Anisotropy in flow and microstructural evolution during superplastic
deformation of a layered-microstructured AA8090 Al Li alloy
/
W. Fan a, B.P. Kashyap b,1, M.C. Chaturvedi a,*
a
b
Department of Mechanical and Industrial Engineering, University of Manitoba, Winnipeg, Canada R3T 5V6
Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology, Mumbai 400 076, India
Abstract
The superplastic forming grade sheets of AA8090 Al /Li alloy were observed to contain layers of different microstructure and
microtexture across their cross-section along the normal to the rolling direction (RD). The surface layer (SL) material contained
coarse equiaxed grains and the dominance of S {1 2 3}[6 3 4] texture whereas the center layer (CL) material contained fine
elongated grains and the dominance of Bs {1 1 0}[1 1 2] texture. Tensile specimens, machined to represent the SL of 0.6 mm
thickness from the surface towards center (SL), the CL of 0.6 mm thickness, obtained by removing the material of 0.6 mm thickness
from each surface towards center (CL), and full thickness (FL) material of 1.8 mm thick, in a sheet of AA8090 Al /Li alloy, were
deformed at optimum superplastic condition of strain rate /1/10 3 s 1 and temperature /803 K to investigate the effect of
loading direction. In SL material, the specimen parallel to RD exhibited maximum and the specimen perpendicular to RD exhibited
minimum flow stresses. This trend was reversed in CL material. The anisotropy in flow stress could be explained on the basis of
texture in the SL material, but the contribution of grain directionality became important in the CL material. The flow behavior of
FL material was found to consist of the composite-like contributions of SL and CL materials.
Keywords: Microtexture; Anisotropy; Superplasticity; Grain morphology; Al /Li alloy
1. Introduction
The presence of fine and stable equiaxed grains is a
microstructural requirement for the materials to exhibit
exceptionally large ductility or superplasticity, in a
uniaxial tensile test at high temperatures and intermediate strain rates. However, several materials, when
processed for superplasticity, initially contain elongated
or banded microstructures [1 /5]. The flow behavior of
the materials with such microstructures is reported [6] to
be quite different from those with equiaxed grains of
uniform distribution. The directionality in the elongated/banded microstructure contributes to anisotropy
in superplastic flow of the materials. In addition to the
contribution of grain morphology and distribution,
texture also contributes to the anisotropy in flow
behavior. The contributions of these two factors to
anisotropy during superplastic flow were evaluated in a
systematic study [7] on the Pb /Sn eutectic alloy.
With the advent of new thermo-mechanical processing methods, it has become possible to develop superplasticity in the multi-component alloy systems of
commercial importance. However, the microstructures
so produced are invariably found to be more complex.
For example, a thin sheet of AA8090 Al /Li alloy was
reported [8 /10] to contain coarse and nearly equiaxed
grains in the surface layer (SL), up to /1/3 of the full
thickness (FL) of the sheet [10], whereas the microstructure in the mid-thickness layer contained fine
elongated grains. In addition, a variation in microtexture from SL material to the center layer (CL) material
was also observed [9 /14]. Such gradients in microstructure and microtexture, from the surface to the midthickness layers, may influence the superplastic flow
behavior and the nature of anisotropy differently from
167
the situation where microstructure and microtexture
throughout the thickness of the tensile specimen were
identical. While some comparison in the superplastic
behavior of CL and FL material was reported [13], there
does not appear to be any attempt made towards
investigating the anisotropy in superplastic flow of
surface and CL material of distinct microstructures
and microtextures. Therefore, the aim of the present
work was to investigate the superplastic flow behavior
of and the microstructural evolution in the SLs, the midthickness layers and the FL material of superplastic
forming (SPF) grade AA8090 Al /Li alloy at different
orientations, with respect to the rolling direction (RD).
2. Experimental procedures
SPF grade AA8090 Al /Li alloy of composition
(wt.%): Al /2.5Li /1.4Cu/1.2Mg /0.11Zr was obtained
from Alcan in the form of 1.8 mm thick sheet. For
obvious proprietorial reasons, no details of the rolling
and heat treatment schedule were made available by the
manufacturer. However, in general the development of
fine grains for superplasticity in such materials involves
a suitable combination of hot/cold rolling and age
hardening treatment. Mechanical working and heat
treatment are coupled in such a way that the stored
energy increases but not to the extent that recrystallization could occur, at the numerous stress concentration
sites, in the course of thermo-mechanical processing.
