Substructure and Texture Evolution in an Annealed AluminumAlloy at Medium Strains

QUANCANG MA, WEIMIN MAO, BIN LI, PAUL T. WANG,and MARK F. HORSTEMEYER

Substructures are widely observed in plastic deformation of aluminum alloys and are ofpractical significance, but studies on characterization of substructure in a scale much larger thanthe grain size and how it affects texture evolution are still lacking. In the current study, weperformed channel die compression on an annealed AA1100 aluminum sheet along the normaldirection (ND) at medium strains at room temperature. The microstructure and texture werecharacterized by electron backscattered diffraction. Texture evolution was simulated by incor-porating octahedral {111}h011i and nonoctahedral {hkl}h011i slip systems in the visco-plasticself-consistent model. The rotation axes and the misorientation angles for the deformationtexture variants were calculated. The results show that substructuring proceeded in all thetexture components but in a heterogeneous manner. The h001i||ND texture presents high-angleboundaries (HABs) of 15 deg through 30 deg without rotation axis clustering and almost noextra high-angle boundaries (EHABs) of 30 deg through 60 deg; while the HABs and theEHABs coexisted in the h011i||ND and the h112i||ND textures. The rotation axes of the EHABspreferentially clustered at h011i and h111i. Under plain strain compression, multiple defor-mation texture variants created by substructuring interwove with each other, resulting in theEHABs with rotation axes clustering. In contrast, the HABs generated by substructuring viadislocation mechanisms showed no rotation axes clustering. Substructuring moderated thetexture component intensity and randomized orientations, and resulted in fluctuation of thea-fiber texture and hindered the increase in the rate of the copper texture development.

DOI: 10.1007/s11661-013-1776-x� The Minerals, Metals & Materials Society and ASM International 2013

I. INTRODUCTION

DURING straining of high stacking fault energy(SFE) face-centered cubic (fcc) metals, e.g., aluminumand its alloys, significant change in microstructureknown as substructuring or grain subdivision is com-monly observed.[1–4] Substructuring mechanisms includedislocation boundaries, noted as dense dislocation walls(DDWs), microbands (MBs), geometrically necessaryboundaries (GNBs), or incidental dislocation bound-aries (IDBs); between these boundaries are dislocationcells inside the original grains.[1–8] The misorientationbetween substructural boundaries increases with strain,and eventually these boundaries evolve into high-angleboundaries (HABs) (>15 deg) that subdivide the origi-nal grains. Extensive studies have been carried out ondislocation boundary alignment on active slip planes

and/or the planes with maximum stress and the dislo-cation configurations with respect to original grainorientation.[1–6,9,10] Although whether DDWs/MBs lie inactive {111} slip plane or the planes with maximumstress still remains a debatable issue, it is clear thatdislocation boundaries subdivide the original grain, andthe misorientation across the dislocation boundaryincreases with the increasing strain.[5,6,11] Hughes andHansen[12] investigated the HAB mechanism by trans-mission electron microscopy (TEM) in heavily rollednickel and aluminum alloys (rolled to 70 pct through 98pct). They pointed out that HABs developed at highstrains were mainly induced by two mechanisms: themicrostructure mechanism and the texture mecha-nism.[12] Also using TEM, Sun et al.[13] studied forma-tion of HABs by substructuring in severely deformedaluminum via equal channel angular extrusion (ECAE).Delannay et al.[14] investigated substructuring by elec-tron backscattered diffraction (EBSD) in a rolled,commercial purity aluminum at 40 pct strain, and theyfound that the misorientation distributions are differentin differently oriented parent grains.Annealed commercial purity aluminum usually devel-

ops strong cube texture before cold-rolling is performed.Practically several reductions in rolling are needed,before the final thickness is reached. Hence, under-standing how microstructure and texture evolutionduring cold-rolling of strong cube-textured aluminum

QUANCANG MA, Postdoctoral Associate, BIN LI, AssistantResearch Professor, and PAUL T. WANG, Deputy Director, are withthe Center for Advanced Vehicular Systems (CAVS), Mississippi StateUniversity, Starkville,MS39762.Contact e-mail: qma@cavs.msstate.eduWEIMINMAO, Professor, is with the Department ofMaterials Science,UniversityofScienceandTechnologyBeijing,Beijing100083,P.R.China.MARK F. HORSTEMEYER, Chair Professor, is with the Center forAdvanced Vehicular Systems (CAVS), Mississippi State University, andProfessor with the Department ofMechanical Engineering atMississippiState University, Starkville, MS 39762.

