nr

. A

Gm

ik, R

Abstract

1. Introduction

These two metals are positioned close to each other in theperiodic table and are quite similar in a number of basic

properties, such as atomic weight, elasticity, strength and

accommodate deformation along the c-axis of the magne-sium hcp unit cell, whereas aluminum with its four slipplanes, each with three slip directions, has 12 slip systems,and thus, unconstrained formability. To satisfy a generalstrain tensor, magnesium has to add additional but notreadily available glide systems and/or mechanical twinning,

Corresponding authors.E-mail addresses: Juergen.Hirsch@hydro.com (J. Hirsch), alsamma-

n@imm.rwth-aachen.de (T. Al-Samman).

Available online at www.sciencedirect.com

Acta Materialia 61 (2013) 81884The use of light structural materials has become inevita-ble in the modern world. The increasing need for improvedfuel economy has created a huge interest in lightweight auto-motive structures. Automotive structures can be best light-ened through innovative design strategies directed towardweight saving (e.g. thin-walled components) and by employ-ing lightweight materials, such as aluminum and magne-sium, which have the lightest density of all commonstructural materials (qAl = 2.7 g cm

3, qMg = 1.7 g cm3).

melting point. Bothmetals have close-packed atomic latticesbut dier in their possible variants, i.e. face-centered cubic(fcc) and hexagonal close-packed (hcp), which explains theirfundamental dierences in forming behavior and single-crystal plastic anisotropy.

While aluminum has already established a leading rolein automotive applications, the use of magnesium for partsof automobile structures is still rather limited. Magnesiumshows limited ambient temperature formability due to ashortage of independent deformation modes that wouldAluminum and magnesium are two highly important lightweight metals used in automotive applications to reduce vehicle weight.Crystallographic texture engineering through a combination of intelligent processing and alloying is a powerful and eective tool toobtain superior aluminum and magnesium alloys with optimized strength and ductility for automotive applications. In the present articlethe basic mechanisms of texture formation of aluminum and magnesium alloys during wrought processing are described and the majoraspects and dierences in deformation and recrystallization mechanisms are discussed. In addition to the crystal structure, the resultingproperties can vary signicantly, depending on the alloy composition and processing conditions, which can cause drastic texture andmicrostructure changes. The elementary mechanisms of plastic deformation and recrystallization comprising nucleation and growthand their orientation dependence, either within the homogeneously formed microstructure or due to inhomogeneous deformation, aredescribed along with their impact on texture formation, and the resulting forming behavior. The typical face-centered cubic and hexag-onal close-packed rolling and recrystallization textures, and related mechanical anisotropy and forming conditions are analyzed andcompared for standard aluminum and magnesium alloys. New aspects for their modication and advanced strategies of alloy designand microstructure to improve material properties are derived. 2012 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Aluminum; Magnesium; Texture; Deformation; RecrystallizationSuperior light metals by texture eand magnesium alloys fo

J. Hirsch a,, TaHydro Aluminum Rolled Products

b Institut fur Metallkunde und Metallphys1359-6454/$36.00 2012 Acta Materialia Inc. Published by Elsevier Ltd. Allhttp://dx.doi.org/10.1016/j.actamat.2012.10.044gineering: Optimized aluminumautomotive applications

l-Samman b,

bH R&D, 53117 Bonn, Germany

WTH Aachen, 52056 Aachen, Germany

www.elsevier.com/locate/actamat

3rights reserved.

ta Mwhich makes it prone to orientation eects caused by itspronounced texture development during thermomechani-cal processing.

Another major dierence between the two materials isthe texture variation during recrystallization of thedeformed structure. In aluminum, depending on alloy com-position and processing conditions, dramatic texturechanges are observed, due to strong selective recrystalliza-tion mechanisms during nucleation and nucleus growth.

Such dramatic but characteristic orientation changes aremissing in conventional magnesium and its alloys, where inmost cases the deformation texture is preserved duringannealing. These texture eects can be best analyzed incold-rolled (i.e. plane strain deformed) and annealedsheet, where stress and strain conditions are well denedand characteristic textures develop. For aluminum, typicalfcc rolling and recrystallization textures occur that aectsheet formability mainly by the related anisotropy ofstrength, ductility and related phenomena (yield locus, form-ing limit, earing). Inmagnesium a strong basal texture formsrapidly during (hot) rolling and remains virtually unchangedduring recrystallization annealing, which aects subsequentsheet formability. To improve magnesium sheet formabilityone needs to depart from standard basal textures.

The following highlights the role of texture engineeringby intelligent processing and alloying as an eective toolfor optimizing aluminum and magnesium alloys for auto-motive applications. The texture evolution and its impacton properties will be analyzed and compared for aluminumalloys (i.e. AA5xxx AlMgMn alloys and AA6xxx AlMgSi alloys) and most common magnesium alloys(containing Al and Zn) that are predominantly used forautomotive lightweight engineering. Special attention ispaid to recent magnesium alloy development through mic-roalloying with rare earth elements. The basic mechanismsinvolved in deformation and recrystallization are analyzedand compared. The role of alloy additions, solute segrega-tion and second-phase precipitates, initial textures, defor-mation and annealing conditions, microstructureheterogeneities and related recrystallization mechanismsare considered, and possibilities are derived for their mod-ication to guide potential strategies of alloy design andmicrostructure engineering for optimal material properties.

2. Magnesium alloys

The tendency of magnesium and its alloys to developstrong single-ber textures during thermomechanical pro-cessing creates major problems for forming at ambienttemperature. The key features of ductile magnesium alloysare a homogeneous microstructure free of brittle interme-tallic particles, and both a grain size and crystallographictexture that promote uniform plastic deformation. Thusto begin with, we will review the recent progress in obtain-ing ductile magnesium alloys via texture optimization that

J. Hirsch, T. Al-Samman /Acserves to delay failure by permitting stable plastic deforma-tion under low stresses. Particular focus is placed on thefavorable eects that rare earth alloying addition can haveon texture sharpness and in generating non-conventionalrare earth textures during conventional deformationprocesses.

2.1. Advantages and limitations of magnesium alloys

Since the 1990s the forgotten materialmagnesium andits alloys have been receiving widespread attention as play-ing an increasingly important role in almost every sector ofthe metal consuming industry. In particular, the past dec-ade has seen rising demand for magnesium alloy develop-ment for structural and automotive applications due totheir large potential for weight saving and, thereby, forimproved fuel economy and decreased exhaust emissions.The range of applications of magnesium alloys has beenbased on liquid and semi-solid molding processes andsheet-forming processes. The vast majority of all magne-sium components (automobile, housings and electric) areproduced by die-casting. Products take advantage of thenumerous appealing properties of magnesium, such aslightness, high specic strength, good machinability, highdamping capacity, good castability, and low melting tem-perature and melting energy [1]. The die-casting processstands out due to its excellent productivity, and the result-ing potential to lower material cost through design strate-gies. Die-casting magnesium alloys can produceacceptable strength values but the room temperature duc-tility remains fairly low. In addition, the problem of gasporosity due to splashing during mold lling remains acomplex issue, which renders heat treatment to improvethe properties dicult [2].

For high-performance structural applications of magne-sium alloys it is desirable to develop wrought magnesiumproducts, such as extruded proles, rolled sheets and for-gings, which in contrast to die-cast products possess higherstrength and ductility. A number of technical issues, how-ever, limit the wider application of existing wrought mag-nesium alloys. The most signicant problems are thelimited formability and strong anisotropy in mechanicalbehavior, particularly at temperatures below 200 C, andthe low corrosion resistance. The corrosion performanceof Mg alloys depends strongly on the microstructure andalloying elements, particularly on the impurity level ofheavy metals, such as Cu, Ni, Co and Fe. Optimizing thecorrosion resistance of Mg alloys requires knowledge ofthe phases present in the alloy, their fraction and distribu-tion, and most importantly their electrochemical compati-bility with the Mg matrix. In MgAlZn (AZ) alloys, ithas been established that the b-phase (Mg17Al12) is catho-dic to the matrix, and can either enhance or degrade thecorrosion resistance of the alloy depending on its distribu-tion [3]. If it is uniformly distributed at the grain bound-aries, forming a dense network, it functions as acorrosion barrier and improves the corrosion resistance

aterialia 61 (2013) 818843 819of the alloy signicantly, but also decreases the ductility.On the other hand, if the amount of b-phase in the matrix

a Mis not sucient, individual b-phase precipitates can causelocalized galvanic corrosion and increase the corrosionrate. For MgAlMn (AM) alloys, the phase Al8Mn5 isused to reduce the Fe and Ni content during casting, whichgives this alloy a reasonable corrosion performance. In Al-free Mg alloys, the corrosion behavior can be improved byadditions of rare earth elements and yttrium.

2.2. Ambient temperature formability characteristics of

conventional magnesium alloys

The limited formability of magnesium alloys at ambientconditions represents a major challenge that prevents themore widespread use of these materials. This limitationleads directly to high production costs and low productionrates. For instance, conventional magnesium alloys aregenerally extruded more slowly than aluminum alloysbecause the temperature window where the material isworkable, yet does not suer incipient melting, is quite nar-row. For the same reason, magnesium alloy sheet produc-tion has so far been restricted to elevated temperaturerolling. In contrast to aluminum and steel, which can benished cold, much of the reduction to nal gauges in mag-nesium has to be performed at elevated temperaturesbetween 300 and 450 C.

