Direct observation of the behavior of grain boundaries during continuous dynamic recrystallization in an Al–4Mg–0.3Sc alloy

Acta Materialia 51 (2003) 4367–4378 www.actamat-journals.com

Direct observation of the behavior of grain boundaries during continuous dynamic recrystallization in an Al–4Mg–0.3Sc alloy L.M. Dougherty a, I.M. Robertson a,∗, J.S. Vetrano b a

Department of Materials Science and Engineering, University of Illinois at Urbana-Champaign, 1304 W. Green St., Urbana, IL 61801, USA b Pacific Northwest National Laboratory, Richland, WA 99352, USA Received 25 February 2003; received in revised form 1 May 2003; accepted 10 May 2003

Abstract The micromechanisms operating during dynamic continuous recrystallization of an Al–4Mg–0.3Sc alloy have been studied in real time by reloading pre-deformed samples at temperature in situ in the transmission electron microscope. The processes responsible for the evolution of the microstructure include the migration and breakup of subgrain boundaries, the destruction of triple junctions, and large-scale volume rotations. These observations are considered in terms of the micromechanisms proposed previously to account for dynamic continuous recrystallization.  2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Electron microscopy; Aluminum alloy; Dislocation dynamics; Dynamic recrystallization

1. Introduction Lower processing temperatures and faster forming rates for superplastic processing can be achieved in some alloy classes by beginning with an unrecrystallized rather than a fine-grained microstructure. In such systems, dynamic continuous recrystallization [1]1 during the early stages pro∗ Corresponding author. Tel.: +1-217-333-6776; fax: +1217-333-2736. E-mail address: [email protected] (I.M. Robertson). 1 Continuous dynamic recrystallization is defined as the homogeneous evolution of the grain structure that involves the migration of high-angle grain boundaries, which is important in hot deformation and in superplasticity of some aluminum alloys. It should be noted that there is some debate surrounding

duces the fine-grained microstructure and permits use of higher strain rates for at least this portion of the process. The microstructural processes responsible for this evolution, however, are not well understood and those that have been proposed have been deduced from static snapshots of the microstructure obtained from interrupted superplastic deformation tests. The proposed mechanisms include: the use of the term continuous dynamic recrystallization to describe these processes, as it is unclear that activity of highangle grain boundaries occurs. If no motion of the high-angle boundaries occurs the process should be correctly identified as extended recovery. For detailed discussion on this topic see Ref. [1].

1359-6454/03/$30.00  2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/S1359-6454(03)00262-3

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앫 The migration and coalescence of subgrain boundaries to cause either annihilation or development of high-angle grain boundaries [2]. This is the same mechanism as proposed for continuous static recrystallization although, in this case, the microstructure develops faster or at lower temperature because of the additional driving force of the applied stress. 앫 The migration of subgrain boundaries only. The misorientation change is directly related to the migration distance because it relieves the lattice curvature introduced by rolling. Nes proposed that the rapid increase in grain boundary misorientation was attributable to the fine particles dissolving ahead of migrating grain boundaries; an effect termed strain-induced particle reversion [3]. The driving force for particle reversion was the coherency energy. The particles reform immediately the grain boundary passes, making the process difficult to verify by experiment. 앫 The migration, interaction, and incorporation of lattice dislocations into subgrain boundaries. The dislocation source is either the original grain boundary triple points [4] or the dissolution of the lowest angle grain boundaries [5]. 앫 The sliding of high-angle grain boundaries with the accommodation mechanism being the rotation of subgrain boundaries attached to them. As the degree of sliding increases, subgrains further from the grain boundary become involved in the accommodation process [6–8]. Hales and McNelley [9] extended this concept by suggesting that sliding was not restricted to high-angle grain boundaries but also occurred on low-angle grain boundaries and along original grain boundaries [10]. Continuity between the sliding subgrain boundaries was achieved by the generation, migration, and absorption of dislocations. In this paper, the first time-resolved observations of the fundamental processes operating during continuous dynamic recrystallization are reported. The processes were observed by reloading predeformed specimens of a superplastic Al–4Mg– 0.3Sc alloy in situ in the transmission electron microscope at nominally the superplastic forming temperature.