Manuscript submitted January 30, 2013.Article published online May 9, 2013

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alloy is of practical significance for mechanical propertycontrol. At medium strains, substructuring proceeds anddominates microstructure and texture evolution.[7] How-ever, there are only limited data in the literatureregarding how substructuring proceeds in a strongcube-textured polycrystalline aluminum alloy and howit affects texture evolution at medium strains, typically 40pct through 70 pct strain in cold-rolling. Chowdhury[15]

studied the texture evolution in cube-oriented polycrys-talline aluminum during cold rolling, but details ofsubstructuring were not reported.

Previous TEM studies have provided valuable insightsinto the nature of the substructure development, but thestatistics is poor because of the tiny sampling volume.Particularly, the correlation between substructuring andtexture evolution needs further characterizations on ascale that is much larger than the average grain size. Themain purpose of the current study is to investigate thesubstructuring mechanisms at medium strains in anannealed aluminum alloy deformed under plane straincompression, and its effect on texture evolution bystatistical measurements and calculations using EBSD.Texture evolution was also simulated and the predic-tions were compared with the measured results.

II. EXPERIMENTAL

A commercial aluminum AA1100 sheet (0.15 Si, 0.53Fe, 0.07 Cu, 99.2 Al, wt pct) with a thickness of 6 mmand H14 status was selected as the experimentalmaterial. Samples with dimensions 10(RD) 9 6(ND) 94.5(TD) mm were cut from the sheet (RD—rollingdirection, ND—normal direction, and TD—transversedirection). The as-cut sample was heated up to 673 K(400 �C) in a furnace and held for 1 hour. The samplewas immediately quenched in water to preserve themicrostructure and served as the initial state. Planestrain compression in a channel die at room temperaturewas then conducted on INSTRON 5869 at a strain rate10�3 s�1 to three intermediate strains: 40 pct, 51 pct,and 67 pct, respectively. A TEFLON lubricant wasapplied to the sample surfaces to mitigate friction duringdeformation. The loading direction was along the ND,the flow direction was along the RD, and flow in the TDwas constrained.

The compressed sample was cold mounted and theRD–ND cross section (the longitudinal plane) of thesample was polished with a series of SiC papers down togrit number 4000. Then the sample was polished with analumina suspension (0.05 lm particle size mixed withethylene glycol). Finally, the sample was polished in avibratory polisher using a colloidal silica suspension(0.02 lm) for 4 hours. EBSD scans were then performedon the polished sample without etching in a high-resolution field emission gun scanning electron micro-scope (ZEISS SUPRA-40 FEG-SEM) with TSLOIM 5.3EBSD data collection and analysis software package. Foracquisition of texture, a large area of 1200 9 1200 lmwith a step size of 2.0 lm in the center of RD–ND crosssection was scanned by EBSD; for microstructuralanalysis, a fine step size 0.2 lm was used in EBSD scan.

III. RESULTS

Figure 1 presents the EBSD inverse pole figure (IPF)maps of the AA1100 samples at strains 0 pct, 40 pct, 51pct, and 67 pct. The initial microstructure exhibitsroughly equiaxed grain structure with a strong cubetexture (Figure 1(a)). Owing to the strong cube texture,a considerable portion of the grain boundaries (GBs) arelow angle GBs (<15 deg) (Figure 1(a)). The equiaxedgrains evolve into pancakes at 40 pct strain. Inside thepancakes, substructure and subboundaries quicklydevelop, especially at strains of 51 pct and 67 pct(Figures 1(b) through (d)). Grain size distributions atdifferent strains were obtained based on EBSD scans(Figure 1). The grain size was measured as the diameterof an equivalent circular area using the EBSD analysissoftware. The grain size distribution at different strainsis presented in Figure 2. In the initial sample, largegrains greater than 50 lm predominate. At strains of 40pct and 51 pct, the grain structure is significantly refined,and the area fraction of the small grains (<20 lm)increases. Both small and large grains coexist, but thearea fraction of the grains larger than 50 lm decreases.At 67 pct strain, small grains dominate the microstruc-ture with few large grains left. From Figures 1 and 2,substructuring can be clearly observed.The orientation distribution function (ODF) of the

