Grain Boundary Mediated Displacive Diffusional Formation of s-phase MnAl

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1 Grain Boundary Mediated Displacive Diffusional Formation of s-phase MnAl JORG M.K. WIEZOREK, ANDREAS K. KULOVITS, CAGATAY YANAR, and WILLIAM A. SOFFA Dynamic in situ heating and postmortem transmission electron microscopy studies, including high-resolution electron microscopy, have been performed for a near equiatomic composition of a Mn-Al-base permanent magnet alloy to investigate the mechanisms of the transformation of hexagonal close-packed e-phase to a chemically ordered tetragonal s-phase with a face-centered cubic-related L1 0 structure. Although a massive mode dominates the kinetics of this composition invariant transformation, we also observed the genesis of a morphologically plate-like s-phase by partial dislocation glide, described as Shockley-type partials with respect to the L1 0 structure, and a transformation mode exhibiting displacive characteristics. The sources for the transformation dislocations facilitating the formation of the morphologically plate-like s-phase were associated with the interfaces between the parent e- and product s-phase produced by the massive transformation mode. Our experiments revealed diffusional and displacive features of the mechanism accomplishing the formation of s-mnal with plate-like morphology. Hence, a hybrid displacive diffusional mechanism has been identified, and the synergistic role of the nucleation interfaces of the massively transformed s-phase has been discussed. DOI: /s Ó The Minerals, Metals & Materials Society and ASM International 2010 I. INTRODUCTION MNAL-BASE alloys of near equiatomic composition derive their attractive ferromagnetic properties from the formation of s-phase, which exhibits large uniaxial magnetocrystalline anisotropy. [1,2] Figure 1 depicts the binary phase diagram for the Mn-Al system. During the appropriate thermal or thermomechanical treatments, the s-phase forms from the chemically disordered, hexagonal close-packed (HCP) e-phase (A3 in Strukturbericht notation) and is stable up to ~1073 K (800 C) (Figure 1(a)). A competing chemical ordering transformation of e-phase to form the metastable orthorhombic e -phase (B19 in Strukturbericht notation) occurs at temperatures below approximately 850 K (~580 C). Therefore, at high temperatures, the s-phase can form without prior e(a3) fi e (B19) ordering. The relationships between the crystal structures of the parent e-phase and the possible chemically ordered products of the e - and s-phase are illustrated in Figure 2. Figure 2(d) shows the conventional (tp4 Pearson symbol) unit cell of the s-phase (L1 0 in Strukturbericht notation with space group P4/mmm). The chemically ordered tetragonal crystal structure of the s-phase JORG M.K. WIEZOREK, Professor, and ANDREAS K. KULOVITS, Post Doctorate, are with the Swanson School of Engineering, Department of Mechanical Engineering and Materials Science, University of Pittsburgh, Pittsburgh, PA Contact akk8@pitt.edu CAGATAY YANAR, Senior Development Scientist, is with the Alcoa Technology Center, Alcoa Center, PA WILLIAM A. SOFFA, Professor, is with the Department of Materials Science and Engineering, University of Virginia, Charlottesville, VA Manuscript submitted December 21, Article published online May 22, 2010 exhibits close-packed octahedral planes {111} s as well as close-packed directions, h110i s and h101i s ; and it is related to the crystal structure of face-centered cubic (FCC) metals. However, the lattice parameter ratio (c/a) of the conventional unit cell of s-phase deviates significantly from unity: (c/a) = (Figure 2(d)). A simple shear of the e -phase with a B19 structure produces the s-phase with an L1 0 -structure (Figures 2(b) and (c)). [1,2] Hence, a displacive mode or shear mechanism within the parent high-temperature e-phase or following chemical ordering to e -phase at low temperatures has been considered for s-phase formation. Early in situ transmission electron microscopy (TEM) studies [3] provided support for such a sequential reaction of chemical ordering by diffusion followed by a displacive shear step e (A3) fi e (B19) fi s (L1o). [3] However, metallographic studies in the 1980s [4] and 1990s [5,6] indicated that s-phase formation in permanent magnet MnAl-base alloys occurred primarily through a compositionally invariant nucleation and growth mechanism akin