed strain path,crystallographic factors governing the mechanism ofdeformation (slip, twinning etc.) and the initial microstructureand texture. In the following subsections, the evolution oftexture has been described on the basis of strain path selectedfor each type of materials. ODF sections showing the locationof ideal orientations during shear based and plane-strainbased deformation processing are shown in Fig. 1, and thecorresponding components are listed in Table 1.2.1.1 FCC materials

The ideal orientations formed under simple sheardeformation in an FCC crystal has been widely studied andreported.1,2) The texture in ECAP materials is described withrespect to the shear plane (SP) and the shear direction (SD).The texture typically consists of fibers, designated as the Afiber with {111} ¬ SP and the B fiber with ©110ª ¬ SD. TheA fiber contains the A, �A, A�1, A

�2 texture components while

the B fiber contains the A, �A, B, �B and C components. Mostof the texture components lie along these fibers. Since therelative intensities of these components are dependent on thesymmetry of the test being carried out, in the case of simpleshear deformation, which exhibits a two-fold symmetryaround the axis perpendicular to both shear plane normal(SPN) and shear direction (SD), the A/ �A, and B/ �Bcomponents have the same intensities. A�1 and A

�2

components exhibit different intensities as they are notsymmetric to the two fold symmetry operation of the SPNand SD. C components are self-symmetric.2.1.2 BCC materials

The ideal orientations formed during simple sheardeformation of BCC materials were first calculated byMontheillet and Jonas.1) During ECAP of BCC materials,

Materials Transactions, Vol. 60, No. 8 (2019) pp. 1457 to 1471Special Issue on Severe Plastic Deformation for Nanomaterials with Advanced Functionality©2019 The Japan Institute of Metals and Materials REVIEW

https://doi.org/10.2320/matertrans.MF201933

the most commonly observed fibers are {110} ¬ SPcontaining the ideal components F, J, �J , E and �E, and©111ª ¬ SD containing the ideal components D1, D2, E and

�E.2) Additionally two more ideal orientations were alsoreported which appeared with the activation of slip on the{123} family of planes.3)

2.1.3 HCP materialsThe ideal orientations formed in HCP materials with near

ideal c/a ratio are fibers B with {0001} ¬ SP (basal fiber) andP ©11.0ª ¬ SD (prismatic fiber), Y, C1 and C2 fibers. However,the stability of the orientation is purely dependent on the c/aratio of the metal which dictates the relative activation of theslip/twin families. Texture after simple shear deformation isrepresented in terms of shear direction SD, normal direction(ND) and shear plane normal (SPN) in pole figures (PFs) andinverse pole figures (IPFs).

A point to be mentioned here is that initially the twoleading groups, led by Tóth and Tomé have representedECAP texture components with additional suffix ‘E’ and ‘ª’added to the shear texture components, respectively. In lateryears, it was mutually decided to discard these additional“suffix” while representing the ECAP textures in order tobring uniformity of representation.

3. Evolution of Texture during Shear Based SPDProcesses

3.1 Equal channel angular pressing (ECAP)Equal channel angular pressing (ECAP) is one of the

extensively studied SPD processing techniques. In ECAP,specimens with square or a circular cross section are passedthrough two dies of equal cross-sectional area inclined atan angle, known as inter-channel angle typically rangingfrom 90°150°. During consecutive ECAP passes, thesample is re-inserted multiple times with or without arotation along the longitudinal axis of the billet in-betweenthe passes. The rotation given between consecutive passesalong the billet longitudinal axis are 0° in Route A,alternating between «90° in Route BA, 90° in the samedirection in Route BC and 180° in Route C. The deformationand velocity gradient for ECAP processing by the routesmentioned above can be found in an extensive review byBeyerlein and Tóth.4)

Owing to the complexity of deformation path in ECAP,texture varies significantly depending upon crystal structure,stacking fault energies and initial texture. However, in allthe processes, textures are unique to the crystal structure.Although ECAP textures have been primarily described asshear textures, the shear plane is identified as the intersectionplane of the inlet and exit channels.5) This causes the idealtexture components to rotate about the transverse direction(TD) by an angle of 45° for an ECAP die with inter-channelangle 90°. Further shift occurs along the TD (i.e. ¤1) with anincrease in the die angle. The die angle has a great influencein the texture formed after each pass. Smaller the die angle,higher is the strain imposed per pass. As a result, rotation ofgrains to ideal orientations will be faster and the textureshould appear stronger.

Besides the die angle, other parameters that affect thetexture are the crosshead speed of the punch, friction betweenthe channel and sample, and the application of a backpressure. It has been concluded from flow line analysis thatthe material flow from top to bottom of the channel is not

Fig. 1 Schematics of SPD processing techniques along with ODF sectionsthe showing position of ideal orientations.33,59,83,117)

S. Suwas and S. Mondal1458

identical. The material closer to the top surface of the channelundergoes deformation more severely than the one at thebottom.6) Therefore, the texture formed will also beinhomogeneous from top to bottom of the sample.