With respect to crystal structure, magnesium diers sig-nicantly from its metallic lightweight competitor, alumi-num, in having a hexagonal crystal structure. At ambienttemperature, plastic deformation of magnesium is limitedto two main deformation mechanisms, basal(0001)h1 1 2 0i slip and f1 0 1 2gh1 0 1 1i mechanical twin-ning. There are three dierent h1 1 2 0i directions in the(0001) plane, and thus the magnesium crystal has onlythree geometrical and two independent slip systems. Onthe other hand, the aluminum crystal has 12 {111}h1 1 0igeometrical and many combinations of ve independentslip systems, and thus easily complies with the von Misescriterion for homogeneous shape change [4]. Other slip sys-tems, such as f1 0 1 1gh1 1 2 0i prismatic slip andf1 0 1 1gh1 1 2 0i pyramidal slip, have the same hai-slipdirection in common but they require a larger criticalresolved shear stress (CRSS) for activation. Therefore, theyare usually harder to operate at room temperature, and atthe same time do not result in deformation out of the basalplane.

Typical magnesium alloy sheet exhibits strong basal-type textures with grain orientations having basal planesparallel to the sheet plane. Under loading in directionseither parallel or perpendicular to the sheet plane, basal,prismatic and pyramidal hai-slip systems fail to accommo-date any deformation because they all have a slip directionparallel to the basal plane and the resulting shear strain inall slip systems is zero. For such cases, slip vectors with acomponent out of the basal plane are required. This canbe accommodated by slip on the f1 1 2 2g-pyramidal plane

820 J. Hirsch, T. Al-Samman /Actin the h1 1 2 3i-direction, which is referred to as pyramidalhc + ai-slip [5]. This slip mode oers ve independent slipsystems, thus satisfying the von Mises criterion for an arbi-trary shape change. On the other hand it has a substantiallylarger slip vector compared to hai-slip, and, hence, a mark-edly higher CRSS. Therefore, pyramidal hc + ai-slip inmagnesium usually requires substantial thermal activation,i.e. high temperatures. With the absence of hc + ai-slip atlow deformation temperatures, the hexagonal crystals haveno means to accommodate the imposed strain along thesheet normal direction by crystallographic slip. This causeshigh stresses, modest work hardening, and, most impor-tantly, premature brittle fracture.

In addition to hc + ai-slip, mechanical twinning on thepyramidal f1 0 1 2g and f1 0 1 1g planes can provide ashear component parallel to the c-axis. The most commontwinning mode in magnesium is the f1 0 1 2g extensiontwinning. This twinning mode accommodates extensionalong the c-axis, and hence for rolling deformation and abasal texture this type of twinning obviously cannot be uti-lized. By contrast, f1 0 1 1g contraction twinning producesa favorable compression strain along the c-axis; however,in practice, while compression twins exist in conventionalmagnesium alloys, they are scarce and do not contributeto plastic deformation in a quantitative manner.

Like cold processing of rolled sheet, cold extrusion ofmagnesium alloys faces similar diculties. During extru-sion processing, the basal planes in the deformed materialare usually aligned parallel to the extrusion direction(ED), forming a strong and homogeneous h1 0 1 0i || EDber texture. Subsequent mechanical loading in compres-sion along the ED will impose a tensile strain componentparallel to the c-axis of crystals and will, thus, massivelyactivate f1 0 1 2g extension twinning throughout thestrongly textured microstructure. As f1 0 1 2gh1 0 1 1icauses an orientation change of 86.3 h1 1 2 0i, the basalplanes of twinned grains will rotate by 86.3 from their ori-ginal orientation, so that their c-axes become closelyaligned with the compression direction. This microstruc-tural evolution normally occurs during the rst 6% of plas-tic deformation and causes a huge and unfavorable texturechange. Subsequent deformation needs to be accommo-dated by unavailable c-axis deformation mechanisms,which severely limits the ductility of magnesium extrusionsfor ED compression. For the reverse case of tensile loadingin the ED, the c-axis of crystals will experience compres-sion, which triggers, however, only little f1 0 1 1g contrac-tion twinning, and thus does not change the overall initialtexture, where the vast majority of grains are favorably ori-ented for prismatic hai-slip. However, the CRSS for pris-matic hai-slip at ambient temperature is relatively highcompared to basal slip, which renders its activation some-what dicult. This again contributes to the serious limita-tion of formability. In essence, the rapid texture formationduring deformation processing severely impacts the ductil-ity and renders the material brittle.

One interesting approach to improve the cold formabil-

aterialia 61 (2013) 818843ity of magnesium is to modify its hexagonal crystalstructure, either by reducing its c/a axial ratio or by

ta Mcompletely transforming it into a cubic structure (e.g.body-centered cubic (bcc)). This can be achieved by alloy-ing with lithium, which also lowers the weight of thealready light magnesium even further (qLi = 0.58 g cm

3).With increasing lithium content the c/a ratio decreasesfrom 1.624 (for pure Mg) to 1.607 (for Mg17 at.% Li,close to the solid-solubility limit) [6]. Between 17 and30 at.% Li content MgLi alloys comprise a two-phasestructure consisting of the a-Mg rich (hcp) and the b-Lirich (bcc) phases. The highly ductile b-single phase struc-ture exists for Li contents greater than 30 at.% [7].Although cubic MgLi alloys do not suer from ductilityproblems, these alloys have high production costs, exhibitlow corrosion resistance and sometimes show low-temper-ature instability. Hexagonal MgLi alloys are also highlyductile at ambient temperatures. This is attributed to a sub-stantial increase in non-basal slip activity over entiregrains, particularly prismatic and pyramidal hc + ai-slip.This is in contrast to pure magnesium, where the activityof these slip systems is negligible at room temperature,and at most limited to regions with large stress concentra-tions in the vicinity of grain boundaries. It is noted that theenhanced non-basal slip activity is not only attributed tothe decrease in the c/a ratio resulting from lithium additionbut also to solid-solution softening eects that promote slipon multiple planes [8]. By comparison to industrial Mgalloys, binary MgLi alloys exhibit modest work hardeningbehavior at room and elevated temperatures. They are alsovery liable to grain growth at temperatures above 300 C,which results in coarse microstructures and sharp recrystal-lization textures [9].

Since there is a strong correlation between the crystallo-graphic texture evolution during thermomechanical pro-cessing, the operating slip and twinning mechanisms, andthe resulting mechanical behavior, it is obvious that con-trolling the microstructure (texture and grain size) hasbroad potential to enhance the mechanical behavior ofwrought magnesium alloys. The principal goal of texturemodication for enhanced formability in magnesium is toachieve a favorable alignment of the basal planes withthe deformation direction, whereas grain renement servesto reduce the activity of twinning and promote additionaldeformation mechanisms, such as grain boundary sliding.A wide range of wrought magnesium alloy compositionsand processing schemes have been investigated, and it hasbeen reported that texture weakening/randomization waspossible in several cases with various alloy and processingconditions. In the following we will review recent attemptsand discuss the underlying mechanisms, pinpointingimportant questions that still need to be answered. We willfocus on conventional deformation processes, such as roll-ing and extrusion, and leave out non-conventional severeplastic deformation methods, such as equal-channel angu-lar pressing or high-pressure torsion, which despite theirgreat potential for grain renement, texture modication

J. Hirsch, T. Al-Samman /Acand superplasticity [1012], have still to nd widespreadapplication and will be addressed in another contributionof this issue. We will certainly not be able to give a compre-hensive overview of all the recent developments in improv-ing the formability of magnesium alloys, rather we willfocus on the critical issues, where substantial progress isbeing made via intelligent texture modication. Finallywe will provide the reader with new perspectives that willguide future magnesium alloy research and development.

2.3. Improvement of conventional magnesium alloys by

intelligent thermomechanical treatment

2.3.1. The role of random texture on ductility

The MgAlZn system is probably the most thoroughlyinvestigated of all magnesium alloy systems because it rep-resents the most common commercial AZ alloys. Withrespect to texture weakening and the improvement of form-ability of these alloys, several reports have been publishedthat focused on the key factors controlling the texture dur-ing deformation. For example, for the widely availablemagnesium alloy AZ31, containing nominally 3 wt.% Aland 1 wt.% Zn, a combination of extrusion and uniaxialcompression at 400 C and a constant strain rate of 104 s1 generated a random texture at strains >1 [13]. Incontrast, lower deformation temperatures and higher strainrates resulted in sharp textures. Further investigationsshowed that not only were the deformation parametersimportant, but the choice of the starting orientation wasalso crucial.

The material investigated was initially hot extruded anddepicted a typical h1 0 1 0i || ED ber orientation. Threedeformation orientations were explored, with the compres-sion direction (CD) being set parallel (Fig. 1a), perpendic-ular (Fig. 1c) and aligned at 45 (Fig. 1e) to the ber axis ofthe extrusion texture. This rendered the activation of slipon basal and non-basal planes, as well as the activationof extension twinning, easier for one starting orientationthan the other. When the material was compressed alongthe ber axis of the extrusion texture, i.e. along the ED,deformation was axisymmetrical for all grains and theresulting nal texture at e = 1.4 was completely random-ized, i.e. the texture intensity was less than 2 times random(Fig. 1b). For the other orientations, though texture weak-ening was evident, the resulting textures were still welldened and characterized by clustering of basal polesaround the CD (Euler angle, U, spread within 30, Bungenotation) (Fig. 1d and f). Subsequent tensile testing ofthe randomly oriented specimen at room temperature anda constant strain rate of 104 s1 showed appreciableenhancement of formability (ef 0.3) over the initialextruded state (ef 0.18) with no decrease in strength(Fig. 2). This shows that texture randomization is a prom-ising avenue for the production of ductile magnesiumalloys.

