an experimental study of the recrystallization mechanism during hot deformation of aluminium
TRANSCRIPT
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Materials Science and Engineering A283 (2000) 274288
An experimental study of the recrystallization mechanism duringhot deformation of aluminium
S. Gourdet, F. Montheillet *
Ecole Nationale Superieure des Mines de Saint-Etienne, Centre Science des Materiaux et des Structures, URA CNRS 1884, 158 Cours Fauriel,
42023 Saint-Etienne Cedex 2, France
Received 10 February 1999; received in revised form 22 October 1999
Abstract
Discontinuous dynamic recrystallization (involving nucleation and grain growth) is rarely observed in metals with high stacking
fault energies, such as aluminium. In this metal, two other types of recrystallization have been observed: continuous dynamic
recrystallization (CDRX, i.e. the transformation of subgrains into grains); and geometric dynamic recrystallization (due to the
evolution of the initial grains). The main purpose of this work was to bring clearly into evidence and to better characterize CDRX.
Uniaxial compression tests were carried out at 0.7 Tm and 102 s1 on three types of polycrystalline aluminium: a pure
aluminium (1199), a commercial purity aluminium (1200) and an Al-2.5wt.%Mg alloy (5052), and also on single crystals of pure
aluminium. In addition, 1200 aluminium specimens were strained in torsion. The deformed microstructures were investigated at
various strains using X-ray diffraction, optical microscopy, scanning electron microscopy and electron back-scattered diffraction.
Observations of the single crystalline samples confirm that subgrain boundaries can effectively transform into grain boundaries,
especially when the initial orientation is unstable. In the case of polycrystalline specimens, after separating the effects of the initial
and new grain boundaries, it turns out that CDRX operates faster in the 1200 aluminium compared to the two other grades.
Moreover, it appears that the strain path does not alter noticeably the CDRX kinetics. 2000 Elsevier Science S.A. All rights
reserved.
Keywords: Hot deformation; Aluminium; Dynamic recrystallization; Single crystals; Subgrain boundaries; Grain boundaries; Misorientations
www.elsevier.com/locate/msea
1. Introduction
Aluminium and its alloys exhibit very high rates of
dynamic recovery, which is generally expected to com-
pletely inhibit dynamic recrystallization. However, the
formation of new grains during hot deformation of
aluminium has been frequently reported. Three types of
dynamic recrystallization are likely to produce such a
microstructure: (i) discontinuous dynamic recrystalliza-
tion (DDRX), i.e. the classical recrystallization, whichoperates by nucleation and grain growth; (ii) continu-
ous dynamic recrystallization (CDRX), which involves
the transformation of low angle boundaries into high
angle boundaries; and (iii) geometric dynamic recrystal-
lization (GDRX), generated by the fragmentation of
the initial grains.
Discontinuous dynamic recrystallization, which is
commonly observed in low stacking fault energy
metals, remains exceptional in aluminium and alu-
minium alloys. Nevertheless, it seems to occur in two
specific cases, viz., in high purity aluminium and in
aluminium alloys containing large particles:
1. Single crystals and polycrystals of 99.999 wt.% alu-
minium have been subjected to hot deformation in
compression by Yamagata [15]. Stress-strain
curves exhibit strong oscillations, typical of DDRX,although more irregular. The associated microstruc-
tures generally display new grains without substruc-
ture. It is therefore not excluded that these grains
have grown after deformation, all the more as
aluminium of such high purity recrystallizes stati-
cally very rapidly, even at room temperature. How-
ever, one micrograph [3] clearly displays the
presence of several grains containing a substructure,
in an initially monocrystalline sample. Moreover,
* Corresponding author. Tel.: +33-4-77420026; fax: +33-4-
77420157.
E-mail address: [email protected] (F. Montheillet)
0921-5093/00/$ - see front matter 2000 Elsevier Science S.A. All rights reserved.
PII: S 0 9 2 1 - 5 0 9 3 ( 0 0 ) 0 0 7 3 3 - 4
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the occurrence of DDRX in high purity aluminium
has been recently confirmed by Ponge et al. [6].
High purity produces two opposite effects. On the
one hand, it favors DDRX by increasing grain
boundary mobility. On the other hand, it can inhibit
DDRX since the high level of recovery prevents
dislocation accumulation, thus reducing the driving
force. Experimental results indicate that the first
effect prevails over the second [7].2. There is some evidence that DDRX can occur dur-
ing hot deformation of AlMgMn alloys, because
a high Mg solute addition raises the dislocation
density and thus the driving force for DDRX, while
large Al6Mn particles (\1 mm) stimulate nucle-
ation. The presence of small new grains adjacent to
large particles has been reported for instance after
plane strain compression of Al-1wt.% Mg-1wt.%
Mn [8] and extrusion of Al-5wt.% Mg-0.8wt.%Mn
[9]. However, the volume fractions of recrystallized
grains remained small and no effect on the shapes of
the stress-strain curves was detected. This means
that DDRX is a possible but limited restorationmechanism in aluminium alloys.
Continuous dynamic recrystallization occurs in turn
by the progressive accumulation of dislocations in low
angle boundaries, leading to the increase of their mis-
orientation and the formation of large angle grain
boundaries when their misorientation angles reach a
critical value qc (qc:15). This mechanism has been
observed in several high stacking fault energy metals,
such as aluminium and aluminium alloys [1013], b
titanium alloys [1416], and ferritic steels [1720]. The
microstructure of a commercial purity (1050 grade)
aluminium strained in torsion has been investigated by
Perdrix et al. [10], and Montheillet [11]. These authorsfound that, at small and medium strains (m:1), the
microstructure consists of the deformed initial grains
containing subgrains, which is typical of a recovered
state. By contrast, strongly strained samples (m:40)
exhibit a completely different microstructure: it is no
longer possible to distinguish the initial grains, and the
former subgrains now appear as crystallites bounded
partly by low and partly by high angle boundaries.