to the so-called massive transformation occurring in metallic and ceramic systems, essentially in agreement with the early observations of Ko ster and Wachtel. [7] The massive mode of transformation has been shown to dominate kinetically during isothermal heat treatments up to 973 K (700 C) and involves nucleation almost exclusively at prior parent e-phase grain boundaries followed by growth via motion of mostly incoherent heterophase interface segments. [8] Other recent work reported and discussed the coexistence of the diffusional massive and the displacive shear modes during the formation of s-phase in MnAl-base alloys. [9,10] Solugubenko et al. [11,12] have emphasized the synergistic relationship of B19 ordering to s-phase formation and have studied 594 VOLUME 42A, MARCH 2011

2 Fig. 1 (a) Binary phase diagram for Mn-Al [13] ; the approximate phase fields for the s-phase are marked by dashed borders together with the temperature range for the in situ TEM heating experiments, 813 K (540 C) T 923 K (650 C). Fig. 2 (a) through (c): Orientation relationships between the parent phase and the possible product phases of the competing ordering transformations for nearly stoichiometric composition MnAl-base alloys [3] ;(a) e-phase (A3), (b) e -phase (B19), and (c) s-phase (L1 0 ). The basal planes (0002) e and (001) e are parallel to the octahedral plane (111) s. The shear vector relating the B19 and L10-structures of the e -phase (b) and s-phase (c) is marked in (b). (d) Conventional unit cell (tp4 Pearson symbol) of the L1 0 -ordered s-phase with lattice parameters. the temperature and composition dependence of the alternative modes of transformation. [9,11,12] Importantly, the displacive shear mechanism seems to involve an interplay between short-range diffusion and chemical ordering or reordering when the displacive transformation front propagates through an array of the up to three different e -phase orientation variants, which can form in each of the parent e-phase grains [e.g., see References 8 12]. Also, shear of the HCP parent phase to FCC stacking might occur with concomitant or sequential atomic rearrangement by diffusion, which produces the L1 0 -ordered s-phase. This mechanism, in which a transformation front involves the migration of arrays of transformation dislocations that produce a characteristic shear and change in stacking sequence in concert with diffusional processes leading to chemical VOLUME 42A, MARCH

3 ordering or reordering, might be called a hybrid displacive diffusional mode of transformation. Here, we report on results of in situ heating TEM studies of s-phase formation, which have been complemented with detailed postmortem conventional TEM and high-resolution TEM (HREM) experiments of an MnAl-base alloy. We identify a synergistic relationship between the massive ordering mode and the displacive diffusional hybrid mechanism of transformation from e-phase to s-phase. We present a new hybrid displacive diffusional mechanism that can exploit the presence of the geometrically necessary interfacial dislocations located in the nucleation interfaces of the massively formed s-phase grains and other suitable heterophase interface segments. This work represents a more detailed exposition of the hypothesis presented in Reference 13. II. EXPERIMENTAL PROCEDURE The MnAl-base alloy used in this study was prepared with vacuum induction melting under an argon atmosphere using high-purity elemental starting materials (99.98 pct Mn, pct Al, and pct C) and had a composition of 54.4 at. pct Mn, 44 at. pct Al, and 1.7 at. pct C. Small amounts of C are added to increase the stability of the s-phase. More details of the alloy preparation are described in a previous report. [8] The electron transparent foils for TEM experiments have been prepared by electropolishing with a twin-jet polisher at room temperature, applying a voltage of 16 V, and using a solution of 8-pct perchloric acid and 92-pct acetic acid by volume. TEM was performed at 200 kv using a Jeol JEM-200 CX (Peabody, MA) and a Philips CM200/FEG instrument and at 300 kv using a Jeol JEM-3010 for the in situ TEM and an FEI Tecnai F30 (Hillsboro, OR) for the HREM, respectively. Single-tilt heating holders from Gatan Inc. (Pleasanton, CA) have been used with the JEM-200CX and JEM-3010 in the heating experiments. III. RESULTS AND DISCUSSION Typically, the transformation in bulk MnAl-base alloys of the disordered parent