AA1100 based on the EBSD scans (Figure 1), and allthe ODFs obtained from modeling in the current studywere calculated using the texture software MTEX[16]

with a Kernel half-width of 5 deg. Since u2 = 45 deg,ODF section is important for fcc alloys where the maintexture components can be observed; the u2 = 45 degODF sections of the AA1100 samples at different strainsare presented in Figure 3. The initial, annealed AA1100indeed has a strong cube texture. In addition, weakcopper and brass texture components can also beobserved (Figures 3(a) and (e)). At 40 pct strain, brass,Goss, and copper textures intensify, while the cubecomponent quickly weakens (Figure 3(b)). At 51 pctstrain, the brass and the Goss textures almost disappear;meanwhile, the copper texture slightly increases, and thecube texture continues to decrease (Figure 3(c)). At 67pct strain, the copper texture appears to be steady, butthe Goss and the brass textures increase again, and theh011i||ND texture (a-fiber) spreads widely, whereas thecube texture keeps decreasing slowly (Figure 3(d)).Owing to the presence of the substructures, the

original grain boundaries have lost the original, identi-fiable morphology. Rolled aluminum usually developstypical texture components, e.g., brass, copper, and Scomponents, while each component may have a fewvariants, for instance, brass component may have twovariants, i.e., brass 1 and brass 2. In order to investigatethe substructure and texture variants evolution, somefiber textures are highlighted because different texturevariants belong to the same fiber texture. To investigatethe behavior of the cube-oriented substructures, wehighlight the h001i||ND fiber texture component(denoted as h001i texture hereafter) with a tolerance15 deg such that other texture components are excluded,as shown in Figure 4. The corresponding correlated

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misorientation (misorientation between one orientationand its neighbors) angles and rotation axis–angle pairsare displayed in Figure 4(d). Since substructuring ismainly caused by HABs, we only show the misorienta-tion >15 deg, and an example of the rotation axisdistribution at rotation angle of 20 ± 2.5 deg, 50 ±2.5 deg, and 60 ± 2.5 deg at the strain of 67 pct isshown in Figure 4(d).

In the h001i texture, the number fraction of the highangle misorientation (>15 deg) increases with theincreasing strain (Figure 4(d)). Meanwhile, the h001i

texture decomposes, and the fraction of remaining h001itexture decreases (Figures 4(a) through (c)). The HABsin the h001i texture are mainly located in the misorien-tation range of 15 deg through 30 deg, and almost noextra high-angle boundaries (EHABs) >30 deg exist inthe h001i texture. The rotation axis distribution at20 deg shows a nearly uniform distribution with nopreferential clustering (Figure 4(d)).Since the brass {011}h211i and the copper {112}h111i

textures are two important orientations in deformedaluminum,[17] investigation of the behavior of the

Fig. 1—EBSD IPF maps of the AA1100 samples under plain strain compression at different strains: (a) at 0 pct strain, a strong cube texture isobserved in the initial state; (b) at 40 pct strain, pancake grains appear, and the cube texture partially transforms to other orientations; (c) at 51pct strain, substructuring proceeds in all grains; and (d) at 67 pct strain, significant substructuring can be observed. The black solid lines repre-sent the boundaries with misorientation angle>15 deg. The typical color triangle represents the ND. EBSD step size: 2 lm (Color figure online).

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h011i||ND and h112i||ND fiber texture components(designated as h011i and h112i textures, respectively,hereafter) is critical for the relationship between sub-

structuring and texture. The h011i and the h112i textureswere highlighted similarly, and the corresponding highangle misorientations and the rotation axis–angle dis-tributions are plotted in Figures 5 and 6, respectively.Similar to the h001i texture, the number fraction of

the HABs of 15 deg through 30 deg increases in theh011i and the h112i textures with increasing strain(Figures 5(d) and 6(d)). The rotation axes in the rangeof 15 deg through 30 deg show no preferential clusteringeither (Figures 5(d) and 6(d)). However, in contrast tothe h001i texture, the EHABs of 30 deg through 60 degappear and its number fraction increases with increasingstrain and the rotation axes preferentially cluster ath111i and h011i (Figures 5(d) and 6(d)).

IV. DISCUSSION

A. Substructuring Mechanisms and Heterogeneity

From the EBSD data (Figures 4 to 6), we observe thatthe average misorientation increases with increasingstrain, which is consistent with the TEM results by Liuet al.[1] The number fraction of the HAB of 15 degthrough 30 deg increases with increasing strain in theh001i, h011i and h112i textures (Figures 4(d), 5(d), and6(d)). Since the HABs of 15 deg through 30 deg increasewith strain, it is consistent with the scaling hypothesis

Grain size in diameter (microns)

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50 100 150 2000

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Fig. 2—Grain size distribution of the AA1100 samples at differentstrains under plane strain compression based on EBSD scans inFig. 1. The fraction of the large grains (>50 lm) decreases, whereasthe fraction of the relatively fine grains (<20 lm) increases withstrain increasing from 40 pct to 67 pct. At 67 pct strain, thereare few grains >50 lm, indicating that significant substructuringoccurred from 40 pct to 67 pct strain (Color figure online).

