Materials Science and Engineering, A 145 ( 1991 ) 21-35
21
Influence of repeated shock loading on the substructure evolution of
99.99 wt.% aluminum
G. T. Gray III
Materials Science and Technology Division, Los Alarnos National Laboratory, Los Alamos, NM 87545 (U.S.A.)
J. C. Huang
Institute of Materials Science and Engineering, National Sun Yat-Sen Universio', Kaohsiung, Taiwan (China)
(Received September 17, 1990; in revised form December 17, 1990)
Abstract
An electron microscopy study has been conducted of the substructure evolution of large-grained
(3 mm) polycrystalline 99.99 wt.% A1 subjected to single- and repeated-shock loading excursions at
-180 °C. Single-shock loading of aluminum is seen to form dislocation cells with a high density of
vacancy loops. The substructure evolution in aluminum with repeated-shock loading is observed to be
progressive in nature and similar to the dislocation arrangements in f.c.c, single crystals and
polycrystals with increasing strain. The substructure evolution, from dislocation cells to planar slip
bands to microbands, is found to be particularly evident adjacent to grain boundaries. The substructure
evolution in high purity aluminum subjected to repeated-shock loading at low temperature is discussed
in terms of the deformation mechanisms, in particular vacancy loop and microband formation, and
compared with previous studies on shock-deformed aluminum.
1. Introduction
The propagation of shock waves through
metals and alloys is well known to induce structure changes such as the formation of dislocations, deformation twins, vacancies or phase
transformation products [1-7]. The deformation
substructures resulting at moderate shock pressures (0-40 GPa) from these defects are generally
observed to be very uniformly distributed on a
grain-to-grain scale. The specific local type of
shock substructure developed in a given metal or
alloy, i.e. dislocation tangles or cells, deformation
twins, stacking faults etc., has been shown to
depend on a number of material factors including
the crystal structure, the relevant strengthening
mechanisms (such as alloying additions, grain
size, second phases, interstitial content etc.),
temperature, stacking fault energy (SFE), and the
imposed shock-loading parameters [1-7]. The
overall observed substructure, while macroscopically uniform, within individual grains may vary
from homogeneously distributed dislocation
tangles or cells to coarse planar slip, microbands,
stacking faults, or twins which reflect locally
heterogeneous deformation. The particular
0921-51)93/91/$3.50
subs(ructure
observed
therefore
critically
depends on the operative deformation mechanisms in the specific material under the specific
shock conditions imposed.
These shock-induced microstructurai changes
have in turn been correlated with variations in the
post-shock mechanical properties, with increasing peak shock pressure in particular, leading to
increases in both the hardness and reload yield
strength [1-7]. Shock loading in some materials
produces greater hardening than quasi-static
deformation to the same total strain, particularly
if the metal undergoes a polymorphic phase
transition, such as pure iron [l, 2]. This phenomenon has been generally attributed to the very
high strain rates associated with shock loading
and the subsonic restriction on dislocation velocity requiring the generation and storage of a
larger dislocation density than for quasi-static
processes [1-3]. In other metals, however, such as
606 l-T6 aluminum or Ti-6A1-4V, shock deformation is observed to cause minimum strengthening with the post-shock yield strength equivalent
to the quasi-statically deformed material yield at
the equivalent strain [8, 9].
© Elsevier Sequoia/Printed in The Netherlands
22
The majority of shock recovery studies have
concentrated on studying the response of
annealed or stress-relieved metals and alloys. Few
studies have probed the shock response of
materials possessing a pre-existing dislocation
substructure formed via heat treatment, quasistatic deformation (such as cold rolling) or prior
shock loading (shock prestraining) [1, 10-16].
The question revolves around whether significant
numbers of additional dislocations can be generated and stored by shock loading in a material
already work hardened by prior cold work and/
or shock prestraining or whether shock-induced
softening will instead occur.
Shock loading experiments on Armco iron
cold rolled to reductions of 20% and 40%
revealed that the cold work did not have an effect
on the post-shock hardness compared with
annealed samples, although it did appreciably
increase the original hardness [1, 10]. However,
prestraining in the ingot iron was observed to
decrease the velocity of the elastic I wave while
increasing its range to 0.6 GPa [10]. Substructurally, prestraining iron by cold rolling [1, 10] or
quasi-static compression [11] was found to inhibit
deformation twinning during subsequent shock
loading. The suppression of deformation twinning was postulated to result from the reduced
incidence of twin nucleation because fewer twins
are necessary to complement slip in accommodating the imposed strain rate in the iron
grains [l l]. This concept is consistent with the
crystallographic observations of Zukas and
Fowler [10] who observed that the degree of twin
suppression was influenced by whether the shock
direction was parallel or perpendicular to the
prior cold rolling direction [10]. The number of
grains exhibiting twin markings was highest when
the shock direction was perpendicular to the
rolling direction; lowest where the shock direction was parallel to the rolling direction [10].
Finally, a study on quenched-hardness high
carbon steels also showed that shock loading
could either soften or harden a previously
hardened structure [1].
In addition to quasi-static prestraining effects,
several studies have investigated the effect of
repeated-shock loading on the substructure
evolution and mechanical response [12-16].