After the 1st pass, texture evolution depends on the ECAProute.4) The direction of shear is reversed in every alternatingpass in route BC and every consecutive pass in route C. Inroute A, the shear planes meet at an angle of 90° and at anangle of 120° in route BA. The monoclinic sample symmetryis lost after the 1st pass when the material is processed byroutes BA and BC, while it is geometrically maintained whenprocessed in routes A and C since the shear plane remains onthe texture-symmetry-axis of the process.7)

A noteworthy observation in most of the work pertainingto ECAP texture is that, there is always a systematicdeviation of experimentally measured texture componentfrom their expected ideal location.5,8,9) This is primarilyattributed to the mechanics of the process which involvesadditional deformation modes operative during the process,in addition to the primary deformation mode as simple shear.A list of texture components that evolve during ECAP isgiven in Table 2.3.1.1 Texture evolution in ECAP processed FCC

materialsIn FCC materials, the components of texture developed

after ECAP are dependent on the deformation mechanism,which in turn depends on stacking fault energy. Althoughsimilar pattern in texture evolution is observed in most of theFCC materials, the key difference manifests in terms of therelative intensities of the components which arise out ofthe differences in the stacking fault energies of the differentmaterials. In the following sections, texture evolution hasbeen discussed for materials with each of the specified rangeof stacking fault energy, for example, Al and Ni, the highSFE metals (³160 and ³128mJ/m2, respectively), Cu, amedium SFE metal (³78mJ/m2) and Ag, a low SFE metal(³22mJ/m2).10)

The most complete study on ECAP texture in aluminium isdue to Suwas et al.,11,12) who investigated textures formed asfunction of processing routes as well as number of passes(Fig. 2). The strong cube texture of the initial materialprevailed after one pass, and the presence of cube texturecomponents was clearly seen along ©001ª ¬ TD fiber. TheODF plots revealed very weak intensity along A�1 and Ctexture components, while the other components, A/ �A andB/ �B, were yet to form. The characteristic texture for theECAP route developed from the second pass. Further passesled to a change in the relative intensities and a slightdeviation in the position of the maxima of the texturecomponents. In routes A and C, the strongest texturecomponents were A�1, and C, respectively, whereas B/ �Bcomponent was strong in route BC. Also, when deformedthrough route A, texture started weakening with subsequentpasses, in route BC it strengthened continuously, in route C,the relative intensities of texture components oscillated. Witheach pass, the influence of initial texture diminished, aphenomena observed in all the ECAP routes.1115)

With decreasing SFE from Al to Ni9,16) to Cu,5,6,8,17) it wasobserved that for there was a change from the dominance ofC component to B/ �B components when deformed byroute A. In Ni, intensity of the C component was the highest,followed by A�1 and B/ �B having nearly equal intensities.

14)

In Cu, the A�1 component has the highest intensity, followedby the C component and then the B/ �B components.6) Thedevelopment of texture as a function of SFE is shown inFig. 3. Texture development in Cu via different processing

Table 1 Texture components formed in various FCC, BCC and HCP materials during Simple Shear Deformation.

Table 2 Texture components formed in various FCC and BCC materialsduring ECAP processing.

Fig. 2 Texture evolution during ECAP processing of Al as a function ofprocessing routes.12)

Texture Evolution in Severe Plastic Deformation Processes 1459

routes is shown in Fig. 4. Interestingly, in the case of copper,Suwas et al.8) investigated the effect of material purity onthe texture evolution and found that the volume fraction oftexture components was dependent on the material variableslike impurity content and starting microstructures. The Cualloy, Cu0.3%Cr, which has SFE close to pure Cu, exhibitsstrong B/ �B components when deformed by route A and astrong A�1 component with some amount of B/ �B componentswhen deformed by route BC.18)

In the case of Ag, texture was dominated by the B/ �Bcomponents, which was attributed to twinning as a resultof its low SFE19) (Fig. 3). After the 1st pass, Ag exhibitedtexture strengthening when deformed in subsequent passesby route A. The B/ �B components strengthened at theexpense of A�2 component. The A

�1 component remained

weak but stable, even up to 3 passes. Microscopicobservations confirmed that twinning was an active modeof deformation even in the third pass.2022)

Skrotzki et al. investigated the dependence of textureevolution on starting textures of another low SFE AlMgalloy, AA5109.23) Starting with a strong cube texture, thematerial showed overall texture weakening after the 1st pass,with the evolution of texture components C, A�1, B/ �B, A and

A�2, in decreasing order of relative intensities. On the otherhand, with a starting rotated cube texture, after the 1st pass, arather strong texture formed with A�2, A

�1, B/ �B, A and C

components, in decreasing order of relative intensities. Withincrease in the number of passes, the effect of starting texturediminished completely. After 3rd pass, the B/ �B componentswere the strongest, a characteristic signature of low SFEmaterial.23)

Studies were also carried out to investigate the evolution ofECAP textures in engineering alloys of aluminium, namelythe AlCu alloy (AA2014)24,25) and it was revealed that theevolution of texture was strongly dependent on the startingmicrostructures and processing routes. The starting materialin aged condition developed weaker textures compared to thesolution treated samples, after 5 passes of ECAP, which wasattributed to higher scattering of strain during deformation bythe fine precipitates in the aged samples.25) This study alsoreported that ECAP by route A led to the strongest texturesand highest heterogeneity from top to bottom of the sample.The weakest texture and lowest heterogeneity was observedpost deformation by route C as a result of reversed strainpath. Texture evolution during deformation via routes BA andBC were found to be nearly similar. On an average, texturesmeasured at the mid plane of ECAP billets, showed thatthe components B/ �B and A�1 showed the highest intensitieswhen deformed by routes A, BA and BC, while A�2 and Ccomponents showed highest intensities when deformed byroute C.24) In another Al alloy AA2195, Suresh et al.26)

observed that up to 3rd pass by route A, the componentsA�1, A

�2, and B/ �B exhibited highest intensities while after the

4th pass A�2 and C components exhibited highest intensities.Skrotzki et al. have extensively studied the heterogeneity

of ECAP textures in Al, Ni, Cu and Ag6,14,16,17,22,27) (Fig. 5).While most of these studies were carried out on samplesprocessed through route A, the samples processed byroute BC also led to texture gradient along the transversedirection, in addition to gradient from top to bottom. In thiscase, the components B and �B, display opposite trends with Bcomponent increasing in intensity from left to right, half wayfrom the top and �B exhibiting the exact opposite trend.27)

