Varieties and Limitations of Dynamic Recrystallization Mechanisms in Al Alloys
H.J. McQueen
Mech.Eng., Concordia University
Montreal, Canada H3G 1M8
ABSTRACT
Aluminum and its alloys are generally noted for dynamic recovery (DRV) during hot working to
such a high level that they do not usually undergo dynamic recrystallization (DRX); although if held
at or above the working temperature they do statically recrystallize. Among the exceptions to be
discussed here is particle stimulated nucleation (PSN) of DRX in Al-Mg-Mn alloys due to large
particles and the dense substructure. In alloys, such as Al-Cu-Zr, which have been processed to
develop medium particles pinning a substructure with high misorientation walls, continuous DRX
takes place during the initial stages of superplastic forming at low strain rates. In Al 99.999 the
grain boundaries are highly mobile so that discontinuous DRX takes place in the classical manner.
At very high strains where the elongated grains have a thickness of only 2 or 3 subgrain diameters,
the opposite serrated boundaries begin to meet, pinching off the grains into short segments; this
gives the appearance of DRX so is called geometric DRX. Finally in the warm end of the hot range,
a noticeable density of boundaries misoriented over 10° develop inside the grains; this could be
attributed to the gradual transition from hot to cold working; however some researchers maintain
this is continuous DRX in which all subboundaries are continually rising in misorientation.
INTRODUCTION
Recently Ponge, Bredehoft and Gottstein [1] published a clear demonstration of discontinuous
dynamic recrystallization (dDRX) in Al of extremely high purity (99.9995%). The behavior of
these single crystal specimens was very similar to those of low SFE fcc single crystals well
-4 -1
documented in the past [2-4]. These results confirm under new conditions (260°C, 10 s )
behavior of 99.999+Al at both (400°C) [5] and 20°C [5,6]. As pointed out by all these authors the
occurrence of dDRX in such pure Al is the result of the very high mobility of the grain boundaries
(GB) which permits DRX, or after working (static) SRX, to take place at 20°C [5-7]. The reasons
why the discontinuous form is generally not observed under greatly varied conditions are explained
[8-18]. The various forms of DRX that have been observed in Al alloys are described and their
critical requirements assessed: particle stimulated nucleation (PSN-DRX) in AlMgMn; continuous
cDRX in Al-Cu-Mg in the early stages of super-plastic deformation and geometric gDRX at high 
where the grain thickness is close to the subgrain size.
HIGH LEVEL OF DYNAMIC RECOVERY
The commercial mechanical shaping of Al alloys with air cooling regularly results in material with
elongated grains, serrated GB and subgrains dependent inversely on Z (=  exp Q/8.3T,  strain
rate, Q activation energy  140kJ/mole), T temperature). (Figures 1,2) [8,9,12,15-18]. During
extensive straining, the subgrains remain equiaxed and constant in size as a result of
repolygonization, that is subboundaries (SGB) migrate or decompose and reknit [8,12-22]; in
association, the serrations in GB which result from interaction with SGB must continually rearrange
as the GBs lengthen in the elongation direction (Figure 2) [23-25]. Under equivalent circumstances,
Cu and Ni of similar purity exhibit equiaxed grains formed by SRX during cooling; where DRX has
occurred, exceptional quenching is required to prevent metadynamic recrystallization (static growth
of DRX nuclei) [9-11]. In high speed rolling at equivalent T to 90% reduction, a high frequency of
highly misoriented potential nuclei in Ni, Cu and Cu-30Zn caused dSRX in a few seconds whereas
in Al a high degree of recovery required over an hour annealing for dSRX [26]. Surveys of the
literature shows that only DRV occurred during varied experiments over a wide range of T,  and
 [14,19-30] and moreover, that analysis by energy dissipation efficiency wrongly affirmed the
presence of DRX [14].
Experiments have been conducted on Al and Al-Mg alloys to high  to induce DRX but a new
phenomenon was discovered when strains were increased to values between 16 and 60 [19-22].