e-phase to the uniaxially ferromagnetic s-phase is dominated kinetically by a massive ordering mode for isothermal treatments between approximately 773 K (500 C) and 973 K (700 C). [8] The massive s-product nucleates almost exclusively at prior parent e-phase grain boundaries (e.g., Figure 3(a)), obeying an approximate orientation relationship with one of the parent e-phase grains into which little or no subsequent growth occurs. [8,14] The orientation relationship of the massively transformed s-phase s m with the e-phase parent grain, into which little or no growth occurs (i.e., with respect to the nucleation interface), can be described as {111} s (0002) e and h-110i s h-2110i e or as h10-1i s h-2110i e. [8] The growth of the massive s-product s m is accomplished with the migration of incoherent heterophase interfaces by essentially random atomic attachment across the growth interface and is associated with the genesis of characteristic defect content (e.g., arrays of stacking faults (SFs), microtwins, and antiphase boundaries in the s m product). [8,15] The bright-field TEM micrograph of Figure 3(a) has been obtained after a partial s-phase transformation at 813 K (540 C) and illustrates some characteristic features of the massive s-product s m. Thus, the semicoherent nucleation interface and the incoherent growth interface are marked in the example TEM micrograph of the early stages of this massive ordering mode of s-phase formation. Furthermore, the contrast in the two parent phase grains (labeled A and B in Figure 3(a)) is consistent with the formation of chemically ordered e -phase precipitates within the e-phase matrix, which is associated with initially elastically accommodated anisotropic misfit strain. A careful inspection of the grain boundary between the e-phase grains labeled A and B revealed the formation of a thin layer of s m -product. The multiple bright-field TEM Fig. 3 (a) TEM bright-field micrograph depicting massive s-phase forming preferentially at prior e-phase grain boundary. The two parent phase grains, labels A and B, have undergone some chemical ordering and represent (e/e )-phase mixtures. The typically incoherent growth interface and the nucleation interface, across which an approximate orientation relationship exists (semicoherent), are marked. (b) Multibeam bright-field TEM micrograph showing massively transformed s-phase (label m), shear-mode transformed s-phase (label s s ), and the parent phase mixture (label e/e ). 596 VOLUME 42A, MARCH 2011

4 Fig. 4 Examples of the plate-like or shear-mode s-product formed during in situ heating TEM experiments at temperatures between about 813 K (540 C) and 923 K (650 C). (a) Bright-field TEM of inclined planar fault packages with inset SAD pattern confirming the presence of e/e -phase mixture with streaks originating from planar faults. (b) HREM for beam direction of e s the plate-like s-phase formation by packages/arrays of planar faults in the e/e -parent phase. Some regions of the product and parent phase are marked by labels s and e, respectively. micrograph of Figure 3(b) shows a representative example of the microstructure in the late stages of the e fi s transformation. Both massively transformed s-phase (labeled m in Figure 3(b)) and regions of morphologically plate-like s-phase, which formed via shear-mode (labeled s s in Figure 3(b)), as well as some remaining e- and e -phase mixture (labeled e/e in Figure 3(b)), are observed. The formation of the morphologically platelike s-phase s s has been observed after most of the prior e-phase volume of the specimens had been transformed to s-phase by the massive mode during in situ heating TEM experiments conducted at temperatures between about 813 K (540 C) and 923 K (650 C). The platelike morphology is depicted in the bright-field TEM and HREM micrographs of Figure 4. The selected area diffraction (SAD) pattern inset in Figure 4(a) indicates that the transforming parent phase is a mixture of e-phase and e -phase. Furthermore, streaking parallel to the basal plane normal of the e-phase in the SAD, is consistent with the presence of the dense arrays of planar faults in the close-packed planes (0002) e. The HREM micrograph (Figure 4(b) shows the orientation relationship between the HCP parent phase (labeled e in Figure 4(b)) and the product phase (labeled s in Figure 4(b)) and confirms the basal plane faulting associated with the early stages of formation of