Multiple-shock loading of 304 stainless steel
increased the amount of martensite formed, the
volume fraction of deformation twins and the
dislocation density observed [12-14]. Successive
shock loading of nickel was seen to produce slight
reductions in the dislocation cell size, increase the
dislocation density in the cell walls and increase
the nickel cell wall widths [13]. In both the nickel
and 304 stainless steel repeated-shock samples,
the reload yield strength and hardness were
observed to saturate with repeated-shock events
[13].
Recent studies on repeated-shock loading
effects have investigated the structure-property
response and thermal stability [15, 16]. Bulk
samples were chosen for these studies because of
recent results revealing differences in foil vs. bulk
sample substructure response due to the effects
of surface contact stresses and the minimum
sample thickness necessary for a stationary shock
front in the foil samples [16, 17]. In addition to
utilizing bulk samples, these recent studies have
carefully documented the residual deformation
strain after shock loading, being no more than 2%
after one or two loadings, 2-5% after five loadings, and 5-7% after ten loadings [15, 16].
Control of the shock loading parameters determining sample residual plastic strain is
considered to be crucially important because of
recent studies revealing the significant influence
of residual strain on post-shock structure-property material behavior [18-21]. Experimental
conditions need to be controlled to achieve a low
residual strain in shock-loaded samples so that
the substructure obtained can be prescribed to
the shock process itself rather than from recovery
processes [21].
Shock studies on A1-2%Mg, AI-6%Mg, and
on a "compound-alloyed high strength A I - Z n Mg-Cu material" showed that the shock hardening response and thermal stability in these
materials was alloy specific [15]. The reload yield
strength of the shock-loaded A1-2%Mg alloy
after a single shock event to 8.5 GPa showed a
strength increase of 16%; however, after annealing for 100 h at room temperature the alloy
properties had reverted to the preshock values
[15]. Successive loading of an AI-6%Mg alloy at
6.5 GPa from one to five times resulted in approximately a 40% increase in the reload yield
strength after the first shock which did not
increase further with subsequent loading and
was thermally stable at 50 °C [15]. The "AI-ZnMg-Cu alloy" displayed no shock strengthening
upon single-event shock loading to 8 GPa [15].
Bulk samples of copper, stainless steel and
AMI-6-aluminum were also repeatedly shock
23
loaded to investigate the mechanical response,
thermal stability and substructure response of
these materials [16]. The reload yield strengths of
the three materials were observed to saturate
after two to five successive shock events, depending on the material [16]. The substructure evolution in the repeatedly shock-loaded materials
revealed the appearance of '~different inhomogeneities" with each additional loading cycle [16].
In the case of copper, this additive evolution
produced a more pronounced dislocation cell
substructure with additional loadings and the
observation of fine deformation twins after five
successive loadings at 9 GPa which were not
present following a single 9 GPa shock [16].
Annealing studies of the multiply loaded samples
showed that at elevated temperatures the strength
characteristics of the specimens repeatedly
loaded by weak shock pulses are higher than
those of specimens strengthened by a single more
powerful pulse [ 16].
Recently, as part of an investigation of deformation twinning in aluminum-base alloys bulk
samples of 99.99% AI were repeatedly shock
loaded at - 180 °C to ascertain whether successive shock loading cycles would twin pure aluminium; they did not [22]. The deformation
substructure instead was cursorily characterized
to consist of dislocation cells containing a high
density of dislocation loops [22]. Quantitative
characterization of the dislocation loop type and
analysis of the substructure evolution due to
repeated loading in the previous study were not
performed. General observations of the substructure in shock-loaded aluminum in previous
studies have not defined a self-consistent pattern
[22-29]. In several studies dislocation tangles [23,
24] have been observed while in other studies
dislocation cells [25-29] have been seen. The
absence of a dislocation cell structure in some of
the experimental studies on shock-loaded aluminum, contrary to the observation of a cell
substructure following quasi-static deformation
in aluminum, is inconsistent with the high SFE of
aluminum as well as a high homologous temperature during room temperature shock experiments
which should enhance both cross-slip and
thereby cell formation.
The purpose of this paper is to report results
on the substructure evolution and deformation
modes in repeatedly shock-loaded bulk high
purity polycrystalline aluminum samples, including quantitative dislocation loop analysis. These
results will be discussed in light of previous
experiments on shock-loaded aluminum and the
documented effects of repeated shock loading
events on material substructure evolution.
2. Experimental procedures
The material studied in this investigation was
99.99 wt.% AI (hereafter referred to as 99.99 AI)
with an analyzed chemical composition (wt.%) of
15 ppm Mg, 7 ppm Cu, 10 ppm Fe, 10 ppm Si,
5 ppm Zn, and balance aluminum. The 99.99 A1
was sectioned from as-received cast material with
a grain size of about 3 mm. Shock recovery
experiments were performed using a 40 mm
single-stage gas gun. The specimen shock assembly consisted of a sample 4.76 mm thick and 12
mm in diameter fitting tightly into a similarly
sized bored recess in the inner momentum ring or
spall plate (25.4mm diameter). This central
cylinder was in turn surrounded by two concentric momentum trapping rings with outside
diameters of 31.7 and 44.5 mm. The surface was
protected from impact and the entire sample
from spallation by a close-fitting cover plate (2.54
mm) and spall plate support (12 mm), respectively. All assembly components were made of
aluminum to ensure impedance matching during
shock loading. The sample assembly was placed
in a steel impact cylinder that allowed the passage
of the sample or inner momentum ring through a
central hole but stopped the projectile. Samples
were soft recovered and simultaneously cooled
by decelerating the sample or inner momentum
ring in a water catch chamber positioned behind
the impact area. Shock loading was done at
approximately liquid nitrogen temperature; the
sample was routinely cooled at - 1 8 0 ° C , as
described previously [22]. The shock experiments
were conducted at - 1 8 0 °C to minimize shockinduced thermal recovery effects.