To summarise, ECAP textures of FCC materials have beenfound to be strongly dependent on stacking fault energy,

(a)

(b)

(c)

Fig. 3 Variation in texture evolution during ECAP processing of FCCmaterials like (a) Al, (b) Cu and (c) Ag as a function of SFE.19)

Fig. 4 Texture evolution during ECAP processing of Cu as a function ofprocessing routes (Unpublished).

Fig. 5 Texture gradient measured by synchrotron radiation after 1 ECAPpass of Cu as a function of distance from top of the billet.27)

S. Suwas and S. Mondal1460

initial texture, processing route, number of passes anddistance from top to bottom of the sample. With decreasein stacking fault energy, the strengthening of the B/ �Bcomponents occur either due to a propensity for twinningor hindrance in the slip activity due to the presence ofprecipitates. The effect of initial texture diminishes withincrease in number of passes. Evolution of texture by route Cis strongly dependent on the number of passes, while inroutes A and BA/BC texture strengthening and weakeningoccurs, respectively. Lastly, the severity of texture is afunction of its distance from top to bottom of the sample withstrongest texture at the top. There is a shift in the location oftexture components from top to bottom, from the correspond-ing ideal locations of shear textures.3.1.2 Texture evolution in ECAP processed BCC

materialsA few BCC metals and alloys have been investigated for

texture evolution during ECAP.2,2831) The early studies dueto Li et al.2,28) showed a good agreement of ECAP texturewith the ideal simple shear texture for BCC materials.1) Liet al. observed the presence of partial ©111ª-fiber and {110}-fiber rotated about TD by 45°, which was furthercorroborated by Bhowmik et al. for IF steels.3) After the1st pass, the material exhibited strong texture with{112}©111ª components. Rotation of the ED-ND plane by45° in counter clock-wise (CCW) direction about the TD axisgave rise to this texture with the components D1 and D2. Theformation of J and F components take place by the alignmentof ©112ª and ©001ª directions along the shear direction,respectively, on the {110} planes which align along the shearplane.3) Additionally this paper also reported two new idealpositions R1 and R2 formed by slip on the (123) planes.These components were observed after the 4th pass by routesA and BC.3) In general, the coincidence of slip plane ordirection or both with the macroscopic shear plane anddirection of the sample was observed in most cases whendeformed by both the routes. Figure 6 shows textureevolution in IF steels processed through routes A and BCof ECAP.3.1.3 Texture evolution in ECAP processed HCP

materialsAs mentioned earlier, the mode of deformation in HCP

materials is dictated primarily by c/a ratio. Mg and Co havetheir c/a close to ideal, as theoretically predicted for a closepacked HCP structure (³1.633). In materials with higher c/aratio, deformation is only accommodated by slip on the basalplane and twinning, whereas with lower than ideal c/a ratio,it is possible to activate slip on other planes also in additionto basal slip and twinning.32) In the following sections,texture evolution in HCP materials with different c/a ratiohas been discussed.3.1.3.1 Texture evolution in materials with near ideal c/aratio

Magnesium and its alloys are the most investigatedmaterials in this category. The comprehensive study byBeausir et al. on texture evolution during ECAP of Mg withstrong basal fiber texture showed the formation of differentfinal textures when processing through routes A, BC and C33)

(Fig. 7). Texture after the 1st pass consists of a weak B fiberformed by rotation of initial orientations by about 105° in the

clockwise direction around TD. It was also observed thatwhen the textures were rotated by +90° before reinsertion ofthe sample for the second pass, the texture appeared verysimilar to the undeformed sample with just a small tilt of¹15° along the TD axis of the sample. The same observationwas made in the texture evolution after each pass in route A.This phenomenon was attributed to rotation of the sample by90° before reinsertion in the real space. Thus, the second passonwards, the rotation of textures before reinsertion positionsthe textures to that of the initial sample. Biswas et al. alsoreported a constant formation of the B fiber on ECAP of pureMg by route A upto 8 passes34) and that with an increase in

Fig. 6 Texture evolution during ECAP processing of IF steel after 1 pass,and after 2 and 4 passes via processing routes: A and BC.3)

Fig. 7 Texture development in Magnesium33) and Titanium41) after 3passes of ECAP via different processing routes.

