An Experimental Study of the Recrystallization Mechanism During Hot Deformation of Aluminium

7/29/2019 An Experimental Study of the Recrystallization Mechanism During Hot Deformation of Aluminium

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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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S. Gourdet, F. Montheillet /Materials Science and Engineering A283 (2000) 274288 275

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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S. Gourdet, F. Montheillet /Materials Science and Engineering A283 (2000) 274288276

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


12/15


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


13/15


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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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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An Experimental Study of the Recrystallization Mechanism During Hot Deformation of Aluminium

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Transcript of An Experimental Study of the Recrystallization Mechanism During Hot Deformation of Aluminium