the morphologically plate-like s-phase. In situ TEM imaging and diffraction showed that the planar defect arrays become increasingly denser as this latter transformation mode continues to operate, which eventually results in the formation of s-phase that contains a high density of twin and SFs on the closepacked plane of the FCC-related L1 0 structure, which is parallel to the basal plane of the HCP e-parent phase. The postmortem TEM data compiled in Figure 5 summarizes this growth scenario for the plate-like s-phase. The parent grain of the (e + e ) phase mixture contains a low-angle boundary (LAGB). The upper part of the bright field (BF)-micrograph in Figure 5 (above the LAGB) exhibits a contrast similar to that observed in Figure 4. The corresponding SAD pattern shows the presence of e-phase and two different variants of e -phase, namely e 0 1 and e0 3 : It also shows streaking along g = [0001] e. However, the lower part of the BF-micrograph in Figure 5 (below the LAGB) is more densely populated with planar defects (i.e., a high density of SFs), and the corresponding diffraction pattern no longer exhibits the (100) superlattice reflection of the e 0 1-phase. The 1101; 1100; and 1101 reflections of the hexagonal e-phase also no longer are present. This pattern can be indexed as s-phase in [101] orientation with a 60 deg rotated twin and with remnants of e 0 3-phase in a [010] orientation. Therefore, it is reasonable to conclude that the e- and e 0 1-phase of this TEM section below the boundary in Figure 5 essentially have transformed to s-phase. The required change in stacking sequences to transform an HCP structure into an FCC structure can be accomplished by introducing partial dislocations with a Burgers vector 1= in the HCP parent phase, which corresponds to Shockley partials of Burger vector 1/6h112i in the FCC structure, on every other closepacked plane of the HCP phase. If we consider the transformation from B19 (e ) to L1 0 (s) rather than from HCP to FCC, then the atomic ordering becomes important. In the HCP structure, three equivalent partial dislocations (1= ; 1= ; and 1= ) can accomplish the transformation of HCP to FCC as long as they move on every other closepacked plane (i.e., the partials do not necessarily have to be of the same type to result in FCC structure). Because the e -phase forms by ordering atoms along the 1120 e (the direction of the HCP matrix), e -phase precipitates can exhibit three orientation variants within the e-phase. The orientation relationships between the e-phase and the three variants of e precipitates are as follows: Variant 1: (0001)e // (001)e and 1120 e // [010]e Variant 2: (0001)e // (001)e and 1210 e // [010]e Variant 3: (0001)e // (001)e and 2110 e // [010]e VOLUME 42A, MARCH

5 Fig. 5 Bright-field TEM and accompanying SAD showing the transformation from the faulted HCP and HCP-related parent phase mixture of the e- and e -phase at lower densities of planar faults to the faulted s-phase as the planar fault density increased sufficiently. Fig. 6 Schematics of (a) the orientation relationships between the e-phase and the three possible orientation variants of the e -phase and (b) the possible shear transformation of the correct e -phase variant to form s-phase in a given prior e-phase grain. Figure 6(a) shows the schematics of the e variants with respect to the hexagonal e phase projected along the [0001] e or the [100] e zone axes. If we consider Variant 1, e 0 1, with respect to matrix (e), then we can write the following orientation relationships (ORs): 1100 e== 001 e e== 031 e e== ½031Še VOLUME 42A, MARCH 2011

6 Because the directions [001] and [031] are crystallographically different in the orthorhombic unit cell, only one type of partial dislocation can shear one of the three possible e -phase variants that may be produced in a given e-grain into the L1 0 structure with the required correct ordering of atoms. For example, to shear the e -variant 1 e 0 1 into the L1 0 structure, the partial dislocation with a Burgers vector of 1=31100 (or 1/ 3[001] with respect to the e -phase and 1/6[11-2] with respect to the s-phase) should glide on the close-packed plane and would be a suitable transformation dislocation (Figure 6(b)). Therefore, if several partial dislocations glide on every other close-packed plane of the parent phase mixture (e + e ) in which there are two or three variants of e, then the resulting structure will not