Instead, the desired microstructure for superplasticity is
obtained subsequently by static or dynamic recrystallization, but prior to superplastic deformation. Details of
metallographic and orientation imaging microscopy
(OIM) techniques have been described in an earlier
publication [10]. For OIM, the strain-free electrolytically polished longitudinal section of the specimen was
scanned in a JEOL 840 scanning electron microscope
(SEM) in a beam-controlled mode. The selected area
included at least 200 grains with the scan grid being
smaller than the grain size. The automated electron back
scatter diffraction pattern analysis, with TSL software
2.0, was used to obtain crystallographic data on a point
by point basis.
Tensile specimens of 10 mm gage length and 5 mm
gage width were machined. The orientations of tensile
axes with respect to the RD (u ) were kept at 0, 30, 45, 60
and 908, as illustrated in Fig. 1. Based on the microstructural and microtextural variation (to be described
in the next section) along the thickness direction, three
groups of tensile specimens were prepared for each
orientation to represent */(i) near surface zone (surface
material, SL), (ii) mid-thickness zone (center material,
CL), and (iii) FL material. In the first two cases, the
thickness of tensile specimens was kept at the 1/3 (0.6
mm) of the sheet thickness. As detailed elsewhere [10],
this classification into SL and CL, with two SLs on
either side of the CL, are based on the appearance of
distinctly different microstructures in the SL, of /1/3 of
1.8 mm FL sheet, and the CL of the same thickness; the
microstructures in the two opposite SLs being similar.
Blanks for tensile specimens of the surface and center
materials were obtained by chemical milling with a
proprietary etchant called Turon 4181. In order to avoid
the pin-holes of the shoulder sections getting deformed,
the shoulders were sandwiched by spot welding two
aluminum pieces of the same size. Constant true strain
rate tensile tests were conducted with a computerized
Instron universal testing machine. All the tests were
conducted at strain rate (o)
˙ of 1 /10 3 s 1 and
temperature (T ) of 8039/1 K, which were determined
to be the optimum conditions for superplastic deformation of this alloy. Heating and soaking time of 20 min
was allowed prior to deformation.
Metallographic samples to observe the cavities were
examined in as-polished condition by an optical microscope as well as in a JEOL SEM. Fracture surfaces of
the tensile specimens were examined in SEM.
3. Results
3.1. The initial microstructures
3.1.1. As-received condition
Due to the difficulty in etching, the microstructure in
the as-received condition of the sheet could not be
revealed by conventional optical metallographic techniques. Therefore, the microstructural characterization
was carried out by OIM technique. Fig. 2(a /c) shows
the microstructures in the rolling surface (a), in the SL
but in the longitudinal section SL (b) and in the CL of
the longitudinal section (c). The contrast in the OIM
images is determined by the quality of diffraction
patterns generated from the scanned specimen area. A
good quality pattern can be obtained from stress free
perfect grains, but can not be obtained from locations
close to grain boundaries and deformed regions. The
quality is also influenced by the size of the electron beam
relative to the size of grains. The contrast in Fig. 2 is
poor because of residual stresses present in the rolled
sheet specimens and due to the smaller grains compared
with the size of the electron beam. Although it was not
possible to measure and quantify grain size and grain
morphology with certainty, the OIM images qualitatively show that the microstructures consist of relatively
coarse and nearly equiaxed grains in the rolling surface
as well as in the SL of the longitudinal section. Fine
elongated and banded grains are seen in the CL of the
longitudinal section.
The texture in a rolled sheet can vary in different
locations of its thickness section. In order to determine if
168
Fig. 1. Schematic representation of tensile axis oriented at different angles (u ) to the RD of the sheet. Also shown are the planes and directions of the
as-received sheet.
Fig. 2. OIM images obtained from different sections of the as-received sheet material.
169
a texture gradient was present in the as-received
material, textures were measured by OIM technique at
several locations on the longitudinal section of the sheet
sample as a function of thickness. Fig. 3 schematically
illustrates the locations from where the OIM images and
pole figures (PF) were obtained. The location A is next
to the surface and E is at the center whereas B, C and D
are located between them. In each location (A /E), the
instrument measured the orientations of all the grains in
the entire area defined by Fig. 3. Each area was 80 mm
along the RD and 200 mm along the normal direction.