Samples were single shocked to 13 GPa for 1
kts pulse duration through the impact of a tungsten flyer plate 2.3 mm thick fixed to a projectile
filled with low impedance glass microballoons
and traveling at 915 m s i. Shock loading to 16
GPa was accomplished using a molybdenum flyer
plate 3.0 mm thick traveling at 1.2 km s-~. The
influence of sequential shock loading on substructure evolution in 99.99 AI was studied by
shock loading a 99.99 AI sample to 16 GPa, reshocking to 16 GPa, and finally shock loading to
14 GPa, all at - 1 8 0 °C. The impacts in the
24
aluminum produce a transient strain (calculated
as 4/3 (In(V/Vo) ) where V and V0 are the initial
and final volumes) for the 13 GPa, 14 GPa and
16 GPa shock levels of 16.7%, 17.6% and 19.5%,
respectively. The sample orientation with the
shock front for all three successive shock events
was kept constant. No precautions were taken to
reduce room temperature annealing between
shocks because of the experimental difficulties
involved. The samples were stored and remachined at room temperature between shocks,
with care taken, i.e. sufficient water cooling
during machining, to minimize sample heating
during machining. The residual plastic strain in
the samples (defined here as the change in sample
thickness divided by the starting sample thickness) was measured to be about 5% following the
combined 16 GPa-16 GPa-14 Gpa repeatedly
shocked series.
Specimens for optical and transmission
electron microscopy (TEM) were sectioned,
through thickness, from the bulk shock-recovered samples. The TEM specimens were purposely not sectioned in plane to the sample
surfaces because of the known influence of
contact shear stresses on the dislocation substructure directly adjoining the surface [19]. TEM foils
were prepared in a solution of 33% nitric acid
and 75% methanol at - 4 0 °C with 20 V using a
Struers electropolisher. Observation of the foils
was performed at room temperature or at
- 1 8 0 ° C on a JEOL 2000 EX at 2 0 0 k V
equipped with both conventional and liquid
nitrogen double-tilt stages.
3. Experimental results
Investigation of the substructure evolution in
99.99 A1 following single-event and repeatedshock loading at - 1 8 0 ° C revealed that the
deformation substructure observed varied with
successive loading events. The substructure of
99.99 AI shock loaded once to 13 GPa at
- 180 °C was observed to consist of a dislocation
cell structure, typically 1/~m in diameter. Figure
1 shows an example of the overall cellular substructure and associated selected area diffraction
pattern (SADP), possessing dense cell walls and a
high density of dislocation loops within the cells.
More loops were generally observed in the
central portions of the cells with a reduced
number directly adjacent to the cell walls. The
substructure observed in the TEM, particularly
Fig. 1. Bright field electron micrograph and SADP (inset) of
99.99 AI shock loaded to 13 GPa at -180 °C showing a
cellular dislocation substructure.
the dislocation loops, was found to be very
metastable under the action of beam heating
requiring use of the liquid-nitrogen specimen
holder. Higher magnification imaging of the loops
using a [111] zone axis, as shown in Fig. 2, reveals
the dislocation loops are hexagonal shaped with
their edges lying along the (110) direction. The
loop size (diameter) ranged from 5 to 60 nm with
an average of about 25 nm. The approximate
loop density, based on calculations from TEM
images, is 2 × 1021 m -3.
As a comparison, TEM characterization of
aluminum explosively shock loaded to 65 GPa at
liquid nitrogen temperature in 1963 by Zukas
[30], stored in the interim at room temperature,
was recently conducted. Figures 3(a) and 3(b)
show a dislocation cell substructure and high
density of dislocation loops similar to the current
results. The similarity of the substructure
morphologies in the two single-event shocks (13
and 65 GPa at about - 1 8 0 ° C ) suggests that a
cellular structure with a high dislocation loop
density is characteristic of aluminum shock
loaded over a large pressure range at low temperatures and appears to be reasonably stable to
room temperature recovery after the shock
process.
The character of the prismatic loops observed
in the current study was analyzed to determine
25
Fig. 2. Bright field electron micrograph of 99.99 AI shock
loaded to 13 GPa at - 18(I °C showing details of dislocation
loops (I110] zone axis).
whether they were vacancy or interstitial type
using the well-documented inside-outside contrast technique [31, 32]. In brief, the loop character was determined as follows [32]:(1) determine
the sample tilting sense and 180 ° inversion of the
TEM image at the magnification; (2) choose a
region and tilt the sample to a low index zone
axis, e.g. [110];(3) determine the loop inclination
sense for the larger loops through tilting, the
positive direction around the loops defined as
clockwise looking down from the electron gun
and b set using the criterion given by Edington
[32]; (4) achieve a diffraction condition of a
strong g and positive s; (5) use the loop determination criteria listed by Edington [32], i.e. if
(g.b)s<O for inside contrast and (g'b)s>O for
outside contrast then the loop is of a vacancy
type; (6) check the result by reversing the g (s still
positive) and see if it is consistent.