Turning to the mechanism(s) responsible for randomiz-ing the texture of AZ31 alloy, it is important to bear in mind

aterialia 61 (2013) 818843 821that the texture given in Fig. 1 is a result of high-tempera-ture deformation comprising concomitant crystallographic

a M822 J. Hirsch, T. Al-Samman /Actslip, dynamic recrystallization (DRX) and possibly graingrowth. Given the fact that the same deformation

Fig. 1. Schematic of the three tested orientations in uniaxial compression atinitial extrusion direction. (a), (c) and (e) represent the starting textures of each(b), (d) and (f) are resulting compression textures at e = 1.4. Texture represconvention using the MTEX tool box [14]. Texture intensity in this and all ot

Fig. 2. Stressstrain curves of tension test at room temperature and104 s1 strain rate comparing the maximum tensile elongation of anextruded AZ31 specimen (Fig. 1a) and a uniaxially deformed specimenwith a randomized texture (Fig. 1b).aterialia 61 (2013) 818843parameters were applied in all three cases, there ought tobe discernible dierences in the deformation and recrystalli-zation behavior that make the achievement of a randomlytextured material quite sensitive to the choice of startingorientation. If we rst consider plastic deformation, it isobvious that the activation scenarios of deformation mech-anisms were dierent for each starting orientation. Forexample, for the orientation where the texture was com-pletely randomized, i.e. the applied force was parallel tothe h1 0 1 0i direction, prismatic slip, f1 0 1 2g extensiontwinning and pyramidal hc + ai-slip were all potentialdeformation mechanisms (thermally activated and geomet-rically favored). However, prismatic slip and mechanicaltwinning can be readily excluded from playing an importantrole since prevalent prismatic slip activity would only retainthe starting orientation, and twinning would cause a suddenand large orientation change from Euler angle / valuesclose to 90 to values close to 0, which was not observedin the recorded texture development as a function of strain,reported elsewhere [13,15]. In this case, only pyramidalhc + ai-slip is expected to dominate the deformation.Despite its great advantage of providing ve independent

400 C and 104 s1 strain rate. CD, uniaxial compression direction; ED,tested orientation with the CD parallel, perpendicular and 45 to the ED.entation in terms of orientation distribution function (ODF) with Bungeher gures is in multiples of a random distribution (MRD).

ta Mslip systems, the activity of this slip nevertheless has to pro-duce some texture component. The absence of pronouncedtexture components, as shown in Fig. 1b, cannot beachieved by any slip condition.

If we add DRX to the scenario, some degree of textureweakening may certainly be expected. As seen, DRX wasevident in all three deformed orientations. The resultingoptical microstructures at e = 1.4 were fully recrystal-lized, and the average DRX grain size was comparablefor all three orientations [13]. It is acknowledged that somerecrystallization mechanisms, such as particle-stimulatedrecrystallization, can generate a large spread of crystallo-graphic orientations, and thus randomize the nal bulk tex-ture by recrystallization and grain growth. However, forthe investigated alloy AZ31, only a few second-phase par-ticles were present, so that any possible role of particles wasconsidered insignicant. Microstructure observations atsmall deformation strains indicated that incipient nucle-ation of DRX occurred mainly through conventional serra-tions and bulging mechanisms of pre-existing high-anglegrain boundaries [1618]. In the course of DRX, withincreasing strain the deformed original grains were pro-gressively consumed by the expansion of dynamicallyrecrystallized regions forming multiple layers of what isknown as necklace structure [19]. The rst necklace grainsformed by the bulging mechanism usually have orienta-tions close to the orientation of the adjacent parent grains[20]. Misorientations of 5 have been predicted to be com-mon, and in some cases misorientations as high as 20 werereported [21]. The expansion of the necklace structure wasshown not to proceed by means of repeated bulging on therecrystallization front, but rather by other mechanisms thatlead to the loss of orientation coherency of the recrystal-lized volume with the deformed matrix [20]. In the caseof Mg, this phenomenon can be pictured as a continuedchange in orientation of the DRX grains as deformationprogresses, up to large misorientations and a tendency totexture randomization. Such a mechanism was proposedin 1982 by Ion et al. [22] for Mg0.8% Al alloy. Theauthors referred to the mechanism as rotation recrystalli-zation owing to lattice rotations in the grain boundaryregions, which were also nucleation sites of recrystalliza-tion. They argued that lattice rotation during deformationarises from the formation of subgrains along prior grainboundaries, and there is a continuous increase of misorien-tation due to absorption of dislocations until a high-angleboundary is obtained.

The question remaining here, however, is why the mech-anism of rotation recrystallization contributed to full tex-ture randomization in one orientation and failed to do soin the other orientations. In a recent in situ study onhigh-purity Al bicrystals subjected to external mechanicalstress at elevated temperatures [23], stress-induced grainrotation concurrent with boundary migration was recordedfor a grain boundary with a mixed tilttwist character. In

J. Hirsch, T. Al-Samman /Acpolycrystals, most grain boundaries are of mixed type,and depending on the stress direction with respect to thegrain boundary plane, the misorientation angle resultingfrom stress-induced grain rotation may increase ordecrease. Therefore, there is good reason to believe thatthe contribution of DRX-grain rotation to texture random-ization is orientation dependent. For a favorable orienta-tion (as in Fig. 1a) the fraction of suitably oriented grainboundaries for stress-induced grain rotation during DRXis most likely large enough to result in a complete loss oftexture. Additional investigations are necessary to providefurther support for this hypothesis.

In another texture study by Backx [24] on extrudedAZ31 combining heat treatment with uniaxial compressionat 250 C and 102 s1 up to dierent strains, texture ran-domization was recorded for 70% deformation strainrecrystallization annealing at 350 C for 30 min. This spec-imen demonstrated a remarkable ductility increase up toe = 0.4 during subsequent uniaxial compression at roomtemperature. The author attributed the enhanced mechan-ical properties to an optimized texture and microstructureobtained by intelligent thermomechanical treatments.Although the underlying mechanisms were not elucidatedin detail, it can be concluded from that work that staticrecrystallization (SRX) can also randomize the texture inAZ31. In essence, the degree of randomization will dependon parameters such as the deformation strain, the grain sizeand the texture of the annealing microstructure; hence theattainment of optimal mechanical properties requires adeep understanding of microstructure evolution duringprocessing.

2.3.2. The role of particle-stimulated recrystallization on

texture formation

Other investigations considered the eect of second-phase particles on the texture development in theMgAlZn alloy system. For dierent aluminum contents,texture formation during plane-strain compression (PSC)at various values of Z was investigated [25], where Z isthe temperature-corrected strain rate, also termed theZenerHollomon parameter; Z _e exp QRT

, where Q is

the activation energy and R is the gas constant. In additionto varying Z, two initial rolling orientations were investi-gated, with the channel-die compression direction, CD,being applied parallel to the previous rolling direction,RD (frequently referred to as in-plane compression). Theextension direction was set in one sample parallel to theprevious sheet normal direction, ND, and in another sam-ple parallel to the previous transverse direction, TD. Thisrendered the activation of f1 0 1 2g extension twinningeasy for the former case and dicult for the latter. Owingto the low Al content in the AZ31 alloy, the amount of b-phase was negligible. In the AZ61 and AZ91 alloys, withnominal Al contents of 6 and 9 wt.%, respectively, Mg17-Al12 precipitates were primarily distributed at the grainboundaries, and within the grain interior had average par-ticle sizes of 510 lm (for AZ61) and 2030 lm (for AZ91).

aterialia 61 (2013) 818843 823During PSC testing at 200 C the volume fractions of thesecond phase in the AZ61 and AZ91 alloys amounted to

2.9% and 5.1%, respectively. At 400 C, the b-phase waspartially dissolved and the respective volume fractions were1.1% and 1.7% [25].

The results showed that the impact of the b-phase pre-cipitates on texture weakening during PSC at 200 C wasmost pronounced for the AZ91 alloy relative to AZ61and AZ31 (Fig. 3). This suggests that the precipitate den-sity, and most likely also the precipitate size, are importantfactors in magnifying the impact of particles on texturedevelopment. The eect of Mg17Al12 precipitates on weak-ening the texture was qualitatively explained by enhancednucleation of DRX at second-phase particles. A sucientlylarge volume fraction of the b-phase causes nucleation ofgrains with diverse orientations (Fig. 4). This view was cor-roborated by deformation experiments at 400 C and 104 s1 up to the same strain of e = 1, which yielded muchsharper textures (Fig. 3). At that temperature the b-phase

was mostly dissolved in the magnesium matrix and hadthus played little part during DRX.

2.4. Improvement of magnesium and its alloys by

microalloying with rare earth (RE) elements

2.4.1. Formation of soft textures in Mg-RE sheet and

extrusions

Magnesium alloys with rare earth (RE) elements havebeen thoroughly investigated in Russia [26], and many indi-cations were long ago found for the improved properties ofthis class of alloys including grain renement, better form-ability at low temperatures, and enhanced strength andcreep resistance at elevated temperatures. In recent years,wrought Mg-RE alloys have again attracted much scienticattention since their deformation and recrystallization tex-tures were found to be much weaker and less common than

nd

824 J. Hirsch, T. Al-Samman /Acta Materialia 61 (2013) 818843Fig. 3. Schematic of the investigated sample orientation for the PSC tests a

terms of (0002) basal pole gures. Top row corresponds to PSC textures atindicates texture randomization. Bottom row corresponds to PSC textures atthe corresponding starting textures of the three AZ alloys (middle row) in

200 C, 104 s1 strain rate and a true strain of 1. Outlined pole gure400 C and same strain rate and nal strain as in the top row.

uree g

ta Mtypical textures observed in conventional Mg alloy sheet orextrusions.

In an attempt to eliminate the yield asymmetry in mag-nesium extrusions, Ball and Prangnell [27] were the rst tonote that a commercial WE43 alloy containing additions ofY and RE elements can develop more random-type tex-tures during extrusion, as opposed to conventional Mgalloys. At that time, the randomized texture in WE43 wasattributed to particle-stimulated nucleation (PSN) ofrecrystallization.