(a) (b) (c)

(d) (e)

Fig. 3—u2 = 45 deg ODF sections of the AA1100 samples at different strains under plain strain compression. (a) at 0 pct strain, a strong cubetexture and weak copper and brass texture coexist in the initial state; (b) at 40 pct strain, a strong brass, copper, and Goss texture appear; (c) at51 pct strain, the brass and Goss texture abnormally weaken again, and the cube texture continues to decrease; (d) at 67 pct strain, the brass,Goss, and a-fiber texture intensify again, while the copper texture remains nearly constant and the cube texture still decreases slowly; and (e) themain texture components positions in u2 = 45 deg ODF section. Levels:1,2,4,6,8,10, and 12 (Color figure online).

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(average misorientation angle increases with increasingstrain) for misorientation distributions of dislocationboundaries in deformed fcc alloys based on extensiveTEM studies.[18,19] It was observed that the h001i grainsmainly exhibit cell structures and perhaps some planardislocation walls, but other grains with different orien-tations show both cells and extended planar boundariesin compressed aluminum.[10] The misorientation angleinduced by the dislocation mechanism may reach up to15 deg through 30 deg.[12] Activation of multiple slipsystems and grain interaction may partially stabilize theh001i texture during channel die compression.[20–22] Thisexplains why the h001i texture only shows the HABs of15 deg through 30 deg.

The h011i and the h112i textures present both HABs of15 deg through 30 deg and EHABs of 30 deg through

60 deg (Figures 5 and 6), indicating that the substructur-ing is inhomogeneous from the perspective of texture. Theh011i and the h112i textures are located in the orientationspace from a- to b-fiber in fcc aluminum. It is known thatorientations tend to transform from brass via S to copperin aluminum during rolling or plane strain compres-sion.[23,24] The h011i and the h112i textures experiencesubstructuring as well, but can only transform to otherdeformation texture variants with random orientations.Thus, different deformation texture variants are involvedin the original texture component and impinge on eachother where the EHABs are present (see analysis belowfromFigure 7). The orientation dependence of the EHABdistribution could be one reason why the misorientationscaling hypothesis shows a marked deviation when alldislocation boundaries are considered.[12,19] Thus, the

Fig. 4—The highlighted h001i||ND fiber texture component with a tolerance of 15 deg at three strains of (a) 40 pct; (b) 51 pct; (c) 67 pct; and (d)the corresponding correlated large misorientation angle (>15 deg) distribution and the rotation axis distribution at 20 ± 2.5 deg, 50 ± 2.5 deg,and 60 ± 2.5 deg at 67 pct strain. The h001i texture decomposes with increasing strain, and cube-band clearly appears at 51 pct and 67 pctstrains. The fraction of the boundaries of 15 deg through 30 deg increases with strains and almost no 30 deg through 60 deg boundaries can beobserved. The rotation axes at 20 deg at 67 pct strain show no preferred clustering. Other orientations in (a) through (c) are excluded (blackareas). EBSD step size: 2 lm (Color figure online).

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preferential occurrence of the EHABs in the h011i andh112i textures with preferred rotation axes (Figures 5(d)and 6(d)) is attributed to the multiple subdivided defor-mation texture variants via substructuring.

In Figure 7 of the EBSD results, substructures in thesubdivided grains and point-to-point misorientationprofiles along four lines at 40 pct and 67 pct strainsare presented. At 40 pct strain, substructuring wasconfirmed by the presence of many subboundaries>15 deg (Figures 7(a) and (b)). Meanwhile, few sub-boundaries with misorientation angle>30 deg are pres-ent at 40 pct strain, but more subboundaries withmisorientation angle >30 deg are seen at 67 pct strain(Figure 7). In general, the substructures and subboun-daries are heterogeneous for differently oriented grains.For example, more cells can be observed in the h001ioriented grains, but cell boundaries and more MBs

coexist in the h011i oriented grain; in contrast, othergrains with different orientations show less substructures(Figure 7(a)). Inside the individual grains, the subdivi-sion is also heterogeneous. For example, along the grainboundary (Line B), dense substructures appear, but notin the interior of the grain (Figure 7(a)). At 67 pct strain,most grains have undergone subdivision to a largeextent, and fine substructures and more HABs andEHABs appear in both ND and RD directions(Figures 7(c) and (d)). Figures 7(b) and (d) confirm thepresence of the HABs and the EHABs mainly created bythe substructuring, not by the original grain bound-aries.[25] The subdivided structure appears inhomoge-neous as well at 67 pct strain. For instance, thecube-band remains, among other types of substructures,either align ~40 deg with RD or parallel to RD(Figure 7(c)). The heterogeneity of substructuring could

Fig. 5—The highlighted h011i||ND fiber texture with a tolerance of 15 deg at three strains of (a) 40 pct; (b) 51 pct; (c) 67 pct; and (d) the corre-sponding correlated large misorientation angle (>15 deg) distribution and the rotation axis distribution at 20 ± 2.5 deg, 50 ± 2.5 deg, and60 ± 2.5 deg at 67 pct strain. Both HABs of 15 deg through 30 deg and EHABs of 30 deg through 60 deg increase with strain. The rotation axisdistribution shows no preferred clustering at 20 deg, but clusterings at h111i and h011i at misorientation angles of 50 deg and 60 deg. Other ori-entations in (a) through (c) are excluded (black areas). EBSD step size: 2 lm (Color figure online).