Figures 4(a) and 4(b) show an inside-outside
contrast tilting experiment, where z=[011] and
g--- _+[ i 1 ] ] showing the nature of loops tested in
this area (identified by the arrows) to be of
vacancy type. The relative inclination of the loop
b, relative to the top and bottom of the foil, and g
used for this area are as shown. Only loops larger
than 20 nm were analyzed because: (1) it was
much easier to determine the loop inclination
Fig. 3. (a) Cellular substructure of aluminum shock loaded
to 65 GPa at liquid nitrogen temperature [26]. (b) Detailed
view of dislocation loops in aluminum shock loaded to
65 GPa.
sense when the loops were larger, and (2) according to Edington [32] the loop image for a loop
smaller than 10 nm is a function of the exact
position of the loop location in the T E M foil, i.e.
whether the loop is located near the top, bottom,
or center of the foil. Analysis of 24 separate
dislocation loops located in several different cells
26
Fig. 5. Bright field electron micrograph and SADP (inset) of
99.99 A1 repeatedly shock loaded (16-16 GPa) showing
planar dislocation arrays and remnant cells.
Fig. 4. Bright field micrograph pair showing vacancy dislocation loops determined by inside-outside contrast [32]: (a)
(g.b)s<0, inside contrast; (b)(g.b)s>0, outside contrast.
The loop inclination and g directions are as labeled. (Small
filled arrows point to loops with the inclination shown while
the open arrow identifies a loop with an opposite inclination.)
was conducted. This analysis showed that all 24
loops were vacancy-type dislocation loops.
The cellular substructure with the surprisingly
high density of vacancy loops in the 13 GPa
single-shock event sample was also seen to affect
the post-shock reload compression yield strength
of the 99.99 A1. Assuming an average loop size of
10- s m, an increment in the critical resolved flow
stress of 30 MPa or approximately 80-100 MPa
in compression yield strength (utilizing a Taylor
factor of 3) may be attributed to the high loop
density using the formulation of Kroupa [33].
These values are not unreasonable compared
with the 150 MPa increase in the compression
yield strength above the 30 MPa yield strength of
the annealed starting material to 180 MPa following the single-event 13 GPa shock. The remaining
strengthening is thought to be due to the dislocation cells and tangles.
The substructure of the 99.99 AI subjected to
repeated-shock events was observed to become
more locally heterogeneous with each successive
shock. Figures 5 and 6 show the substructure of
99.99 A1 shocked twice (hereafter referred to as a
16-16 GPa shock). The substructure after the
initial single-event 16 GPa shock was virtually
identical to that after the 13 GPa shock described
above. The substructure of the 16-16 GPa
shock-loaded sample is observed to consist of
"coarse" [34] planar dislocation arrays (slip
bands) lying near or along {111 } planes, remnants
of dislocation cells containing a limited number
of dislocation loops, and microbands (Fig. 5). In
the repeatedly shocked samples some of the dislocation ceils were also observed to exhibit interface dislocation structures and increased
misorientation, as seen in the streaking in the
SADP of Fig. 5, suggesting local reorganization
into subgrains in response to the increasing local
strain. The overall breakup of the cellular structure and appearance of the coarse slip bands,
27
Fig. 6. Bright field electron micrograph showingdetails of
dislocation arrays and dislocation debris in a repeatedly
shock-loaded 16-16 GPa 99.99 AI sample.
microbands and subgrains were most evident in
regions adjacent to grain boundaries. The dense
dislocation arrays or slip bands appear as a group
of closely spaced slip planes which exhibit variable local misorientation with the matrix from
zero to several degrees depending on the individual band. Several coarse slip bands are seen to
subdivide portions of the substructure into a
checker-board-like array containing dislocation
debris and occasional dislocation loops (Fig. 6).
The dislocation motion associated with the
second 16 GPa shock in the 16-16 GPa sample is
observed to have annihilated the high density of
dislocation loops, throughout the grains, present
following the single 16 GPa shock, leaving a few
isolated loops.
In addition to the planar dislocation arrays or
slip bands, numerous microbands were seen in
the 16-16 GPa sample. In this paper we follow
the definition of microbands as substructural
features consisting of double dislocation walls,
approximately 0.1-0.4/~m (mostly 0.2 ~m) apart
when viewed edge on, lying within a few degrees
along crystallographic slip planes and enclosing a
volume possessing 1°-3 ° misorientation with
respect to the surrounding matrix [35-37]. The
morphology and density of microbands were
Fig. 7. Bright field electron micrograph showing isolated
microbands mixed with remnant ceils in 16-16 GPa shockloaded 99.99 AI.
observed to vary between the grain interiors and
directly adjacent to the grain boundaries and/or
grain boundary triple points. Figure 7 shows a
moderate density of well-formed microbands
interspersed with remnant dislocation cells and
dislocation debris from a grain interior area in the
double-shocked 16-16 GPa sample. The propensity for microband formation and the relative
local density of microbands adjacent to grain
boundaries were seen to be considerably higher
than observed in the grain interiors. Figure 8
shows a region exhibiting a higher density of
microbands adjacent to a grain boundary area.
Some of the microbands observed in the repeatedly shocked 16-16 GPa sample additionally
appeared to have a close relationship with the
coarse planar slip bands, with both of these bandtype structures lying on active slip planes.