The OIM images of A /E layers separately (not presented here) were shown by colour code, provided by
software, to represent the specific orientations of the
grains. The PFs representing these locations are shown
in Fig. 4(a? /e?). A comparison of these PFs with those of
the main components of the rolling texture in fcc
materials [15] suggests the following. At locations A
and B, which correspond to the SL, the texture is copper
(Cu): {1 1 2}Ž1 1 1 type with a mixture of S
{1 2 3}Ž6 3 4 and Cu components. The texture in the
CL at D and E is primarily brass (Bs) {1 1 0}Ž1 1 2
type with a mixture of Bs and S. The transition from Cu
type to Bs type texture is seen at C, which is located
approximately between the surface and the CLs.
3.1.2. Just before tensile deformation
Fig. 5(a /c) shows the microstructure and prominent
microtexture components present in the surface and CLs
of the tensile specimens after heating to and soaking at
the test temperature, and just before the start of
deformation. It is seen that, in this initial microstructure, the grains in the surface and its adjoining layers,
along the thickness direction, are coarse and nearly
equiaxed whereas they are fine elongated in the mid-
thickness (center) zone. In these two layers, the dislocation structure and the prominent texture components
were also different [16]. The mid-thickness layers contained nearly 40% low angle ( B/158) boundaries [10,17]
whereas the SLs contained mostly the well-developed
grains of high angle ( /158) boundaries. As shown by
the histograms of various texture components and their
volume fractions in Fig. 5(b and c), the SL contains the
following texture components in the order of their
decreasing
volume
fractions */S{1 2 3}Ž6 3 4 /
Cu{1 1 2}Ž1 1 1 /C{0 0 1}Ž1 0 0 /
Bs{0 1 1}Ž2 1 1/G{1 1 0}Ž1 0 0. In contrast to this,
the CLs are dominated by the Bs texture component.
The volume fractions of various texture components
present in these layers follow a sequence in the
descending order of Bs /S /Cu /C /G. Thus, the
FL of the material consists of three distinct layers*/two
SLs and one mid-thickness layer with the characteristic
microstructure and microtexture. Therefore, the flow
behavior, anisotropy in flow and concurrent microstructural and microtextural evolution in these layers
were investigated separately and the results were compared with those of the parent FL material.
Static annealing during heating to and soaking at the
test temperature, prior to superplastic deformation, does
not seem to have much influence on the texture present
in the as-received material. Thus, the texture present
prior to tensile testing, but after static annealing during
heating to and soaking at the test temperature, is
introduced by the thermo-mechanical processing, which
is used to produce microstructure that makes the
material superplastic. Such observations have been
also reported by other investigators [11,12,18 /20].
However, static annealing during heating and soaking
at the test temperature does facilitate the etching process
to delineate the grain boundaries.
3.2. Anisotropy in tensile flow
Fig. 3. Five equally divided locations (A /E) in half thickness of
longitudinal section of the as-received sheet, at which OIM measurements were made and the PFs were obtained.
Tensile specimens of SL, CL and FL materials of
different orientations, as defined in Fig. 1, were
deformed to failure at the optimum superplastic condition of o˙1103 s1 and T /803 K. True stress (s)/
true strain (o ) curves obtained are presented in Fig. 6(a /
c).
The stress /strain curves for the SL material, Fig. 6(a),
exhibits the maximum flow stress for the specimen of
u /08 orientation and the minimum flow stress for the
specimen of u /908. Between these limiting cases, the
flow stress is found to decrease with an increase in the u value. Also noticeable in this figure is a gradual decrease
in the slope of s /o curves with the increase in u . That is,
the flow hardening rate decreases as the orientation of
tensile axis deviates more from the RD.
The stress /strain curves for the CL material, Fig.
6(b), exhibit the maximum flow stress for the specimen
170
Fig. 4. PFs obtained from five equally divided longitudinal sections along the half thickness of the as-received sheet material.
of u /908 orientation and the minimum flow stress for
the specimen of u /08. However, the s /o curves for the
intermediate orientations do not follow a similar dependence of flow stress on u. The slope of the s /o curve for
u /08 is noted to be less than that of u/908.
For FL material, Fig. 6(c), the s /o curves for
different orientations lie within an envelope whose
upper bound is the s /o curve corresponding to 308
orientation and the lower bound is that corresponding
to 08 orientation. A large number of inter-weavings of
the s /o curves suggest no systematic effect of orientation.
The stress /strain curves in Fig. 6 suggest that the
difference in flow stress amongst the specimens of
different orientations in general increases with increasing strain. However, it may be pointed out here that this
widening in-plane anisotropy is in variance with the
anisotropy reported earlier [21] from the strain measurements in the width and thickness directions of the
deformed specimens. The anisotropy in out-of-plane
directions decreased dramatically by superplastic deformation.
