Texture evolution of ﬁve wrought magnesium alloys during route...

Acta Materialia 53 (2005) 3135–3146

www.actamat-journals.com

Texture evolution of five wrought magnesium alloys during routeA equal channel angular extrusion: Experiments and simulations

S.R. Agnew a,*, P. Mehrotra a, T.M. Lillo b, G.M. Stoica c, P.K. Liaw c

a Department of Materials Science and Engineering, University of Virginia, Charlottesville, VA 24590, USAb Environmental and Energy Sciences Division, Idaho National Engineering and Environmental Laboratory, Idaho Falls, ID 83415, USA

c Department of Materials Science and Engineering, University of Tennessee, Knoxville, TN 37996, USA

Received 20 January 2005; received in revised form 9 February 2005; accepted 10 February 2005

Abstract

Equal channel angular extrusion (ECAE) has been demonstrated to induce unusual deformation textures and resulting properties

in magnesium alloys, such as the remarkably enhanced room temperature ductility first reported by Mukai et al. [Mukai T, Yamanoi

M, Watanabe H, Higashi K. Scr Mater 2001;45:89]. This paper documents a wide range of textures which evolve during ECAE of

magnesium alloys. The fact that different alloys exhibit different texture evolutions is an indication of distinctions in the balance of

deformation mechanisms which operate within the different alloys. Polycrystal plasticity modeling is used to develop explanations

for these texture distinctions in terms of the relative activities of non-basal secondary slip modes, involving Æaæ and Æc + aæ type dis-locations. AZ alloys appear to exhibit balanced secondary slip of non-basal Æaæ and Æc + aæ dislocations, while ZK60 and WE43

appear to favor non-basal Æc + aæ slip. A binary Mg–Li alloy exhibits a radically distinct texture evolution, which is associated with

large-scale strain accommodation by non-basal Æaæ slip.� 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

Keywords: Magnesium; Texture; ECAE; ECAP; Crystal plasticity; Ductility

1. Introduction

Texture can result in strong anisotropy and asymme-

try, particularly in the case of the plastic behavior ofnon-cubic materials, such as hexagonal close packed

(hcp) magnesium and its alloys. One example is the ten-

sion/compression strength asymmetry exhibited by mag-

nesium extrusions, where the yield strength in

compression along the prior extrusion axis may be as lit-

tle as half that of the tensile yield strength along the

same direction (e.g., Refs. [1,2]).

If the flow stresses and hardening behaviors of theindividual deformation mechanisms are known, poly-

crystal-plasticity models provide a means of predicting

1359-6454/$30.00 � 2005 Acta Materialia Inc. Published by Elsevier Ltd. A

doi:10.1016/j.actamat.2005.02.019

* Corresponding author. Tel.: +1 434 924 0605; fax: +1 434 982 5660.

E-mail address: [email protected] (S.R. Agnew).

the evolution of texture given boundary conditions

appropriate for the prescribed deformation [3]. Con-

versely, if the texture evolution and strain path are well

characterized, an investigation of the underlying defor-mation mechanisms is possible [4]. Both avenues are

examined in the context of equal channel angular extru-

sion of magnesium alloys in the present paper.

1.1. Equal channel angular extrusion

A novel metal forming process, which has come to be

known as equal channel angular extrusion (ECAE) (e.g.,

Ref. [5]) or pressing (ECAP) (e.g., Ref. [6]), can impart

enormous levels of strain without significantly changing

the overall dimensions of the sample. In this regard, it is

similar to torsion, which has also been used extensivelyin the study of strain hardening [7,8] and hot working

ll rights reserved.

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3136 S.R. Agnew et al. / Acta Materialia 53 (2005) 3135–3146

[9,10]. Some researchers have also noted a similarity in

the geometry of the ECAE process and metal cutting

[11]. The main interest in such techniques of severe plas-

tic deformation (SPD) has been the potential to reduce

the grain-size of the material [12]. The concept of reduc-

ing the microstructure scale using SPD is not new (e.g.,Refs. [13,14]). In fact, it is a core tenet denoted ‘‘simili-

tude’’ in a leading theory of strain hardening [15]. The

interest in grain-size reduction is driven by the possibil-

ity to produce ultrahigh strength metals [16] and high

strain rate superplasticity (e.g., Refs. [17,18]). Other

methods of manufacturing nanocrystalline or ultrafine-

grained metals have been plagued by extrinsic defects

[19,20], such as cracks and pores, limiting their abilityto push the limits of grain-size strengthening or high rate

superplasticity. SPD has proven capable of increasing

both the strength and ductility of metals [21] as well as

consistently producing materials that exhibit high strain

rate superplasticity [17].