Furthermore, the misorientation angles, which display a
bimodal distribution at small strains (with subgrain
boundaries less than 15 and initial grain boundaries
between 30 and 63) become uniformly distributed be-
tween 0 and 63 at large strains. Perdrix et al. [10] have
explained these results by assuming a progressive trans-
formation of subgrain boundaries into grain
boundaries. However, this mechanism remains contro-
versial and some authors have suggested that the in-
creased fraction of high angle boundaries could result
from GDRX.
Nevertheless, there is a general agreement to consider
that the transformation of low angle boundaries into
high angle boundaries effectively takes place when the
boundaries are pinned by small particles. This mecha-
nism has been used to promote superplasticity in Zr
bearing or high Mg aluminium alloys (Al-6wt.%Cu-
0.4wt.%Zr [21,22], Al-0.25wt.% Zr-0.1wt.% Si [23,24],
Al-10wt.% Mg-0.1wt.% Zr [25,26], Al-10wt.% Mg-
0.5wt.% Mn [27,28]). These alloys are generally cold or
warm rolled, to increase their dislocation density. Un-
der these conditions, subgrains form very quickly dur-ing the subsequent hot tension testing. Since the
boundaries are pinned by small particles (Al3Zr,
(Al8Mg5)b or Al6Mn, according to the alloy composi-
tion) and continuously absorb dislocations, the sub-
grains transform into grains without growing. A fine
and equiaxed grain structure is thus obtained in the
early stages of superplastic deformation.
Finally, geometric dynamic recrystallization has first
been described by McQueen et al. [29] in a commercial
purity aluminium. During deformation, the original
grains flatten (compression) or elongate (tension, tor-
sion), and their boundaries become progressively ser-
rated while subgrains form. Consequently, the grainboundary area per unit volume grows strongly and an
increasing fraction of subgrain facets is made of those
initial grain boundaries. Ultimately, when the original
grain thickness is reduced to about two subgrain sizes,
the grain boundaries begin locally to come into contact
with each other, causing the grains to pinch-off [29
32]. In the case of Al-5Mg alloys, the tendency to the
serration of boundaries is stronger and a secondary
process of GDRX by pinching off of the serrations has
been reported [33,34].
The first purpose of this work was to bring forward
further evidence for CDRX. The main objection to theresults of Perdrix et al. [10] was that the high fractions
of large angle boundaries did not necessarily result
from CDRX, but could also be due to the evolution of
the initial grain boundaries. In order to exclude the
latter possibility, single crystalline samples were used.
The second objective was to better characterize CDRX.
The influence of the following parameters was therefore
investigated:
crystalline orientation, by using single crystals of
various orientations;
purity, by testing three grades of polycrystalline alu-
minium, ranging from 99.99 to 97 wt.% Al. Impuri-
ties and alloying elements reduce the recoverycapacity of the material, since they decrease its stack-
ing fault energy (although this effect seems to be
relatively weak in aluminium alloys) and solutes
as well as precipitates reduce the dislocation mobil-
ity;
strain path, by comparing the microstructures
obtained from uniaxial compression and torsion
tests.
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Table 1
Chemical compositions of the three aluminium grades
Mg Si Cr Mn Fe Cu Zn
12 1 1 101199 (mg/g) 457 4
1500 62 829 57001200 (mg/g) 59 200
0.115052 (wt.%) 0.22.49 0.08 0.3 0.01
2. Experimental procedure
The three selected grades of aluminium, a pure 1199
aluminium, a commercial purity 1200 aluminium and
an AlMg 5052 alloy (Table 1), were provided by the
Centre de Recherches de Voreppe (Pechiney) in the
form of hot rolled plates. The initial structures were
fully recrystallized, exhibiting equiaxed grains of 220
mm in the 1199 aluminium, flattened grains of average
size 200 mm in the 1200 aluminium, and again equiaxed
grains of 80 mm in the 5052 alloy (Table 2). The
textures displayed a strong cube component, especially
in the 1200 grade. Cylindrical compression specimens
were machined with their axes parallel to the rolling
direction. In addition, torsion specimens were prepared
from the 1200 aluminium with the torsion axis parallel
to the rolling direction. Finally, single crystals were
grown from the 1199 aluminium and sectioned to ob-
tain cubic specimens with a 001, 011 or 111
direction parallel to the compression axis (Table 3).
Hot compression tests were carried out at a constant
strain rate m;=102 s1 and 0.7 Tm (i.e. 380, 368 and
333C for the 1199, 1200, and 5052 grades, respec-
tively). The specimens were lubricated with graphite to
minimize strain inhomogeneities, and water quenched
within 1 s after deformation. The torsion specimenswere strained under the same strain rate and tempera-
ture conditions, and water quenched within 5 s. The
compression specimens were then sectioned parallel to
the axis, ground and electropolished (5 ml HClO4, 95
ml C2H5OH, 20 V, 0C, 60 s); some of them were
subsequently anodized (10 ml HBF4, 90 ml H2O, 30 V,
20C, 120 s). Since deformation was not uniform, local
strains were estimated using a mechanical model [35].