The substructural stability of the 16-16 GPa
sample to subsequent shock loading was probed
by additionally shock loading at 14 GPa; this
sample was termed 16-16-14 GPa. The substructure of the 16-16-14 G Pa aluminum sample
was similar to that of the 16-16 GPa sample
although the number of microbands, both in the
grain interiors and in the grain boundary areas,
was observed to increase and the number of
28
Fig. 8. Bright field electron micrograph showing a high
density of microbandsadjacent to a grain boundary.
Fig. 9. Brightfield electron micrographof a 99.99 AI sample
followingrepeated-shock loading at 16, 16 and 14 GPa.
visible dislocation loops was reduced to almost
zero. Overall, the substructures seen in the grain
interiors and adjacent to the grain boundaries
were more similar for the 1 6 - 1 6 - 1 4 GPa sample.
The "checker-board" subdivision of the substructure by the prior planar dislocation arrays was
noted to coexist with microbands; in some cases
the microbands appeared to continue directly
from a planar slip band (Fig. 9). In Fig. 10 several
small microband segments are seen to be directly
associated with portions of planar dislocation
lines. The concentration of microbands was
observed to be particularly high directly adjacent
to grain boundaries as seen in the 16-16 GPa
sample. Finally, the regions between the microbands in the 1 6 - 1 6 - 1 4 GPa sample contained a
reduced amount of dislocation debris suggesting
considerable sweep of the dislocations into the
microband and cell boundaries.
4. Discussion
The substructure developed in 99.99 A1 subjected to multiple-shock loading excursions
provides a detailed view of the effects of repetitive high-strain-rate deformation on the substructure evolution in a high SFE f.c.c, metal. Due to
the intrinsic nature of the shock-loading process
the substructure evolution in a shock excursion in
Fig. 10. Bright field electron micrograph of a 99.99 AI
sample followingthe 16-16-14 GPa shock series showing
the coexistenceof planar dislocationarrays and microbands.
actuality is composed of a
regime occurring at a very
s -l) strain rate (termed the
reasonable stable stress,
compressive loading
high (about 105-108
shock rise), a time of
i.e. a hold regime
29
(termed the pulse duration), and finally a tensile
release of the applied compressive load returning
the sample to ambient pressure at a lower strain
rate (one to three orders of magnitude lower than
the rise, termed the rarefaction). Collectively the
loading sequence in a shock amounts to a singlecycle loading excursion with stress-strain change
and a strain-rate jump with elastic and plastic
deformation operative in two directions at two
different strain rates. In this regard the shock
process may be compared with a single high
amplitude "fatigue-type" cycle with a dwell time
representing the pulse duration [38]. Overall,
investigation of the deformation substructure
evolution during post-shock microscopic examination of shock-recovered samples must therefore consider the total path history of
stress-strain and strain-rate change comprising
the shock.
The substructure evolution in the repeatedly
shock-loaded 99.99 A1 is a reflection of the progression of the deformation mechanisms and
resultant morphologies in the framework of the
shock cycle compared with the influence of
amount of strain, temperature and strain rate on
the dislocation arrangements in conventionally
deformed f.c.c, metals. The discussion of the
substructure evolution in aluminum subjected to
single- and repeated-impact loading will be
addressed separately.
4.1. Single shock
Although the permanent plastic strain in the
repeatedly '~soft" shock-recovered 99.99 AI
samples is quite small, the deformation microstructure in the shock-loaded aluminum due to
the collective shock strain exhibits a sequence of
substructure development consistent with many
of the quasi-static observations as a function of
strain and temperature. The substructure evolution in quasi-statically deformed single crystals of
copper as a function of strain is well known to
exhibit three stages with increasing strain up to
fracture [39]. The effect of increasing strain on
the dislocation cell substructure of aluminum has
shown that the cell size is diminished sharply at
first and then reaches a limiting value; the absolute value of the limiting cell size is dependent on
the type of deformation (tensile elongation or
cold rolling) employed [39]. Studies of the substructure evolution in high purity aluminum as a
function of temperature have found that the
dislocation density for the same strain is
increased as the deformation temperature is
decreased, resulting in a reduced cell size; this
indicates that dislocations move shorter distances
at lower temperatures requiring the generation of
a higher density of dislocations to accommodate
the applied stress [39].
The substructure ew)lution in shock-loaded
aluminum has not defined a consistent pattern
[23-30]. Examination of the experimental details
in the previous studies observing cells compared
with those exhibiting dislocation tangles or a
"polygonal" structure [27, 29] suggests that differences in the starting grain size, sample configuration and temperature, the shock-loading
technique utilized (including the residual strain)
and the peak pressure may partially account for
the differences.
The initial study of Rose and Berger [23]
observed a substructure of dislocation tangles in
shock-loaded stress-relieved aluminum foils
0.0254 mm (0.001 in) thick with an unspecified
starting grain size. A direct comparison of their
results with the current study is complicated by
the known differences between bulk and foil
sample response to shock loading and the potential for high residual strains in foil samples [16,
17]. Late-time radial loading effects can lead to
significant reorganization of the substructure of
shock-loaded metals [21[. The question of starting grain size may also influence the substructure
developed in aluminum. Tensile studies on
99.995% AI deformed to 15% strain at room
temperature found that in coarse-grained aluminum (290 j~m) a well-defined cell structure was
formed whereas in fine-grained (30 ~m)
aluminum a tangled substructure was developed
[40]. This observation suggests that in the coarsegrained aluminum, upon removal of the applied
stress, recovery processes are more operative
leading to reorganization of the tangled dislocation substructure into a low energy cellular dislocation configuration [41 ]. This observation is in
agreement with the observation in one study of a
cellular-type dislocation substructure in aluminum with a starting grain size of 0.9 mm [25]
shock loaded at room temperature.