From the tensile test data, the values of peak stress
(sp), the strain at which it occurs (o p), and the elonga-
tion-to-failure, expressed in the form of true strain
(o max), are listed in Table 1 to show the effects of
orientation in SL, CL and FL materials. A close
observation of the data in this table suggests the
following, with few exceptions. (i) Maximum ductility
is exhibited by FL material whereas CL material
exhibits the minimum but, there is no effect of orientation. (ii) Except for u /08, SL material shows the
minimum values of sp and the same decrease with
increasing u , whereas sp increases in CL material with
u . (iii) Except for u /08, o p has the lowest values for the
SL material, but no influence of orientation is evident.
3.3. Effect of specimen location and orientation on
microstructural evolution
In order to examine the variation in the nature and
extent of microstructural evolution, including cavitation,
grain structure and texture, the SL and CL tensile
specimens of different orientations (u /0, 30, 45 and
908) were deformed to a selected strain of 1.0. For
metallography, the samples were prepared from the
longitudinal and transverse sections of the gage portion
of all the tensile specimens of different orientations.
171
Fig. 5. Initial microstructure (a) and microtexture in the surface (b) and center (c) layers of tensile specimens upon heating to and soaking at the test
temperature of 803 K for 20 min, prior to deformation. The grain structure and texture components in the surface and CLs are noted to be different.
The texture components are: S {1 2 2}Ž6 3 4; Cu {1 1 2}Ž1 1 1; Bs {0 1 1}Ž2 1 1; C {0 0 1}Ž1 0 0 and G {1 1 0}Ž1 0 0.
Microstructures were examined at various magnifications and finally photographs were taken at low and
high magnifications to explore the effects of specimen
location and orientation.
3.3.1. Cavitation
The specimens of different orientations were examined in as-polished condition upon superplastic deformation to o /1.0. In SL material, the size and number of
cavities were found to be similar in the longitudinal and
transverse sections. Also, no noticeable difference in
cavitation behavior was seen for different orientations
of tensile specimens. In CL materials, cavities were
found to be maximum for the orientation of u /08; and
they decreased in number and size for the larger values
of u . The minimum cavities were observed for u /45
and 908 orientations in the longitudinal and transverse
sections of the tensile specimens, respectively. A comparison of the cavity levels in the SL and CL materials
for all orientations revealed that the cavities were more
prevalent in the CL material. This is illustrated in Fig.
7(a and b) by the micrographs taken from the SL and
CL tensile specimens oriented at u /08.
3.3.2. Grain structure
Typical microstructures obtained upon superplastic
deformation to o /1.0 are presented in Fig. 8(a /e). In
the course of superplastic deformation, grain growth
and transformation from elongated grains to equiaxed
grains occurred. As compared to the initial microstructure in Fig. 5(a), the grain morphology and grain sizes
are seen to be uniform and homogeneous in Fig. 8(a).
Owing to the clear microstructure obtained by optical
metallography at this stage no OIM images are presented here. The microstructures in the longitudinal and
transverse sections, of each SL and CL materials, were
similar. However, the grain boundaries in the CL
material were not sharply etched and a trace of initial
banding was found to remain in the case of intermediate
orientations of tensile loading, viz. u /30 and 458, Fig.
8(c and d). The microstructure of the tensile specimen of
u /08 orientation, on the other hand, revealed equiaxed
grains with well-developed grain boundaries, Fig. 8(b).
172
3 1
Fig. 6. Stress(s ) /strain(o ) curves for (a) SL, (b) CL and (c) FL materials deformed at various orientations (u ) with respect to RD. (/o110
˙
s ;
T /803 K). The s /o curves for SL material show a systematic effect of orientation, with u /08 exhibiting maximum and u/908 the minimum flow
stress. The trend is somewhat reversed in CL material, and not so systematic in FL material.
Also, worth noting is the appearance of intergranular
precipitate free zones, the extent of which was prominent in the specimen of u /458 orientation, Fig. 8(d). In
this case, the precipitates also appeared to be somewhat
finer. The precipitate free zone appeared less prominent
in the SL material, Fig. 8(e).
3.3.3. Microtexture
Microtexture as a function of orientation was measured in such a way that the tensile axis for each
specimen orientation (u , Fig. 1) was aligned to the
default RD in OIM. Then the measured texture data
were rotated by an angle, which was the one between the
tensile axis and the sheet RD, viz. u /30, 45, 60 and 908,
because the present material has fcc crystal structure.