In what follows, the strain values correspond to the
center part of each specimen, where all observations
were carried out. However, even at large strains, the
strain gradient was quite low, e.g. in the compression
direction, Dm/Dz:0.2 mm1 at m=1.5 [35]. The tor-sion cylinders were ground to obtain a flat surface, and
then prepared the same way. The strain and strain rate
undergone by these samples were estimated to approxi-
mately 80% of their nominal (surface) values.
The deformed microstructures and textures were in-
vestigated using polarized optical microscopy (POM)
on the anodized specimens, scanning electron mi-
croscopy (SEM) in the channeling contrast mode, elec-
tron back-scattered diffraction (EBSD) and X-raydiffraction on the electropolished specimens. POM dis-
plays both the initial grains and the new crystallites
(subgrains or new grains), whereas only the crystallites
are generally revealed by SEM. This is due to the fact
that POM colors are related to the crystalline orienta-
tions, and the original grains are associated with re-
gions of similar colors. This is not the case in SEM,
where crystallites of similar orientations can display
quite different grey levels, and conversely. However,
POM is known to overestimate subgrain sizes
[31,36,37], and is therefore inadequate for quantitative
analyses of the substructure. For this purpose, SEMmicrographs were therefore used. Moreover, the color
contrast in POM is not precise enough to distinguish a
new grain from a subgrain and even less to estimate the
misorientation between two crystallites. The orientation
of each crystallite was thus determined using EBSD and
the misorientation associated with each boundary sub-
sequently calculated. On account of detection limitation
in hot worked structures, boundaries with a misorienta-
tion angle of less than 1 were not taken into consider-
ation. This omission is likely to be of no consequence,
however, since the study is focused on the transition
from low angle to high angle boundaries around 15. In
addition to these local texture measurements, global
textures were investigated by X-ray diffraction.
Table 2
Mean intercept lengths of the grains in the hot rolled plates along the
rolling (RD), transverse (TD), and normal (ND) directions
RD TD ND
2261199 (mm) 204 227
1200 (mm) 285 215 98
8678 595052 (mm)
Table 3Crystallographic directions parallel to the compression axis (CA) and
perpendicular to the lateral faces (TD1 and TD2) of the cubic single
crystals
TD1 TD2CA
[100][001] [010]
[100] [011(][011]
[112(][11(0][111]
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Fig. 1. Stress-strain curves of the 001, 011, 111 single crystals
and polycrystals of 1199 aluminium.
3. Experimental results
3.1. Pure aluminium single crystals
During the compression tests, the cross-sections of the
001 specimens remain square, as expected since the
001 crystallographic axis and the geometric axis
(square cross-section) are both fourfold, and the trans-
verse directions TD1 and TD2 are crystallographicallyequivalent (Table 3). The 011 (twofold crystallo-
graphic axis) specimens lengthen in the 100 direction.
It has been shown [38] that this shape change can be
explained by the activation of the four octahedral slip
systems (111)[101(], (111)[11(0], (1(11)[1(01(], and (1(11)[1(1(0].
Finally, the cross-sections of the 111 specimens be-
come irregular, which is due to the combination of their
threefold crystallographic axis and four-fold geometric
axis.
Fig. 1 shows that the flow stresses reach a steady state
value of about 10.5 MPa in the case of the 001 and
111 specimens (note that the polycrystal flow stress
tends to the same value); by contrast, it is much higherfor the 011 specimens (about 14 MPa). It will be
shown below that these values are closely related to the
subgrain sizes. The flow stress drop associated with the
111 crystals can be mainly attributed to the evolution
of the Taylor factor: indeed, at m=0, M=33/2:3.67,while at m=1.5 (101 orientation, see below) M=6:2.45.
Global texture measurements show that the 011
orientation is perfectly stable (Fig. 2(b)). The 001
orientation is also fairly stable, since its decomposition
only starts at m=1.5 (Fig. 2(a)). On the other hand, the
111 orientation is very unstable (Fig. 2(c)). At m=0.3,
the compression axis is roughly parallel to 112 andeventually reaches the 011 stable orientation at m=
1.5. These results are in good general agreement with the
literature, although the rotation amplitudes are larger
than those observed by Mecif et al. [39], after uniaxial
compression of aluminium single crystals at the same
temperature. This difference can be attributed, to a large
extent, to the higher strain (m=1.5 vs 0.35) and strain
rate (102 vs. 2104 s1) applied in the present work.
The aspect of the sections perpendicular to the com-
pression axis does not change significantly with increas-
ing strain for the 100 specimens observed by POM
(Fig. 3(a, b)). However, the band structure observed on
the lateral sections at m=0.3 transforms into a subgrain
structure at larger strains. It should be noted that SEM
reveals the presence of many subgrains within these
bands. The 011 specimens exhibit symmetrical mosaic
patterns on their (011() lateral sections (Fig. 3(c, d)), with
dislocation walls parallel to the planes of the
activated slip systems. SEM also reveals the formation
of small subgrains inside the cells at m=1.5. Such a
mosaic microstructure of the 011 specimens (which
Fig. 2. Global textures of the monocrystalline specimens strained to
m=1.5. (a) 001; (b) 011; (c) 111. The horizontal and vertical
axes are associated with the TD1 and TD2 directions, respectively
(Table 3).
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elongate only in the [100] direction) has already been
observed in the case of single crystals of various orien-
tations deformed by channel die compression [40]. This
suggests that such a particular structure is due to the
geometry of the slip systems, which allows a unidirec-
tional strain to take place. By contrast to the previous
orientations, the 111 specimens exhibit an inhomoge-
neous microstructure. Horizontal and tilted bands are
displayed on the TD1 lateral sections at m=0.3. How-ever, this inhomogeneity tends to vanish at larger
strains.