A final area of ECAE research has focused on

modeling the plastic deformation of metals and the cor-

responding texture evolution, in particular. Rather thanenforcing primarily pure shear characteristic of conven-

tional rolling or extrusion, ECAE essentially enforces

simple shear at the intersection of two channels

(Fig. 1). In addition to the unique deformation geome-

try, the work-piece may be passed through the die

multiple times for further deformation. A nomenclature

has been developed in the literature [6] to describe multi-

ple pass deformation routes, where route A denotes nosample rotations about the billet axis between passes,

route BA denotes 90� back-and-forth rotations, route

BC (simplified to route B in the current work) corre-

sponds to continuous 90� rotations, and route C denotes

180� rotations. It has been noted that these routes have

varying degrees of redundancy in the strain they impart,

with route A being the least redundant, and route C

essentially ‘‘undoing’’ the strain each second pass [22].Thus, the ECAE process offers a distinct possibility to

generate unique textures [23]. Comparisons of

experimental textures and those simulated using poly-

+

90˚

==

y'x' y

xzz'

(a) (b)

Fig. 1. Schematic of the ECAE processing showing (a) an idealized

case where a simple shear is localized within an infinitesimal band at

the intersection of the channels and (b) where the simple shear

accumulates along flow lines, which simultaneously rotate.

crystal-plasticity simulations have been used to test the

underlying assumptions about the strain path during

ECAE [24–27]. Researchers have also explored the use

of polycrystal models as a foundation for predicting

the microstructure refinement, which will take place dur-

ing deformation, and ECAE has been used as a test case[28]. The vast majority of ECAE texture studies have fo-

cused on cubic metals, such as Al (e.g., Ref. [27]), Cu

(e.g., Refs. [24,26]), Fe [25,29], and their alloys. There

have been comparatively few studies of ECAE textures

of non-cubic metals, such as Be [30], Mg [Error! Book-

mark not defined, 31, 32], Ti [33], and Zr [34].

1.2. Wrought magnesium alloys

Wrought magnesium alloys exhibit marginal cold

formability and the accepted explanation is the intrinsic

plastic anisotropy of the magnesium single-crystal [35].

However, recent modeling and transmission electron

microscopy (TEM) of deformed polycrystals has

stressed that magnesium grains may be less anisotropic

than single-crystal measurements have suggested[36,37]. It has also been reported that the effect of

grain-size upon the strength of magnesium is potent

due a high apparent Taylor factors [38,39], and the effect

upon ductility may also be dramatic [40]. The current

study of ECAE processing of magnesium alloys was ini-

tiated due to the potential to produce simultaneously

high-strength, high-ductility, and low-density alloys

through grain refinement [41–43].The range of textures observed in conventional mag-

nesium wrought products is limited and the potentially

positive impact on the mechanical behavior of magne-

sium alloys by developing a unique texture using ECAE

processing was first reported by Mukai et al. [75]. Alloy

AZ31 exhibited a 2–3 times improvement in the tensile

elongation after ECAE processing and annealing (in or-

der to restore the original grain-size of �15 lm). It wassuggested that the ductility improvement was due to the

randomization of the originally strong extrusion texture.

Subsequent work verified the incredible improvement in

ductility for both AZ61 [31] and AZ31 [32] and showed

that it was due to a change in the texture. However, the

resulting texture was shown to be stronger than the ori-

ginal extrusion texture, not weaker [32].

The objective of the current study is to determine thedeformation texture evolution which is enforced by

ECAE processing of five magnesium alloys (Table 1)

subjected to various processing routes. These texture

data could be useful in designing processing treatments

to explore the effect of texture on various behaviors

in magnesium (and some other hcp alloys as well).

The present article focuses on route A processing

and the possibility to make inferences regarding therelative activities of deformation mechanisms (slip

modes) within the different magnesium alloys using

Table 1

Alloys examined in this study, their nominal compositions, and typical applications

Alloy Al Zn RE Zr Mn Other Applications

AZ31 3 1 – – 0.3 Fe, Ni, Cu low Forgings, extrusions, photoengraving plate

AZ80 8 0.5 – – 0.3 High-strength forgings

ZK60 – 6 0.5 – High-strength forgings

WE43 – – 4a 0.5 – 3Y High-temperature castings, forgings

Mg4L – – – – – 4Li Experimental alloy

a RE mischmetal, rich in Nd in this case.

S.R. Agnew et al. / Acta Materialia 53 (2005) 3135–3146 3137

polycrystal-plasticity simulations [4]. A companion pa-

per [44] documents the texture evolutions of some of

the alloys subjected to route B and C processing.

2. Procedures

2.1. Material

The selected alloys are representative of the two most

common wrought alloy families, AZ (aluminum–zinc)

and ZK (zinc–zirconium); a new class of commercial al-

loys designed for use in high-temperature applications,

WE (yttrium-rare earth) [45]; and, finally, an experimen-

tal binary alloy with 4 weight percent (wt%) Li (near the

solid solubility limit of the a-hcp phase) that has beenthe focus of a number of prior studies [4,46–49]. The al-

loys used in this study were received from commercial

vendors listed in Table 2, with the exception of the

experimental lithium containing alloy, which was pro-

cessed at the Oak Ridge National Laboratory (ORNL).