Texture Evolution in Severe Plastic Deformation Processes 1461

the deformation temperature, the activity of the pyramidal©c + aª II type slip system increased, resulting in theappearance of C2 fiber along with B fiber.35)

In route BC too, formation of B fiber texture was observed.This B fiber forms at a position which is 20° rotated in theclockwise direction position around the simple sheardirection axis. In route C, similar B fiber texture formed,with the only difference being a split in the fiber from the 3rdpass onwards, which is best visible in (0002) pole figure. Thefindings were consistent with studies carried out later.36)

Studies have also been carried out on effect of alloyingadditions on the texture evolution during ECAP of Mg alloyssuch as AZ31, AZ80, ZK60, WE43, Mg4L.3740) Based onthese studies it could be inferred that deformation of the AZgroup of alloys took place predominantly by basal slip with alow activity of non-basal ©aª and ©c + aª slip, and the B fibersshowed an inclination of ³55° to the ED. For the ZK and WEgroup of alloys B fiber aligned along ND, caused by a higheractivity of non-basal ©c + aª slip systems in conjunction withbasal slip. MgLi alloys too develop similar textures withbasal fiber aligned along ND, but do so because of higheractivity of prismatic ©aª slip.3.1.3.2 Texture evolution in materials with c/a < 1.633

Studies carried out on ECAP processed Ti, showed thatafter the 1st pass, texture resembled the characteristic simpleshear textures of HCP materials (Fig. 7). A small deviation of5° away from the ideal position was noticed. It was alsoobserved that the basal poles tend to align in orientationsbetween the normal direction (ND) and shear plane normal(SPN) direction.41,42) This was clearly different from the caseof Mg, where the B fiber aligned along the SPN. Textureformed in subsequent passes in route A exhibit a similartrend observed in the case of Mg. In another study with thesame die geometry by Jager et al.,43) the texture formed after4th pass of ECAP by route A corroborated the results of thisinvestigation. The only difference was due to the applicationof a back pressure between 270 and 590MPa. The texturesdeveloped were found to be almost similar to those reportedin a previous study by Beyerlein et al. on the texturedevelopment in pure Be processed through routes A and C ofECAP for 2 passes.44) They had previously reported that thetextures developed in Be were tilted by ³15° after 1st and2nd pass of ECAP by route A. Yapici et al. reported similartrend in texture evolution for pure Be, Zr, and alloys Ti64 andAZ31, which leads to the conclusion that the trend ofevolution of ECAP texture is more or less similar for all HCPmetals with c/a ratio 1.633ECAP textures of HCP materials with a c/a ratio greater

than ideal has not been studied extensively because of theirlimited application. The paper due to Bednarczyk et al.47)

reported texture evolution in single phase Zn alloy, Zn0.5Cu. The samples were subjected to 4 passes of ECAP ina 90° die by route BC, and the resulting texture wascharacterized by the presence of prismatic P and pyramidal©aª - Y fibers with a deviation of ³15° from the idealorientation.

To summarize, it was observed that after ECAP process-ing, a basal B fiber texture formed in materials with c/a ratiosclose to and lower than ideal, while formation of prismaticP and pyramidal Y fiber textures were observed in hcpmaterials with c/a ratio greater than ideal. It was alsoobserved in both Mg and Ti that multi-pass ECAP by route Cled to the formation of a split basal fiber texture.

3.2 High pressure torsion (HPT)High pressure torsion is the perhaps the oldest SPD

processing technique.48) This technique has been usedextensively over the last couple of decades owing to thepossibility of subjecting the material to extremely highamount of strains. Some review papers on the mechanics ofHPT, and the physical and mechanical responses of thematerial after processing by HPT can be easily found.49,50)

Texture evolution during HPT of different types of materialsis discussed in the following sub-sections.3.2.1 Texture evolution in HPT processed FCC materi-

alsTextures developed after HPT of FCC materials very

closely resemble the simple shear textures as shown for thecase of high purity Cu in Fig. 8.51) Naghdy et al.52) reportedthat commercially pure recrystallized aluminium exhibitedsimple shear type texture after being subjected to HPT toa strain of only 0.75, however without the characteristicmonoclinic symmetry. Symmetrically equivalent compo-nents, such as B/ �B and A/ �A, showed different relativeintensities, which was attributed to the domination of textureweakening effect resulting from grain fragmentation and

S. Suwas and S. Mondal1462

rotation over the texture strengthening. At higher strains,texture weakening further dominates. With increasing strains,£ > 4.0, continuous strengthening of the C and Bcomponents was observed at the expense of rotation ofgrains away from A�1 and A

�2 ideal positions.

52) Korneva et al.studied texture evolution in a CuCr alloy under HPT afterannealing at 550°C/1000°C and found that the position andmaxima of shear components are nearly identical in both thesamples.53) The position of texture components was deviatedfrom the ideal simple shear texture components by a smallangle along ¤1. Similar texture evolution was observed in Cuphase of a nanolayered CuNb composite post HPT.54) Studyon HPT processed pure Al57) indicated texture evolution fromA�1, A

�2 to C components up to a strain of 4, followed by an

overall texture weaking. Pure Ni58) too shows a strengtheningof C and B/ �B components, similar to pure Cu and 70/30Brass.

The study by Skrotzki et al. on a randomly texturednanocrystalline Pd10 at%Au sample, also showed thepronounced dominance of B/ �B components, which attaineda maximum relative intensity at a shear strain of ³6 andtexture evolution got saturated.55) Weak A�2 and Ccomponents were also observed beyond this deformationlevel. Another important observation in this study was theshifting of all the texture components along ¤1 by ³10°. Themechanism of formation of strong B fiber was attributed to aneasier ©112ª slip, which was found from texture simulationsusing Taylor model.