2. Experimental procedures The Al–4Mg–0.3Sc system was selected for this study as Mg remains in solid solution and only Al3Sc particles, which are stable against coarsening at moderate temperatures [11,12], are present to influence the behavior. The composition of the alloy was 95.37Al–4.35Mg–0.28Sc–0.02Si– 0.03Fe. It was cast in 203.2 × 254 × 50.8 mm book molds and the surfaces scalped after cooling. The castings were warm-rolled at 573 K to a thickness of 6.35 mm, solution treated for 4 h at 848 K, and then aged for 8 h at 553 K. This aging treatment corresponds to the peak-hardness condition [13]. Following the aging treatment, the Al3Sc particles were distributed homogeneously as spherical precipitates with a bimodal size distribution. Large particles, which existed in the as-received state, had diameters in the range from 125 to 225 nm and had a semi-coherent interface with the matrix. Small coherent particles with diameters from 12 to 25 nm were distributed homogeneously throughout the grains, and these inhibit static recrystallization [11]. The material was cold rolled to 0.19 mm (70% reduction) and dog-boned shaped specimens, with a 2.5 cm × 0.625 cm gauge and 2.5 cm × 1.875 cm grip, were produced and subsequently deformed at 733 K at a constant strain rate of 10⫺3 s⫺1. In the unrecrystallized condition, the alloy exhibited an engineering strain of 376% compared to 181% for a fine-grained microstructure. To follow the evolution of the microstructure during superplastic deformation, tests were interrupted at true strains of 0.1, 0.2, 0.4, and 0.8. The microstructure was retained by quenching the samples using a liquid nitrogen spray, which results in a cooling rate of 75°/s. Electron backscattering, orientation imaging microscopy, and post-mortem TEM analysis established that continuous dynamic recrystallization processes were complete by a true strain of 0.4 [13]. Therefore, samples for the in situ TEM deformation experiments were prepared from specimens deformed to a true strain of 0.2. The samples were cut from the gauge section parallel to the rolling direction as opposed to the transverse direction, ground to 150µm thickness, and the central section electropolished to electron transparency using a twinjet-pol-

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ishing unit with an electrolyte of 5% perchloric acid in methanol. The electrolyte temperature was held between 234 and 228 K and the polishing voltage was 15–25 V. The in situ deformation experiments were performed at nominally the superplastic deformation temperature in either a Hitachi 9000 or a JEOL 4000 TEM using Gatan high temperature straining stages. It should be noted that the stage design precludes control and monitoring of the sample temperature during straining, as heating depends on the degree of contact between the sample and the heater. Calibration of the stage for the JEOL 4000 TEM showed a temperature difference between the sample and the heater of upwards of 150°. The difference was, however, not constant and could vary during the course of a straining experiment. To ascertain that the combination of straining and heating was responsible for the observations presented in this paper some samples were heated without straining and others were strained at different temperatures. The heating experiments established the static recrystallization temperature for electron transparent samples. This was 733 K, which is 50° lower than in a bulk specimen. Straining at different temperatures confirmed that the observed behavior was temperature dependent and at temperatures lower than 683 K, 50° below the bulk superplastic temperature, rapid crack nucleation and propagation occurred. Thus, despite the lack of temperature control it was possible to distinguish experimentally the appropriate conditions for performing the experiments. Static and dynamic recrystallization processes were also kinetically very different. Static recrystallization occurred rapidly and dynamic recrystallization slowly, making them easy to distinguish. Results from the dynamic experiments were recorded on S-VHS videocassettes via a Gatan TV-rate camera system, giving a time resolution of 1/30th s.

3. Experimental results Although the emphasis of this paper is on the dynamic mechanisms operating during dynamic continuous recrystallization, the microstructural changes revealed by post-mortem analysis of inter-