For this alloy, the pure metals (99.99 wt% Mg, and

99.9% Li) were alloyed in a mild steel crucible within a

resistively heated (Pt-wire) furnace (�800 �C) andpoured into a heated (�150–200 �C) Cu mold (74 mm

diameter). The alloying and casting operations were

conducted within a glove-box, which was evacuated

and back-filled with a mixture of Ar, SF6 and O2 as a

protective cover gas. The ingot was, then, directly ex-

truded (4:1) into a 35 mm diameter bar. Conventional

metallography, including a picral-acetic etch commonly

used for magnesium alloys [50], revealed an essentiallyequiaxed grain structure in all of the as-received alloys

Table 2

Vendor and as-received condition of alloys examined

Alloy Vendor Initial condition

AZ31 Mark Metals Plate (25 mm)

AZ31 Timminco Ltd. Cast extrusion

AZ31 Mark Metals Extrusion (32 m

AZ80 Intercontinental Manuf. Extrusion (178

ZK60 Mark Metals Extrusion (76 m

WE43 Magnesium Elektron Extrusion (50 m

Mg4L n/a Extrusion (35 m

a O – fully annealed, F – as deformed.

with the exception ZK60, which had some larger elon-

gated grains dispersed among a matrix of fine equiaxed

grains [51].

2.2. Texture measurements and analyses

Crystallographic texture measurements were con-

ducted using X-ray diffraction in the reflection geometrywith a four circle goniometer and Cu Ka radiation.

Experimental f10�10g, (0002), and f10�11g, pole figureswere collected on a 5� · 5� grid for sample tilts,

a = 0�80�, and azimuthal rotations, u = 0�355�. Defo-

cusing corrections were made using experimentally

determined defocusing curves from random powder

samples. Complete orientation distributions and recal-

culated full pole figures were determined using theWIMV algorithm within the popLA (preferred orienta-

tion package of Los Alamos) code [52]. The texture sam-

ples were sections from the flow plane (or x–y plane in

Fig. 1). The X-ray spot (�2 mm) was focused on the

more uniformly deformed central portion of the billets.

2.3. ECAE processing

Equal channel angular extrusion was performed at

the Idaho National Engineering and Environmental

Laboratory (INEEL). The die utilized has an included

angle u = 90� between the square (22 mm) entrance

and exit channels (Fig. 1). There is no relief radius in

the die corner, rather a sliding bottom wall on the exit

channel after the patented design by Segal et al. [53].

Such a design has been shown to reduce the strain inho-mogeneity discussed above. The die was operated within

Tempera Grain size (lm)

O 50 [51]

billet n/a Large (mm)

m B) F 6 [32]

mm B) F 140

m plate) F Bimodal � 10, 50 [51]

m B) F 6.5

m B) F 38

Table 3

Processing conditions explored during this investigation

Alloy/condition Temp. (�C) Passes/route

AZ31/plate 300 1, 2, 4, 8A

AZ31/cast 300, 200BPa 1A and 2, 8B

AZ31/extr. 200BP 1, 2A and 2, 8B

AZ80/extr. 200BP 1A and 2, 8B

ZK60/extr. 260 1, 2, 4, 8A and 4B and 4C

WE43/extr. 325 2, 4A

Mg4L/extr. 260 1, 2, 4A

BP, back-pressure.a Initial pass at 300 �C, subsequent passes at 200 �C.


a 445 kN universal testing machine. Billets were ma-

chined from the as-received materials with dimensions

of 22 · 22 · 100–150 mm for ECAE processing. It was

necessary to process all of these alloys at elevated-tem-

peratures ranging from 175 to 325 �C (see Table 3) in or-

der to avoid fracturing the billets. Depending upon the

alloy, a temperature between 200 and 300 �C was first

attempted in view of hot-workability data from the liter-ature [1]. The temperature adopted was the lowest at

which successful extrusions were possible, in order to

promote grain refinement (see Table 3). The whole

ECAE die is heated to the desired temperature using

built-in cartridge heaters. The billets were lubricated

with MoS2 paste, inserted into the die, held until they

reached the processing temperature (approximately

25 min), and, then, pressed at a rate of 25 mm/min (or12 mm/min in the case of alloy WE43). Samples sub-

jected to multi-pass processing were machined slightly

to allow them to fit back in the entrance channel.

Some of the ECAE runs were performed with a back-

pressure enforced by a second ram and hydraulic cylin-

der on the exit channel. The back-pressure superimposes

a hydrostatic compressive stress on top of the primarily

shear stress at the intersection of the channels in order toprevent shear failures. For example, an initial attempt to

process alloy AZ31 at 200 �C or below resulted in cata-

strophic shear failures (Fig. 2). However, processing the

alloy at 300 �C yielded only marginal grain refinement

Fig. 2. Shear failures exhibited by the alloy AZ31 ECAE processed at (a) 20

Use of a backpressure during 200 �C processing prevented such failures. (Th

[51]. Enforcing a back-pressure of 58–77 MPa (as com-

pared with the �300–350 MPa entrance pressure) pre-

vented the shear failures mentioned above and enabled

processing at a temperature of 200 �C, which resulted

in substantial grain refinement [32]. Attempts to process

at lower temperatures (175 �C) were unsuccessful evenwith a back-pressure (Fig. 2(b)).