be L1 0. The shearing of e precipitates other than those belonging to the singular correct variant would result in complex planar faults. Hence, in a given e-phase grain, two obstacle variants and one correct variant of the e -phase exist. A significant local rearrangement of the Mn and Al atoms would be necessary to produce the ordering along the [001] direction characteristic of the s-phase. [9,11,12] An example is observed in Figure 5. The e-phase and, in this case, the crystallographically correctly oriented e 0 1 variant are transformed into s-phase below the LAGB. Reflections of both e- and e 0 1-phases disappear and are replaced by reflections corresponding to s-phase. The e 0 3-variant, however, acts as an obstacle variant to the shear-type transformation in this particular grain. Reflections corresponding to the e 0 3-variant remain visible with emerging reflections of the newly formed s-phase. Figures 7 and 8 present individual still frames extracted from video recordings of dynamic sequences of s-phase transformation during the in situ TEM heating experiments. Figure 7 shows extraordinary micrographs extracted from a dynamic video sequence that reveal the injection of transformation dislocations from the nucleation interface of the massive product (labeled s massive in Figure 7(a)). These gliding transformation dislocations generate s-phase either by direct shear from the appropriate B19-ordered regions of e -phase and/or by creating SFs and precursory FCC regions in the HCP e-phase, which order by short-range diffusion concomitantly or subsequently. The nucleation interface formed by the massively transformed s-phase at segments of prior e-phase grain boundaries is, in most cases, a semicoherent interface and can be a coherent interface with respect to the massive s-phase and the e-phase grain into which little or no growth occurs. Fig. 7 Still frames extracted from a dynamic video sequence of an in situ heating TEM experiment depicted the nucleation of the planar faults at the nucleation interface for the previously formed massive s-product, s-massive. The original parent phase grain boundary, the massive product, and the parent phase e/e -phase that transforms to a plate-like s-product are marked. The arrow indicates the nucleation of a new planar fault at the nucleation interface of s-massive and the growth facilitated by a transformation dislocation into the e/e -parent phase. VOLUME 42A, MARCH

7 Fig. 8 Growth of morphologically plate-like s-phase into the mixture of e- and e -phase. The s-phase plate formed by the shear -mode marked by the arrow in (a) grows to a position in (b) where its growth temporarily slows and stagnates (c) and (d). Although compositionally invariant, the massive transformation accomplishes a change in the stacking sequence from that of HCP for the parent phase to an FCC-related product phase. Hence, it is either a crystallographic or a geometric requirement that transformation dislocations reside at the nucleation interface to accommodate the stacking sequence change. Such transformation dislocations encountering the singular correct orientation variant of the three possible orientation variants of the e -phase can produce the s-phase directly by shear (Figure 6(b)), [9] whereas the other two e -phase orientation variants will be effective obstacles to the transformation dislocations and the s-phase transformation in this mode, unless local reordering occurs as recently was discussed in detail by other investigators. [9,11,12] The progression or growth of the plate-like s-phase into the e-e phase mixture exhibited periods of rapid progress interrupted intermittently by periods of stagnation or rest during our in situ hot-stage TEM observations (e.g., Figure 8). This would be consistent with individual transformation dislocations propagating into the parent phase mixture of e- and e -phase by gliding on the close-packed planes until they encounter an obstacle (e.g., one of the two incorrect e -phase orientation variants that formed in the prior pure e-phase grain during in situ TEM heating experiments). In the as-quenched and annealed bulk samples, all three orientation variants of e exist within the e-matrix. However, an important observation was made in the samples that were exposed to in situ heating TEM experiments concerning the e -variants. The multibeam BF-micrograph of Figure 9 shows the faulted s-phase (dark contrast) and the