PFs of SL samples of different orientations are presented in Fig. 9(a /d) along with the axis system defining
the tensile axis orientation with respect to the RD. A
noticeable effect of tensile loading direction in the PFs is
observed, with the maximum intensity increasing with
an increase in u. This is also shown by the general
increase in volume fractions of the dominant texture
components S and Cu in Fig. 9(e).
PFs of CL samples of different orientations are
presented in Fig. 10(a /d). However, in this case, there
does not appear any noticeable effect of tensile loading
direction in the PFs, with the maximum intensity
remaining within a narrow band of 5.53 /7.95. This is
also illustrated by nearly constant (but within a large
scatter band) volume fractions of the dominant texture
173
Fig. 6 (Continued)
Table 1
Effect of tensile loading direction on maximum strain to failure (o max),
peak stress (sp, MPa) and the corresponding strain (o p) levels for the
surface (SL), center (CL) and FL materials
Layer
Property
Orientation of tensile axis wrt RD (u )
08
308
458
908
SL
o max
sp
op
1.344
9.99
1.203
1.523
9.11
0.897
1.475
8.53
0.870
1.383
8.27
0.937
CL
o max
sp
op
1.387
9.76
1.066
1.311
9.44
1.060
1.455
9.68
1.075
1.354
9.99
1.003
FL
o max
sp
op
1.664
9.27
0.866
1.636
9.68
0.938
1.637
9.55
1.133
1.647
9.66
1.055
components Bs and S plotted as a function of u in Fig.
10(e).
3.4. Effect of specimen location and orientation on
fracture behavior
Fracture surfaces of the SL, CL and FL specimens of
different orientations were examined at various magnifications but the fractographs were taken at three
magnifications of 50 /, 200/ and 600/. In general
the failure was of pseudo-brittle type, with the voids
being separated by the different proportions of ductile
(tearings) and nearly flat zones. The whole fracture
surfaces of the SL, CL, and FL specimens of orientations u /30, 45 and 908 are presented in Fig. 11(a /c). A
comparison of the fractographs of the SL (denoted by s
in Fig. 11(a /c)) and CL (denoted by c in Fig. 11(a /c))
Fig. 7. Micrographs illustrating less cavities in SL (a) and more
cavities in CL (b) materials upon deformation to o /1.0. (u /08; o
˙
1103 s1 ; T/803 K).
specimens reveals that the voids are more numerous in
the CL specimens than in the SL specimens for all the
orientations. In the SL samples, there appears to be a
predominance of non-cavitated but extensively plasti-
174
Fig. 8. Micrographs illustrating grain growth and transformation of elongated grains into equiaxed grains (compare with Fig. 5(a), the same axis
3 1
system being employed for both figures) upon superplastic deformation to o /1.0. (/o110
˙
s ; T /803 K): (a) FL material (u/08), (b /d) CL
material at orientations of (b) u /08, (c) u /308, and (d) ? u/458, (e) SL material at u/458. There are regions of banding and pronounced
precipitate free zones in c and d.
cally deformed regions. As a function of orientation, no
noticeable effect is seen in the fractographs of SL
material. However, the fractographs of the CL material
(denoted by c in Fig. 11(a /c)) suggest that the proportion of the void area decreases whereas the proportion
of the plastically tearing region increases as the orientation shifts to the greater value of u. The fractographs of
the FL material (denoted by F in Fig. 11(a /c))
predominantly exhibited tearings, except that the voids
were prominently present at the orientation of 308.
A comparison of the fractographs of the SL and CL
materials at a high magnification is illustrated in Fig.
12(a and b), for the samples of 458 orientation. In the
fractograph of SL, there appears to be a smaller number
of voids, which are surrounded by the large flat or
tearing zones. In the case of CL, there appear a large
number of voids, which were developed amidst a group
of grains. Such groups of grains, containing the voids,
are partitioned by a network, which has undergone
extensive plastic deformation/tearing.
4. Discussion
4.1. Flow anisotropy and role of microstructure
Although superplasticity is a high temperature deformation phenomenon where slip and related events in
general become unimportant, the analysis of anisotropy
in a duplex stainless steel [22] led to the suggestion that
the crystallographic texture plays significant role in
explaining the anisotropy during superplastic deforma-
tion as well. The anisotropy in tensile properties of
AA8090 Al /Li alloy was extensively studied [9,11,23 /
28] at room temperature, and it was explained by the
variation in Taylor factor (M /sy/tcrss) as a function of
u . M is the ratio of yield stress (sy) to critical resolved
shear stress (tcrss), which is a polycrystalline equivalent
to the inverse of the Schmid factor. In spite of some
evidence of anisotropy during superplastic deformation
of this material, no systematic attempt was made to
understand its origin. Presented below is an attempt to
examine whether the observed anisotropy in the stress /
strain curves (Fig. 6) can be explained through texture.