Among the low SFE FCC materials, the most notable studyis due to Skrotzki et al.,56) who investigated texture evolutionin HPT processed high-entropy alloy CrMnFeCoNi up toshear strain ³170. A weak, yet a characteristic shear typetexture with dominant B/B components was observed, whichwas attributed to deformation mainly by partial dislocationslip accompanied by twinning and grain boundary sliding.3.2.2 Texture evolution in HPT processed BCC materi-

alsRelatively fewer studies have been performed on texture

evolution during HPT of BCC materials. Zhao et al. studiedthe texture of HPT compacted Fe powder.59,60) The [110] polefigures showed the strengthening of BCC texture componentswith increase in deformation, i.e., transition from top tobottom surface of the sample. A strengthening of typicalsimple shear texture components of BCC materials wasexhibited from the undeformed to transition to severelydeformed zones. Texture evolved in the severely deformedzone is shown in Fig. 8. The texture however, manifested asignificant rotation of nearly 10° from the ideal positions oftexture components and also a noteworthy deviation from itscentro-symmetric peaks. Another observation was the non-uniform development the shear texture components; therelative intensities of texture components in decreasing orderwere J/ �J , D1/D2 and F. The E/ �E components showed arelative weaker intensity.59)

Huang et al.61) and Renk et al.62) reported the textureevolution of pure vanadium and tantalum, respectively, whensubjected to HPT. They found that after subjecting the sampleto a rotation of 180°, the material exhibited a ©212ª ¬ NDfiber near the center and ©101ª ¬ ND texture near the edge.However, upon subjecting the material to 10 turns, ahomogeneous texture ©101ª ¬ ND was observed. This wasattributed to the occurrence of slip on the (110) planes ofBCC materials. It was also observed that Ta developedstrongest shear texture when processed at room temperaturefollowed by annealing. Other BCC materials such as Nb63)

and ¢-Ti6466) have been subjected to HPT, but textureevolution has not been studied.3.2.3 Texture evolution in HPT processed cubic inter-

metallicsStudies have also been done on various other alloys having

a cubic based structure, such as Ni50Mn29Ga21 magneticSMA,67) NiAl and YCu intermetallics having a B2 structureand TiAl intermetallic having an L12 structure.68,69) Chulistet al. reported the appearance of strong oblique cubecomponent (001)©100ª and a weak (110)©001ª F component.In NiAl, YCu and TiAl intermetallics also, the oblique cubecomponent formed in addition to weak shear texturecomponents. All the texture components were found to beslightly deviated from their ideal positions. Except for TiAl,the relative strength of texture components weaken with anincrease in deformation temperature while that of the cubecomponent strengthens. This strengthening of oblique cubecomponent was attributed to the occurrence of discontinuousdynamic recrystallization. In TiAl, neither did the texturestrengthened, nor did it exhibit any clear trend in the volumefraction of components with varying temperatures.69) Overall,all the intermetallics showed a tendency of forming obliquecube component.

Fig. 8 Texture evolution during HPT processing of FCC,51) BCC59) andHCP70) materials.

Texture Evolution in Severe Plastic Deformation Processes 1463

3.2.4 Texture evolution in HPT processed HCP materi-als

3.2.4.1 Texture evolution in materials with c/a ratio closeto ideal

Most of the studies in this category of HCP materials havebeen carried out on Mg. Bonarski et al. studied the textureevolution during HPT under the application of differenthydrostatic pressures,70) and found that under all conditions,a B fibre formed post HPT. However, higher hydrostaticpressures led to a corresponding decrease in texture intensityand also introduced deviation from the ideal position of theB fiber. Torbati-sarraf et al. studied the texture evolutionduring HPT in a ZK60 Mg alloy71) having an initial texturewith strong intensity of basal poles aligned perpendicular tothe extrusion axis, a texture termed as “cylindrical texture”, atypical extrusion texture of HCP materials.32) The basal polesreoriented to align along ND after 5 rotations, while theprismatic and pyramidal poles aligned perpendicular toND.71,72) Texture developed in Mg after a shear strain of³55 under 4GPa pressure is shown in Fig. 8.3.2.4.2 Texture evolution in materials with c/a ratio1.633

Unlike in the case of ECAP, there are quite a few studiescarried out in HCP materials with a c/a ratio greater thanideal. These studies have mostly been carried out at roomtemperature with a few reports also available on HPT atcryogenic as well as high temperature.63,8082) This waspossible because of high hydrostatic pressure associated withthe HPT. Srinivasarao et al. studied texture evolution in CPZn and found that under pure compression, a weak splitbasal texture parallel to ND formed. However, with increasein the number of rotations, the maximum intensity of basalpoles along ND increased. In fact, a proportional increase inthe volume fraction of basal fiber with deformation wasnoticed.82)

3.3 Friction stir processing (FSP)Friction stir processing is a relatively new SPD technique,

originating from friction stir welding, wherein a material issubjected to intense localized plastic deformation to yield aUFG microstructure.83) The details of the process can befound in Ref. 84. The biggest advantage of this techniqueis the spatial tailoring of microstructure and texture in thesample by optimising the processing parameters, such asthe tool rotating and traversing speeds.85) The process isassociated by intense shear deformation, hence developscharacteristic shear textures. Textures in FSP materials aremeasured with respect to the working direction (WD), normaldirection (ND) and transverse direction (TD).