surrounding matrix consisting of e + e. In the SAD pattern taken from the faulted region of the s-phase, the superlattice reflections are present, despite being very weak, which indicates that this faulted region is not transformed completely to s-phase but contains several e -phase particles. A series of tilting TEM experiments have been performed for the e + e matrix adjacent to the band of the displacively transformed faulty s-phase shown in Figure 9 and for regions of the same matrix grain away from the band of s-phase. In Figure 10, the SAD patterns of the matrix near the sheared region (Figure 10(b)) and the patterns that were taken away from thatregion (Figure 10(a)) are shown. The zone axes are 1120 ; 0110 ; and 1010 with respect to the e-phase. SAD patterns in Figure 10(a) confirm that all three variants of the e -phase exist in the matrix grain region away from the shear-transformed s-phase band. The variants responsible for the superlattice reflections in Figure 10(a) are labeled as e 0 1 ; e0 2 ; and e0 3 (other reflections from all three variants overlap). In Figure 10(b), the reflections for e 0 1 and e0 2 are missing 600 VOLUME 42A, MARCH 2011

8 Fig. 9 (a) Multibeam BF TEM micrograph and (b) corresponding SAD (see text for details). Fig. 10 SAD patterns obtained from a series of tilting experiments in postmortem TEM after transformation during in situ TEM heating experiments (see text for details). from the 1120 and 0110 patterns, and only the e 0 3 superlattice reflection is present in the 1010 pattern (i.e., only one variant of e -phase is present in the matrix grain near the transformed s-phase band of Figure 9). This solves the problem of partial dislocations cutting through different variants of e -phase. Although it is currently unclear what drives the local orientation variant selection for the e -phase if only one e -variant is present as a result of local variant selection, the formation of s-phase by a displacive or shear mode can be envisaged by Shockley partials of the same kind gliding on close-packed planes of the matrix. It is necessary to have these partial dislocations on every other close-packed plane to complete the e + e fis transformation. However, if the partials do not glide on every other plane, then SFs would be introduced into the product phase. The TEM observations revealed that the plate-like s-phase actually forms from the overlapping of individual SFs introduced in a random manner into the e- + e -phase mixture. The relationship between the parentandthe product phases by the glide of appropriate 1=6 112 dislocations with respect to the s-phase can be envisaged with reference to Figures 2 and 6(b). The glide of partial dislocations of the opposite sign (e.g., 1=6 112 in this case) also could glide on close-packed planes, and they can lead to the formation of twin variants in the product s-phase. These scenarios are consistent with the in situ and postmortem TEM and HREM observations reported here. VOLUME 42A, MARCH

9 Fig. 11 BF TEM micrographs of plate-like s-phase and s S, revealing an APB structure; (a) multibeam BF TEM and (b) two-beam BF of APB-related s-phase domains. The formation of s-phase by the glide of partial dislocations of the same kind and sign would result in the accumulation of stresses ahead of the growth front. However, if an equal number of partials of opposite sign glide to form the s-phase, then the overall transformation stress would be zero. Therefore, the formation of microtwins during the displacive mode to produce s may be proposed to be a stress-relaxation mechanism. Because the matrix contains ordered (e -phase) and disordered regions (e-phase), the product of the dislocation shear-mitigated displacive mode would consist of L1 0 and FCC structures. The subsequent ordering of the FCC regions would transform the whole product of the shear mechanism into L1 0. The impingement of ordered domains eventually would form antiphase domain boundaries (APBs). Figure 11 shows micrographs of displacively formed s-phase that contains fine-scale APBs. The multibeam bright-field micrograph of Figure 11(a) clearly depicts the presence of planar faults on the 111 plane with respect to the s-phase. The micrograph of Figure 11(b) is a BF image of the same area shown in Figure11(a) but is taken with a two-beam condition with g ¼ 110 near the [112] zone axis of s-phase and confirms the presence of small APB-related L1 0 domains. Note