Misorientation maps were plotted from EBSD mea-
surements for each single crystal strained to m=1.5. In
the 001 specimen (Fig. 4(a)), a large number of
boundaries exhibit a misorientation greater than 15,
which indicates that subgrain boundaries have trans-
formed into grain boundaries. By contrast, the
boundaries in the 011 specimen (Fig. 4(b)) exhibit
low misorientations, generally smaller than 6, so that
no high angle boundaries have formed. The behavior of
the 111 specimen (Fig. 4(c)) is intermediate, since
only a small fraction of the interfaces consists of high
angle boundaries.
The evolutions of the misorientation distributions
with increasing strain are compared for the three orien-
tations in Fig. 5. In the case of the 001 crystal, the
average misorientation strongly increases with strain.
At m=0.9, a significant fraction of misorientations al-
ready exceeds 15, which clearly means that part of the
low angle boundaries have transformed into large angle
boundaries. At m=1.5, this trend is more pronounced,
since almost 20% of the interfaces are now grainboundaries. However, the microstructural steady state
is not yet attained, which suggests that a more recrys-
tallized microstructure (with crystallites bounded
mainly by large angle boundaries) could form at larger
strains. By contrast, for the 011 orientation, all the
measured misorientations are lower than 15, even at
m=1.5. Furthermore, no evolution is noticeable be-
tween m=0.9 and 1.5, and thus, the formation of grain
boundaries is unlikely, even at larger strains. The be-
havior of the 111 specimens is more complex. In
particular, the specimen strained to m=0.9 has devel-
oped a large amount of very high angle boundaries
(30 60). This is due to the splitting of the initial
orientation into two components (Fig. 6), which was
Fig. 3. POM microstructures of the monocrystalline specimens. (a) 001, m=0.3; (b) 001, m=1.5; (c) 011, m=0.3; (d) 011, m=1.5.
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Fig. 4. SEM micrographs and associated misorientation maps of the monocrystalline specimens strained to m=1.5. (a) 001; (b) 011; (c) 111.
The vertical and horizontal axes are associated with the CA and TD 1 (001 or 011) or TD2 (111) directions, respectively (Table 3).
observed in only one specimen of this orientation. It
should also be pointed out that, in the case of this
unstable orientation, large angle boundaries are formed
in the early stages of the deformation: at m=0.3, they
already represent more than 8% of the boundaries.
3.2. Polycrystalline specimens
The mechanical behavior of the various specimens
obtained from the compression and torsion tests was
investigated (see Fig. 1 for the 1199 grade). In all cases,
the flow stress seems to reach a plateau at m:0.3
(although a small decrease of the torsion stress is
expected at very large strains [10,11]). The flow stresses
reach steady state levels of 10.5, 21 and 77 MPa for the
1199, 1200 and 5052 grades, respectively. Global tex-
ture measurements were carried out on the compression
specimens. The pole figures of the three grades (Fig. 7
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Fig. 5. Strain dependence of the misorientation distributions for the three single crystals (N: number of measurements, q(: average misorientation
angle).
Fig. 6. Formation of deformation bands in the 111 specimen strained to m=0.9.
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Fig. 7. Global texture of the 1199 polycrystalline specimen strained to
m=1.5. The horizontal and vertical axes are associated with the TD
and ND directions, respectively (Table 2).
All the microstructures are quite homogeneous. At
m=0.3, subgrains are well formed in the 1199 and 1200
grades (Fig. 8(a)), whereas they are not visible in the
5052 grade, because their formation is delayed by the
strong solute content of this alloy. Note also that at
m=1.5, the grain thickness is still much larger than the
subgrain size (Fig. 8(b)), which indicates that the initial
grains are far from decomposing into smaller grains by
a GDRX process. In Fig. 9, the mean intercept lengthsD of the subgrains measured from the SEM micro-
graphs are compared with a theoretical value calculated
from the flow stress |, using the relationship |/G=hb/
D. The value of h=28 was chosen here, since it is
commonly used for aluminium [34]. In the case of the
1199 and 1200 grades, the calculated subgrain sizes are
consistent with the measured ones, although slightly
larger. The coefficient 28 seems therefore overestimated,
and a better agreement between the two sets of data is
obtained by using a value of 24. However, according to
Blum et al. [34], this difference is due to the condensa-
tion of free dislocations into additional subgrain
boundaries between the end of deformation andquenching, which causes a decrease of the average size
D. By contrast, the calculated subgrain size is much
smaller than the experimental value for the 5052 alloy.
This is due to the slow formation of the subgrains in
that case, since the presence of regions without well-
formed subgrains leads to an overestimation of the
mean intercept length.
Fig. 10 displays misorientation maps of the polycrys-
tals strained to m=1.5. In the 1199 specimen (Fig.
10(a)), two initial grain boundaries can still be iden-
tified because they form a continuous chain of high
misorientation segments (\30), whereas isolated high
angle boundaries with lower misorientations (15 30)are very probably new grain boundaries. In the case of
the 1200 aluminium strained in compression (Fig.
10(b)), the map shows the presence of a large fraction
Fig. 8. POM microstructures of the 1200 polycrystalline specimens.
(a) m=0.3; (b) m=1.5. The compression axis is vertical.