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rupted tests are shown to establish the changes that must be explained by the in situ experiments. The electron micrographs in Figs. 1 and 2 show the evolution of the grain and dislocation structure as a function of strain to 0.4 true strain. After 1000 s at 733 K, the pre-deformation hold state, a mostly recovered microstructure with a small, slightly elongated subgrain structure was obtained, Fig. 1(a). The average subgrain size, as determined using ASTM standard E112-88, was 0.5 µm. Recovery was complete and the subgrains were equiaxed with an average size of 1.7 µm at 0.1 true strain. The subgrain size continued to increase with increasing strain to 0.2 but the size distribution became less uniform, Fig. 1(b). The smaller subgrains were equiaxed and the larger ones were elongated suggesting subgrain coalescence. After 0.4 true strain, the microstructure was dominated by high-angle grain boundaries, the grain structure was coarser, and subgrains and subgrain boundaries were rarely found, Fig. 1(c). These observations, which are consistent with electron backscattering studies, suggest dynamic continuous recrystallization was complete by 0.4 true strain. At the dislocation level, changes also occur with increasing strain, Fig. 2. The lattice dislocation and subgrain boundary order and density decrease with increasing strain. Dislocations were commonly anchored at large Al3Sc particles with the other end attached to a nearby grain boundary. The dislocation structure at the particle matrix interface increased in complexity with increasing strain. Determining the mechanisms by which these structures evolve is difficult from the snapshot micrographs shown in Figs. 1 and 2, but observations of this type have been the basis for current models. To observe the reactions, specimens deformed to 0.2 true strain were reloaded at nominally the superplastic forming temperature in situ in the transmission electron microscope. Montages showing the beginning and final microstructures following such an experiment are presented in Fig. 3. The changes include movement of arrays of dislocations, removal of subgrain boundaries, and dislocation emission from a triple junction. For ease of referencing, static features are indicated by a letter and dynamic changes by a number. The par-

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Fig. 1. Comparison of the grain size as a function of strain. (a) No strain, but held at 733 K for 1000 s, (b) 0.2 true strain, and (c) 0.4 true strain.

Fig. 2. Weak-beam dark-field micrographs comparing the dislocation structure in samples deformed to true strains of (a) 0, (b) 0.2, and (c) 0.4.

tial subgrain boundary 1 has broken free at one end, presumably from the junction between grain boundaries 3 and 5. With increasing load at temperature, the dislocations comprising subgrain boundary 1 move collectively in the direction indicated, breaking contact with grain boundary 5 and 2 in the process. This reaction is shown at higher magnification in Fig. 4; for point of reference the initial position of subgrain boundaries 1 and 2 and some of the larger Al3Sc particles are marked. In association with the breakup of this junction, new dislocations are created and ejected from the junction into the grain. This was a common observation and shows dislocation ejection from grain boundary triple junctions but this is a responsive rather than a causative effect. Subgrain boundary 2 is now unconstrained at both end points and it breaks free from its interaction with the cluster of large Al3Sc particles a, and the component dislocations move toward and are incorporated into grain boundary 5. Grain boundary 5 also changes during this time, showing more cusps (marked by

arrowheads) than it had initially. These cusps are indicative of the grain boundary being pinned locally as it tries to rotate perpendicular to the surface to minimize its area. A small grain x exists between grain boundaries 4 and 5. It is unaffected by the deformation at the elevated temperature. This observation illustrates the usefulness of the in situ technique as a snapshot such as Fig. 4(b) could be interpreted as the formation of a new grain at the intersection of pre-existing grain boundaries. The duration of the experiment was at least 60 min, which is sufficient for grain growth to occur. Subgrain boundary 3 also moves, consuming grain boundary 7 in the process. This movement may be better seen in Fig. 5, which shows the initial and final positions, with the initial positions in outline form on both images. The motion of the grain boundary is in the direction indicated by the black arrow. Comparing the positions of the same feature shows the grain boundary has moved, and in the process shortened the length of grain boundary 7 by 0.5 µm. This corresponds to the grain

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Fig. 4. Disintegration of a subgrain boundary 2. (a) Initial and (b) final configuration.

Fig. 3. Comparison of the same region following straining in situ in the TEM at nominally the superplastic temperature. (a) Initial configuration and (b) final configuration. Time between (a) and (b) was 60 min.

boundary sweeping out an area of the order of 1 µm2. There is also a shift in the position of grain boundary 6 as indicated by the white arrow. This movement makes the angle at the junction between

boundaries 4, 6, and 8 less acute. Another grain boundary has appeared in one of the grains and is labeled 9. The origin of this boundary is unknown but it is likely due to a migrating boundary being trapped by large Al3Sc particles. The disintegration of a grain boundary in association with the migration of a subgrain boundary is shown in the series of images captured from videotape and presented in Figs. 6–8. The schematic shown in Fig. 6 represents the initial configuration and the labeling used to describe the events leading to the destruction of grain boundaries 3 and 5. The dotted line represents the position of the subgrain boundary at a later stage. The first event is the destruction of segments of grain boundary 5 by the sequential release of arrays of dislocations. Two dislocation arrays are identified in Fig. 7 and are labeled 1 and 2. While, the creation of array 2 was not observed, array 1 is associated with grain boundary 5 and was created as sub-boundary 3 migrated towards boundary 4, leaving portions of grain boundary 5 unattached. The unattached por-

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Fig. 6. Initial configuration of the grain boundaries. The numbers indicate features involved in the reactions, the letters refer to reference points.