that the g vector used in this micrograph is perpendicular to the displacement vector R ¼ 1=3 111 ; and therefore, the SF contrast is not visible in the micrograph (g 9 R = 0) of Figure 11(b). The subsequent L1 0 ordering of the FCC regions, produced by the shearing of prior HCP e-phase, requires short-range diffusion. The overall interplay between structural shear and atomic diffusion to produce a new, more stable phase essentially defines what the current authors call a hybrid displacive diffusional transformation. [16] The in situ heating TEM observations showed that the partial dislocations responsible for the transformation shears generally originated at the grain boundaries of the e-phase or at the stagnated and immobilized massive transformation interfaces and propagated into the e + e -phase mixture. Figure 12 is a set of schematics of the hybrid or shear mode and shows the emergence of a plate-like morphology of s-phase synergistically with the massive product consistent with the many observations made during this in situ TEM study. Given that the systematic orientation relationship obeyed across the nucleation interface between massive s-phase s m and one of the parent e-phase grains (e.g., e 2 in Figure 12(a)), which can be described as h11-20ie h101is or h110is and (0001)e {111}s, [8,14] an array of dislocations must reside at the nucleation interface to accommodate the change in the stacking sequence of the close-packed planes (Figure 12(b)). The Burgers vectors of these interfacial dislocations can be described with respect to the s-phase as b = 1/6h112i. These partial dislocations are glissile on the close-packed planes in both s- and e-phases that are continuous across the nucleation interface. Once these partial dislocations move into the less stable e-phase changes in the HCP stacking sequence are accomplished, which initially results in basal plane SFs, and if accomplished on every second (0002) e the plane produces FCC stacking. Together with some short-range diffusion, the motion of these partial dislocations, therefore, can accomplish the transformation from e-phase to s-phase with plate-like morphology (i.e., this is the shear-mode s-phase s S ) in Figure 12. If the moving partial dislocations, which act as transformation dislocations (e.g., Figures 4 through 8), encounter orientation variants of the e -phase that cannot be sheared directly to s-phase, then, again, short-range diffusion is required to accomplish s-phase transformation, which may be the rate-limiting step for this transformation mode. Because this mechanism implies concomitant displacive shear and short-range diffusion, it represents a possible hybrid displacive diffusional transformation mode. It must be emphasized that this new hybrid mechanism of transformation does not preclude a primary role of B19-ordering. Instead, it is suggested that both mechanisms together are operative under most transformation conditions and together can explain what have seemed to be rather disparate results or mechanisms reported previously. This electron microscopy study of s-phase formation in a Mn-Al-C permanent magnet alloy clearly shows that both massive and hybrid displacive diffusional modes coexist even within the 602 VOLUME 42A, MARCH 2011

10 Fig. 12 Schematics (a) through (d) describing the synergistic role of the massive mode of s-formation s m for the hybrid displacive diffusional mode of plate-like or shear-mode s s s-formation, which is consistent with experimental data (see text for details). range where the massive transformation is kinetically dominant. Although the shear mode has been related to B19-ordering in the parent HCP phase, our results and analysis suggest that a hybrid displacive diffusional mode may occur synergistically in conjunction with the nucleation interface of the massive product. This synergism with the massive product together with the mechanism related directly to B19-ordering are consistent with experimental results, indicating faulting in the HCP phase as well as the formation of a metastable FCC precursor to s-phase formation. However, it is recognized that the transformation dislocations, which are incorporated within the nucleation interface as partials accommodating the stacking sequence change, can interact with e -particles in close proximity. In this article, we refer to a hybrid displacive diffusional transformation as a phase transformation wherein the