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be ascribed to the combination of original grainorientation and grain interaction.[4,26,27] The EHABmechanism caused by the subdivided deformationtexture variants will be further discussed in Sec-tion IV–C.

B. Substructuring Effect on Texture

During substructuring, texture evolves accordingly. At40 pct strain, a significant part of the cube texturetransforms to other components including brass, copper,andGoss (Figures 3(a) and (b)). The quick decrease in thecube texture at 40 pct strain may originate from threereasons: (1) the initial large grain size facilitates substruc-turing[7]; (2) high density of low-angle GBs among thecube-oriented initial grains favors the subdivision of thecube texture (Figure 1(a)); and (3) the most favorable slip

systemmay dominate the plastic deformation at the earlystage (<40 pct strain) and result in quick rotation awayfrom the cube texture. The Goss texture is a transitioncomponent between the cube and the brass texture, thus itis also enhanced at 40 pct strain (Figure 3(b)). However,as the strain increases to 51 pct, the brass and the Gossdisappear, whereas the copper slightly increases and thecube slightly decreases (Figures 3(b) and (c)). Strikingly,the brass and the Goss slightly enhance again at 67 pctstrain (Figure 3(d)). Hence, a fluctuation in the a-fibertexture is observed (Figure 3).Subdivision continues in the deformation texture

components. A single brass crystal was proven stable inchannel die compression, because the two slip systemshave the same maximum Schmid factor, and are simul-taneously activated and cancel out the brass crystalrotation.[28,29] However, the activated two slip systems

Fig. 6—The highlighted h112i||ND fiber texture with a tolerance of 15 deg at three strains of (a) 40 pct; (b) 51 pct; (c) 67 pct, and (d) the corre-sponding correlated large misorientation angle (>15 deg) distribution and the rotation axis distribution at 20 ± 2.5 deg, 50 ± 2.5 deg, and60 ± 2.5 deg at 67 pct strain. Both HABs of 15 deg through 30 deg and EHABs of 30 deg through 60 deg increase with strain. The rotation axisdistribution shows no preferred axes clustering at 20 deg but strong clusterings at h111i and h011i at misorientation angles of 50 deg and 60 deg.Other orientations in (a) through (c) are excluded (black areas). EBSD step size: 2 lm (Color figure online).

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lead to strong shear strain eRD–TD rendering the brass asunstable.[14,28–30] Thus, in the polycrystalline case, acti-vation of only two slip systems in brass is actuallyimpossible. Under plane strain compression, if parts of abrass 1 variant are subdivided into brass 2 or S variants[29]

or vice versa, the signs of shear strain eRD–TD of brass 1and brass 2 or S variants are opposite, and the total eRD–

TD is minimized. Similar to the subdivision of the brass, acopper texture would subdivide into copper 1, copper 2,

or S variants to eliminate the large shear strain eRD–ND

due to the activation of four slip systems in the copper-oriented grain.[30–32] Since both brass 1 and brass 2 belongto the brass texture, the brass texture still has a strongintensity at 40 pct strain (Figure 3(b)). However, withincreasing strain, brass 1 and brass 2 continue tosubdivide. During the breakup of the brass into nonbrassvariants (e.g., S texture variants),[29] the brass texturequickly weakens because of the substructuring

Distance (micron)

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Fig. 7—Substructuring heterogeneity in the AA1100 samples in EBSD IPF maps. (a) at 40 pct strain, substructuring occurs and propagates inboth intragranular and intergranular heterogeneous styles; (b) the point-to-point misorientation profile of Line A and B shows >15 deg bound-ary within the h011i||ND grain and the h001i||ND grain; (c) at 67 pct strain, extensive substructuring proceeds, and the thin cube-bands sand-wich other substructures; (d) the point-to-point misorientation profile along Line C and D shows >30 deg EHABs inside an original grain. Thethin solid black lines in (a) and (c) represent the boundaries with misorientation angle>5 deg. The color triangle represents the ND. EBSD stepsize: 0.2 lm (Color figure online).