Fig. 9. |/G versus b/D plot for the single crystals and polycrystals
strained to 1.5 (the shear moduli G=20.6, 20.8, 21.2 GPa for the
1199, 1200, 5052 grades, respectively, at 0.7 Tm, and the Burgers
vector length b=2.861010 m).
illustrates the texture of the 1199 aluminium) show that
the texture consists mainly of a 011 fiber component,
which is especially strong for the 5052 alloy. Since a
strong cube component was observed in the initial
state, this means that the crystallites have strongly
rotated during compression.
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Fig. 10. SEM microstructures and associated misorientation maps of the polycrystalline specimens strained to m=1.5. (a) 1199 specimen; (b) 1200
compression specimen; (c) 1200 torsion specimen. The compression axis is vertical.
of high angle boundaries, but the original grainboundaries are no longer recognizable. It is likely that
subgrains have rotated more in the 1200 than in the
1199 grade, thus leading to a fragmentation of the
initial grain boundaries. In the torsion specimen of Fig.
10(c), the fraction of high angle boundaries is very
similar, although some chains of segments exhibiting
very large misorientations at the top of the map look
like initial grain boundaries. The misorientation distri-
butions of the 1199 aluminium are depicted in Fig. 11.For this aluminium, as for the two other grades, there
is a progressive shift of the small misorientations to-
wards the larger ones, which leads to an increase of the
average misorientation of about 8 between m=0.3 and
m=1.5. The fraction of subgrain boundaries with small
misorientations (B6) is strongly reduced; the number
of subgrain boundaries with larger misorientations in-
creases at first, and then decreases, whereas the fraction
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of high angle boundaries continuously increases, reach-
ing more than 30% at m=1.5. This indicates that the
very low angle boundaries (B6) are continuously con-
verted into medium angle boundaries (615), which in
turn transform into high angle boundaries. Such an
evolution may be referred to as continuous
recrystallization.
3.3. Comparison of the geometric and continuous
dynamic recrystallization kinetics
In the polycrystalline specimens, the increase of the
high angle boundary fraction is due not only to the
formation of new grain boundaries, but also to the
expansion of the initial grain boundary area. In order
to get a quantitative estimation of the respective effects
of CDRX and GDRX, a geometrical model was used.
The evolutions of the lengths per unit area of initial
grain boundaries, new high angle boundaries, and low
angle boundaries were determined from a simple
derivation detailed in the Appendix A.Fig. 12(ae) illustrate the GDRX and CDRX kinet-
ics. The strain dependence of the initial grain boundary
fraction is not monotonic since it grows due to the
flattening of the initial grains, whereas the formation of
subgrain boundaries makes it decrease. At m=1.5, the
initial grain boundaries represent about 8% of the
interfaces for the 1200 and 5052 grades. The 1199
aluminium differs by a much higher fraction, about
17%, which is due to a smaller initial grain size/sub-
grain size ratio. By comparing the thickness h of the
initial grains at m=1.5 (the latter were estimated by
calculation to 68, 80, and 33 mm for the 1199, 1200, and
5052 grades, respectively) and the subgrain sizes D atthe same strain (14.3, 6.6 and 3.4 mm, respectively), it
appears that the h/D ratios are about 5 for the 1199
grade, and 1012 for the 1200 and 5052 grades. This
confirms that GDRX is more developed in the pure
aluminium. But this ratio still remains far from 2,
which means that the strain achieved by compression is
not large enough to obtain a complete GDRX
structure.
With regard to CDRX, the 001 single crystal (Fig.
12(d)) and the polycrystal of same purity (and strong
initial cube texture, Fig. 12(a)) display similar behav-
iors: at m=1.5, roughly 1518% of their interfaces
consist of new high angle boundaries. When deforma-
tion bands form in the 111 crystal, the fraction of
new grain boundaries rises very quickly; if not, it
remains rather small, about 5 8%. Among the poly-
crystalline specimens, the highest fraction of new grainboundaries is observed in the 1200 specimens (35% in
compression, 39% in torsion at m=1.5), then in the
5052 alloy (about 24%), and last in the pure aluminium
(only 15%). This can be explained by the very high
recovery rate in pure aluminium, which lowers the
accumulation rate of dislocations in the subgrain
boundaries. On the other hand, in the Al-Mg alloy, the
solute atmospheres impede dislocation movements,
thereby delaying the formation of subgrain boundaries
[7]. It should be noted that the grain boundaries present
in the 1199 aluminium originate in equal parts from
GDRX and CDRX, while those present in the 1200
and 5052 grades have mainly developed by CDRX. Thepresent results also indicate that the compression and
torsion specimens display quantitatively the same be-
havior. Therefore, the strain path does not seem to
modify the CDRX kinetics significantly.
4. Discussion
4.1. New grain boundaries and deformation bands
Experiments carried out on single crystals confirm
that the subgrain boundary misorientations stronglyincrease with strain: a maximum of at least 15 is
reached for the three investigated orientations at m=
0.9. Moreover, the transformation of low angle
boundaries into high angle boundaries is clearly demon-
strated for the two unstable orientations, as well as in
the polycrystalline specimens where the fraction of high
angle boundaries is much larger than expected for the
deformed initial grain boundaries.
Fig. 11. Strain dependence of the misorientation distributions for the 1199 polycrystals ( N: number of measurements, q(: average misorientation).
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Fig. 12. Strain dependence of the surface fractions pertaining to the
various kinds of interfaces, viz. low angle boundaries (LAB), new
high angle boundaries (nHAB) [including the deformation band
boundaries (DB)], and old initial high angle boundaries (oHAB). (a)
1199 polycrystals; (b) 1200; (c) 5052; (d) 1199 001; (e) 1199 111.