In two other room temperature shock-loading
studies tangled substructures were seen in aluminum, with a starting grain size of 1100 /~m,
shocked to pressures up to 8 GPa at two pulse
durations [24] and in 99.5 wt.% AI shocked to 5.8
GPa [28]. While it is impossible to resolve unambiguously the differences in the previous
30
studies, it is postulated that due to the approximately doubling of peak pressure in the current
study, and corresponding increase in the dynamic
strain, a substantially higher density of point
defects and dislocation debris were generated
compared with the previous studies. This higher
defect density may perhaps also be influenced by
the low temperature of the shocks [39]. Upon
release of the shock and subsequent warming of
the samples to room temperature the local debris
substructure then recovered. The pronounced
metastability of the high density dislocation loop
and dislocation structure under the action of
beam heating in the transmission electron microscope supports the contention that the substructure of the large-grained aluminum shocked at
- 180 °C in this study has undergone some local
stress relaxation, particularly in the presence of a
high density of point defects, forming the
observed dislocation loops and cells.
Within the cell interiors of the cellular dislocation substructure in the 13 or 16 GPa shockloaded samples a surprisingly high density of
vacancy loops was observed. Measurements by
Kressel and Brown [3] of point defects in shockloaded nickel samples revealed that five-sixths of
the point defects were vacancies and one-sixth
were interstitials. Murr and Ross [42] and Murr et
al. [43] also determined that the shock-induced
dislocation loops in 304 stainless steel and
molybdenum were predominantly vacancy-type
loops. Both previous studies indicated that at
least 80% of the shock-induced point defects
were vacancies. Assuming a similar percentage of
vacancies produced in the present study, it
appears likely that most of the dislocation loops
are vacancy type. The interstitial concentration of
about 20% or less is considered too low to allow
for local clustering of sufficient interstitials to
produce a large loop, i.e. greater than 20 nm.
This hypothesis is consistent with the identification of only vacancy loops in the current study.
The point defect concentration generated
during the shock may be estimated using the
model developed by Meyers and Murr [44] who
applied the idea that point defects are formed via
a dislocation jog mechanism. The point defect
concentration C in their model is given by the
relation
C =Kp:L2/2nb
(1)
where K is a ratio equal to about 0.25; p is the
dislocation density, about 1015 m-2 in this study;
L is the dislocation travel distance, about 0.4/~m
in the shock-loaded aluminum samples (where
the dislocation cell size is about 1/~m and using
the statistical traveling distance which is 7tR/4
with R being 0.5 ~m in the current case); n is the
number of atoms per unit volume (four atoms per
unit cell), about 6 × 10 28 m-3; and b is Burger's
vector for perfect unfaulted prismatic loops
(b = a(110)/2), about 0.286 nm.
Using this data in eqn. (1) the value of C is
1.16 × 10-3. Based on the examination of TEM
micrographs, the observed loop density N is
2 × 1021 m -3. However, a sufficient number of
loops that are near the top and bottom surfaces of
the foil will have disappeared because of the
truncation by the foil surfaces or local vacancy
diffusion to surfaces or the cell walls. Therefore
we assume for calculation purposes that the
actual starting loop density N' is about 3 × 10 2~
m- 3. We will assume that half of the point defects
aggregate by diffusion to form loops (thought to
be a reasonable assumption, excluding the point
defects that diffuse to the dislocation cell walls or
remain in the matrix); therefore the effective
vacancy concentration for loop formation is C/2.
By equating the volume fraction of vacancy loops
calculated from both the loop measurement and
evaluation of the vacancy concentration, we
obtain the equation
N'Jrd2 b
Cnv i
4
2
or
2Cnvi I ~/2
a=IN'G j
in which the volume for one loop is assumed to be
jrd2b/4 and v~ is the volume occupied by one
vacancy (equal to 1.66 × 10-29 m 3, i.e. one-fourth
of the volume of the unit cell). The average loop
size d, calculated from the above equation, is
about 29 nm, which is in reasonable agreement
with the experimental value of about 25 nm. This
calculation supports the idea that a jog mechanism is most probably responsible for the high
density of loops in shock-loaded aluminum.
Our thinking regarding the jog mechanism is
presented as follows. As a result of extensive
dislocation interactions during shock loading, a
significant number of jogs (and also kinks, which
are not considered here because they can glide
31
without the involvement of non-conservative
movement) will be formed. Depending on the
sign of the two interacting dislocations, a jog can
be viewed either as being associated with an extra
half plane (as if it were a row of interstitials) or as
being associated with an empty half plane, as if it
was composed of a row of vacancies. Under the
conditions of quasi-static deformation, such jogs
may not be able to move, such that there are not
extensive point defects generated and hence there
will not be a high number of dislocation loops.
However, during shock loading it is much more
effective for a dislocation with a jog to move [44].