2.4. Modeling texture evolution

It will be shown that ECAE processing results in dif-

ferent texture evolutions for the different classes of

alloys. The viscoplastic self-consistent (VPSC) polycrys-

tal modeling scheme originally proposed by Hill [54] andimplemented later by Hutchinson [55] has been shown to

be very effective for modeling the plastic response and

texture evolution of non-cubic metals [56]. The details

of the particular VPSC algorithm employed in the pres-

ent work are explained in the paper by Lebesohn and

Tome [56]. In brief, the texture of the polycrystalline

aggregate is represented by assigning a finite number

(�1000) of discrete crystallographic orientations a vol-ume fraction. The resulting discrete texture approxi-

mates the more continuous orientation distribution

measured using X-ray diffraction. Each orientation is as-

signed an anisotropic viscoplastic constitutive response

characteristic of the single-crystal, e.g. 57, where all

the slip and twinning mechanisms have been assigned

a high stress exponent, n = 20. This exponent is not in-

tended to explicitly reflect the overall rate-sensitivity ofthe material. Rather, it serves to round the vertices of

the single crystal �yield surface�, providing a mathemati-

cally convenient and unambiguous connection between

the stressing direction and the resulting strain rate.

The coupling of the grain-level response with the

aggregate is accomplished using a self-consistent

homogenization scheme based upon Eshelby�s inclusionformalism [58]. The interaction strength is assignedthrough the parameter neff = 10, which is an intermedi-

ate value between the compliant tangent modulus

0 �C without any back-pressure, and (b) 175 �C with a back-pressure.

e billets� widths and thicknesses are 22 mm.)


approach, neff = n = 20, and the more rigid secant mod-

ulus approach, neff = 1. Simulations are performed in an

incremental fashion and after each small straining step

(De � 0.02), the grains� shapes and individual orienta-

tions are updated to account for the crystallographic

rotations due to slip on the specified slip systems, andthe critical resolved shear stresses may be updated to ac-

count for a variety of effects ranging from strain harden-

ing, to latent hardening, Bauschinger-type effects.

Recent implementations have even linked these latter

phenomena to microstructure features, such as disloca-

tion substructures [59,60].

In order to explore the influence of different slip

systems on the texture evolution, different criticalresolved shear stresses (CRSSs) were assigned to

the various slip and twinning modes known to be ac-

tive within magnesium h1�210i(0002) or basal Æaæ,h1�210i f10�10g or prism Æaæ, h11�23if11�2�2g or pyra-

midal Æc + aæ, and h10�1�10if10�12g tensile twinning.

For the current simulations, all the individual defor-

mation mechanisms were assumed to behavior in a

perfectly plastic fashion, which is consistent with theobservation that metals may not significantly harden

during hot deformation. Twinning is treated distinctly

from dislocation slip in that the CRSS in the anti-

twinning direction is set to a very high (impossible)

stress level, and the crystallographic reorientation

associated with twinning is accounted for using a pre-

dominant twin reorientation scheme described else-

where [61]. It is noted that neither dynamic recoverynor recrystallization is explicitly modeled, although

their role in the deformation process of these alloys

at these temperatures is undeniable. Additionally,

changing the stress exponent, n, introduced above

would also affect secondary slip mode activity (lower

n values would promote more secondary slip). How-

ever, rather than allow multiple factors to influence

the activity of slip systems, the present investigationfocuses on the direct connection between a slip mode�sactivity and its CRSS value.

The overall boundary conditions of the ECAE pro-

cess are approximated by a simple shear within an infin-

itesimal layer aligned with the plane of intersection

between the two channels, as suggested by the inventor

[23,53,62]. For the present die with an included angle,

u = 90�, the deformation gradient, F 0, and the corre-sponding velocity gradient, L 0, are:

F0 ¼1 �2 0

0 1 0

0 0 1

0B@

1CA and L0 ¼

0 �2 _c 0

0 0 0

0 0 0

0B@

1CA ð1Þ

expressed within the rotated frame labeled x 0 and y 0 in

Fig. 1 (and _c is the nominal strain rate). A transforma-

tion of basis allows expressing these within the labora-

tory coordinate frame, x and y,

F ¼2 �1 0

1 0 0

0 0 1

0B@

1CA and L ¼

_c � _c 0

_c � _c 0

0 0 0

0B@

1CA: ð2Þ

This approach ignores the fact that the actual deforma-

tion takes place over a finite width, and that there maybe a finite radius at the intersection of the dies by design

or by virtue of a ‘‘dead metal zone’’ in the corner of the

die [63]. Some researchers have painstakingly deter-

mined the specifics of the strain path experienced by

the work-piece as it passes through the die either exper-

imentally [5,64,65] or theoretically [11,66–69]. It is gen-

erally agreed that the overall deformation is well

described by Eqs. (1) and (2). However, some research-ers have concluded that it is necessary to include details

of the strain path, while others have obtained good re-

sults assuming the simplified view. In order to model

multi-pass ECAE, it is emphasized that the material

must undergo a right hand rotation of 180� � u( = 90�for the present die) about the z-axis (shown in Fig. 1)

prior to re-insertion in the die, regardless of the process-

ing route

RA ¼0 1 0

�1 0 0

0 0 1

0B@

1CA: ð3Þ

For route A, this is sufficient and additional rotations

will apply for the other routes. This is a critical issue that

many researchers appear to have over-looked and led

them to suggest that route A processing is nothing morethan sequential simple shear of the same sense in the

same direction. As documented previously [24,70], route

A processing involves shearing on an intersecting planes

and may be modeled by enforcing the following defor-

mation for even numbered passes (2,4,6,8, . . .)