Fig. 5. Motion of grain boundary 3 with the corresponding reduction in length of grain boundary 7. The outlines show the initial grain boundary structure. The circled precipitates serve as a reference point.

tion breaks away at some critical but undetermined length. It migrates as a complete unit, suggesting that it is composed of edge dislocations. The migration of array 1 towards array 2 is shown in the series of captured images presented in Fig. 7. Array 1 moves toward array 2 and the two combine to form a single dislocation array, Fig. 7(d). The new array moves as a unit until it encounters some of the large Al3Sc particles at which it is pinned, Fig. 7(f). It breaks up into segments and then into individual dislocations that move across the grain and are incorporated into a surrounding grain boundary. This was the only condition in which arrays of dislocations (subgrain boundaries) com-

bined and the lifetime of the new structure was relatively short. The migration of subgrain boundary 3 can be seen by comparing the schematic in Fig. 6, in which the change in grain boundary position is indicated by the dotted line, and Fig. 8. In Fig. 8, it is trapped by Al3Sc particles c (visible) and d (not visible). The subgrain boundary does not break free of these particles as a unit but breaks into segments as indicated in the schematic in Fig. 6. The central segment becomes mobile, breaking free first from particle d. This causes section 3b to move and rotate towards the grain boundary as the upper portion is still pinned at c. The dislocations move toward grain boundary 4 at which they are accommodated. On accommodation, the grain boundary structure changes, as expected. However, the change in contrast in the grain boundary is short-lived due to the rapid spreading kinetics of dislocations in the grain boundary at this temperature. During the breakup of subgrain boundary 3, copious numbers of dislocations are generated from an unidentified source at particle d. With further straining the remaining segments break free of their pinning points, reverse direction and move back across the grain. Grain boundary 5 also disintegrates, ejecting dislocations from the location of the original triple junction in the process.

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Fig. 7. Migration, coalescence, trapping and breakup of dislocation arrays ejected from a grain boundary. The time for the reaction is given in seconds.

Fig. 8. Position of subgrain boundary at a later stage. It is now trapped by particles “c” and “d”.

The process of disintegration of a subgrain boundary trapped at a particle, which is marked by an arrowhead in Fig. 9(i), is shown in the series of captured images presented in Fig. 9. The breakup involves the ejection of the component dislocations from the sharp bend as opposed to the generation and emission of new dislocations. The destruction of the subgrain boundary results in it breaking into two arrays, which are apparent in Fig. 9(h) and (i). The segment that has broken free rotates in the direction in which the component dislocation moved. The ejected dislocations move through the grain, their motion impeded by the small Al3Sc particles, as evidenced by their line shape. Although it is clear that the small particles impede dislocation motion, the details of the interaction were not studied in this investigation. The dislo-

cations moved through the grain until their path was blocked by the grain boundary between grains 1 and 2 (1/2), Fig. 10. Also seen in Fig. 10 is the subgrain boundary from which the dislocations were ejected and a triple junction. Initially, grain boundary 1/2 shows no internal structure and it appears to be free of trapped dislocations, which is consistent with the rapid spreading kinetics of grain boundary dislocations at the test temperature. The advancement and intersection of the dislocations ejected from the subgrain boundary is shown in the series of images presented in Fig. 10(b)–(d). Many dislocations were observed interacting with and being accommodated in grain boundary 1/2. This incorporation changes the grain boundary structure and its energy. The grain boundary has several options to respond to this change. It can reconstruct to minimize the energy, it can emit dislocations into either grain 1 or 2 in accordance with the slip transfer rules, it can migrate, or it can break apart [14,15]. In this case, it breaks apart and in the process disintegrates the triple junction as shown in the series of images in Fig. 11. The annihilation of the grain boundary between grains 2 and 3 is followed by the disintegration of the boundary between grains 1 and 2. Grain boundary 2/3 moves as a unit and shows some curvature, Fig. 11(e), presumably because it is still pinned somewhere along its original length. The dislocations from the disintegration of boundary 1/2 move into grain 1, which provided the dislocations that initiated this reaction. The contrast

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Fig. 9.