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(Figure 3(c)). The copper texture still slowly increasesbecause of the transformation from S to copper; mean-while, substructuring proceeds in the copper texture.Thus, substructuring slows down the rate of coppertexture increase. Otherwise, copper texture would rapidlyincrease with strain (Figure 8). The origin of there-enhancement of the a-fiber with a wide spread at 67pct strain (Figure 3(d)) may result from further deforma-tion that causes the substructures to re-aggregate alongthe a-fiber. The a-fiber texture fluctuates at strains 40 pct,51 pct, and 67 pct presumably resulting from the cycle ofsubstructuring with strain: substructuring–texture weak-ening–dislocation accumulation–texture strengthening–substructuring. Thus, substructuring is able to moderatethe deformation texture components’ intensity (e.g.,copper) and results in fluctuation in a-fiber texture. In

fact, the volume fraction of typical deformation textures,including the brass and the Goss textures measured byX-ray diffraction, in a rolled aluminum alloy wasobserved to markedly fluctuate in the reduction from 20pct to 75 pct.[5] Only substructuring can rationalize this a-fiber texture fluctuation. This texture fluctuation wouldsubside when the substructure size reduces down to acritical size at high strains where further substructuringmay no longer occur.[7] The slow decrease in the cubetexture from 40 pct to 67 pct strain (Figures 3(b) through(d)) confirms that the cube texture indeed is partiallystable during compression in aluminum.[20–22,33]

To further understand the effect of substructuring onthe texture, we simulated the texture evolution using thevisco-plastic self-consistent (VPSC) model.[34,35] Nosubstructuring was considered in the modeling, and

(a) (b) (c)

(d) (e) (f)

(g) (h) (i)

Fig. 8—VPSC simulation of texture evolution starting from the initial texture of the AA1100: (a) a {111} pole figure obtained from EBSD scan(Fig. 1(a)). (b) The initial 500 orientations are spatially discretized from (a), serving as the starting orientations. (c) The calculated initial ODFu2 = 45 deg section from (b). (d) through (f) VPSC predictions of textures at (d) 40 pct, (e) 51 pct, and (f) 67 pct strain in u2 = 45 deg ODFsection based on twelve {111}h011i octahedral slip systems. (g) through (i) Simulation results by combining 12 octahedral slip systems{111}h011i, six {100}h011i slip systems, six {110}h011i slip systems, and twelve {112}h011i slip systems at (g) 40 pct, (h) 51 pct, and (i) 67 pctstrain. The ratio of the critical resolved shear stress of {111}h011i and {hkl}h011i is set to 1:3. It can be observed that VPSC overpredicts thecopper and cube textures compared with the measured textures in Fig. 3. Additional nonoctahedral slips only slightly slow down the increaserate of copper texture. ODF intensity levels: 1, 2, 4, 6, 8, 10, 12, 14, and 16.

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the results were compared with the measured texture.The 500 orientations (Figure 8(b)) were spatially dis-cretized from the initial texture (Figures 3(a) and 8(a))and served as the starting orientations. Since the channeldie compression was performed at room temperature,the 12 standard octahedral {111}h011i slip systems maydominate the deformation, while nonoctahedral {100}-,{110}-, and {112}h011i slips may also activate in aminimal way at ambient temperature.[36–39] Thus, weused twelve {111}h011i slip systems and combination of12 octahedral slip systems and six {100}h011i slipsystems, six {110}h011i slip systems, and twelve{112}h011i slip systems with a 3:1 ratio of the criticalresolved shear stress (CRSS) of nonoctahedral andoctahedral slip (since the {111} planes are the mostdensely packed) in the VPSC simulation, respectively. Inthe VPSC modeling, each orientation interacts with auniform matrix represented by all the orientations. Thesimulated textures at strains of 40 pct, 51 pct, and 67 pctare presented in Figures 8(d) through (i). From theVPSC simulations, the typical deformation texturecomponents, such as copper, brass, Goss, and cubecan be successfully predicted based on 12 octahedral{111}h011i slips and combination of octahedral andnonoctahedral {100}h011i, {110}h011i, and {112}h011islips (Figures 8(d) through (i)). The VPSC simulationresults show that the copper texture quickly increases,and the cube texture continuously decreases, while thebrass texture slightly decreases with strain (Figures 8(d)through (i)). The inclusion of the nonoctahedral slipsystems does not change the overall texture, but slightly

slows down the rate of increase in the copper texture andleads to an insignificant decrease in the brass texture(Figures 8(d) through (i), and 3(e)). Compared with themeasured textures (Figures 3(b) through (d)), the VPSCsimulation overpredicts the copper and the cube inten-sities (Figures 8(d) through (i)). No a-fiber fluctuation isobserved in the simulation. The simulation confirms thatsubstructuring plays a vital role in the texture evolutionat intermediate strains, and more refined models[8] areneeded to capture the texture evolution.