F00 ¼0 �1 0

1 2 0

0 0 1

0B@

1CA and L00 ¼

� _c � _c 0

_c _c 0

0 0 0

0B@

1CA

ð4Þand that expressed by Eq. (2) used for odd numbered

passes (1,3,5,7, . . .). For other routes, it may be more

convenient to rotate the material frame of reference

(e.g., for describing the material�s texture), rather thanthat of the deformation, in between each subsequent

pass.

3. Results

3.1. Initial textures

The initial textures of pre-extruded samples (Table 2)

are typical of magnesium extrusions (Fig. 3) having the

Fig. 3. Inverse pole figures showing the dominant h10�10i fiber texture in all the extruded magnesium alloys. ZK60 and MgLi are most strongly

textured, WE43 is most weakly textured, and the large AZ80 forging billet shows a distinct secondary Æ0001æ fiber.


basal planes aligned with the extrusion axis and a ten-

dency for the h10�10i directions to align with the extru-

sion axis. Alloys ZK60 and Mg–Li have the strongest

initial textures, and WE43 is the most weakly textured.

The texture of the large AZ80 extrusion is distinct, in

that there is also a secondary Æ0001æ fiber aligned with

extrusion axis. The effect of initial texture on the texture

evolution during ECAE was examined using two otherinitial conditions of alloy AZ31: a plate with an axis-

symmetric Æ0001æ fiber aligned with the plate normal

with an intensity of four multiples of a random distribu-

tion (MRD) and a weakly textured direct-chill cast

billet.

3.2. The effect of alloying

The initial textures (labeled 0) and the texture ob-

served after 1, 2 and 4 passes by route A are shown in

Fig. 4. All the texture results are presented in terms of

recalculated (0002) and ð10�10Þ pole figures from the

flow plane (i.e., the plane which contains the entrance

and exit channels) with the exit direction (the x-axis in

Fig. 1) to the right. Due to the 2-fold rotational symme-

try of the process (about the flow plane normal, thez-axis in Fig. 1), and the symmetry of the experimental

pole figures, the orientation distributions and full pole

figures were calculated assuming monoclinic symmetry.

The pole figures are, thus, completely represented by

their upper half in order to save space.

After a single pass, alloy AZ31 essentially exhibits a

Æ0001æ fiber texture (Fig. 4(a)) with the dominant c-axis

fiber is oriented approximately 20� aft of vertical (they-axis in Fig. 1). Subsequent passes serve to strengthen

this texture, although it is noted that there are enormous

rigid body rotations (Eqs. (2) and (3)) and plastic rota-

tions inherent to the process. Therefore, this result is

not evidence of texture stabilization, as might occur dur-

ing monotonic straining, such as multi-pass rolling.

Rather, it demonstrates reorientation to a very similar

texture after each subsequent pass. After subsequent

passes, there is also a notable strengthening of the

h10�10i parallel to (or i) z. Finally, there is a trend for

the dominant Æ0001æ texture component to split into

two components, one closer to the z-axis. Alloy AZ80

processed under the same conditions adopted an essen-

tially identical texture, which is not shown in the interest

of space.

The texture evolution of alloy ZK60 (Fig. 4(b)) showsthree major distinctions from that observed for the AZ

alloys after a single pass. (i) The major Æ0001æ fiber is

oriented closer to the y-axis; (ii) there is a secondary fi-

ber Æ0001æ iz; and (iii) h11�20ik z, placing maxima in the

ð10�10Þ pole figure 30� away from the z-axis. During

subsequent passes, the secondary texture component

Æ0001æ iz weakens, and the dominant fiber �5� aft of

the y-axis. The texture after 8 passes by route A is verysimilar to that shown after 4 passes, albeit with higher

texture strength, with peak intensities in the basal pole

figure of �8 multiples of a random distribution

(MRD), in comparison with �6 MRD for 4 passes. Al-

loy WE43 is designed for high temperature strength and,

thus, required the highest processing temperature

(325 �C) of all the alloys examined. Although it exhibits

slightly weaker textures throughout (maximum peakintensities of �4 in the basal pole figure), the texture

evolution of alloy WE43 (Fig. 4(c)) is most similar to

that of the other fine-grained Zr-containing alloy in this

study, ZK60.

The experimental Mg–Li alloy exhibits the most dis-

tinct texture evolution (Fig. 4(d)), which is not surpris-

ing in view of the unique properties and deformation

mechanisms already reported for this alloy (e.g., Refs.[46–49]. The initial extrusion texture is similar to the

other alloys due to the high symmetry of the extrusion

process itself. The lower symmetry of the ECAE process

reveals a strong distinction between Mg–4wt%Li and

the other alloys. The dominant texture component in

all the other alloys (Æ0001æ iy) becomes secondary, while

the Æ0001æ iz dominates the texture by the 4th pass

through the die. This appears consistent with the prior

Fig. 4. (0002) and ð10�10Þ pole figures (upper half) show initial extrusion textures (labeled 0) and ECAE textures after 1, 2, and 4 passes by route A.