The fact that RE elements are relatively expensive andnot readily available throughout the world has promotedconsiderable eorts to examine the potential of RE addi-tions at microalloying levels and to check whether texturemodication is also possible when RE elements are in solu-tion. In their study on sheet textures in MgZn alloys con-taining dilute additions (15) in black and low-angle boundaries (515) in yellow. (b) Inverse polof discrete and contour data.

J. Hirsch, T. Al-Samman /Actimes in the literature that sheet textures of many RE ele-ments containing Mg alloys cause a greater tilt of basalpoles in the sheet transverse direction, TD, rather than inthe rolling direction, RD (Fig. 5), which promotes moreactivation of basal slip during loading in TD than in RD.This causes a mechanical and plastic anisotropy responsein terms of in-plane yield strength, elongation to failureand strain anisotropy (r-value) dierent from thoseobserved in conventional Mg alloy sheet material. TheTD-spread in RE sheet textures was found to reduce theplanar anisotropy in comparison with conventional alloysthat exhibit r-values between 2 and 4. Reduced planaranisotropy, Dr, of rolled sheet was reported to enhancethe forming behavior under straining conditions, which isrequired for sheet thinning [29].

The benecial eect of RE elements in magnesium alloyshas stimulated research on a wide range of alloys andcompositions involving dierent mixtures of RE elementsand Y. In a study to examine the eect of microalloyingaddition of Y- and Nd-based mischmetal (Table 1) onthe extrusion texture and microstructure of MgZnZr(ZK) alloys, appreciable texture and microstructure modi-cation was achieved (Fig. 6). Fig. 6a and b show represen-tative electron backscatter diraction (EBSD) micrographs(inverse pole gure maps) of the extrusion microstructuresof both ZK10 (standard alloy version) and ZWEK1000(microalloyed version with RE/Y) alloys, respectively,and the corresponding misorientation angle distributions(MADs) of the grain boundaries. Clearly, the addition ofmischmetal and Y aects the recrystallization behavior ofthe material and weakens the texture. The large elongatedgrains in the extruded ZK10 alloy were replaced by coarseequiaxed grains in the RE/Y-containing alloy, which wasreected in the MAD by a pronounced reduction of themisorientation peak at 5 and concomitant increase ofthe misorientation peak at 30. The peak intensity at lowmisorientation angles of 5 is an indication of the extentof subgrain structures within the unrecrystallized elongatedgrains. The misorientation peak at high-angle boundariesof 30 is most probably due to a common 30[0001] mis-orientation relationship during growth of recrystallized

in Fig. 3. (a) Kikuchi band contrast (BC) map with high-angle boundariesure (IPF) map of the circled area in (a) and corresponding texture in terms

aterialia 61 (2013) 818843 825grains in hexagonal materials (e.g. [3034]) and its intensitycan be correlated with the recrystallized volume fraction.

As evident from the orientation map in Fig. 6b, therecrystallized grains in the modied alloy comprised abroad spectrum of orientations that were dierent fromthe sharp h1 0 1 0i || ED deformation orientation that dom-inates the traditional extrusion texture in the ZK10 alloy.The wide variety of orientations of recrystallized grainsand their prevalent volume fraction in the microstructureof the ZWEK1000 alloy was found to randomize the bulktexture to a large extent. While this is still a subject ofcurrent research it is proposed that stable second-phaseparticles with a complex mixture of Zn, Nd, Y and Mgare eective nucleation sites for DRX during extrusion at400 C and give rise to randomly oriented nuclei. Fig. 7shows good microstructural evidence of very negrains (d 5 lm) spotted in the immediate vicinity of

Fig. 5. (a) Recalculated (0002) pole gure and (b) basal pole density plot as function of tilt angle about the sheet normal direction, ND, showing anexample of RE sheet texture in a Mg0.9% Zn0.7% La0.2% Zr (wt.) alloy with obvious trend for TD spread opposed to the other sheet directions.

Table 1Chemical composition of the investigated benchmark alloy ZK10 and the modied version containing Y and Nd-rich mischmetal.

Alloys Zn (%) Zr (%) Nd (%) Ce (ppm) Gd (ppm) Y (%) Mg

ZK10 (standard) 1.51 0.41 RestZWEK1000 (RE/Y modied) 1.47 0.38 0.20 31 11 0.19 Rest

Fig. 6. EBSD microstructures (IPF maps) along the extrusion direction, ED, and corresponding misorientation angle distributions for (a) conventionalalloy ZK10 (benchmark), and (b) modied alloy ZWEK1000.

826 J. Hirsch, T. Al-Samman /Acta Materialia 61 (2013) 818843

ta Msecond-phase particles, whose size varied between 1 and10 lm. Other grains, outside of the highlighted area inFig. 7 (where no particles were present), were 10 timeslarger. While the role of particles in generating new recrys-tallized orientations may appear to account for a small vol-ume fraction, recrystallizing grains seemed to readilyundergo grain growth, consuming a larger volume of theextrusion microstructure, and thereby contributing to tex-ture weakening.

From the above ndings it can be concluded that smalladditions of Nd-rich mischmetal and Y to the extrusionalloy ZK10 can signicantly weaken the extrusion texture,yet the position of the texture peak was still close to theh1 0 1 0i ber. This is not necessarily the case for otherextrusion alloys with additions of RE and Y. Stanfordand Barnett [35] reported a shift in the orientation peakof extruded Mg alloy ME10 (1 wt.% Mn and 0.4 wt.%Ce-rich mischmetal) from h1 0 1 0i (seen in AZ31) orh1 0 1 0ih1 1 2 0i (seen in M1) to a new position h1 1 2 1i

Fig. 7. Optical micrograph (ED plane) of the extruded ZWEK1000 alloyshowing one example of numerous ne-grained areas associated withsecond-phase particles.

J. Hirsch, T. Al-Samman /Ac(about 30 from h1 1 2 0i of the inverse pole gure), whichwas termed in later studies the RE texture component. Astudy by Huppmann et al. [36] on a similar alloy ME21(2.1 wt.% Mn and 0.7 Ce-rich mischmetal) examined therole of extrusion parameters (billet temperature, extrusionratio and cooling conditions) on the formation of the REtexture component. In addition to the aforementionedh1 1 2 1i component, they observed two additional RE tex-ture components, h1 1 2 2i and h2 0 2 1i, parallel to theextrusion direction. Moreover, it was shown that the vari-ation of extrusion parameters and the cooling conditionsallows a variation in the extrusion texture. Both studiesconcluded that the RE texture component is well orientedfor basal slip when tested in the appropriate orientation,which results in a substantial gain of ductility and a reduc-tion of the tensioncompression asymmetry typical forconventional wrought Mg. While the focus of this paperis on dilute MgRE alloys, it is noted that unusual texturesdevelop also in heavily alloyed MgRE extrusions withmore than 10 wt.% total RE. For example, in an extrudedMg7Gd6Y0.5Zr (wt.%) alloy [37], an unusual prismatictexture component was observed, in which the c-axis wasaligned parallel to the extrusion direction, perpendicularto the typical basal extrusion texture of magnesium alloys.

2.4.2. Comparison of dierent RE additions with respect totexture modication

In order to gain additional insights into the origin of REtexture modication, recent eorts have focused on theeect of single RE additions on the formation of defor-mation and recrystallization textures. Basically, dierentRE elements can form dierent intermetallic compounds[26,38] that contribute to distinct properties. Additionally,in the form of solutes, individual RE elements may interactdierently with dislocations and grain boundaries. In singleRE alloying studies, individual RE elements that are read-ily soluble (e.g. Gd, Nd and Y) or insoluble (e.g. Ce andLa) in Mg are usually chosen to study the role of RE inthe form of both solute solutions and two-phase alloys.Investigations were mainly concerned with (a) examiningthe potential of individual elements as texture modiers,and (b) determining the amount of RE element requiredto modify the texture. Based on ndings obtained for bothaspects, good progress has been made so far toward high-lighting the underlying mechanisms for RE texturemodication.

2.4.3. The role of REsolute segregation

In a comparative study on microalloying Mg with dier-ent RE elements and Y for texture modication [39], it wasdemonstrated that the minimum level required is dierentfor each RE element. At suciently high concentrationsbetween 300 and 600 ppm, MgLa, MgCe and MgGdproduced the RE texture component (h1 1 2 1i parallel toED). Strangely, Y did not show this texture type at anyof the concentrations examined (0.020.17 at.% or 0.080.62 wt.%). A similar trend was reported for extrudedMgMn alloys with single additions of Ce, Y or Nd [40],where Y was a much weaker texture modier than Ceand Nd.

EBSD investigations have shown the formation of theRE extrusion texture to be associated with a high propen-sity for DRX nucleation at shear bands [41]. Furthermore,it was proposed that the eect of individual RE elements onthe location and intensity of the RE texture component isdue to solute interaction with dislocations and grainboundaries. Ce and La have larger atomic radii than Gdand Y, and were found to show the strongest eect on tex-ture and grain size [39]. In support of their hypothesis thatRE solutes have a strong interaction with dislocations andgrain boundaries, the authors conducted another study onextruded Mg1.5 wt.% Gd alloy that went through a tex-ture transition (h1 0 1 0i || ED ! h1 1 2 1i || ED) withincreasing extrusion temperature [42]. The aim of the studywas to investigate the correlation between solute segrega-

aterialia 61 (2013) 818843 827tion to grain boundaries and the RE texture eect. Bymeans of energy-dispersive X-ray spectroscopy (EDS)

performed in a transmission electron microscope (TEM),the authors reported a signicant increase in the local con-centration of Gd at grain boundaries compared to the bulk.Small amounts of other non-RE elements (Al, Mn and Cu)were also present in the material but showed less segrega-tion than Gd. The level of solute segregation was foundto depend on the extrusion temperature and correlated wellwith the texture transition that occurred at 490 C. Tem-perature plays a role in the competing eects of solutediusivity and grain boundary mobility. Elevated tempera-tures promote solute diusivity but decrease segregation,and at suciently high temperatures the boundary migra-tion rate becomes large enough for the grain boundary tobreak free from its solute cloud and move freely [43].Fig. 8 is a cast microstructure of a ZNK100 alloy (wherethe letter N represents Nd) with 0.1 at.% Nd, showing good

coupled and contribute strongly to the formation of unu-sual RE recrystallization textures. It should also be notedthat RE solute segregation can be coupled with RE soluteclustering in the matrix according to recent atom probetomography measurements [42]. Solute clusters couldprompt plastic strain heterogeneities similar to thoseobserved in systems with shearable particles, and couldthus certainly aect the nucleation and growth process ofrecrystallization.