3.3.1 Texture evolution in FSP processed FCC materialsStudies carried out on texture evolution during FSP of Al

alloys8690) have shown that the nugget zone exhibits typicalshear textures of FCC materials. Field et al. reported that inthe Al alloys 1100, 6061 and C485, the material on the topsurface of the plate and adjacent to the FSP tool shoulder, hadtexture characteristic of shear deformation, whereas texture ofthe material below the tool displayed the characteristictexture for uniaxial deformation. A strong ©111ª A fiber waspresent in the shearing region of the weldment, whereas©100ª fiber was observed in the region compressed under theFSP tool.91) Strong ©111ª ¬ ND fiber prevailed even in themid-plane of the weld nugget. The study also revealed that inthe material under the tool pin, partial A fiber forms. Similartextures were observed in 7xxx series of Al alloys.92,93)

However, in the presence of reinforcements in the matrix,textures formed after FSP appear more diffuse because of thehindrance the reinforcement caused to the material flow. Inanother study on AlMg alloy by Suresh et al.,94) it wasobserved that the nugget region tended to exhibit strongð111Þ½1 �21� and (011)[100] texture components, which wereanalogous to simple shear texture components in FCCmaterials. In the thermo-mechanically affected zone (TMAZ),strong ð111Þ½1 �10� and ð112Þ½11 �1� texture components wereobserved with the recrystallized part of TMAZ showingorientations close to ð011Þ½23 �3�.94) Nadammal et al. alsoreported the formation of a predominant component C in theshear texture near to the top surface of the nugget zone andstrong B/ �B and A�2 components towards the middle of thenugget zone in the Al alloys 2024 and 221986,89) (Fig. 9). Inthe AlMg alloy 5086, the C component was found to havethe highest intensity owing to static annealing after thedeformation.86) In multi-pass FSP, the Al alloy 2024exhibited a strong C component, a texture characteristic ofhigh temperature-high strain deformation.88) Apart from Al,HSLA steel was investigated for the evolution texture postFSP, where it was observed that the material flow duringFSP occurred in the high temperature austenite phase. Thisresulted in the formation of ideal shear components of FCCmaterials, which gave rise to a BCC transformation texturethat followed a Kurdjomov-Sachs (KS) orientation relation-ship84) (Fig. 9).

Fig. 9 Texture evolution during FSP processing of FCC (Al alloys,86)

HSLA 65 steel84)) and HCP (Ti)97) materials.

S. Suwas and S. Mondal1464

To conclude, textures formed in FCC materials post FSP,were found to be similar to torsion, ECAP and HPT textures.The cube component strengthened in Al and its alloys. WhileB/ �B components showed strengthening with a decrease inSFE. The positions of texture components coincided withideal simple shear orientations of FCC materials.3.3.2 Texture evolution in FSP processed BCC materials

Various reports are found on the FSP of BCC materials.Some of them include FSP of a near - ¢ alloy Ti-5111 byKipling et al.,95) and FSW of Al-steel joints by Pattersonet al.96) The material was deformed in the high temperatureBCC phase due to the heat generated during FSP, and hencethe D2 ð �1 �12Þ½111� shear component of the high temperatureBCC phase must have formed as a result, because uponcooling the (0002) basal poles aligned perpendicular to thedirection of shear with a 35° tilt from ND.3.3.3 Texture evolution in FSP processed HCP materials

Bahl et al. have studied the texture evolution in cp-Ti postFSP and observed the formation of two distinct texturecomponents:- one with the basal poled tilted by 5° from NDand the other tilted by 38° from ND in the sample processedusing 600 rpm (Fig. 9). The volume fraction of basal polesin this condition was ³6%. When the rpm was increased to1250, the sample showed the weakest texture with thevolume fraction of the basal poles as 4%.97) Young et al.studied the FSP textures of the Mg alloy AZ31 and observedthat a strong basal peak formed at a tilt of ³45° from theprocessing direction. However, similar to Bahl et al., they tooobserved that with increase in the deformation rate, thereappeared a decrease in the texture strength.98,99)

4. Evolution of Texture during Various Plane StrainBased SPD Processes

4.1 Accumulative roll-bonding (ARB)Accumulative roll bonding (ARB) is an SPD processing

technique that does not require any specially designed die orheavy load cells for acquiring heavy forming loads. Instead,ARB can be performed using a conventional rolling mill. Theprocess was invented for severe plastic deformation by Tsujiand his co-workers.100,101) The final product after ARB is inthe form of rolled sheets with UFG microstructure. It is mostcommon to carry out ARB at room temperature due to theease of operation. However, high temperature ARB process-ing is also practiced. Milner et al.102) assumed that there wasno significant drop from the starting temperature as thesamples were very thin, transferred very rapidly from furnaceto the rolling mill (¯2 s) and deformed at high strain rates(20 s¹1).

The evolution of texture during ARB has been wellestablished and was found to be inhomogeneous throughoutthe sample thickness.103,104) The texture components havebeen generally described in terms of rolling texturecomponents. However, there exist subtle differences betweenrolling and ARB textures.4.1.1 Texture evolution in ARB processed FCC materi-

alsStudies on ARB of pure Al up to 8 passes by Huang

et al.105,106) revealed the weakening of texture with increasein number of passes. Locally, most lamella was characterised

by typical rolling texture components.105) Chekhoninet al.107) carried out texture investigations on ARB process-ing of high purity and commercial purity (CP) Al laminates.They observed that the CP-Al layers preferentially exhibittedorientations close to S, Cu and Bs, while the high purity Allayers exhibit orientations close to Cube and Rotated Goss.Studies on ARB of aluminium alloys by Roy et al. showedthat after subjecting two Al alloys, AA5086, AA2219 to 8pass ARB, the intensities of the texture components indecreasing order are: Bs, S, Dillamore, Cube, Cu andGoss.108110) Nearly similar observations were made in otherstudies on Al alloys, for example, like Al0.2%Sc,111)