The distinctions suggest that different deformation mechanisms operate in the different alloys. The extrusion axis is to the right, and the intensity scale

is the same as in Fig. 3 in this figure and all that follow.


observation of enhanced non-basal slip in this alloy, in

particular, the observation that prismatic slip dominated

the single-crystal deformation of Mg at elevated

temperatures [47].

3.3. The effect of initial texture

The hot-rolled AZ31 plate has a very distinct initial

texture compared to the extrusions discussed above,

however, the resulting ECAE textures (Fig. 5) are simi-

lar. The primary distinction is the complete absence of

any Æ0001æ iz, which most likely results from the factthat the initial texture places essentially all the grains

with their c-axes far from this orientation. A second dis-

tinction is that the dominant Æ0001æ fiber component is

closer to the z-axis after ECAE. Taking a most simplistic

view, the ECAE textures shown in Fig. 5 are similar to

the initial plate texture. However, the initial plate tex-

ture exhibits an orthotropic symmetry characteristic of

the rolling process, while the ECAE textures exhibit alower (monoclinic) symmetry characteristic of ECAE

[25]. Proceeding to 8 passes by route A only succeeds

in strengthening the texture. Beginning with a very

weakly textured, but coarse-grained, as-cast AZ31 mate-

rial resulted in an ECAE texture evolution almost iden-

tical to that shown for the extruded AZ31.

3.4. Modeling the texture evolution during ECAE, route A

In all cases, it has been assumed that the basal Æaæ slipmechanism is the easy slip mechanism and, therefore, it

is shown to accommodate most of the strain and leads to

the overall similarity between the textures of the various

alloys (see Figs. 3–5). In most of the simulations, the

three slip modes explored were basal Æaæ, prismatic Æaæ,and pyramidal Æc + aæ slip. In a few cases, the pyramidalÆaæ slip and f10�12g tension twinning modes were

included.

3.4.1. Modeling the effect of the non-basal slip mode

activity

The activities of the non-basal slip modes are var-

ied by assuming different critical resolved shear stress

(CRSS) values, relative to the CRSS for basal slip(sbasal = 1). A range of CRSS values from 2 to 8 is

explored for both prismatic and Æc + aæ slip modes.

Fig. 6 shows the simulated textures after a single

ECAE pass, where the initial texture was assumed

Fig. 5. (0002) and ð10�10Þ pole figures of AZ31, from initially (a) hot-rolled plate (labeled 0) or (b) cast billet and ECAE textures after 1, 2, and 4

passes by route A.

Fig. 6. (0002) and ð10�10Þ pole figures of simulated single-pass ECAE textures beginning with an initially random texture. Labels indicate relative

CRSS values of the basal, prismatic, and pyramidal Æc + aæ slip modes used in the simulations.


to be random. The contour plots presented have been

smoothed over 10� in keeping with the fact that the

initial discrete textures were determined for a

10� · 10� · 10� grid in Euler space. In all cases, basal

slip dominates the strain accommodation and causes a

dominant basal fiber texture to evolve. The clearest

case of this is labeled 1-8-8, indicating that there isa relatively high CRSS value for both of the non-ba-

sal slip modes sprismatic = sÆc + aæ = 8. The result is a

simple Æ0001æ fiber component tilted �20� aft of

the y-axis. Non-basal slip mechanisms accommodate

no more than �20% of the strain individually, with

the prismatic Æaæ slip mode being about twice as ac-

tive as the Æc + aæ.Increasing the activity of the Æc + aæ slip system up to

�20% by lowering its CRSS value to sÆc + aæ = 4 leads to

a splitting of the dominant Æ0001æ fiber into two compo-

nents, with one closer to the y-axis. There is a corre-

sponding weakening of the intensity in the ð10�10Þpole figure, with some tendency for h10�10ik z-axis. On

the other hand, if the prismatic slip mode becomes more

active (accommodating up to �40% of the strain for the

cases where sprismatic 6 4), orientations close toÆ0001æ iz are stabilized and the band in the ð10�10Þ polefigure splits into peaks corresponding to h11�20ik z.

Simulations incorporating the pyramidal Æaæ slip

mode yield very similar results to those involving

prismatic slip, since a combination of prismatic and

basal slip gives rise to the same shear strains and

crystallographic rotations as pyramidal Æaæ slip.

Therefore, to keep the simulation parameters to a

minimum, the non-basal slip of Æaæ dislocations is

considered collectively and only prismatic slip is mod-

eled in practice [71]. Additionally, while f10�12g ten-sion twinning is known to be a very important

deformation mechanism at low temperatures, it is less

active at elevated temperatures [2]. A few simulations

including tension twinning were performed, however,

the texture components which resulted are not ob-

served experimentally.