Destruction of a subgrain boundary trapped at a large Al3Sc particle. The time is given in seconds.

Fig. 10. (a) Initial grain boundary structure, including the triple junction. (b)–(d) Dislocations entering the grain boundary between grains 1 and 2. The origin of the dislocations was the breakup of the subgrain, which was described in Fig. 9. The time is given in minutes and seconds.

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Fig. 11. Disintegration of a grain boundary triple junction and the accompanying large-scale volume rotation. The arrows indicate the direction of motion of the subgrain boundaries. Times are in minutes and seconds and continue from Fig. 10.

around the junction changes as boundaries 1/2 and 2/3 are destroyed. This is consistent with there being a local volume rotation to a common orientation. This rotation continues as grain boundary 2/3 propagates through grain 3. Continued high temperature straining caused the grain boundary between grains 1 and 3 to disintegrate, resulting in a complete local volume rotation to a common orientation. In several of the in situ heating and straining experiments, a grain boundary triple point migrated along a grain boundary, resulting in the destruction of that boundary. As the triple point migrated, dislocations were ejected from the migrating junction. In Fig. 12, the junction between grain boundaries 1, 2, and 3, migrated along boundary 2, resulting in a shortening of its length. In conjunction with this reduction of length of grain boundary 2, dislocations were emitted from the triple junction and grain boundary 3 disintegrates as shown in Fig. 12(e). The component dislocations from grain boundary 3 are accommodated in grain boundary 4. As the junction migrates in the direction indicated in Fig. 12(b), dislocations were ejected from the triple point, Fig. 12(c) and (d). Eventually, these dislocations are incorporated into grain boundary 4. Both grain boundaries 2 and 3 have been removed by these reactions, and the junction that was originally between grain boundaries 1 and 2 is now between grain boundaries 1 and 5, Fig.

12(f). In annihilating grain boundary 2, a ledge has been created on grain boundary 5, as indicated by the arrowhead in Fig. 12(f)

4. Discussion From post-mortem analysis of superplastic deformation in which continuous dynamic recrystallization occurs, it has been found that there is an increase in subgrain size, a transformation from a low-angle grain boundary structure to a highangle one, and the existence of non-uniform lattice dislocation densities at high strains. These effects must be accounted for by the proposed operative mechanisms, which are subgrain boundary migration and coalescence; subgrain boundary migration, accelerated by particle reversion, through a curved lattice; dislocation generation, migration, and incorporation into low-angle grain boundaries; subgrain rotation to accommodate sliding along grain boundaries inclined with respect to the tensile axis; and subgrain boundary sliding accompanied by the dissolution of low-angle grain boundaries through dislocation emission. These mechanisms are all based on interpretation of postmortem studies of microstructure following superplastic deformation. Each of these mechanisms is considered in regard to the current dynamic observations.

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Fig. 12.

Migration and annihilation of a triple junction. The time is given in seconds.

From the in situ high temperature deformation experiments, subgrain boundary migration and coalescence does not appear to be a dominant mechanism during continuous dynamic recrystallization. Subgrain boundaries do migrate, but their motion is impeded by the large Al3Sc particles, which exist in the microstructure due to processing conditions. Once trapped, the subgrain boundaries disintegrated ejecting individual dislocations that migrated through the grain until they reached a grain boundary into which they were incorporated. The incorporated dislocations are rapidly accommodated within the grain boundary, which is indicative of the high spreading kinetics in grain boundaries at this deformation temperature. The mechanism of subgrain boundary migration, accelerated by particle reversion, through the lattice was not observed. The fine Al3Sc particles do not appear to have much affect on subgrain boundary motion, although they were observed to impede the passage of individual dislocations. Most likely, the small particles are never forced into solution and the dislocations comprising the low-angle boundaries either cut the particles or bypass them in a climb-assisted process made possible by the use of both thermal and strain energy. The migration of subgrain boundaries is hindered by the large Al3Sc particles although not to the extent suggested by the Zener drag mechanism; this process is discussed elsewhere[16]. No evidence supporting the mechanism of dislo-