When the grains (100 m ) become so thin that the serrations on opposite sides touch, the grains
pinch off into short segments containing low  SGB (Figures 2,3). Because of the serrated
boundaries, the microstructure consists of subgrain sized crystallites with about one third of their
perimeters high angle boundaries. After such straining with similar flow curves, specimens with
2000 m grains still exhibited elongated grains with equiaxed subgrains (about 4 to 6 across) and
with the same texture as the 100 m grains; this was quite different from Cu which had undergone
dDRX [8-10,31,32]. This process of DRV which produced grain refinement was called geometric
gDRX; the phenomenon has often been mistakenly reported as dDRX [13,31]. Subsequently it was
shown that gDRX occurred also in Al-11Zn, Al-2Mg-1Si, and Al-5Mg at strains which decreased as
Do was finer and the subgrain size ds larger due to higher T or lower  (Figure 4) [15,24,25,29,30].
In Al-Mg alloys, it was shown that the serrations became much more meandering (billowing) than
in Al and could pinch off at their base creating crystallites which did not grow because of the equal
substructure within and around (Figures 2,4) [22-25,29,33,34]. Such detachment and rotation of
serrated units has been called rotation rDRX [17,34,35].
DISCONTINUOUS DYNAMIC RECRYSTALLIZATION
At high strain rates in commercial purity polycrystalline Cu, Ni,  -Fe and stainless steels, classical
discontinuous dDRX is characterized by a single peak in the flow curve with softening to a steady
state regime [8-11,16,36-39]. The nuclei form as necklaces at the GB and have usually been
preceded by serration formation; growth stops at a size Ds characteristic of Z
independent of initial size Do (Ds << Do: when Ds > Do/2, the formation pattern changes to give
distinct multiple waves of DRX [11,36]). Unlike dSRX for which growth of grains ends due to
impingement, in dDRX it ends due to the re-insertion of the substructure; in steady state continual
distributed nucleation takes place to keep Ds constant and grains almost equiaxed to high strains
(10-40) (Figure 5) [10,11]. In 300 series stainless steels, dDRX nucleates at the GB at only a
fraction (0.1) of the serrations and proceeds according to the Avrami theory analysis to complete
one wave by the start of the steady state [37-39]. Nuclei have been observed forming within the
grains before the peak and also during steady state. The critical strain  cdDRX for initial nucleation
is always higher than that for SRX  cdSRX which occurs spontaneously during holding at T
following the test [10,11,16,37]. As  rises,  cdDRX rises whereas  cdDRX decreases; however, the
rate of DRX rises (tcDRX =  cdDRX/  and t0.99DRX =  s/  ) as expected from the higher substructure
density. The density of high strain energy sites for dDRX (enlarged cells of high misorientation) is
greater than that for dSRX in which static recovery (SRV) has an opportunity to develop the nuclei
[37]. In stable single crystals,  rises to a critical value greater for higher levels of DRV (lower
heterogeneity) and of Z, being much higher than in polycrystals [2-4,40,41]. Generally shown by a
rapid drop in  , a single grain nucleates at a high angle disorientation boundary, as for SRX [40,41]
or for DRX in large grains of polycrystals [42], and rapidly produces twin chains with new
orientations favouring high GB mobility as in annealing of cold deformed material [2-4,40,41].
The observations of dDRX in 99.999 Al have been mentioned in the introduction [1,5,6]. However
in ancient reports, when zone refined 99.9995 Al was worked and held at T between 77 and 293 K,
SRX was observed after longer times for lower T. Either in these cases or when
annealed at T above that of working, there were a little recovery and a higher rate of nucleation (but
not growth) compared to less refined or dilutely re-alloyed metal indicating that it is the formation,
rather than the migration, of the mobile boundary which is enhanced [43-44]. In very high T creep
tests on large grained Al specimens, sudden rises in  were associated with migration of long
segments of GB over some fraction (<0.4) of the grain area without change in number of grains; in
the absence of nucleation it was named dynamic grain growth DGG [46]. During low  ,
deformation at 350-450°C of Al with a high density substructure from thermomechanical
processing, nucleation of dDRX took place when the ratio of  exceeds a critical minimum;
however, growth ceased as the matrix recovered dynamically [47]. Generally, secondary straining
at increased T or decreased  induces rapid conversion to the substructure characteristic of the new
condition [12]. In duplex hot working tests on Cu or  Fe, DRX could be speeded up or delayed by
different levels of reduced stress [10,49,50].