2.4.4. Impact of dierent RE elements on the mechanical

properties

While the addition of RE elements to thermomechani-cally processed magnesium primarily aims at weakeningthe texture, the desired engineering outcome is improveddeformation behavior and enhanced property combinationof strength and room-temperature ductility. The linkbetween texture and formability has been well established,

828 J. Hirsch, T. Al-Samman /Acta Materialia 61 (2013) 818843evidence of preferential segregation of RE solute atoms tograin boundaries. Apparently, the local concentration ofNd at the grain boundary was in fact much higher thanits bulk concentration, leading to the precipitation of veryne second-phase particles (

ta Mbasal poles toward the sheet transverse direction. Goodcorrelation was found between the texture intensity andmaximum elongation to fracture, and also between thehardening exponent n and the texture. The Gd alloy withthe weakest texture showed a remarkable increase ofroom-temperature ductility up to 30% in all three sheetdirections. This was coupled with an average r-value of1.025 and a planar anisotropy Dr of 0.05, which translatesinto enhanced sheet thinning, less earing propensity and auniform strain distribution in the plane of the sheet. Inaddition, the hardening exponent was increased up to animpressive value of 0.45, which was found to eectivelyinhibit the onset of localized deformation described bythe Conside`re criterion [45]. Regarding the recrystallizedgrain size, there was no striking dierence in the dierentalloys. The values were in the range of 2842 lm. The ulti-mate tensile strength was comparable to conventional sheetbut the yield point could denitely benet from someimprovement, obtained potentially by optimizing the alloycomposition and rening the grain size. It is challenging,however, to increase the yield strength without sacrifyingductility.

2.4.5. Eect of RE elements on the relative activities of slip

at elevated temperatures

Although there is general agreement in the literaturethat texture modication by addition of RE elementsoccurs primarily during recrystallization and grain growth,the role of RE elements (in the form of both solute and par-ticles) on the deformation mechanisms and the microstruc-ture which precedes recrystallization remains important.From a recent study [44] aiming at quantifying the defor-mation mechanisms during high-temperature plane-straincompression of magnesium alloy ME20 containing ceriumas the main RE element compared to a benchmark alloyAZ31, it was seen that despite the same deformation condi-tions (Z parameter) and qualitatively comparable initialtextures, the two alloys showed considerably dierent tex-ture development (Fig. 10). This behavior was the resultof signicant dierences in the activation of deformationmodes, which was analyzed on the basis of experimentaltexture development and texture simulations using anadvanced cluster-type Taylor model that accounts for graininteraction [46].

As evident from Fig. 10, the texture development in theconventional alloy AZ31 was very similar at both Z condi-tions, 200 C/102 s1 (high Z) and 400 C/104 (low Z).From previous studies, it is established that the resultingAZ31 textures at e = 1 are the outcome of prevailing pris-matic slip activity with a small contribution from other hai-slip modes operating in some disoriented grains. On theother hand, the RE alloy ME20 showed an interestingtrend of rotating the basal poles from their initial TD ori-entation toward the compression direction, CD, whichindicates a reduced activity of prismatic slip relative to

J. Hirsch, T. Al-Samman /Acother slip modes, particularly during high Z deformationat 200 C. This view was substantiated by simulationresults shown in Fig. 10 revealing that at 200 C pyramidalhai-slip was much more important than prismatic slip; infact, at incipient deformation it was even more importantthan basal slip that was impeded by the initial TD orienta-tion. By increasing the temperature to 400 C, there was astrong competition between prismatic and pyramidal hai-slip to accommodate the deformation. This competitionwas extended to include also basal slip at strains largerthan 20%. Under such a multiple slip condition, pyramidalhai-slip played an important compensating role preventingthe PSC deformation texture from turning into either abasal or prismatic texture, which has positive implicationsfor ductility.

From the above ndings it seems that Ce additionchanges the relative CRSS of the various slip systems com-pared to benchmark AZ31, which points in the direction ofsolid-solution hardening and softening eects on the dier-ent slip/twin mechanisms, even in the presence of particles(micrometer-sized Mg12Ce and nanometer-sized pure Mnparticles). There are several reports in the literature onthe eect of non-RE elements such as Al, Zn and Li onsoftening prismatic slip at the expense of basal slip [4751]. However, a few studies on RE elements (e.g. [52])reported that Gd alloying tends to increase the CRSS forprismatic slip, which is consistent with the results presentedin Fig. 10 that show obvious strengthening of prismatic slipin the high Z regime. On the other hand, when the temper-ature was increased to 400 C, solid-solution strengtheningof the prism plane seemed to diminish, probably due tochanges in the particle/solute concentration. The strength-ening eect of prismatic slip did not seem to be coupledwith strengthening of pyramidal hai-slip.

It is noted that for the ME20 alloy, the addition of Ceand the presence of nanometer-sized Mn particles stronglyretard DRX, and thus prevent it from playing a signicantrole in the texture evolution shown in Fig 10. The eect ofdierent initial grain sizes between AZ31 (40 lm) andME20 (15 lm) was not considered when discussing themechanisms responsible for the dierent texture develop-ment of the two alloys. However, based upon prior knowl-edge regarding the impact of grain size on texturedevelopment, we do not suspect a dierence of 25 lmto disprove the hypothesis of solute-related eects on thedeformation slip activity.

2.4.6. The eect of RE elements on the stacking fault energyand hc+ai-slip activation at room temperature

As noted earlier, the addition of RE elements producesweak recrystallization textures that promote basal slip,which contributes to enhanced room-temperature ductility.There are also reports in the literature that the addition ofRE elements in solid solution increases the ductility duringcold deformation without involvement of recrystallization[5355]. A recent comprehensive study by Sandlobeset al. [55] combined electron density functional theory

aterialia 61 (2013) 818843 829(DFT) calculations with TEM to quantify the fundamen-tal mechanisms responsible for solute-related ductility

830 J. Hirsch, T. Al-Samman /Acta Materialia 61 (2013) 818843enhancement during cold deformation of MgY alloys.The authors reported that both methods (TEM andDFT) revealed a signicant decrease of the intrinsic stack-ing fault energy (SFE) with the addition of Y, and con-cluded that the reduced number of intrinsic stackingfaults enables the formation of dislocation structures onpyramidal planes, and acts as a heterogeneous nucleationsource for hc + ai-dislocations, which correspondingly pro-motes the activation of hc + ai-pyramidal slip at roomtemperature.

2.4.7. The role of non-RE elements on RE texture

modication

Alloying with RE elements is generally done using eitherpure magnesium, resulting a simple binary MgRE alloys,or dilute magnesium-based alloys, such as MgMn, MgZn or MgZnZr, resulting in more complex alloy systems.While this has recently been intensively studied, little atten-tion has been paid to the role of accompanying non-REelements on the RE texture modication. Preliminaryresults indicate that the texture weakening eect of RE ele-ments is substantially magnied by the addition of othernon-RE elements, such as Zn and Zr.

Fig. 10. Comparison of the PSC texture development between a conventional Aand low (400 C/104 s1) Z deformation up to e = 1. Texture prediction resultdeformation modes and their relative activities as a function of strain.Figs. 11 and 12 show interesting comparisons betweenbinary MgRE alloys and their corresponding quaternarysystems with added Zn and Zr. Fig. 11 presents the castmicrostructures of Mg1.07% Gd (wt.) (G1) and Mg0.92% Zn1.04% Gd0.57% Zr (ZGK110) alloys, and theresulting textures after rolling at 400 C and thicknessreduction of 80%, followed by annealing at 400 C for1 h. It is evident that the strongest texture weakeningoccurred in the ZGK110 alloy during recrystallizationannealing (Fig. 11b). The corresponding texture of the bin-ary G1 alloy revealed a considerable reduction in basalpole intensity of one texture component (compared withthe rolling texture) but the total maximum intensity wasstill high. The combination of non-RE elements (Zn andZr) with Gd seemed to be important also for the develop-ment of an RE-sheet texture with TD-spread during roll-ing, whereas the addition of only Gd to magnesiumresulted in a conventional double-peak basal texture withhigh intensity. Although these are preliminary results, itis obvious from Fig. 11 that the formation of RE texturesand the associated texture randomization during recrystal-lization depend strongly on grain size and the nature of thesecond-phase particles, such as composition, size, shape,

Z31 alloy and an RE-containing alloy ME20 during high (200 C/102 s1)s (outlined) of ME20 oer an estimate of the relative CRSS values of active

ta MJ. Hirsch, T. Al-Samman /Acdistribution, etc. Another important aspect, not directlyseen in Fig. 11 is the nature of solutes and their interaction

Fig. 11. Comparison of the cast microstructures and second-phase precipitatio1.07% Gd (wt.) (G1) and (b) Mg0.92% Zn1.04% Gd0.57% Zr (ZGK110) a

Fig. 12. Comparison of the sheet texture development in rolled binary Mgadditionally Zn and Zr (bottom) as a function of annealing temperature.aterialia 61 (2013) 818843 831with each other and with microstructural features, such asdislocations and grain boundaries.

n, and the resulting rolling and annealing textures between (a) binary Mglloys.