The jog theoretically can move in any direction so
that either the row of interstitials extends, generating more interstitials, or the row of vacancies
extends, thereby creating more vacancies. In
reality, a jog tends to travel in the sense that more
vacancies are formed since the formation energy
of an interstitial is typically a factor of 5 larger
than that of a vacancy in common metals [45].
Accordingly, a significantly greater number of
vacancies are usually formed when dislocations
are forced to move and interact. The difference in
formation energy for a vacancy and an interstitial
(about 5:1 ) is consistent with the observed difference in point defect concentration ratio (5:1) of
vacancies and interstitials observed in shockloaded nickel [3].
The observed prismatic dislocation loops are
thought to result from the point defects created
by the motion of jogs. Because of the very limited
time during shock deformation, point defect
diffusion cannot be very long range. Consequently, dislocation loops are formed via the
clustering of very locally distributed point
defects. In addition, because of the short time
available there should not be an appreciable
loop-free zone near the dislocation cell walls
which are considered to be a point-defect sink.
Only those defects in close proximity to the cell
walls will be capable of diffusing to the sinks
during the limited time of the shock. This is
consistent with the experimental observations in
this study. In addition, since vacancy diffusion is
generally much faster than interstitial diffusion,
even if a higher fraction of the generated loops
were interstitial, contrary to that believed, few
interstitial loops would probably be expected.
Taken collectively, the low density of interstitials
generated during shock deformation and their
low diffusion rate should result in the majority, if
not all, of the observed loops being vacancy type.
It has been shown in previous investigations
[46, 47] that unfaulted (prismatic, b = (a/2)(110))
loops are more stable than faulted (Frank, b =
(a/3){ll 1))loops. There is more energy associated with the faulted loops because of the extra
stacking fault energy involved, especially for
materials with a high SFE Y~t such as high purity
aluminum. Accordingly, we would not expect to
observe faulted loops in shock-loaded aluminum.
In addition, the loop density in shock-loaded
high purity aluminum is considerably higher than
that previously observed in other shock-loaded
f.c.c, metals. This is thought to result possibly
from two causes. First, the very high SFE y~ of
aluminum (166 mJ m -') [481 makes it easier to
have extensive dislocation interactions due to the
effectiveness of cross-slip. In metals with lower y~
(such as silver with a y~ of 16 mJ m e) [48], it is
more difficult to move dislocations with extended
jogs because of the possible production of large
stacking faults (details can be found in ref. 49).
Jogs are therefore more readily created and can
subsequently travel more easily in aluminum,
thereby more profusely producing point defects,
than in a lower y~. metal. Secondly, the high
density of dislocation loops observed in aluminum will be influenced by the high point defect diffusion rate in aluminum, which has a relatively
low melting temperature. A higher vacancy rate
in aluminum will greatly increase the propensity
for sizable loop formation. In metal systems
possessing higher melting temperatures and
hence a lower vacancy diffusion rate, coupled
with a lower y,t, extensive numbers of dislocation
loops are probably unlikely. The lack of extensive
numbers of dislocation loops in shock-loaded
copper and nickel (y~. of 45 mJ m e and 125 mJ
m e, respectively)[48] is consistent with this
argument.
4.2. Repeated shock
The substructural features observed in the
repeatedly shock-loaded aluminum provide a
view of the local substructural stability and evolution of the dislocation cells and high dislocation
loop density, formed during the initial low temperature shock excursion, to further plastic deformation. The substructures are seen to reflect both
a reorganization of the pre-existing shock micro,structure and the increasing total sample strain
with repeated loading. The 16-16 GPa and
16-16-14 GPa samples represent increasing
total transient strains of approximately 40°/,, and
32
57%, respectively. The application of the second
shock cycle causes the dislocations comprising
the cells formed during the first shock-loading
cycle to sweep out the majority of the vacancy
loops throughout the grains. This observation is
consistent with experiments on quenched aluminum possessing vacancy loops which when cold
rolled 5% swept out all the pre-existing loops
[50]. In addition to sweeping out the loops, in the
vicinity of the grain boundaries the second shock
is seen to concentrate the dislocations in the cell
walls into planar dislocation boundaries lying
along {111 } planes and into subgrains. The planar
dislocation arrays interspersed with the cells are
similar in appearance to the dislocation structures
seen in f.c.c, single crystals in stage III [39, 51]
and structures termed "deformation" or "matrix"
or "transition" bands or dense dislocation walls in
polycrystals [52-54]. Two sets of orthogonally
oriented "bands" in cold-rolled aluminum [54],
microbands in Ni-Co [55], microbands in nickel
[56], and dense dislocation walls or microbands
in lightly rolled aluminum [57, 58] have also
produced a substructure similar in general
appearance to the substructure shown in Fig. 8.
The similarity of the substructural changes in
the current aluminum polycrystals (grain size,
about 3 m m ) to single-crystal substructural
morphologies is consistent with several experimental observations and modeling predictions
[59, 60]. A study of the influence of grain size on
the substructure evolution in deformed copper
found that large grains (250 /~m) behaved like
single crystals ]59]. The formation of the planar
slip boundaries and then microbands initially
adjacent to the grain boundaries further reflects
this "single-crystal-like" behavior. For largegrained polycrystals the accommodation between
adjacent grains is primarily confined to the region
along the grain boundary where a larger number
of slip systems are activated in contrast to the
grain interior [52, 60]. The activation of significant secondary slip near the grain boundaries
to accommodate the intergrain compatibility
stresses is believed to cause the local breakdown
of the dislocation cells into planar dislocation
bands, microbands and subgrains. Some of the
local rearrangement of the cells into subgrains
with repeated loadings most likely also reflects
some "dynamic recovery" and strain-path reversal processing of the initial shock substructure by the second and third shock excursions
]61, 62].