AA6016 and AA1050112) (Fig. 10). The primary differencebetween the textures formed in Al and Al alloys was thatin the alloys, the strength of the Cu component was thehighest111) and the components exhibited a higher orientationspread from the ideal positions.107,112117) Dominance ofrolling texture components was observed during ARB ofpure Ni by Bhattacharjee et al.,118) with the strength ofcomponents in decreasing order S ¼ Cu ¼ Bs followed byweak intensities of Goss, Cube and RD-rotated cube. Incopper, studies by Takata et al.119) and Suresh et al.103) haveshown that though the textures at the surface and the centerare qualitatively similar, showing components like S, Cu, Bs,Goss, cube, and RD-rotated cube, their relative intensities atboth the sample locations are different. They found that nearto the surface, a strong cube texture formed, with S and Bsorientations being second strongest and Cu being the weakestorientation. On the contrary, in the center of the sample, thedeformation texture components were relatively weaker. Itexhibited a strong cube component with a nearly continuous£-fiber with Bs and Goss components showing similarvolume fractions. The higher intensities of S, Cu and Bs

RD

TDND

Fig. 10 Texture evolution during ARB processing of FCC (Al),115) BCC(Nb)133) and HCP (Mg)136) materials.

Texture Evolution in Severe Plastic Deformation Processes 1465

was attributed to continuous recrystallization (CRX), whichretains the deformation texture components. The high amountof Bs component, as against Cu, was observed because ofimposing large amount of shear. The higher volume fractionsof cube and RD-rotated cube components were observed dueto discontinuous recrystallization (DRX) during inter-passannealing.103)

Other studies were carried out on FCC-BCC nano-lammelar composites, such as the Cu/Nb composites,processed via ARB, which have captured the textureevolution in each phase.120129) Carpenter et al. observedthat in the Cu phase, when the thickness of each layerremained above the nanoscale regime, the C component inrolled FCC materials remained dominant.122) On the otherhand, when the thickness of the layer reduced to

However, ARB of TRC samples developed a noticeable splitbasal texture along RD, while the split basal texture of theARB processed TRC + HT samples was inconspicuous. Thissplit basal texture was attributed to the activation of ©c + aªpyramidal slip.136) Similar observations in Zr were made byCarpenter et al.,132) who carried out ARB processing of ZrNb sheets, where the Zr phase developed a split basal texturealong RD.132) Knezevic et al. revealed that in Zr, large-straindeformation is accommodated by the simultaneous activationof prismatic, basal and pyramidal slip.133) This texture ischaracteristic of rolling textures of HCP materials having ac/a ratio less than ideal.

4.2 Multi-axial forging (MAF)Multi-axial forging is perhaps the easiest of all the SPD

techniques in terms of die design and processing owing to thewide range of processing temperatures (0.10.7 Tm). It is aplane strain deformation process, owing to the constraintsimposed by the die. The process is also called multi-axialcompression, either as a synonym or as a variant. Thesamples to be deformed are always square cross-sectionedbillets having dimensional ratios of either x:y:z = 1:1:1 orx:y:z = 1:1:2. The samples are then inserted into the MAFdie with the longest dimension aligned parallel to thecompression direction, and then compressed to half itsheight. Multi-axial compression (MAC), a variant of MAF,has also been studied in certain cases.137,138) In MAC, no dieis used. Instead the sample is compressed sequentially alongeach axis in a simple forging press. The fundamentaldifference in MAC and MAF lies in the shape of the shearplane and therefore the strain path during processing. Ineach pass of MAF, no deformation occurs in one axialdirection perpendicular to the compression axis due toconstraints imposed by the die. The sample undergoesdeformation under plane strain condition during each pass.The sample is then rotated to realign the longest dimensionparallel to the forging direction (FD). Similarly, deformationis carried out by repeated compression in the x, y and zdirections, which comprises of one cycle. The true strainexperienced by the sample in each press is ³0.8. Thus, theaccumulated true strain imparted to the sample after onecycle is about 2.4. Despite the large strain imparted, thehomogeneity of the strain within the sample is lowercompared to ECAP and HPT.

Textures evolved during MAF are usually compared tothose of rolling textures. Due to the rotation of the sample foraligning the three orthogonal axis along the compression/forging direction, the macroscopic strain path induced ineach pass (comprising of three presses), is different from thecenter of the sample to the surface. As a result, texturesdeveloped in MAF are quite different from those in otherSPD processing.4.2.1 Texture evolution in MAF processed FCC materi-

alsIn spite of a number of investigations carried out on MAF/

MAC of various FCC materials such as Al alloys, Cu alloysand Ni alloys, the evolution of texture during the process hasbeen scarcely documented. Moghaddam et al.139) found thatMAF of 7075 AlZn alloy led to the strengthening of Gossand Cube texture components. Formation of Goss and cube

components was also observed in a 2024 AlCuMgalloy.140142) In another study by Rao et al.,143) it wasobserved from an EBSD micrograph that after 9 cycles ofMAF processing of a 6061 Al alloy, the final texture was(110) ¬ FD. However, due to absence of data along the otherdirections, it was not possible to describe the texturecompletely. The predominance of (110) ¬ FD texture wasalso reported in another study on MAF of CP-Al by Zhao.137)