Ce alloy (top) and a quaternary version of the same alloy containing

a MAnother example is given in Fig. 12 for annealing tex-ture modication as a function of annealing temperaturein a binary Mg alloy containing 1.15% Ce and a similarversion containing in addition 1.05% Zn and 0.57% Zr(wt.). While texture weakening was evident in both alloysupon 1 h annealing at 350 and 400 C, it was much morepronounced in the quaternary alloy containing Zn andZr, where the sharp rolling texture (10 times random)was strongly randomized (2 times random). Notably,for annealing temperatures up to 300 C, both alloys exhib-ited similar annealing textures, which were not much dier-ent from typical basal textures of conventional magnesiumsheet. With this information at hand, future investigationswill focus on annealing temperatures between 300 and350 C to examine recrystallization and detect changes inthe microstructure and phase relations between the twoalloys in order to improve our understanding of the roleof non-RE elements in texture randomization duringrecrystallization.

3. Aluminum alloys

3.1. Advantages and applications

In contrast to magnesium, the slightly heavier elementaluminum has already established a leading role in a widerange of applications, due to its light weight, relatively easyfabrication and attractive mechanical properties. The useof these alloys started with their spectacular developmentin the aerospace industry as soon as high-strength variantswere developed and became available in sucient quanti-ties, more than 100 years ago [56]. In the past few decades,increasing amounts of these alloys have also been used inautomotive applications due to the creation of robustand easily applicable variants with good strength, form-ability, crash and corrosion performance [57].

3.2. Aluminum sheet alloys for automotive applications

The two main alloy systems used in automotive applica-tions are the non-heat-treatable AlMg (AA5xxx) and theage-hardened AlMgSi (AA6xxx) alloys, which show agood combination of sucient strength and good formabil-ity [5762]. For specic applications (e.g. in bumpers andcrush-zone elements) the high-strength AlZnMgCu(AA7xxx) aerospace alloys are also being used mostlyas extruded parts. However, due to limitations in corro-sion, joining and age-hardening characteristics, this andthe other high-strength AlCuZnMg (AA2xxx) alloygroup are less suitable for conventional mass-producedautomotive parts. However, as well as alloy additions, ther-momechanical processing is of major importance [63]. Thisimplies close control of texture, which can be quite strongin Al alloys [64] and thus signicantly inuences anisotropyand related key properties, such as yield locus and forming

832 J. Hirsch, T. Al-Samman /Actbehavior [65,66].The strength and formability of the main group of non-heat-treatable AlMg (5xxx) sheet alloys for automotiveapplications is based on the mechanism of solid-solutionhardening by Mg additions, usually up to 5% [58]. Thisalso causes a high strain-hardening eect that helps to sta-bilize the sheet during stretch forming, avoiding local neck-ing. Thus high-Mg-containing AlMg alloys enhance bothkey properties, i.e. strength and formability.

3.3. 5xxx AlMgMn alloys

The use of 5xxx AlMg alloys is well established in chas-sis and various structural applications, due to their goodforming behavior. They are also used in sheet panels, butseldom as exterior parts due to sheet surface irritation ofLuders lines, caused by the PortevinLe Chatelier eect.At room temperature and slow strain rates an avalanchetype dislocation motion occurs in AlMg alloys due tothe strong dislocation interaction (pinning) with diusingMg atoms.

AlMg alloy strength is mostly dened in terms of asoft-annealed O temper. The signicant strength contri-bution due to strain hardening is seldom used in automo-tive parts (in contrast to beverage cans [67]) due to heattreatments during processing (e.g. paint baking) and thesoftening eects involved in any use phase at elevated tem-peratures. If exposed to elevated temperatures for longtimes medium (max. 3%) Mg-containing alloys are useddue to the eect of intercrystalline corrosion (IC) in corro-sive environments. In such cases for alloys containing(>3%) Mg, special precautions in either processing or inapplication conditions are required, while for predominantmoderate-temperature applications (e.g. in marine environ-ments) high (5%) Mg alloys are common. In recent yearsnew variants of AlMg alloys with higher Mg contents (e.g.3.5%) [69] or with small additions of Cu have been investi-gated showing improved properties, some of which areused for body-in-white applications by Japanese automo-bile companies [68].

3.4. 6xxx AlMgSi alloys

6xxx AlMgSi are the well-established age-hardeningsheet alloys used in many automotive parts, including exte-rior panels, which have high requirements for surfaceappearance. They also provide a high eective strength inthe age-hardened T6 condition, and lower strength andgood formability in the T4 (solution annealed and agedat room temperature) condition [58,70]. Their characteris-tic strength evolution occurs following an additional heattreatment which is applied in car production after forming,in some cases simply during paint baking cycle of the body-in-white. In the temperature range around 185 C a signif-icant increase in strength occurs by precipitation hardeningwhen the sheet was processed by a high solution anneal

aterialia 61 (2013) 818843(>540 C) with fast quench in a continuous annealing line.

ta MThe age-hardening behavior can be further enhanced by aspecial pre-treatment pre-bake (avoiding the specicretrogressive eect observed after long-term room-temperature aging) and by deformation applied duringthe production of parts, thus achieving sucient strengthafter part forming and paint baking [70].

3.5. Other alloying additions

Mn has a benecial inuence on the mechanical proper-ties and is added mainly to control grain size by formingsub-micrometer-sized particles [71]. The microstructuralchanges involved during annealing often include textureeects since particles modify textures by changing therecrystallization mechanisms (enhancing PSN [72]) and/orinhibiting grain boundary mobility. This is especiallyimportant in AA6xxx alloy sheet due to the required highsolution annealing temperatures that otherwise inducegrain growth. The corresponding recrystallization mecha-nisms can most eectively be controlled by choosing theappropriate processing routes, especially the heat treat-ments, starting with ingot pre-annealing and homogeniza-tion [73,74].

Other elements, such as Fe, that must also be consideredin industrially processed alloys, are usually present asimpurities. Due to their low solubility they can have a neg-ative eect on formability by forming large constituentparticles during eutectic solidication. They are the sourceof local cracking during forming operations. Proper ingothomogenization cycles, additions of Mn and large rollingreductions aect their form and size and help to reducetheir negative inuence, as well as aecting texture forma-tion during intermediate and nal annealing treatments[73,74].

3.6. Forming characterization and simulation

In simple characterization of sheet material behaviorformability is often expressed by parameters such asstrain-hardening coecient n-value and anisotropy coef-cient r-values, which are easily measurable in simpletensile tests and were therefore used in early forming simu-lation software. Advanced forming simulation methods areneeded, however, for anisotropic materials (such as alumi-num or advanced steel) which require more detailed (up to16 or even more) parameters to fully describe the materialow under various conditions [66]. These models weredeveloped for aluminum sheet where some anisotropyeects need to be considered due to complex textures. Con-ventional interstitial-free steels, in contrast, show rathersimple ber textures with rotational symmetry in the sheetnormal [75], which is easily described by classical BishopHill continuum models. For new steel grades (e.g. TWIP/TRIP steels), however, the advanced models are now beingapplied since directional eects play a major role in their

J. Hirsch, T. Al-Samman /Acspecic forming behavior. New crystal plasticity-basedcodes have also been developed and applied in order toinclude detailed texture data to describe Taylor-typecrystallographic slip and predict related anisotropy eects[76].

3.7. Textures in aluminum alloy sheet

3.7.1. Rolling textures in aluminum sheet

During cold rolling of any aluminum alloy a typical fccrolling texture develops, which is composed of two bers[77]: (i) the a-bre (= h100i parallel to the normal direc-tion, ND, which connects the {011} h112i B orientationwith the {011} h011i G (Goss) orientation) occurring atlow strains; and (ii) the b-bre which aligns the main threetexture components C (copper), S and B (brass). Bothbers show some characteristic scatterings and peaks,depending on the initial (casting or recrystallization) tex-ture, and thus on the preceding processing history.

In nal customer operations the fully annealed tempersare the most widely used conditions (except for beveragecan production [67]). Hence, the corresponding recrystalli-zation textures are usually more relevant, and aect thecustomer forming processes involved. The recrystallizationtextures of rolled and annealed aluminum alloys have beeninvestigated in detail and fully described (e.g. [78,79,84]) byadvanced texture evaluation methods (X-ray diraction,EBSD).

3.7.2. The Cube texture

The most prominent fcc recrystallization texture is theclassicalCube component (Fig. 13),which appears inmosthot-rolled Al-alloys with characteristic shifts or scatteringsof the orientation peak [78]. After cold rolling and annealingit is dominant in high-purityAl alloys.However, also in non-heat-treatable Al-Mg (AA5xxx) and in age-hardenableAlMgSi (AA6xxx) sheet alloys this component can nucleateand grow signicantly during hot rolling and fast and high-temperature heating of the required nal solution annealingtreatment, aecting strength and formability [80].

The Cube texture originates from a classical nucleationmechanism of metastable orientations existing as smallbands in the highly rolled microstructure. In hot rollingthey gain in stability due to the activation of non-octahe-dral slip systems. These Cube bands are surrounded bydivergent orientation zones which rapidly enhance thelocal orientation gradient needed for sucient grainboundary mobility. Growth is then strongly accelerated,supported by a preferred oriented growth eect due to a40 h111i orientation relationship to the surroundingmain rolling texture (b-bre) orientations, including allof its variants: due to its high symmetry the Cube orien-tation can compensate minor deviations from this rela-tionship and resistant grains can easily be by-passedwhen surrounded. This explains why the Cube texture isthe best compromise orientation according to recrystalli-zation simulations based on statistical orientation rela-

aterialia 61 (2013) 818843 833tionships, including preferred nucleation and orientedgrowth eects.

a M834 J. Hirsch, T. Al-Samman /Act3.7.3. The R texture

Another prominent recrystallization texture componentin rolled and annealed Al sheet is the R component(Fig. 13c), which is very similar to the S rolling texturecomponent, indicating either an in situ recrystallizationor a strain-induced boundary migration (SIBM) nucleationprocess with subsequent growth [81]. It mainly occurs inFe-containing (AA1xxx and AA8xxx) alloys and is usuallynot observed in rolled and recrystallized AlMg (5xxx) andAlMgSi (6xxx) alloys used as automotive sheet. It thus

Fig. 13. Cube texture component in Al alloys. (a) Hot-rolled AA5182 AlMsection only). (c) Cube and R texture in 95% rolled commercial-purity Al (AA(AA6111).aterialia 61 (2013) 818843can be concluded that the recrystallization mechanismsrelated to the R texture are not dominant in these alloys.