The stress /strain curves for the SL and CL materials,
Fig. 6(a and b), exhibit anisotropy in flow behavior but
in an opposite manner. An attempt was made to
understand this anisotropy by employing the variation
in the magnitudes of Schmid factor for different
orientations of tensile axis with respect to the RD, i.e.
as a function of u , and considering the main texture
components, viz., S, Bs, and Cu. As the Schmid factor
increases, the yield stress is known [29] to decrease. In
other words, in polycrystalline materials, yield stress
increases with an increase in Taylor factor M . The
values of M for the S, Bs and Cu texture components
were calculated by using the Taylor/Bishop-Hill model
[30 /32] and plotted as a function of u by Fricke, Jr. and
Przystupa [33]. In the SL material, the volume proportion of S, Cu and Bs texture components are nearly
45:17:10 whereas the relative Schmid factors are about
0.313:0.270:0.323 for u /08, 0.313:0.313:0.417 for u/
308, 0.286:0.270:0.270 for u /458 and 0.385:0.270:0.270
for u /908. On the basis of the variation in Schmid
175
Fig. 8 (Continued)
factors of the individual texture components, as a
function of u (see Fig. 8(a) in [33]), the Cu texture
component suggested a higher flow stress whereas the S
and Bs texture components suggested lower, but mutually comparable, flow stresses at u /08 orientation. At
u /908, Cu and Bs texture suggested mutually equal but
higher stress than that suggested by S texture component. At intermediate orientations, the S component
suggested the maximum stress whereas the other two
texture components suggested the minimum stress.
When several texture components are present, one could
speculate that deformation should take place as soon as
the minimum critical stress required for slip in any one
of these components is reached. Accordingly, Bs and S
texture components become the most favorable for
deformation at u /08, S texture component favors
deformation at u /908 and the Bs texture component
becomes favorable at intermediate orientations. As the
SL material is dominated by S type texture component,
and the u /08 orientation has lower Schmid factor than
that at u /908 orientation, the former orientation is
expected to exhibit higher flow stress. This explains the
relative positions of the s /o curves for the two
orientations, i.e., the maximum flow stress for u/08
orientation and the minimum flow stress for u/908
orientation in Fig. 6(a). The CL is dominated by Bs
texture for which the Schmid factor is lower at u/908
than that at u/08. Therefore, the maximum flow stress
176
3 1
Fig. 9. (a /d) PFs obtained from tensile specimens of SL material deformed to o /1.0 (/o110
˙
s ; T /803 K) for various orientations (u 8): (a)
0; (b) 30; (c) 45 and (d) 90. (e) Plot of the volume fractions (%) of different texture components as a function of orientation.
for u /908 and the minimum flow stress for u /08 (Fig.
6(b)) can be understood. However, this does not explain
the positions of s /o curves for the intermediate
orientations; the variation in Bs texture showed the
maximum values of Schmid factors at u /30 and 458
but the flow stresses for these orientations are not the
minimum ones.
Unlike in single crystals, macrostrain in polycrystalline materials is achieved by the operation of five
independent slip systems in each grain, which maintains
177
3 1
Fig. 10. (a /d) PFs obtained from tensile specimens of CL material deformed to o /1.0 (/o110
˙
s ; T/803 K) for various orientations (u 8):
(a) 0; (b) 30; (c) 45 and (d) 90. (e) Plot of the volume fractions (%) of different texture components as a function of orientation.
178
Fig. 11. Low magnification full view of fracture surfaces of SL, CL and FL materials (marked by s, c and F, respectively) deformed at different
3 1
˙
s ; T/803 K). Irrespective of orientation, more voids are noticed in the CL material.
orientations (u 8): (a) 30, (b) 45 and (c) 90. (/o110
co-ordination between the adjoining grains. In the
presence of texture of different types, the group of
adjoining grains should respond to the compatibility
requirements and contribute to deformation, instead of
individual grains of nearest neighbour doing so. Therefore, it may be important to consider the gross effect of
all the major texture components to explain the anisotropy in flow behavior.