3.4.2. Modeling the effect of the initial texture

Simulations beginning with an initial texture, which

models the AZ31 extrusion or plate texture, exhibit

more non-basal slip mode activity than an initially ran-

dom texture. Even with the relative CRSS values set

quite high, sprismatic = sÆc + aæ = 8, the prismatic slip sys-

tem accommodates �25% of the strain for the case of

the initial extrusion texture. This trend results in a stron-

ger Æ0001æ iz fiber component (Fig. 7) than observed forthe initially random texture. Similarly, decreasing the

CRSS for the Æc + aæ slip mode to sÆc + aæ = 4 results in

�40% of the strain being accommodated by that mode

throughout the deformation.

Fig. 7. Pole figures showing textures simulated using an initial extrusion texture. Again, the labels indicate the relative CRSS values of the basal,

prismatic, and pyramidal Æc + aæ slip modes.

Fig. 8. Pole figures showing textures simulated using an initial plate

texture. Again, the labels indicate the relative CRSS values of the

basal, prismatic, and pyramidal Æc + aæ slip modes.


When the CRSS for the prismatic slip mode is quite

low, sprismatic = 2, the Æ0001æ iz component becomes

the dominant feature in the texture labeled 1-2-8. Corre-spondingly, there are strong texture components within

the flow plane shown in the ð10�10Þ pole figure. The caseinvolving extensive Æc + aæ slip (�35% of the strain

accommodation) labeled 1-8-4 is similar to that simu-

lated from a random texture.

Since only alloy AZ31 was processed from plate mate-

rial, only CRSS conditions appropriate for simulating

the AZ31 texture evolution were explored. Fig. 8 showsthe impact of increasing the contribution of the Æc + aæslip mode. Again, increasing the activity of Æc + aæ causesthe dominant Æ0001æ fiber to split into two components

and with the stronger of the two rotated close to the

y-axis.

3.4.3. Modeling multi-pass ECAE texture evolution

The metallographic investigation of some ECAE pro-cessed magnesium alloy sample [51] reveals that the

grain shapes after ECAE are equiaxed. This observation

combined with the assumption of no slip system harden-

Fig. 9. Pole figures showing textures simulated after a second route A pass

simulations performed with the same CRSS ratios.

ing reduces the problem of multi-pass simulation to an

issue of initial texture. As for the case of the plate tex-

ture above, the focus here will be on simulating the tex-

ture evolution of alloy AZ31, thus the simulations wereall performed with the CRSS combination of 1-8-6 (i.e.,

sprismatic = 8 and sÆc+ aæ = 6).

The four cases in Fig. 9 show the texture evolution

during a second ECAE pass by route A, beginning with

simulated single pass textures with (a) random (Fig. 6),

(b) plate (Fig. 8) and (c) extrusion (Fig. 7) initial tex-

tures. Each of these cases appear rather similar, how-

ever, beginning with the random texture has theultimate result of promoting significant activity of the

prismatic slip mode during the second pass, which re-

sults in components not observed experimentally.

Finally, second pass textures beginning with the experi-

mental textures after a single pass from the extrusion

(Fig. 4) and plate (Fig. 5) textures were simulated. This

allows exploring the possibility that simulating multiple

passes results in compounding errors. The result fromthe experimental plate texture was very similar to that

of the simulated (labeled Plate), while the subtle distinc-

tion between the simulated extrusion (labeled Extrusion)

and the experimental (labeled ECAE) are noteworthy.

4. Discussion

While all five of the examined alloys develop qualita-

tively similar h10�10i fiber texture during conventional

axis symmetric extrusion (see Fig. 3), the lower symme-

try of the ECAE process results in the significant quali-

tative differences between the ECAE textures of the five

alloys shown in Fig. 4.

with different initial textures as labeled and described in the text. All


4.1. Comparisons of experimental and simulated textures

Comparisons of the experimental and simulated tex-

tures yield a number of insights regarding the activities

of the different non-basal slip systems within the various

magnesium alloys. Ultimately, the mechanistic hypothe-ses developed through this sort of inverse approach are

best examined through direct observation (e.g., using

transmission electron microscopy).

The texture evolution of alloy AZ31 does not show

any of the signatures of the extensive non-basal slip of

Æaæ dislocations, such as Æ0001æ and h11�20ik z-axis,which are evidenced by all the other alloys (Fig. 4)

and in the simulations involving significant non-basalÆaæ slip (labeled 1-2-8 and 1-8-8 in Figs. 6 and 7). Rather,

the AZ31 textures are similar to those simulated with

conditions of balanced non-basal Æaæ and Æc + aæ slip (la-

beled 1-8-6 in Figs. 6 and 7). Finally, the texture evolu-

tion of the initially cast (randomly textured) AZ31 is

similar to the pre-extruded material. This is particularly

significant because these samples spanned a wide range

of initial grain-sizes, and the similarities between theECAE textures of cast and extruded suggest grain-size

does not significantly influence the texture evolution.

The texture evolution of the pre-extruded alloy AZ80

is very similar to that of extruded AZ31 (not shown here

in the interest of space). The similarity between the two

alloy�s textures also holds for route B processing [44].

This not surprising since the relative strengths of the dif-

ferent slip modes are most strongly affected by the solidsolution alloying. Alloy AZ31 already contains more

than the equilibrium solid solubility of aluminum. Thus,

alloy AZ80 has a similar solid solution Al content, be-

cause the additional Al is contained in second phase

particles.