However, the misorientation increase of the subgrain
boundaries and their transformation into high angle
boundaries has been controversial for a long time,
principally on the basis of the work by Kassner and
McMahon [30]. These authors studied the microstruc-
tural evolution of high purity (99.999 wt.%) polycrys-
talline aluminium specimens strained in torsion (370C,
5104 s1). The observations were mainly carried
out by transmission electron microscopy, and the mis-orientations were estimated from either dislocation
spacings (subgrain boundaries) or selected area diffrac-
tion patterns (grain boundaries). These authors noticed
only very limited increases in subgrain misorientations
with strain: the mean values varied from 0.5 at m=0.2
to a saturation value of 1.2 from m=1.2 to 16. Most of
the discrepancies between this prior work and the
present investigation can be explained by differences in
the experimental procedures. First of all, the aluminium
used by the authors was purer, and the strain rate,
lower. This means that recovery was more efficient, and
thus the increase in misorientation was slowed down.
Furthermore, by contrast to the SEM-EBSD techniqueused here, TEM allowed to account for boundaries
with very low misorientations (B1), therefore decreas-
ing the average misorientation value. However, some
boundaries with medium (about 10), and large misori-
entations (\30) were also observed by Kassner and
McMahon [30]. Since a limited number of measure-
ments (about 20 for each specimen) were carried out,
the authors misorientation distributions are not
smooth and discontinuous values are observed within
the range 515, instead of the continuous distributions
displayed here (Figs. 5 and 11). This is a reason why
these medium angle boundaries, which were also
present in single crystalline specimens strained undersimilar conditions [41], were interpreted as interfaces
between persistent deformation bands (although such
bands were not clearly identified), and not as former
subgrain boundaries transformed into new grain
boundaries.
It should be noted that the classification of the
various boundary types has been developed in the case
of cold deformation and its application to hot worked
structures in not obvious. Indeed, at temperatures be-
low 0.4 Tm, several kinds of interfaces are observed.
The initial grains decompose into deformation bands
separated by a thin (12 mm) transition band contain-
ing dislocation cells. Inside the deformation bands,
random low misorientation cells are grouped together
in cell blocks separated by dense dislocation walls of
higher misorientations [43]. However, as temperature
increases, dislocation walls become thinner and a much
more homogeneous microstructure develops. Generally,
above 0.6 Tm, only equiaxed subgrains are observable
inside the deformed initial grains. Even when a decom-
position occurs, the deformation bands are not easily
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S. Gourdet, F. Montheillet /Materials Science and Engineering A283 (2000) 274288 285
observed on micrographs (Fig. 6). The reason is that
they appear bounded by true grain boundaries, instead
of the straight above-mentioned transition bands. Dif-
ferentiation between several types of low and high angle
boundaries is therefore difficult in hot worked struc-
tures, and the relevance of such a classification even
questionable.
4.2. Influence of crystalline orientation
Results obtained on single crystals confirm that the
CDRX kinetics strongly depend on the crystallographic
orientation. The fully stable 011 orientation enables
only a limited increase of the misorientations. No high
angle boundary creation was observed, although some
misorientations are close to 15 at m=0.9. In the case of
the quasi-stable 001 orientation, the formation of
new high angle boundaries, with 15 30 misorienta-
tions, was observed in the specimens strained to m=0.9
and 1.5. For the 111 orientation, which is fully
unstable, new grain boundaries in the same misorienta-
tion range were observed, starting even at m=0.3. Suchresults are in total agreement with previous work on
aluminium single crystals. Theyssier et al. [40] observed,
after channel die compression of single crystals of the
same purity, that some orientations (e.g., {112}111
or {421}112) led to the formation of high angle
boundaries, whereas other ones (e.g. {110}112) did
not. More recently, Ponge et al. [6] investigated the
misorientations developed during hot uniaxial compres-
sion (260C, 4104 s1) of a very high purity alu-
minium (99.9995 wt.%), which recrystallized by DDRX.
In the unrecrystallized matrix, these authors observed
the occurrence of smaller misorientations in the 011
single crystal (B9), than in the 112 specimen (up to26). The misorientation ranges were thus very close to
those obtained in the present work.
4.3. Continuous dynamic recrystallization and grain
rotations
In order to clarify the origin of the new high angle
boundaries, it is interesting to look more closely at the
relationship between global texture and local misorien-
tations. In the case of uniaxial compression, the follow-
ing remarks can be made:
1. Stable orientation. During straining, the global tex-
ture of the 011 crystals remains unchanged, ex-
cept that it becomes slightly less sharp. This is due
to strain hardening and dynamic recovery: disloca-
tions accumulate in the subgrain boundaries, thus
altering the initial orientation. Misorientations up to
15 have been reported in pure aluminium single
crystals. The results obtained on polycrystals indi-
cate that in a less pure aluminium, such as the 1200
grade, the misorientations are probably larger, thus
leading to the formation of high angle grain
boundaries. A question still remains open: which
mechanism leads to the increase in misorientation?
If one considers that dislocations of opposite signs
are created in equal densities during straining, each
dislocation wall will absorb dislocations of both
signs, keeping its misorientation at a low level. A
misorientation increase can only occur if a subgrain
boundary absorbs an excess of dislocations of onesign, which supposes that the various types of dislo-
cation are not uniformly distributed in the material.