An exceptional case is particle stimulated nucleation (PSN-dDRX) in Al-5Mg-0.7Mn during both
torsion and extrusion (Figure 5) [12,51]. PSN-dSRX is not only common in Al alloys with
constituent particles (>0.6 m ) but is used industrially for grain size control; PSN arises from the
fine cells formed around the rigid particles due to turbulent plastic flow [52]. However at elevated
T, the high level of DRV would cause the fine cells to decompose in a pure Al matrix whereas this
was not likely in an Al-Mg matrix where Mg atmospheres retard dislocation mobility resulting in
flow stresses and frequency of SGB (reciprocal subgrain size, l/ds) larger by a factor of about 4 [2830].
CONTINUOUS DRX IN AL ALLOYS
Continuous SRX has been observed in dispersoid alloys where fine particles strongly stabilize the
substructure inhibiting the formation of nuclei and the growth of grains even in extended annealing
at high T. However annealing for medium time at intermediate T resulted in most of the dislocation
walls converting into normal but strongly pinned GB which defined grains only slightly larger than
the initial subgrains [53]. In development of fine grained material capable of superplastic
deformation (SPD), TMP at about 300°C in the form of numerous rolling passes have been applied
to Al-Cu-Mg-Zr or Al-10Mg-Zr alloys [54,55]. During subsequent heating at the SPD temperature,
dSRX did not occur due to the particles, nor did cSRX in the short time. However, during the initial
stages of SPD (   100%), the material underwent cDRX at a much higher rate than has been
common for cSRX [55-57]. In particular, the presence of shear stresses induced more rapid
conversion of cell walls with greater than 15° misorientation and the sliding at those boundaries and
additional lattice strain led to the increase in misorientation and conversion of the remaining cells
walls. For less stabilized alloys, deformation under SPD conditions led to DRV and progression
towards the characteristic substructure [48].
In 99.5 Al, Montheillet and co-workers [31,58] have examined hot worked specimens quantitatively
to establish whether SGB were increasing in misorientation with strain and developing into true GB.
Recently Gourdet and Montheillet [59] presented experimental evidence for polycrystals (99.2 and
99.992% Al) of a greater density of high angle boundaries than could be accounted for by plastic
thinning of the grains accompanied by growth of softer grains (at expense of harder ones) leading to
loss of original GB. A marked increase of high angle boundaries in unstable single crystals, and
only a small increase in stable ones (Figure 6) were similar to those of Ponge et al [1]; however as
the crystals were 99.99 Al, no dDRX occurred. Deformation bands were identified only in the
<111> single crystal, although Kassner, McQueen, and Blum [13,15-19,60] have argued that all
new high misorientation boundaries arose as persistant disorientation walls between deformation
bands (as found in the torsion of single crystal 99.99 Al [60].
In the mesoscale model of Gourdet and Montheillet, subgrain boundaries continuously form, some
annihilating and some increasing in misorientation  to reach the critical value (~10°) for
conversion into GB and eventual absorption by migrating ones [31,58,59]. The continual increase
in  of SGB so that at a strain much higher than  s the distribution stabilizes with a noticeable
fraction of  > 5° < 10° is the significant difference from Blum and McQueen [15-18]. In cDRX
theory, the increased density of dislocations in the SGB has little effect on the flow stress since it
depends mainly on the internal dislocation density. Moreover, the steady state is achieved partly
through a slow GB migration by a factor of about 100 less than that for dDRX under similar
conditions [62]; this also serves to increase the average grain size by elimination of some grains.
However, this model does not assume any precise mechanism to cause the increase in  of some
facets. Wert and colleagues [56,57] have shown that an experimental array of subgrains with a 
distribution from earlier straining will not easily develop increased misorientation through Taylor
rotations during additional straining. The continued increase in  of some SGB to produce a few
high  segments is classed by Gourdet and Montheillet as cDRX although being quite different
from cSRX in which almost all the cell walls transform into high  GB. Moreover, the texture of
Al sharpens with increasing strain and does not become more random as in Cu when new grains
appear due to dDRX [32]. Furthermore, the incidence of such high  segments increases as T falls
into the warm range (Al,300-100°C) although optical microsocpy continues to show the elongated
grains and x-rays the same texture; a complete network of such high  block walls in cold working
has never been classed as cDRX.
