Influence of repeated shock loading on the substructure

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 macroscopi- cally 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

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 phenom- enon 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].

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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, quasi- static 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]. Substruc- turally, 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 accom- modating 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 deter- mining 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 AI-Zn- 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-Zn- Mg-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

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 inhomo- geneities" 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 alu- minium; 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

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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 concen- tric 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 shock- induced thermal recovery effects.

Samples were single shocked to 13 GPa for 1 kts pulse duration through the impact of a tung- sten 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, re- shocking to 16 GPa, and finally shock loading to 14 GPa, all at - 1 8 0 °C. The impacts in the

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aluminum produce a transient strain (calculated as 4/3 (In(V/Vo) ) where V and V 0 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 re- machined 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 pur- posely 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 7 5 % 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 200kV 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 repeated- shock 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 -180°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

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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 deter- mination 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 bot tom 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) accord- ing 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

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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 inter- face 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,

Fig. 6. Bright field electron micrograph showing details of dislocation arrays and dislocation debris in a repeatedly shock-loaded 16-16 GPa 99.99 AI sample.

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Fig. 7. Bright field electron micrograph showing isolated microbands mixed with remnant ceils in 16-16 GPa shock- loaded 99.99 AI.

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 vari- able 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

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 band- type 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

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Fig. 8. Bright field electron micrograph showing a high density of microbands adjacent to a grain boundary.

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 16-16-14 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 16-16-14 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 repeti- tive 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. 9. Bright field electron micrograph of a 99.99 AI sample following repeated-shock loading at 16, 16 and 14 GPa.

Fig. 10. Bright field electron micrograph of a 99.99 AI sample following the 16-16-14 GPa shock series showing the coexistence of planar dislocation arrays and microbands.

actuality is composed of a compressive loading regime occurring at a very high (about 105-108 s - l) strain rate (termed the shock rise), a time of reasonable stable stress, i.e. a hold regime

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(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 single- cycle 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 micro- structure 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 abso- lute 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 coarse- grained 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 un- ambiguously the differences in the previous

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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 micro- scope 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 shock- loaded samples a surprisingly high density of vacancy loops was observed. Measurements by Kressel and Brown [3] of point defects in shock- loaded 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 identifica- tion 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

o r

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

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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, gener- ating 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 shock- loaded 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. Conse- quently, 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 3mm) 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 large- grained 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 re- versal 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 experi- encing 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 conti- nuity 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.

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 double- wall structure may then lead to microbands as described in the model proposed by Huang and Gray [37]. The coexistence and apparent connec- tion 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 pancake- shaped 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 large- grained 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

33

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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Influence of repeated shock loading on the substructure