cation generation, migration, and absorption into low-angle grain boundaries, proposed to increase the misorientation of these boundaries, was found during the current in situ high temperature straining experiments. One source of lattice dislocations frequently observed was from the disintegration of low-angle boundaries trapped at large Al3Sc particles. The component dislocations from these boundaries migrated freely as lattice dislocations, interacting with the small Al3Sc particles before being incorporated into a grain boundary. However, they were never observed to incorporate into low-angle boundaries and increase the boundary misorientation. Instead, these lattice dislocations were observed to either pass through the low-angle boundaries or temporarily incorporate into them and then contribute to their disintegration. Dislocations were emitted from grain boundary triple junctions but this was a consequence of other events such as the breakup or migration of the junction rather then being a primary response. The mechanism of subgrain rotation to accommodate grain boundary sliding has been suggested to account for the transformation of low-angle grain boundaries into high-angle grain boundaries. Subgrain rotation was observed to occur during the current real-time experiments, but only the removal of low-angle boundaries by this process was noted. The development of higher boundary misorientations through subgrain rotation was not observed. This could be the result of limitations imposed by the thin foil geometry of the TEM tensile samples.

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Rotations to higher angles would require the subgrains to rotate out of the plane of the foil, which is energetically unfavorable. Therefore, subgrain rotation may be an important mechanism during continuous dynamic recrystallization even though the increase in subgrain boundary misorientation was not observed during these experiments. Recent work by McNelley et al. [17] on an Al– Cu–Zr material using electron backscattering with orientation imaging microscopy, has suggested that regions between adjacent deformation bands formed during the final cold-rolling process become high-angle grain boundaries upon annealing and/or high temperature deformation. These can then begin to rotate during deformation under grain boundary sliding conditions. No evidence for this was found in these experiments but it must be remembered that these observations are at a finer scale and in a different orientation, which would make it difficult to see the deformation band. This does not preclude the mechanism suggested by McNelley et al., it just does not support it. The mechanism of subgrain boundary sliding is based on the premise that sliding can occur along subgrain boundaries with misorientations as low as 5°. However, no observations made during these experiments can be attributed to shear along lowangle grain boundaries. Subgrains did not appear to behave as independent crystals capable of shearing along their boundaries. Rather, they behaved as collections of dislocations that migrated and eventually disintegrated. Migration and disintegration of subgrain boundaries, as observed in these experiments, implies rotation of one or both subgrains but does not imply any type of sliding mechanism. The current investigation has shown that the dislocations produced by the breakup of subgrain boundaries migrate through the grain until they interact with and are accommodated in a grain boundary. This process can result in the annihilation of grain boundaries and their junction, which results in a local volume rotation and creates a new grain orientation that is an average of the prior grains. Hoagland [18], using atomistic modeling, has observed that a triple junction can disintegrate in response to an increase in grain boundary energy associated with incorporation of lattice dislo-

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cations, see Fig. 13. The equilibrium grain boundary structure is shown in Fig. 13(a). A dislocation was then introduced into one of the grain boundaries at 500 K, which caused a change in the grain boundary energy. The configuration of the grain boundaries 4 ps later is shown in Fig. 13b. The boundaries and the triple junction are destroyed, emitting dislocations in the process. This simulation appears, at least qualitatively, to be consistent with the experimental observations.

5. Conclusions In situ elevated temperature deformation experiments have been performed in the electron microscope to reveal the deformation process operating

Fig. 13. Atomistic model of a triple point formed at the junction of ⌺33a, ⌺33c, and ⌺9 boundaries. Atoms are shaded according to their excess energy relative to perfect crystal ranging from the dark (lowest energy) to light (highest energy). (a) The equilibrium structure at 0 K, (b) the same configuration after 4 ps at 500 K in which a grain boundary dislocation has been accommodated by emission of glide dislocations from the grain boundaries [18].

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during dynamic continuous recrystallization. These studies have revealed: 1. Subgrain boundaries migrate, become trapped at large Al3Sc particles, and break apart by emitting component dislocations that move through the lattice until they are accommodated in grain boundaries. 2. A mechanism by which grain boundary junctions can be destroyed and volume rotations can occur has been directly observed for the first time. The destruction is in response to mobile lattice dislocations impinging on and being incorporated in the grain boundary.

Acknowledgements This research was carried out in the Center for Microanalysis of Materials, University of Illinois, which is partially supported by the US Department of Energy under grant DEFG02-91-ER45439. The work at Illinois (LMD and IMR) was supported by a grant from Pacific Northwest National Laboratory (PNNL). The work of JSV at PNNL was supported by the Materials Science Division, Office of Basic Energy Sciences, US Department of Energy (DOE) under Contract DE-AC06-76RLO 1830.

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