elementary mechanisms involved in producing the product phase within the parent phase constitute a characteristic shear component propagated by an array of glissile dislocations coupled with longrange or short-range diffusion to effect a change in composition or degree of order. This says nothing about a lattice invariant shear, which combines with the Bain strain to produce an invariant plane strain in the classic crystallographic theory of martensite formation. It must be remembered that this composite deformation is constituted to minimize the strain energy attendant to the transformation, and therefore, it should not be surprising to find various solid-state transformation mimicking features of martensite. IV. SUMMARY In situ and postmortem TEM investigations of the transformations to s-phase in a Mn-Al-C-base alloy have been conducted. Typical transformation conditions (e.g., isothermal annealing between about 813 K (540 C) and 923 K (650 C) after quenching-in of the high-temperature HCP e-phase) involve both the massive ordering mode and the hybrid displacive diffusional modes. Intrinsic synergistic roles between the heterogeneous nucleation of massive s-phase s m and the morphologically plate-like or shear -mode s-phase s s formation have been identified and discussed based on a hybrid displacive diffusional mode, which involves the motion of partial dislocations that act as transformation dislocations and concomitant short-range diffusion. ACKNOWLEDGMENTS The material presented in this article is based on work supported by the National Science Foundation. Any opinions, findings, conclusions, or recommendations expressed in this material are those of the authors and do not necessarily reflect the views of the National Science Foundation. We are grateful for the assistance of E.A. Stach and V. Radmilovic during in-situ heating TEM experiments, which have been performed with the JEM3010 at the National Center for Electron Microscopy, Lawrence Berkeley National Laboratory, Berkeley, CA, and to H. Heinrich, now at VOLUME 42A, MARCH

11 the University of Central Florida, for assistance with HREM experiments performed with the 300 kv FEG- TEM at the Eidgeno ssiche Technische Hochschule Zu rich, Zu rich, CH. Finally, we acknowledge the use of facilities of the Materials Micro-Characterization Laboratory in the Department of Mechanical Engineering and Materials Science of the University of Pittsburgh. REFERENCES 1. S. Kojima, T. Ohtani, N. Kato, K. Kojima, Y. Sakamoto, I. Konno, M. Tsukahara, and T. Kubo: AIP Conf. Proc, AIP, New York, NY, 1975, p Y.Z. Vintaykin, V.A. Udovenko, I.S. Belyatskaya, N.N. Luarsabishvili, and S. Yu. Makushev: Phys. Met. Metalloved., 1974, vol. 38, p J.J. Van Den Broek, H. Donkersloot, G. Van Tenedloo, and J. Van Landuyt: Acta Metall., 1979, vol. 27, p A.V. Dobromoslov, A.E. Ermakov, N.I. Taluts, and M.A. Uimin: Phys. Stat. Sol. A, 1985, vol. 88, p D.P. Hoydick, R.J. McAfee, and W.A. Soffa: in Boundaries and Interfaces in Materials, A.H. King and D.B. Williams, eds., TMS, Warrendale, PA, 1998, p D.P. Hoydick and W.A. Soffa: Scripta Mater., 1997, vol. 36, p W. Ko ster and E. Wachtel: Z. Metallknd., 1960, vol. 51, p C. Yanar, J.M.K. Wiezorek, V. Radmilovic, and W.A. Soffa: Metall. Mater. Trans. A, 2002, vol. 33A, p P. Mu llner, B.E. Burgler, H. Heinrich, A.S. Sologubenko, and G. Kostorz: Phil. Mag. Lett., 2002, vol. 82, p J.H. Kim, C.T. Lee, and W.K. Choo: in Solid-Solid Phase Transformations 99, M. Koiwa, K. Otsuka, and T. Miyazaki, eds., JIM, Sendai, Japan, 1999, vol. 12, p A.S. Solugubenko, P. Mullner, H. Heinrich, and G. Kostorz: Microsc. Microanal., 2003, vol. 9, p A.S. Solugubenko, P. Mullner, H. Heinrich, M. Wollgarten, and G. Kostorz: J. Phys. IV, 2003, vol. 112, p J.M.K. Wiezorek, C. Yanar, E.A. Stach, V. Radmilovic, and W.A. Soffa: in Hybrid Displacive Diffusional Transformation in Manganese Aluminum-Base Alloys, Solid-to-Solid Phase Transformations in Inorganic Materials, J.M. Howe, D.E. Laughlin, J.K. Lee, U. Dahmen, and W.A. Soffa, eds., TMS, Warrendale, PA, 2005, p C. Yanar, V. Radmilovic, W.A. Soffa, and J.M.K. Wiezorek: Intermetallics, 2001, vol. 9, p T.B. Massalski: Binary Alloy Phase Diagrams, TMS, Materials Park, OH, B.C. Muddle and J.F. Nie: in Characteristics of Diffusional Displacive Transformation Products, Solid-Solid Phase Transformations 99, M. Koiwa, K. Otsuka, and T. Miyazaki, eds., JIM, Sendai, Japan, 1999, p VOLUME 42A, MARCH 2011