MODIFIED DEFORMATON MODEL; WARM DRV
In the model of Blum and McQueen for steady state creep of many fcc metals and hot working of
Al, the subgrains remain equiaxed and constant in size w, wall dislocation density (1/s) and internal
dislocation density i -0.5 being characteristic of the T and  conditions and providing for the steady
state stress [15-18,29,30]. In the composite model, dislocations pass at equal  through the soft
subgrains retarded by a small back stress and through the hard subboundaries aided by a high
forward stress. The substructure is developed during the transient strain  s and is completely
reconstituted during  =  s by migration, annihilation, unknitting and reforming of dislocation
walls. Disorientation boundaries develop during the initial straining between deformation bands
that are following alternative Taylor slip systems and rotate in different manners [15-18]. These DB
include geometrically necessary dislocations, become permanent unlike incidental SGB and
increase in  as straining proceeds. In hot working the frequent complete rearrangement of the
substructure prevents the development of additional high  walls within the deformation bands as
occurs in cold working due to microbands and block walls developing from new slip systems
[18,62-64]; ultimately DBs become equivalent to GBs having little effect on  at high T [1618,62,63].
The critical strain  cdSRX for static recrystallization dSRX is lower than  s for steady state so that
nucleation occurs after deformation stops if held at the working T; SRV causes high misorientation
features to develop into nuclei [10,16,36]. Behavior of SRX after hot working supports the model
above; the rate and grain refinement increase with  only up to steady state, thus indicating no
increase in density of high  features [8,9,12,16]. DRX is inhibited by the DRV being more
uniform than the SRV or by dislocations accumulating in any potential nucleus, thus reducing the
energy differential which would drive growth [10,12,36]. In polycrystals, high  cells develop
preferentially near GB, leading to more intense nucleation there. In single crystals, the DB are the
only source of nucleation of new grains in SRX after hot working [40,41] and of DRX when GB
mobility is unusually high as in 99.999 Al. In Al polycrystals of standard purity, GB develop
serrations on the scale of the subgrains but no new grains develop during straining; in Al-5Mg,
billowing serrations may pinch off at their base (rDRX) forming detached crystallites that do not
grow because of high internal dislocation density and that hasten the pinching-off process
[24,25,28,29]. The serrations not only cause gDRX through pinching off at high  of the long thin
grains but at low  , effectively hasten the migration of triple junctions by closing the acute angle
created by plastic flow; elimination of sharp grain ends provides average thickening of the grains
(Figure 1) [ ]. Moreover, during the rearrangement of serrations as GB lengthen in association with
SGB repolygonization [23], there may be net motion of GB into grains with above average
dislocation density. This feature of the gDRX mechanism appears to contribute to the development
of intense textures with low Taylor factor at high torsional strains [65,66] in Al [19] and in ferrite
[67].
Warm working is the domain between hot (> 0.6 Tm, 300°C Al) and cold working (< 0.35 Tm,
100°C Al) and its characteristics have recently been analysed [16,17]. The mechanical behavior
changes smoothly across the range; steady state can be attained at lower T for progressively higher
strain and stress with the level of polygonization decreasing and the spacings w,  -0.5 and s
decreasing as  /G rises [15,18,29,30]. Thus one can expect that as T decreases, there will be new
features related to the dense dislocation walls, blockwalls and microbands which are observed for
 >1 at 20°C. Thus dislocation walls would reach  >10° only in short segments on descending
from the hot working range and in greater lengths as T moves towards the cold domain. The
presence of such walls to the extent of 10 or 20% should not be classed as continuous cDRX
[58,59] as indeed their presence to a greater extent in cold working is not so classed. Instead of
naming the phenomena after a mechanism to which it attains in such a low degree that the
characteristics are barely similar, it would be better to specify it as a different form of the ideal, hot
hDRV namely warm wDRV. The low level of polygonization attained in cold working at high
strains, nominally Stage III or IV in   plots (   d / d) , could be termed cold cDRV [16-18].