C. Substructuring Effect on Extra High-Angle Boundary

Because substructuring continues during deforma-tion, the deformation texture components evolve todifferent variants. These texture variants in compressedfcc aluminum for the Bunge Euler angles {u1, F, u2}

[40]

are the following: brass 1 {35,45,0}, brass 2 {325,45,0},copper 1 {90,35,45}, copper 2 {270,35,135}, S 1{60,32,65}, S 2 {300,148,245}, S 3 {120,148,245}, S 4{240,32,65}, and Goss {0,45,0}. In a fine EBSD scanningmap at 67 pct strain (Figure 9), these different defor-mation texture variants and the cube texture in fccaluminum were colored by their specific orientations(Euler angles) with a tolerance of 20 deg. Substructureswith other orientations (no deformation texture variantsand the cube) shown in white are considered as randomcomponents (Figure 9). The orientation map at 67 pctstrain (Figure 9) confirms that the different deformationtexture variants are created by substructuring. Thesesubdivided deformation texture variants interweave

Fig. 9—Deformation texture variants interweave during substructuring at 67 pct strain. The colored rectangles are the orientation legends thatrepresent the specific texture components within 20 deg of the ideal orientations. White color represents other orientations, i.e., random. Thesolid black lines are the boundaries with a misorientation angle>5 deg. EBSD step size: 0.2 lm (Color figure online).

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inside an original grain. Moreover, the random texturecolored in white in Figure 9 appears to be significant asa result of substructuring. Orientation randomization bysubstructuring leads to the lower texture intensity thanin the VPSC modeling results (e.g., copper and cubetextures) (Figures 3(b), (c), 8(d) through (i)).

Theoretically, there are 36 different neighboring pairsamong the deformation texture variants: Goss, brass 1,brass 2, copper 1, copper 2, S 1, S 2, S 3, and S 4(Table I). The misorientation and the correspondingrotation axis of the 36 neighboring pairs were calcu-

lated by MTEX[16] (Table I) and are plotted inFigure 10. All of the misorientation angles of the 36texture variants pairs are larger than 15 deg, and manyrotation axes are equivalent and preferentially locate inh011i-h111i in the IPF (Table I and Figure 10). Notethat the misorientation angle (i.e., the disorientationangle) is the smallest misorientation angle among the 24variants in fcc aluminum.[41] There is a large numberfraction of misorientation angles between 30 and60 deg. In addition, the corresponding rotation axespreferentially cluster at h111i-h011i (Table I andFigure 10). The calculated misorientation and rotationaxis distributions are consistent with the measuredEHAB misorientation and the corresponding rotationaxis distribution in the h011i and the h112i textures(Figures 5(d), 6(d), and 10). The number fraction ofthese texture variants increases with strain because offurther substructuring. This is the reason why thefraction of EHAB of 30 deg through 60 deg misorien-tation increases with strain in the h011i and the h112itextures (Figures 5(d) and 6(d)). The calculated and theexperimental results confirm that the EHABs indeedresult from substructuring in the deformation texturecomponents. According to the calculation and experi-mental observations, the HABs of 15 deg through30 deg mainly result from substructuring by dislocationmechanism,[12] whereas the EHABs of 30 deg through60 deg are largely generated by the impingement ofdifferent deformation texture components, i.e., throughmechanism of deformation banding. It is conceivablethat if the deformation mode changes, e.g., simpletension or torsion, then the deformation textures willchange accordingly. Thus, substructuring may proceeddifferently. Moreover, the effect of substructuring maybecome negligible after high strain deformation,because there is little subdivision at high strains.[7]

Therefore, typical deformation textures may normallydevelop and enhance after large strains.

Table I. Misorientation Angles and Rotation Axes of Different 36 Neighbor Pairs Among the Ideal Deformation Texture Variants

in fcc Aluminum Under Plane Strain Compression

No. Neighbor Pair Rotation Angle-Axis d deg [u v w] No. Neighbor Pair Rotation Angle-Axis d deg [u v w]