The R texture is interpreted as a retained rolling textureS component, but often shows distinct dierences in inten-sity and exact position. If Cube nucleation is suppressedearly (e.g. by simultaneous precipitation eects) or runsout of nucleation sites (e.g. at extreme rolling reductionsas in foil rolling) and no other nucleation sites are activethe R-orientation emerges and can grow substantially[82]. It is assumed to nucleate from existing orientations,

g (full ODF). (b) Cold-rolled and annealed AA6016 AlMgSi (u2 = 01145). (d) ND rotated Cube and P texture in age-hardenable AlMgSi

e.g. by a bulging SIBM mechanism at existing grain bound-aries. It can grow preferentially by a 40h111i relationshipinto one of its own symmetrically equivalent variants. Itthen shows systematic deviations from the original S orien-tation, either due to preferred nucleation eects (near the Corientation) or preferred growth eects (near the 40h111irotated S orientation) [81].

3.7.4. PSN recrystallization textures in aluminum

In industrial rolled and recrystallized AlMg (5xxx) andAlMgSi (6xxx) alloys recrystallization nuclei often origi-nate at deformation inhomogeneities, i.e. particles or shearbands [80]. The latter is the Q component (near{013}h231i) which was observed to originate at shearbands as observed in highly cold-rolled AlMg alloys

illustrated in Fig. 14ac where the nal recrystallizationtexture is inuenced by the homogenization during the ini-tial ingot preheating treatment, with other process param-eters kept constant [79,80]. In Fig. 14 the exact cubeorientation together with its RD scatter increase at theexpense of the PSN (e.g. P) orientations with increasinghomogenization treatment. The recrystallization texturesin AlMgSi alloys consist of the Cube orientation withsome RD and/or ND scatter and/or the P orientation(and some R orientation) in dierent proportions depend-ing on the particles formed during early processing. Herealso the supersaturation level plays a role which can inhibitearly growth of preferred (Cube) nuclei by simultaneousne precipitation. The decrease of cubeND and P (the R ori-entationnot shown in Fig. 14appeared to scale with

J. Hirsch, T. Al-Samman /Acta Materialia 61 (2013) 818843 835[83]. The rst is the P (PSN [72]) component (near {011}h122i) originating from local deformation inhomogeneitiesaround second-phase particles. Both mechanisms lead tosignicantly lower texture intensities, as compared to Cubeand R textures (Fig. 13c and d), but certain texture eectscan be observed and can help to reveal the underlyingmechanism.

PSN is an important recrystallization mechanism inindustrial aluminum sheet, inuencing textures mainly bygenerating a nearly random orientation distribution. How-ever, some characteristic texture eects can be observed inthe weak recrystallization texture (as shown, e.g., inFig. 13d) revealing a peak in {011} h122i P and in{001} h310i ND rotated cube; the main peaks are veryclose to orientation accumulations predicted from the roll-ing texture by a complete 40 h111i transformation. Thistexture is interpreted as a preferred growth selection pro-cess taking place out of the large spectrum of potentialnuclei around particles, e.g. in early stages of recrystalliza-tion (microscale growth selection).

In heat-treatable AlMgSi AA6xxx alloys recrystalliza-tion is governed by PSN but this also competes with othernucleation mechanisms, in particular with the Cube nuclei.Hence the nal texture can vary and depends on the rela-tive amounts of nucleation at either orientation. This isFig. 14. Particle-stimulated nucleation (PSN) texture in cold-rolled and annprocesses.the Cube orientation) reveals a decrease in the eciencyof PSN with increasing precipitation of Mg2Si particlesduring hot rolling and/or nal annealing. These nely dis-persed particles tend to impede PSN more strongly thannucleation at the Cube bands. Hence, nuclei emerging fromthe Cube bands (i.e. Cube- and cubeRD-oriented grains)and the grain boundaries (i.e. R-oriented grains) will pre-vail in the recrystallization texture [80].

3.8. Texture eects on mechanical properties

Fig. 15 illustrates the resulting plastic anisotropy interms of the variation in r(a) (Fig. 15a) and the (symme-trized) earing proles h(a) of deep-drawn cups (Fig. 15b)for the three dierent materials. The texture variations havea strong impact on the resulting plastic anisotropy since thePSN textures (CubeND and P orientations, e.g. homogeni-zation states C-1 and C-2, see Fig. 14a and b) generate max-ima for both earing and r-values at an angle of about 45 tothe RD and minima at 0/90. Theoretical r-values werepredicted from the texture data [73] to give the completer(a)-curves tted to the experimental points in Fig. 15a.Texture control can balance the pronounced anisotropicproperties of the prevalent Cube-recrystallization texturewith the opposite 0/90 ears and minimum r-values atealed AA 6010 AlMgSi alloy sheet with dierent ingot pre-annealing

s (s

a M45. The resulting texture variation plays an important rolein the material properties since their opposite eect onanisotropy can be used to improve the formability of AlMgSi autobody sheet.

3.8.1. StrengthAny variation in crystallographic texture has an imme-

diate impact on material strength due to the changes ingeometry of the dierent slip systems activated upon ongo-ing or subsequent deformation. The simplest explanationfor this can be given by assuming single slip only on the slipsystem with the highest orientation factor m (= 1/Schmidfactor l = cos k cos v, with k and v being the anglebetween the tensile and slip plane normal and the slip direc-tion, respectively) in a tensile test where the variation canbe up to 84% (m varies from 3.67 for the h111i directionto 2 for a central direction with k = v = 45), neglectingany orientation-dependent strain-hardening eects duringdeformation! For the other extreme of multiple slip (a min-imum of ve slip systems) as dened by the Taylor modelthe variation still can be as high as 61%. Any slip system

Fig. 15. Plastic anisotropy of the textures shown in Fig. 14: (a) r-value(experimental).

836 J. Hirsch, T. Al-Samman /Actactivation will occur within the limits of these two models,and hence the variation range in strength can be estimatedaccordingly.

3.8.2. Formability

The Cube texture is typical for all aluminum alloysrecrystallizing during hot rolling. In a fully recrystallizedAlMgMn hot strip the Cube texture causes a character-istic anisotropic behavior, e.g. as strong 0/90 earing in acup deep drawn from a circular blank as shown in Fig. 15b.It is technically quite important, e.g. for AlMgMn canbody stock, where a strong hot strip Cube texture isrequired for low earing and good formability of the cold-rolled nish gauge sheet [67].

Full or signicant partial recrystallization in the pre-ceding rolling stepseither during hot rolling or by inter-mediate annealingconsiderably reduces the nal Cubetexture strength. This means that a large recrystallizedfraction during hot-rolling passes reduces Cube texturestrength [71], while a short interpass times (as in multi-stand hot-rolling lines) or reduced temperatures increaseit. However, minimum Cube texture is needed when hotstrips are used for forming operations and a more isotro-pic deformation behavior is required, e.g. for the deepdrawing of wheel disk fabricated from an 8 mm AlMg3hot strip. Here optimum formability is provided by con-trolled partial recrystallization where some of the hot-roll-ing texture is preserved to eectively balance the oppositeanisotropy of the Cube recrystallization texture. Analmost even cup rim prole can be produced which other-wise may lead to severe deviations in the rotationalsymmetry.

To describe and compare sheet formability properties,forming limit diagrams (FLDs) are used [65]. The maxi-mum strain achieved under all biaxial strain statesdepends on the strain hardening (e.g. n-value) and localow pattern, including the anisotropic material ow. Itis evaluated from tests under dierent loading/strainingconditions. In recent years methods have been developedto generate FLDs from their principal determining factors

ymbols: experiments; lines: simulated from textures); (b) earing proles

aterialia 61 (2013) 818843[85]. The simulation shows how strain hardening aectsthe level of the FLD curves and texture modies theirshape.

Fig. 16 shows the results of an industrial AA6016 sheet,processed in two dierent (proprietary) ways, revealing apronounced formation of a PSN texture with a signicantincrease in formability (6016-A30+) [86]. Anotherimportant aspect of the formability of AlMgSi AA6016sheet (also encountered in steel sheet forming) is the forma-tion of uneven surfaces during sheet forming (called rop-ing). As indicated in Fig. 17, a localized band-likeformation of many ne-grained Cube orientations isresponsible for this eect. Systematic process variationsare applied to control the texture in order to reduce thevolume and uneven distribution of Cube grain orienta-tions and thus avoid roping. Both cases demonstrate theimportance of texture control in Al sheet processing toensure quality and enhance performance for automotiveapplications.

ta MJ. Hirsch, T. Al-Samman /Ac4. Summary and concluding remarks

The common features and signicant dierences of tex-ture formation during sheet rolling and annealing havebeen analyzed for the two most important lightweight met-als used in automotive applications. Conventional thermo-mechanical processing of aluminum and magnesium alloys

Fig. 16. Textures of AA6016 autobody sheet with processing variations.(a) Conventionally processed AA6016. (b) Advanced processed AA6016.

