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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
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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
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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
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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
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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
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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
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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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