The experimental texture evolution of alloy ZK60 is

not precisely modeled by any of the simulated textures

shown in Figs. 6 and 7. Nevertheless, the case labeled1-8-4 in Fig. 7 shows the most similarity. The experi-

mental texture after 1 pass exhibits some of the signa-

tures of the prismatic slip mode�s activity, such as

Æ0001æ and h11�20i components iz-axis). However, sub-

sequent passes show more tendency for the splitting of

the Æ0001æ fiber close to the vertical axis, which is a sig-

nature of extensive Æc + aæ slip, and an absence of any

h10�10i components within the flow plane, which is asignatures of extensive prismatic slip. Alloy WE43

exhibits a qualitatively similar texture evolution, how-

ever, it is weaker at every stage. This trend may be

due to a randomizing effect related to particle stimulated

nucleation during recrystallization [72].

The final alloy examined in this study, Mg–4wt%Li

shows the most similarity to the case labeled 1-2-8 in

Fig. 7, in particular the sharp texture componentsshown within the flow plane of the ð10�10Þ pole figure.

Thus, it appears that extensive prismatic Æaæ slip occurs

in this alloy, as suggested in a number of previous stud-

ies, specifically in the study of elevated temperature

deformation of Mg–Li single-crystals which showed

only non-basal slip Æaæ above a certain temperature

[47], as opposed to the dominant role played by the ba-

sal slip mode in all other alloy and temperaturecombinations.

Finally, it is well known that these alloys do undergo

dynamic recrystallization (DRX) at the observed tem-

peratures and strain rates employed in these studies

[72,73]. However, the effect of DRX on the texture evo-

lution of magnesium is not definitively known. In fact,

the study of static recrystallization textures in magne-

sium alloys is only in its infancy [74]. Nevertheless, com-parisons of the textures induced by cold-rolling and by

hot-rolling suggest that there is a randomizing effect

associated with hot deformation. Even in cases where

the texture was qualitatively well-modeled by the poly-

crystal simulation, the magnitude of the texture strength

was exaggerated. Additionally, there were cases where

the simulated texture after a single ECAE pass exhibited

characteristics shown experimentally only in the multi-pass results. This suggests that the texture evolution is

slowed by unaccounted for processes, such as recovery

and DRX.

5. Conclusions

Equal channel angular extrusion offers the potentialto induce previously unobserved crystallographic tex-

tures, which may be of significant interest for modifying

the properties. This paper documents some of the vari-

ety of textures which are generated by route A ECAE

processing, for a number of magnesium alloys, with a

range of initial textures and grain-sizes. For a given al-

loy, route A processing appears to result in similar tex-

tures from one pass to the next. However, it isemphasized that the material undergoes large rigid-body

rotations during the course of loading into the die and

during the process itself, so it is not really the same ori-

entations persisting, rather each pass through the die in-

volves a texture evolution that results in a final texture

very similar in appearance to the previous pass. Initial

texture is shown to have a pronounced effect, particu-

larly in dictating the presence or absence of certain tex-ture components which arise due to their rotational

stability.

The lower symmetry of the ECAE process as com-

pared with conventional deformation processes reveals

distinctions in the texture evolutions of the major classes

of wrought magnesium alloys, despite the similarity of

their conventional extrusion textures. In turn, these tex-

ture evolutions are used to extract information aboutthe relative activities of the deformation mechanisms,

which accommodate the plastic deformation, using an


inverse approach. Alloys AZ31 and AZ80 demonstrate

basal slip dominated deformation with a small balanced

contribution of the non-basal Æaæ and Æc + aæ slip modes.

Alloys ZK60 and WE43 also show similarities with each

other, and a significant contribution from the non-basal

Æc + aæ slip mechanism (in addition to basal slip) appearsto explain most of the observed texture components. Fi-

nally, the unique texture evolution of the Mg–Li solid

solution exhibits a number of characteristics of simu-

lated textures resulting from extensive activity of the

prismatic Æaæ slip mode. Although the simulated textures

mentioned above qualitatively reproduce many of the

experimental observations, there is a tendency to over-

predict the strength of the textures. A more detailedunderstanding of texture evolution during the hot defor-

mation of magnesium alloys awaits a better understand-

ing of the effect of dynamic recovery and

recrystallization processes on the texture evolution.

Acknowledgments

This research was supported by a grant from the US

Department of Energy, Office of FreedomCAR at

ORNL and INEEL. The encouragement of our pro-

gram managers, P.S. Sklad and S. Diamond, is grate-

fully acknowledged. National Science Foundation

Grants EEC-0203415 and DMR-0231320 to the Univer-

sity of Tennessee are also acknowledged. Thanks to

C.A. Carmichael and K. Blakeley of ORNL, for theirhelp with casting and extrusion of experimental alloys,

and to B. MacDonald (Timminco), P. Chaudhury (for-

merly with Intercontinental Manuf.), and K. Clark

(Reade Manuf.) for providing the commercial alloys

examined in this study. M. Olson is acknowledged for

performing some of the texture measurements.

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