Fig. 17. EBSD of AA6016 surface showing aligned recrystallized grains incube texture.produces sharp ber textures, which in fcc aluminum accu-mulates in the b-ber with characteristic peaks, dependingon the preceding processes and initial textures, whereas formagnesium the basal plane rapidly aligns parallel to thedirection of primary material ow. While aluminumdeforms by classical slip on {111}h110i slip planes, easilyaccommodating strain incompatibility in any orientation,mechanical twinning in magnesium involves only the c-axisdeformation mode at room temperature with drastic conse-quences for texture and resulting mechanical performance.For standard Mg sheet or extrusion textures, subsequenttensile or compressive loading parallel or perpendicularto the predominant ow direction requires deformationalong the c-axis, which under ambient conditions meansaccommodation by mechanical twinning. For c-axis con-traction, f1 0 1 1g twinning only activates at very highstresses at which void nucleation occurs. For c-axis exten-sion, f1 0 1 2g twinning is not very helpful, as it hardensthe texture by abruptly rotating it into contraction alongthe c-axis. The diculty of forming magnesium at temper-atures below 150 C is, therefore, attributed not only to alack of respective deformation modes but also to the con-sequence of activating twinning, particularly under c-axiscontraction.

As highlighted in this paper, attempts to enhance theambient formability of magnesium alloys have aimed atlowering the stress for non-basal slip activity, by alloyingwith Li for example, or by combining texture optimizationand grain renement obtained by conventional ornon-conventional processing. Additions of RE elementshave shown promise in this regard. Yttrium addition, forexample, has a ductilizing eect on ambient formabilitythat was attributed to enhanced pyramidal hc + ai-slipactivity associated with a signicant decrease of the intrin-sic stacking fault energy. Another favorable eect of REelements is on texture modication that seems to hold formischmetal additions or single additions of Ce, La, Nd,Gd, etc. The level of texture modication depends stronglyon the RE concentration, the processing parameters andsometimes on other non-RE alloying elements, such asZn. Modied textures can show a decrease in texture inten-sity up to random level or exhibit RE-texture compo-nents that depart from standard sharp sheet andextrusion textures. The weaker textures produced in thesecases are better aligned for basal slip, which results inenhanced ductility and reduced mechanical anisotropy.As a secondary eect, addition of RE elements also renesthe grain size of the alloy, which confers an additional duc-tility benet.

In rolled and annealed aluminum alloys the eect ofrecrystallization generates a spectrum of new orientationsand texture components. Their interpretation and themethods derived to control them for achieving advancedproperties in practical applications are highly developedand industrially applied. The classical Cube component

aterialia 61 (2013) 818843 837occurs in many hot-rolled, as well as cold-rolled andannealed Al alloys, in some cases with characteristic shifts

a Mand/or scatterings. It is a major inuence on many aniso-tropic properties, including forming behavior and relatedsurface eects. The R-component is very similar to themain S rolling texture component indicating either anin situ recrystallization (i.e. extended recovery) or SIBMnucleation processes with subsequent growth. In industrialalloys second-phase particles play a major role, modifyingthe Cube texture and generating the P-component. Withvariations in particle size and distribution and in the solutecontent of ne precipitating elements suppressing the earlyCube nucleation dominance, the textures in AlMgSialloys can be systematically varied and properties modiedand formability signicantly improved. Other weak com-ponents originate at deformation inhomogeneities such asshear bands, typical for weak recrystallization textures ofhighly cold-rolled and annealed high-Mg-containing AlMg AA5xxx alloys.

For magnesium alloys containing RE elements, recrys-tallization plays a signicant role in modifying the defor-mation texture. With the employment of moderncharacterization techniques, such as electron backscatterdiraction, atom probe tomography, high-resolutionenergy dispersive X-ray spectroscopy and rst-principlesmodeling, the mechanistic understanding of deformationand dynamic recrystallization in correlation with RE ele-ments has been markedly improved. Numerous mecha-nisms based on solid-solution and particle eects havebeen suggested to be primarily responsible for the RE tex-ture modication. The most important solute-relatedmechanism involves solute segregation to grain boundariesthat serves to retard recrystallization, which has crucialeects on the nucleation characteristics (e.g. the spectrumof orientations that can nucleate or the type of nucleationsites) and the relative mobility of dierent boundaries.

PSN is also observed in magnesium alloys, and seems tobe particularly promoted in the those containing RE addi-tions. This mechanism provides more randomly orientednuclei that weaken the sharp rolling or extrusion textures.It may appear to involve small volume fractions of themicrostructure, and be hence considered less important,but in fact the orientations concerned may dominate thebulk texture following recrystallization and grain growth.

In many cases, where both solutes and particles are pres-ent, the RE-induced texture change can be expected to beimpacted by several mechanisms at once. While the roleof recrystallization has been suciently highlighted bymany authors, the role of deformation in nucleating somemechanisms for texture modication prior to recrystalliza-tion should also be considered important. In particular,whether the deformed microstructure allows better devel-opment of deformation heterogeneities, such as shearbands or other strain localizations that contain larger crys-tal rotations, or whether the twinning behavior shows anydistinctions from traditional Mg alloys regarding certaintwin modes (e.g. compression twins becoming equally

838 J. Hirsch, T. Al-Samman /Actfavorable with tension twins when loaded accordingly).The favorable eect of some RE elements in providing mul-tiple slip conditions can also have an impact on the nucle-ation of grains with distinct orientations.

It is anticipated that microalloying of RE elements com-bined with optimized processing for grain renement andsoft textures are going to help future conventionally pro-cessed wroughtmagnesium alloys become an important partof automotive applications. It is, however, noteworthy thatin some cases, and after a careful choice ofmany parameters,soft textures can be achieved with traditional (non-RE)alloys, such as AZ31 or AM50. We showed that a combina-tion of extrusion and uniaxial compression at 400 C in theextrusion direction succeeds in randomizing the texture toa large extent, which demonstrated positive implicationsfor subsequent ambient forming for this class of alloys.

In both magnesium and aluminum the formation ofrecrystallization texture is controlled by the limited nucle-ation events and/or by a subsequent process of growthselection out of the corresponding orientation spectrum.In both materials specic alloying elements can be addedand thermomechanical processing modied so that theyaect the basic recrystallization mechanisms and allowthe texture, and thus the nal properties, to be varied.The principal texture eects on product properties showvariations in:

strength, either by the orientation dependence of slip orindirect by changing slip modes, e.g. during agehardening;

formability, where deformation geometry is aected, e.g.in the form of the r-value in tensile deformation or cupheight and ange thickness in deep-drawing operations.

Since property control is a key issue in many practicalapplications, textures and related process variables areaccordingly being industrially controlled. Furthermore, ifanalyzed carefully, useful information can also be obtainedabout the underlying physical processes involved in indus-trially processed material. These two aspects combine boththe fundamental and the practical aspects of textureresearch for material control and process optimization inindustrial applications.

Acknowledgments

T.A.S. would like to acknowledge nancial supportfrom the Deutsche Forschungsgemeinschaft (DFG), GrantNo. (AL 1343/1-1), and D.A. Molodov for helpful discus-sions. T.A.S. also thanks Indranil Basu, Feng Jiao, Xiao-hui Li and Konstantin Molodov for their eorts in thereported results. The assistance of Arndt Ziemons in pro-ducing magnesium alloys is sincerely appreciated. J.H. isgrateful to Olaf Engler and other former coworkers atthe Institute of Physical Metallurgy and Metal Physics(IMM) at RWTH Aachen University and current

aterialia 61 (2013) 818843colleagues at Hydro Aluminium R&D Center, Bonn for

ta Mmany years of fruitful co-operation and many helpfuldiscussions.

Appendix A. Representation of texture in cubic and

hexagonal materials

Texture is a collective term for a non-uniform distribu-tion of crystallographic orientations in a polycrystallineaggregate [87]. A sample in which crystallographic orien-tations are fully random is called a random texture or isoften said to have no texture. If the crystallographicorientations are not random, but have some preferred ori-entation, then the sample has a pronounced texture. Intexture measurement using X-ray or electron beam dif-fraction, the volume fraction of a particular family ofcrystallographic planes with respect to a sample coordi-nate system is quantied. The axes of a sample coordinatesystem are normally chosen according to important direc-tions associated with the external shape change of thespecimen. For rolling deformation those are usually therolling direction (RD), the normal direction (ND) andthe transverse direction (TD). The texture obtained in thisway is an average value of the whole sample volume (pro-vided that that the measured sample volume is statisticallylarge enough).

In order to understand how crystallographic textures areformed and how they aect material properties, it is essen-tial to have a method of representing and characterizingthem. Data obtained from texture measurements can bepresented graphically in dierent ways. The most importantones are pole gures (PFs), inverse pole gures (IPFs) andorientation distribution function (ODFs) [87,88].

1. Pole gure. PFs are two-dimensional stereographic plotsthat describe the relation between the orientation of thecrystallographic axes of a crystal with respect to a coor-dinate system of a sample. An (hk l) pole gure showsthe distribution of the {hkl} poles in the sample.

2. Inverse pole gure. IPFs are two-dimensional stereo-graphic plots that use the crystal axes as the referenceframe instead of the sample axes. They show the distri-bution of crystallographic axes parallel to a specic sam-ple direction. In the case of uniaxial deformationgeometry (extrusion, uniaxial compression, etc.) IPFsare very convenient for describing the texture, since incontrast to PFs they require only one well-dened sam-ple direction.

3. Orientation distribution function. The full three-dimen-sional representation of crystallographic texture is givenby the ODF, which is not measured directly but rathercalculated by combining the data from several experi-mentally measured pole gures. An ODF describes theorientation of each crystal relative to three Euler angles,i.e. u1, U and u2 (Bunge notation) that, in a simpliedsense, dene the dierence in orientation between the

J. Hirsch, T. Al-Samman /Accrystallographic axes and the sample