Similar structural modifications sometimes called cDRX [68] have been observed in the warm
working of ferrite [68-70], eg. walls of about 13° forming blocks enclosing low  cells in Fe-19/23
Cr at 1050°C,  = 2.7 [69]. In the ferrite of a duplex stainless steel (Fe-21Cr-10Ni-3Mo, 1200°C,
0.7 s-1,   1.3), subgrains with mean  = 4.8° had 10% of the walls evenly distributed between 10
and 20°; this was classed as an altered form of DRV extended to less perfection at lower T [70]. In
line with that reasoning, wDRV is proposed as a mechanism with reduced repolygonization and a
modest proportion of high  walls yet still maintaining equiaxed subgrains of constant size. All
the significant mechanisms have been summarized in Table. 1.
CONCLUSIONS
During deformation above 0.5 Tm in Al, and even in Al-5Mg, dynamic recovery leads to a steady
state regime during which equiaxed cells maintain constant spacing of walls and also of wall and
interior dislocation densities inversely related to stress. Because of this high level of DRV, DRX
takes place only under special circumstances. At purities above 99.999% grain boundary mobility is
so enhanced that discontinuous dDRX can take place. Even with particle stimulation, dDRX will
not take place in 99.99 Al but it will in Al-5Mg with 0.7 Mn. In Al alloys with suitable pinning
particles and after TMP to produce a dense substructure, continuous cDRX takes place at a higher
temperature and lower strain rate suitable for superplastic deformation. In Al and most alloys,
geometric DRX (essentially DRV) takes place at a strain which reduces grain thickness to about 3
times the subgrain size. After discounting transition boundaries between deformation bands, the
increased incidence of 5-10° segments in Al over the range 300-100°C can be explained by the
diminished level of warm DRV leading to the gradual transition to the cold-work block wall
structure and associated augmented strain hardening.
Acknowledgements: The author is indebted to Sophie Gourdet and F. Montheillet for provision of
research results before publication and long discussions about their theory of continuous DRX. He
is grateful to S. Gourdet for the graphs of SGB misorientations. He has benefited greatly from
extended collaboration with Wolfgang Blum, University of Erlangen-Nurenberg, that helped clarify
themes in this paper.
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TABLE 1 RESTORATION MECHANISMS IN Al (SUMMARY)
MECHANISMS
(STRAIN HARDENING)
(Dislocation Slip)
RECOVERY
RECRYSTALLIZATION
DISCONTINUOUS
Particle Stimulated
(> 1 m )
CONTINUOUS
Fine Particle (<0.2
m ) Stabilized
Non Stabilized**
Partial**
GEOMETRIC
DYNAMIC (During  )
Taylor Constraints, deformation
bands, transition boundaries
rising  with  , all T.
Texture Formation
DRV #1
 monotonically rises under applied
strain to become constant in
steady state defined by Z.
DRX
dDRX (Classical)
 rise to peak, soften to steady state only
in 99:999+Al at 400°, 280°, 20°C
In prior cold worked Al
In Al-7Mg
STATIC (After  )
SRV
Gradual decrease in  , falling
substructure stress field
declining rate
SRX
dSRX (Annealing)
After incubation, rapid soften to
completion, rates higher grains
finer as Tdef down  ,  up, not in
steady state
Texture: preferred nucleation,
PSN-dSRX
growth
Alloys with constituent particles
cSRX
Very slow almost no growth of
cells
PSN-dDRX
Only in Al-Mg + particles
cDRX
After low T TMP, during SPD in supral,
Al-l0 Mg,Al-Li
Warm work, some high  walls #2
Not cold work, many high  walls #3
gDRX
Pinching off of very thin elongated
serrated grains
Serration Detachment
rDRX
and Rotation
In Al-Mg alloys meandering serrations
pinch-off at base
GRAIN GROWTH
DGG (  accelerates)
Normal or secondary grain
In high T, low  creep
growth
#1 DRV (creep, hot working > 0.6 Tm) equiaxed subgrains (repolygonization) with constant spacing of
walls and of wall and interior dislocation
#2 DRV (warm, 0.5-0.4 Tm, 300-100°C) mainly as in #1 but with some high  segments
#3 DRV (cold working) small cells ~ 0.5  m, blocks (~ 5  m) with dense dislocation walls, rising  ,
no repolygonization during strain hardening transient
** This classification under debate in this paper