Studies by Basinski and Jackson [63-65],
Basinski and Basinski [66] and Sharp and Makin
[67] have shown that planar slip bands are formed
in single crystals when a second single-slip deformation strain path follows a preliminary deformation path on another system. The nature of the
shock process, in the absence of significant radial
release effects [21], results in the sample experiencing a high-strain-rate uniaxial-strain path in
compression followed by tensile release. The
uniformly distributed cellular dislocation substructure observed after the single-shock excursion suggests that a low number of operative slip
systems may have been activated to respond to
this predominantly uniaxial-strain cycle and may
have initially dominated, especially in the grain
interiors. Upon repeated-shock loading the substructure evolves in a sequential manner with
increasing total strain. This process, however,
appears to saturate throughout the grains after
the second shock excursion as the total transient
strain is about 40%. The progressive nature of
repeated-shock cycles is most evident in the
substructural rearrangement adjacent to the grain
boundaries. This observation is consistent with
the general concept that the overall strain continuity in polycrystals is maintained primarily at the
grain boundaries and not simultaneously across
the whole grain [52]. An overall additive nature to
the shock-produced substructures and the transition to a more inhomogeneous substructure with
increasing total transient strain is consistent with
the observations of Gubareva et al. [16] on repeatedly shock-loaded copper.
As in the single-crystal case, the formation of
planar dislocation arrays along {111} planes is
thought to reflect an increasing local misorientation, such as seen in the SADP shown in Fig. 5,
between the cells with increasing local strain. In
addition, while small in magnitude (about 5% in
the 16-16-14 GPa sample) the increasing value
of the residual strain in the repeatedly shocked
samples adds a late-time radial tensile loading
component to the uniaxial-strain shock excursion.
Because the radial stresses produce a small
stress-path change, accommodation of these
stresses primarily near the grain boundaries may
also help with the breakdown of the cells by the
mechanism postulated by Sharp and Makin [67].
Application of the local stresses plus the
opposite-signed dislocations composing the cell
walls is postulated to reach a critical point where
a portion of a cell wall will locally break down.
33
Once disrupted, further dislocation activity will
localize within such a path forming a planar dislocation array. The development of more misorientation across the planar dislocation arrays is
then thought to drive locally microband formation. Further development of the slip bands by the
generation of polarized dislocations, annihilation
of the primary dislocations in the central band
region, and secondary slip between the doublewall structure may then lead to microbands as
described in the model proposed by Huang and
Gray [37]. The coexistence and apparent connection of the planar dislocation arrays and cellular
dislocation arrangements with the microbands as
seen in Fig. 10 suggest a potential substructural
development link between cells and microbands.
Observations of Barlow et al. [54] on cold-rolled
aluminum and Bay et al. on lightly rolled pure
aluminum [58] have also covered in detail the
coexistence between dense dislocation walls,
small elongated subgrains or small pancakeshaped cells, and "bands" or microbands which
appear similar to what we have described as
microbands.
5. Summary
An electron microscopy study has been conducted on the substructure evolution of largegrained polycrystalline 99.99 wt.% AI subjected
to single- and repeated-shock loading at
- 180 °C. The shock sequence is considered to
be composed of a single-cycle stress-strain path
jump loading excursion with elastic and plastic
deformation operative in compression followed
by tensile directions. Accordingly, the shock
process may be idealized as a single high amplitude "fatigue-type" cycle. Single-shock loading of
aluminum is seen to produce a substructure of
dislocation cells containing a high density of
dislocation loops. TEM analysis of the dislocation loops revealed that the loops were all
vacancy type. The high measured loop density
(about 2 x 1021 m -3) is believed to be consistent
with the inherent large number of point defects
generated by the motion of jogs during shock
deformation, the high ~'~f(166 mJ m -e) [48] of
aluminum, and the high point defect diffusion
rate in aluminum. The observation of dislocation
cells in the current study, compared with random
dislocation tangles in some previous studies, is
believed to be linked to variations in the starting
grain size, sample configuration and temperature,
the shock loading technique utilized (including
residual strain) and the peak shock pressure.
With repeated-shock loading, the substructure
evolution in pure aluminum is found to be
sequential in nature. This overall progression of
the substructure is consistent with several previous repeated-shock studies. The substructure
evolution in shock-loaded aluminum is observed
to progress from dislocation cells to planar slip
bands to microbands, both lying on {1 11 } planes,
and subgrains. This dislocation morphology
transition from cells to planar dislocation arrays
is similar to the dislocation arrangement progression in f.c.c, single crystals and polycrystals as a
function of strain. The breakdown of the cells,
development of concentrated cell walls along
primary slip traces, and the development of
increasing local misorientation in the planar slip
bands is postulated to be a potential link in the
progression between dislocation cells, subgrains
and microbands.
Acknowledgments
The authors wish to acknowledge the many
contributions of C. E. Frantz to the design of
these recovery experiments and B. Jacquez for his
assistance in operating the gun facility. This work
was performed under the auspices of the U.S.
Department of Energy.
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