The magnitudes of Taylor factor associated with S, Bs
and Cu texture components were obtained at different
values of u from Fig. 8(a) in [33]. These values were then
used to calculate the average value of M according to
rule of mixture, on the basis of volume fractions of these
texture components. The values of Taylor factor so
obtained (Mg) are plotted in Fig. 13 for SL and CL
materials. Also included is a plot of Taylor factor for FL
material, whose values were determined from the mean
Mg values of SL and CL materials, in the ratio of 2:1
(two surfaces and one mid-thickness layer in a FL
material). Now, the following observations can be made
by comparing the stress /strain curves of SL, CL and FL
materials of different orientations in Fig. 6 with the
respective Mg values in Fig. 13. (i) For SL material (Fig.
6(a)), the highest flow stress for u /08 orientation can
be ascribed to the highest value of Mg; and the variation
in flow stress as a function of u follows the same trend
179
Fig. 12. High magnification fractographs of tensile specimens from SL
(a) and CL (b) materials tested at an orientation of u/458.
as does the plot of Mg vs. u . Thus, the texture has
similar influence on anisotropy as reported at low
temperatures. (ii) For CL material (Fig. 6(b)), the
highest flow stress is observed at u /908, corresponding
to the maximum value of Mg, and the minimum stress
observed at u /08 correspond to a lower value of Mg.
However, a similar dependence of stress on Mg is not
obeyed for intermediate orientations. The s /o curves
for intermediate orientations lie above that for u/08
and below that for u /908. On the basis of the
maximum values of Schmid factor or the minimum
value of Mg, the lowest flow stress would be expected at
some intermediate orientation of the tensile specimens,
which is not the case. (iii) For FL material, the values of
Mg varies over a narrow range and so does the flow
stress for different orientations (Fig. 6(c)). In this case,
the minimum flow stress of the s /o curve corresponding to u /08 and the maximum flow stress of the s /o
curve corresponding to u /308, and so also for the
remaining orientations, can be understood. On the basis
of this semi-quantitative comparison of s /o curves and
Taylor or Schmid factors for different orientations, the
anisotropy can be attributed to texture, when the
microstructure contains equiaxed grains.
The failure to account for anisotropy through texture
in CL material may be related to the presence of
elongated grains. In the Pb /Sn eutectic alloy, anisotropy was attributed [7] to the presence of such
microstructure rather than texture because the anisotropy was eliminated once the grains became equiaxed
at a strain of /300%. What is still puzzling from the
nature of stress /strain curves in Fig. 6 is that, in spite of
the reduction in texture and evolution from elongated
grains towards equiaxed grains with increasing strain
[10,17], the anisotropy in flow stress is seen to increase.
This requires further work to follow the microstructure,
Fig. 13. Plot of gross Taylor factor (Mg, determined by using the Taylor factors for the S, Cu and Bs components from Ref. [33] according to their
volume fractions in rule of mixture) for SL, CL and FL materials. The mean values representing Mg for the FL material were determined from that of
the SL and CL materials in a proportion of 2:1 while using rule of mixture.
180
microtexture and substructure evolution and to analyze
their independent and synergistic contributions to
anisotropy. The microstructure (Fig. 8) and microtexture (Figs. 9 and 10) examined at o /1.0 in the present
study do not support the existence of anisotropy. It
remains to be seen whether these basic sources of
anisotropy evolving in the early stages of deformation
can influence the mechanisms for superplastic deformation to cause anisotropy even after their own disappearances. For instance, as reported in earlier publications
[10,17], the CL material was highly textured and it
contained a large proportion of low angle boundaries,
about 40% in the beginning, which rapidly decreased to
10% along with a reduction in texture upon superplastic
deformation to o /1.0 for u /08 orientation. Recently
Huang and Humphreys reported [34] that such a change
in misorientation angle has a strong effect on the
kinetics of subgrain growth.
The observations of cavities at o /1.0 and the
fractographs (Fig. 11), at even larger strains, did not
reveal significant effect of orientation. A reason for the
absence of orientation effect may be the fact that these
observations were made at large strains, whereby the
texture had already diminished and the microstructure
had become more or less equiaxed and homogeneous.