This assumption does not seem unrealistic, however,
all the more as misorientations between adjacent
subgrains will increase such inhomogeneities. In-
deed, slip system activities can be affected by small
orientation changes. Furthermore, among the vari-
ous orientations introduced by strain, those belong-
ing to the 011 fiber are fully stable and will
probably not disappear, thus leading to increased
and permanent misorientations. A slight trend to-
wards fiber formation in the pole figure of the 011
crystal strained to m=1.5 can be noticed (Fig. 2(b)).If compression specimens could be deformed to very
large strains, a 011 fiber texture would certainly
be observed.
2. Unstable orientations. New interfaces can be intro-
duced by both strain hardening and lattice rota-
tions. Two cases must be distinguished, according to
whether deformation bands occur or not. Let us first
consider the case in which the whole crystal rotates
towards the same orientation. In the 111 single
crystal strained to m=0.3, the EBSD local pole
figure clearly shows that some orientations still re-
main close to the initial one, whereas others are
already located near the final one. That is why highangle boundaries are observed at such a low strain.
However, at m=1.5, all the crystallites have reached
their final 011 orientation, which explains the
lower fraction of high angle boundaries. The subse-
quent behavior has been described in (i). In the
second case, when the initial orientation splits into
symmetrical components, very large and permanent
misorientations are rapidly built up. Deformation
bands are expected to occur only for specific initial
orientations, e.g. 001, 111, or the intermediate
uuw orientations, and they are more likely to
occur in single crystals than in polycrystals (except
for very large initial grains).
Lyttle and Wert [23] have formulated three models
based on dislocation glide, boundary sliding, and neigh-
bor switching to account for the increased misorienta-
tions during straining of superplastic alloys. They
concluded that combination of the boundary sliding
model and the neighbor switching model most closely
reproduced the misorientations measured experimen-
tally. In the case of the dislocation glide model, the
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S. Gourdet, F. Montheillet /Materials Science and Engineering A283 (2000) 274288286
misorientations tended to decrease rather than increase,
therefore reflecting texture development. However, it
can be inferred from the above discussion that the
convergence of the crystallite orientations and the cre-
ation of large misorientations are not incompatible. A
reason is that strain hardening tends to move the
crystallites away from their ideal orientations. This
effect was probably very pronounced in this case be-
cause of the low recovery level associated with highalloying element content. Moreover, it was experimen-
tally found that, for some reason, the lattice rotation
rate varied from one crystallite to another, thus intro-
ducing high misorientations during the transient. All
these disturbing factors were not taken into account in
the model. The increase in subgrain boundary misorien-
tation by accumulation of dislocations was probably
predominant in the early steps of straining. However,
since the crystallite size was very small, grain boundary
sliding and grain switching probably prevailed as soon
as large enough misorientations were built up.
4.4. Kinetics of the continuous dynamic recrystallization
It is worth noting that the various mechanical and
structural parameters tend to their steady state values
at different rates. The flow stress remains approxi-
mately constant from m=0.3, except for the 111
crystal. In this case, the crystallographic rotation delays
the steady state up to m=0.9. The subgrain size is
generally well established at m=0.3. However, a sub-
grain refinement is still observed in the 5052 alloy at
m=1.5. By contrast, the misorientations are not yet
stabilized at m=1.5, with the only exception of the
011 crystal. This is due to the combined effects of the
increasing misorientations of the subgrains and newlyformed grain boundaries on the one hand, and the
increasing surfaces of the original grain boundaries on
the other hand. Apart from some minor cases, the flow
stress and the average subgrain size thus reach their
steady state values quite rapidly (m=0.3), while the
average misorientation is still increasing at m=1.5.
These experimental results are thus in contradiction
with the similitude principle, which stipulates that the
various microstructural spacings are inversely propor-
tional to the flow stress. For instance, McQueen and
Blum [42,44] proposed that there is a unique relation-
ship between the average dislocation spacing s in the
subgrain boundaries, (or, equivalently, the average sub-
grain boundary misorientation q:b/s) and the flow
stress. The fact that the average subgrain boundary
misorientation increases without affecting the flow
stress significantly can be explained if one considers, as
established by several authors [45,46], that the strength-
ening associated with dislocations inside the subgrains
is larger than that due to dislocations in the subgrain
boundaries. The former is thus the main controlling
parameter of the flow stress. The assumption of Mc-
Queen and Blum was based mostly on creep data from
specimens deformed to low or moderate strains, which
can explain why these authors did not observe any
misorientation increase of the subgrain boundaries.
4.5. Elementary mechanisms of the continuous dynamic
recrystallization
Continuous dynamic recrystallization has sometimes
been labeled extended dynamic recovery, either because
some authors restrict the term recrystallization to the
classical discontinuous recrystallization, or because the
microstructures generally consist of crystallites only
partially bounded by high angle boundaries. This termi-
nology is somewhat misleading, however, since dynamic
recovery can have detrimental as well as beneficial
effects on CDRX. Indeed, the above experimental ob-
servations indicate that the CDRX process results from
the combination of three elementary mechanisms:
1. The formation of subgrain boundaries. These
boundaries are created with a very low misorienta-tion angle (about 1), as a result of dynamic
recovery.
2. The transformation of subgrain boundaries into
grain boundaries. The increase in misorientation of
the subgrain boundaries is more or less rapid, ac-
cording to the material and the experimental condi-
tions. It was shown that it is accelerated by medium
recovery levels and when the initial orientation is
unstable. If such favorable conditions are brought
together, the subgrain boundaries can be gradually
transformed into grain boundaries.