4.2.2 Texture evolution in MAF processed BCC materi-als

Studies by Bhowmik et al.144) and Gurao et al.145) haveshown that in the texture of MAF processed IF steels, theintensities were concentrated at particular values of ¤1 whenmeasured near to the centre, predominantly having{011}©211ª and {011}©111ª components, while near to thesurface, (001) fiber along ¯ = 0/90° and (111) fibre along¯ = 54.7° in ¤2 = 0° and 45° sections were observed.Additionally intensities were found at {110}©111ª and{110}©011ª locations as well144,145) (Fig. 11). Then again,observations were made in pearlitic steels deformed via twodifferent routes which gave different major texture compo-nents. When ¢-Ti alloy, with a chemical composition of Ti10V3Fe2Al was subjected to 1 pass MAF, post solutiontreatment, it was found that a weak (001)[100] cube typetexture and ð2 �12Þ½ �324� component were formed, whereas,uniaxial compression resulted in the formation of ©111ª ¬ND-£-fiber texture.146) This was observed due to theactivation of multiple slip systems, which resulted from thealteration of strain path in each pass, which led to aweakening of texture.147)

4.2.3 Texture evolution in MAF processed HCP materi-als

A study was done on MAF of magnesium above itsrecrystallization temperature (T > 100°C). The startingmaterial exhibited an axisymmetric texture as can be seenfrom the (0002) pole figure where the basal poles occupy theperipheral positions.148) It was observed that MAF led to theevolution of a split basal texture after 2 passes (Fig. 11).After the 1st MAF cycle, a two-fold split basal componentwas observed along the forging direction in the (0002) polefigure. After the second pass, a four-fold split in the basalcomponents was observed. A gradual weakening of texture

Fig. 11 Texture evolution during MAF processing of BCC (IF steel)145)

and HCP (Mg)148) materials.

Texture Evolution in Severe Plastic Deformation Processes 1467

was noted in the pole figures indicating that the alteration ofcompression axis with each pass resulted in the activation ofdifferent strain paths.

5. Summary

In general, it was observed that deformation by ECAP lead

Table 3 Texture components formed in various FCC, BCC and HCP materials during various SPD techniques.

S. Suwas and S. Mondal1468

to texture evolution dominated by processing route. Unlikeroutes A and C, the absence of a dyad symmetry about TDaxis lead to the development of asymmetric textures in allmaterials deformed by routes BA and BC. Strengthening oftexture components was observed in most cases whendeformed by route A, while strength of shear componentswhen processed by route C was dependent on the passnumber. In FCC metals and alloys, with a decrease in theSFE, a shift from A/ �A to B/ �B components was observed. InBCC materials, predominance of D1, D2 and F componentswas observed. While in HCP materials, strong basal texturewas observed in materials having c/a ratio less than and nearto ideal, with deviation in fiber position appearing as afunction of c/a ratio. In HCP materials with c/a > 1.633 thepredominance of prismatic and pyramidal fibers wereobserved. The shift in texture components away from theirideal orientation position was a function in the internalchannel angle of the ECAP die.

In HPT and FSP processing, position of texturecomponents post deformation of all FCC, BCC and HCPmaterials closely coincided with the torsion textures of therespective materials. In FCC materials, strength of thecomponents was SFE dependent. In HCP materials, texturedevelopment was primarily governed by the c/a ratio. Thec/a ratio dictates the deformation mode. When basal planeswere closely spaced, multiple slip and twinning systems canget activated which yields a B-fiber, whereas, when basalplanes are less closely spaced, (the case of materials withc/a > 1.633), only basal and twinning systems are respon-sible for accommodating the applied deformation, leading tothe formation of prismatic and pyramidal fibers.

In ARB and MAF, textures were found to be similar torolling textures of FCC, BCC and HCP materials. In cases ofeasier recrystallization, FCC and BCC materials, predom-inantly exhibit cube textures, whereas, in other cases, rollingtexture components are exhibited. HCP materials, reveal atexture analogous to HCP rolling textures with a split basalfiber texture, which is characteristic of the c/a ratio.

In conclusion, it was established that the texture evolutionwas path dependent, and that the conventional SPDprocessing techniques gave rise to a texture either similarto simple shear textures or rolling textures. Also, there stillappears to be a wide scope for studying texture evolutionduring MAF of FCC materials, and FSP and MAF of BCCand HCP materials. Texture components formed in variousFCC, BCC and HCP materials during various SPD process-ing techniques are compiled in Table 3.

Acknowledgement

SS gratefully acknowledges the interactions with Profs.Laszlo Tóth (University of Lorraine, France), WernerSkrotzki (TU Dresden, Germany), Günter Gottstein (RWTHAachen, Germany), H.-G. Brokmeier (TU Clausthal,Germany), Drs. B. Beausir and J.-J. Fundenberger (Uni-versity of Lorraine, France). A large part of the SPD workwas carried out during the stay of one of the authors (SS) inFrance and Germany, for which he acknowledges thefunding/fellowship from Labex DAMAS, University ofLorraine, France and Alexander von Humboldt Foundation,

Germany. The review work was initiated during author’s (SS)stay at TU Dresden as Frederich Wilhem Bessel ResearchAwardee from Alexander von Humboldt foundation,Germany for which he is grateful to the foundation. Theauthors would like to thank K.U. Yazar, and R.J. Vikram,their help in manuscript preparation.

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