4.2. Composite-like flow behavior
The stress /strain curves in Fig. 6(a /c) were replotted
to compare the flow curves of SL, CL and FL materials
at each orientation. As shown in Fig. 14(a /d), at all the
orientations, the stress /strain curves of CL material are
at higher level than that of SL material. This difference
in the stress levels of the two materials can not be
explained on the basis of the difference in the texture
components because Mg for SL material is estimated to
be greater than that for CL material over the wider
range of orientations (Fig. 13). Then the difference in
flow stress may be a result of the difference in microstructures in the two layers (Fig. 5). The elongated
grains, like the one present in CL, are known to be
unfavorable for high temperature deformation whereas
the material with equiaxed grains, like the one present in
SL, is known to deform at lower stress [2].
Also included in Fig. 14(a /d), are the two stress /
strain curves for FL material*/one obtained experimentally and the other calculated from the data of SL
and CL materials, according to rule of mixture and by
considering that the FL consists of two SLs and one
mid-thickness layer. The experimental stress /strain
curves for all the orientations are seen to lie between
those of the SL and CL materials up to large strain
levels, supporting the composite-like flow behavior. The
experimental and calculated stress /strain curves are
seen to be entirely comparable at u /08 whereas, at
other orientations, the two curves are noted to match
over the early part of deformation and then deviate
subsequently. At larger strains, the experimental stress /
strain curves exhibit higher flow stress than that of the
calculated ones. Since microstructural evolution was not
investigated as a function of strain for all the orientations, it is difficult to pin point the source of the higher
experimental flow stress. However, it is worth considering the proportion of the grain boundaries falling in a
specific direction with respect to tensile loading direction, because those grain boundaries, which are inclined
to tensile axis, undergo sliding and migration more
preferably [35]. Considering the microstructure in Fig. 5,
which contains elongated grains in the mid-thickness
layer, the loading directions other than u /0 and 908,
will have more inclined boundaries. Such inclined
boundaries may not initially facilitate sliding mechanism
because of the dominance of texture (dislocation activity), but upon reduction of texture at some strain level,
sliding and migration could take place in these boundaries more preferably. These could also result in an
enhanced grain growth during superplastic deformation
at such orientations. This is suggested to be a probable
reason for the deviation between the experimental and
calculated stress /strain curves at larger strains.
The comparison of cavities in the SL and CL
materials, deformed to the fixed strain of 1.0 (Fig. 7)
and to fracture (Figs. 11 and 12), revealed the latter
material to be more prone to cavitation. This, in
elongated grains (CL material), can be ascribed to the
difficulty encountered in the mutual accommodation
and deformation processes of grain boundary sliding
and diffusion. Further, grain rotation and rearrangement, which otherwise relieve the stress concentration
built-up by grain boundary sliding during superplastic
deformation, is also delayed till the grains become
equiaxed.
5. Conclusions
A study on superplastic deformation in different
directions of SL, mid-thickness layer, CL and FL sheet
of AA8090 Al /Li alloy, with the variation in microstructure and microtexture, led to the following conclusions.
(1) The SL material, which was dominated by S:
{1 2 3}[6 3 4] texture and contained nearly equiaxed
grains, exhibited maximum flow stress in the RD and
minimum flow stress in perpendicular to RD. The flow
anisotropy could be ascribed to the presence of texture.
(2) The CL material, which was dominated by Bs:
{0 1 1}[2 1 1] texture and contained fine elongated
grains, exhibited maximum flow stress in a direction
perpendicular to RD and minimum flow stress along the
RD. The anisotropy in flow was attributed to the
combined effects of texture and grain directionality.
181
Fig. 14. Comparison of stress /strain curves for SL, CL and FL materials at various orientations (u8 ): (a) 0, (b) 30, (c) 45 and (d) 90. Also included
are the stress /strain curves calculated (shown by symbols 0 /0) from the data of SL and CL materials according to rule of mixture. In each of the
figures (a /d), the top curves are for CL, the bottom curves for SL, and the FL curves experimental (solid curve) and calculated according to rule of
mixture (0 /0) are between the two. The stress /strain curves derived from the data of SL and CL materials suggest a composite-like behavior of FL
material for u /08 and for other orientations in the early part of deformation.
(3) Irrespective of loading direction, the flow stress for
CL material was found to be greater than that for SL
material. The experimental stress/strain curves for FL
material were found to lie between that for the SL and
CL materials.
(4) The stress /strain curves for FL material could be
explained by the composite-like contributions of the
layered SL and CL materials. However, at large strains,
the effects of concurrent microstructure and microtexture should be incorporated in the rule of mixture.
Acknowledgements
The authors would like to thank the consortium of
Manitoba aerospace industries and the Natural Sciences
182
and Engineering Research Council of Canada for
financial support. Technical assistance by D. Mardis
and J. Van Dorp is very much appreciated.
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