3. The elimination of subgrain and grain boundaries.
Measurements carried out on a b titanium alloy[13,47] and some aluminium alloys have recently
shown that the grain boundaries migrate, even in
the absence of classical DDRX, although at much
lower rates. All the interfaces present in the volume
swept by the migrating boundaries disappear, which
certainly plays a major role in the establishment of
the steady state, for both GDRX and CDRX. It has
been shown for instance that, in compression, the
initial grains can reach a quasi-steady state thick-
ness, which is an increasing function of the
boundary velocity [13,47]. Moreover, the elimina-
tion of interfaces allows the subgrain size and the
misorientation distribution to stabilize after a tran-
sient period.
From the previous experimental observations, a
model of CDRX has recently been proposed, in which
the above three mechanisms are combined in order to
predict the microstructural evolutions [38,48]. The main
features are well reproduced: during the transient, the
subgrain size decreases while the misorientation grows,
and the subgrain boundaries are gradually converted
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S. Gourdet, F. Montheillet /Materials Science and Engineering A283 (2000) 274288 287
into grain boundaries; ultimately, the crystallite size
reaches a steady state value, as well as the misorienta-
tion distribution, although at larger strains.
5. Conclusions
Various aluminium specimens were submitted to uni-axial compression and torsion testing at 0.7 Tm and
102 s1, up to m=1.5. The main results obtained after
investigation of the hot worked structures are the
following.
(1) Experiments carried out on 001, 011 and
111 single crystals confirm that the subgrain
boundary misorientations strongly increase with strain.
For the three investigated orientations, the largest val-
ues are close to or beyond 15. In the case of the
stable 011 orientation, the increase in misorien-
tation is not sufficient to transform low angle into
high angle boundaries. Among the two other orienta-
tions, the conversion occurs earlier for the most un-stable one: it starts at m=0.3 in the unstable 111
crystal, but only at m=0.9 in the metastable 001
crystal.
(2) Most of the new high angle boundaries have
1530 misorientations. They originate either from dif-
ferences in the rotation rate towards the final orienta-
tion, or from fluctuations around the average
orientation introduced by strain hardening and dy-
namic recovery. In some cases, the initial orientation
splits into symmetrical components, causing the occur-
rence of deformation bands. In this case, very high
misorientations (\30) are rapidly built up.
(3) In the case of polycrystals, the high angleboundary fractions strongly increase with strain. Since
both continuous and geometric dynamic recrystalliza-
tion are likely to occur, calculations were carried
out in order to estimate the surface fraction of the
initial grain boundaries. It turns out that the high angle
boundary fraction is much larger than could be ac-
counted for by the expanded initial boundaries, thus
confirming the presence of many new high angle
boundaries.
(4) Continuous dynamic recrystallization is more effi-
cient in the commercial purity aluminium than in the
pure aluminium and the Al-Mg alloy. This
suggests that the transformation of low angle
boundaries into high angle boundaries is faster when
the recovery level is neither too high (the accumulation
rate of dislocations in the subgrain boundaries de-
creases), nor too low (the formation of subgrains is very
slow or they do not form at all). Comparisons between
compression and torsion specimens also indicate that
the strain path does not alter CDRX kinetics notice-
ably.
Acknowledgements
The authors are indebted to Professor J.J. Jonas,
McGill University, Montreal, for providing access to
the torsion facility. They are also grateful to Professor
H.J. McQueen, Concordia University, Montreal, for
many fruitful discussions. The work of S. Gourdet was
supported in part by the Region Rhone-Alpes, France,
through a scientific fellowship (Avenir program).
Appendix A
The evolution of the fraction of the various interface
types (i.e. low angle boundaries, initial and new high
angle boundaries) is addressed here in the case of
compression. Similar equations apply for torsion (see
Ref. [38] for more details). Since experimental measure-
ments provide interface lengths per unit area, a two-di-
mensional approach was chosen. Initial grains and
subgrains are approximated by ellipses with semiaxes
a=y/4 DCA and b=y/4 DTD, where DCA and DTDdenote the measured intercept lengths parallel to the
compression axis and the transverse direction,
respectively.
The mean intercept lengths of the initial grains can
be accurately measured only at m=0. Their evolutions
with strain are evaluated by taking into account the
flattening of the grains and the migration of the initial
boundaries, which leads to an increase of the average
thickness [13,38]. Along the direction parallel to the
compression axis:
D:CA=DCAm;+26 (1)
where the migration rate 6=0.1 mm/s at m;=0.01 s1
[13,38]. This yields after integration:
DCA=(D0DS) exp(m)+DS (2)
where D0 is the initial length and DS=26/m;, or, for the
semiaxis:
ag=(a0aS) exp(m)+aS (3)
where aS=y6/2m;. Moreover, perpendicular to the com-
pression axis:
bg=b0 exp(m/2) (4)
(in the above equations, a0 and b0 are the semiaxes
lengths at m=0).
In addition, the initial grain boundary length is af-
fected by the presence of serrations. Measurements
from POM micrographs showed that the length ratio k
between a serrated and a straight boundary first in-
creases with strain (as subgrains form) and then
remains approximately constant. This evolution can be
accurately described by the equation k=1+
0.25 [1 exp( 4m)] . The current length per unit
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area of the initial grain boundaries is then given by
kPg/2yagbg, where the ellipse perimeter Pg is calculated
numerically from the lengths of the semiaxes ag and bg.
Since the subgrain semiaxes asg and bsg can be mea-
sured at various strains, the total interface length per
unit area is simply given by Psg/2yasgbsg. The total
(old+new) high angle grain boundary length is then
estimated by multiplying the total interface length by
the fraction of high angle boundaries measured byEBSD (see the misorientation distributions). Finally,
the length per unit area of the new high angle
boundaries is obtained by difference.
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