1 Goss–brass 1 35.0 [0 1 1] 19 brass 2-S 2 20.7 [7 9 �3]2 Goss–brass 2 35.0 [0 �1 �1] 20 brass 2-S 3 20.7 [7 9 �3]3 Goss–copper 1 55.0 [�1 1 0] 21 brass 2-S 4 51.2 [96 100 17]4 Goss–copper 2 55.0 [1 1 0] 22 copper 1–copper 2 60.0 [1 �1 �1]5 Goss–S 1 44.7 [7 4 5] 23 copper 1–S 1 17.5 [�9 �4 �3]6 Goss–S 2 44.7 [7 4 5] 24 copper 1–S 2 17.5 [�9 �4 �3]7 Goss–S 3 44.7 [7 4 5] 25 copper 1–S 3 52.0 [46 95 100]8 Goss–S 4 44.7 [7 4 5] 26 copper 1–S 4 52.0 [46 95 100]9 brass 1–brass 2 60.0 [1 �1 1] 27 copper 2-S 1 52.0 [46 95 100]10 brass 1–copper 1 35.7 [100 76 �5] 28 copper 2-S 2 52.0 [46 95 100]11 brass 1–copper 2 35.7 [76 100 �5] 29 copper 2-S 3 17.5 [�9 �4 �3]12 brass 1–S 1 20.7 [7 9 �3] 30 copper 2-S 4 17.5 [�9 �4 �3]13 brass 1–S 2 51.2 [96 100 17] 31 S 1-S 2 33.0 [�9 �9 �5]14 brass 1–S 3 51.2 [96 100 17] 32 S 1-S 3 60.4 [�2 3 3]15 brass 1–S 4 20.7 [7 9 �3] 33 S 1-S 4 40.1 [3 5 �3]16 brass 2-copper 1 35.7 [76 100 5] 34 S 2-S 3 40.1 [3 5 �3]17 brass 2-copper 2 35.7 [100 76 5] 35 S 2-S 4 60.4 [�2 3 3]18 brass 2-S 1 51.2 [96 100 17] 36 S 3-S 4 40.1 [�9 �9 �5]

Fig. 10—The distribution of the misorientation angle and rotationaxis of the 36 neighbor pairs among the ideal deformation texturevariants in fcc aluminum. The texture variants include: Goss{0,45,0}, brass 1 {35,45,0}, brass 2 {325,45,0}, copper 1 {90,35,45},copper 2 {270,35,135}, S 1 {60,32,65}, S 2 {300,148,245}, S 3{120,148,245}, and S 4 {240,32,65}. The results show that most ofmisorientation angles are located in 30 deg through 60 deg, and therotation axes preferentially cluster at h111i and h011i.

4414—VOLUME 44A, SEPTEMBER 2013 METALLURGICAL AND MATERIALS TRANSACTIONS A

V. CONCLUSIONS

Channel die compression on an annealed cube-tex-tured AA1100 aluminum alloy was conducted at strainsof 40 pct, 51 pct, and 67 pct. The microstructure and thetexture of the compressed samples were characterized byEBSD. The texture evolution was also simulated bymeans of VPSC model incorporating 12 standardoctahedral {111}h011i slips and a combination ofoctahedral and nonoctahedral {hkl}h011i slips, respec-tively. The misorientation angle and the rotation axis forthe deformation texture variants were also calculated.The following conclusions can be drawn:

1) Substructures developed in the annealed AA1100 alu-minum alloy in the strain range of 40 pct through 67pct. Grain refinement and orientation randomizationtook place by substructuring. Both stable and unstableorientations experienced substructuring during defor-mation. Substructuring proceeded in a heterogeneousmanner because of original grain orientation and graininteraction. The HABs of 15 deg through 30 deg andthe EHABs of 30 deg through 60 deg at mediumstrains were mainly created by substructuring.

2) The HABs of 15 deg through 30 deg did not haverotation axes clustering in h001i, h011i, and h112ifiber textures and the number fraction of HABsincreased with strain. It was mainly caused by sub-structuring through dislocation mechanisms.[12]

3) The EHABs of 30 deg through 60 deg with rotationaxis preferentially clustering at h111i-h011i in theh011i and h112i fiber textures resulted from defor-mation texture variants that interwove with eachother because of substructuring. Few EHABs wereobserved in the h001i texture because of the partialstability of the cube texture as more cell structuresexisted in the cube-textured grains.[10]

4) The texture simulation based on octahedral{111}h011i and nonoctahedral {hkl}h011i slips over-estimated the copper and the cube texture intensity.One reason could be that no substructuring effectwas considered in the modeling. Substructuring ran-domized orientations and moderated the textureintensity, and resulted in fluctuation in a-fiber tex-ture, e.g., the brass and the Goss.

ACKNOWLEDGMENTS

The authors gratefully acknowledge the financialsupport from the Department of Energy, Contract No.DE-FC-26-06NT42755, and the Center for AdvancedVehicular Systems (CAVS) at Mississippi State Uni-versity. The authors are also grateful to Robert Malleyand Stephen Horstemeyer for their assistance with theexperiments at CAVS.

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