RECRYSTALLIZATION NUCLEATION AND GRAIN ...

Proceedings of the 6th International Conference on Recrystallization and Grain Growth (ReX&GG 2016) Edited by: Elizabeth A. Holm, Susan Farjami, Priyadarshan Manohar, Gregory S. Rohrer, Anthony D. Rollett, David Srolovitz, and Hasso Weiland TMS (The Minerals, Metals & Materials Society), 2016

RECRYSTALLIZATION NUCLEATION AND GRAIN GROWTH IN Al-1%wt.Mn SINGLE CRYSTALS OF STABLE ORIENTATIONS Magdalena M. Miszczyk1, Henryk Paul1, Julian H. Driver2 1

Polish Academy of Sciences; Institute of Metallurgy and Materials Science; 25 Reymonta St.; 30-059 Krakow, Poland 2 Ecole des Mines de Saint Etienne; Centre SMS; 158 Cours Fauriel; 42000 Saint Etienne, France Keywords: Nucleation and grain growth, Orientation relations, Al-Mn alloy, Single crystal, SEM/EBSD Abstract The crystallographic aspects of recrystallization nucleation have been characterized in Al1%wt.Mn single crystals of Goss{110} and brass{110} orientations. Samples were compressed 40% and then lightly annealed to partial recrystallization. SEM/EBSD orientation maps revealed that boundaries with misorientation angles above 8o are not present in the deformed structures. During annealing high angle boundary segments (>15o) were created at an advanced stage of recovery preceding grain nucleation, i.e. some of the low-angle boundaries evolve into high-angle boundaries. As the dislocation-free subgrains are now partly or completely surrounded by high-angle boundaries, the orientations of nuclei are different from the orientations of their deformed/recovered neighbours. There is also a strong relation between asdeformed orientations and the limited number of recrystallized grain orientations. The results are discussed in terms the creation of twist or tilt grain boundaries of high misorientation by thermally activated, co-ordinated movement of dislocations during recovery. Introduction The structure and texture transformations during annealing of deformed face centred cubic metals has been examined for several decades, e.g. [1,2]. There is a widely accepted opinion that the orientations of recrystallized grains originate from the orientations of the as-deformed areas from which they grow, e.g. [3-5]. However, there is still no agreement for the mechanism by which the orientations of the deformed structure are transformed into primary nuclei orientations. In particular, it is still an open question whether the nuclei/grain orientations are present in the deformed structure or the orientations of the deformed state evolve towards new grain orientations. The details of post-nucleation grain growth are also still not clear. Barrett [6] first suggested that the growth potential of a given nucleus/grain depends on the orientation difference across the migrating recrystallization front. This implies that near-equiaxed grains should be formed during annealing of deformed single crystals of a uniform microstructure/texture. However, in the early stages of recrystallization of grains without significant deformation gradients a considerable number of recrystallized grains exhibit a strongly elongated shape [7-9], despite a near-uniform misorientation relation in all directions. This anisotropic growth suggests that it is not only the misorientation across the recrystallization front but the type of migrating grain boundary, e.g. their twist or tilt character, which is an important factor controlling migration velocity, e.g. [7-9]. In order to clarify this problem, it is

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advisable to perform experimental research under conditions, where the number of 'free parameters' is minimized. Therefore, the current investigations have focused on an analysis of transformations which occur in model single crystals with stable orientations. The objective of this paper is to characterize the ‘crystallographic parameters’ of preferred misorientation relationships and the type of grain boundary appearing between a new grain and the deformed/recovered region during the early stages of recrystallization of high SFE metals. The choice of Al-1%wt.Mn single crystals of stable orientations simplifies the current microstructural and texture analyses and improved its clarity: no initial grain boundaries and deformation twins, absence of deformation-induced high-angle (>15o) boundary segments, and limited recrystallization twinning. Moreover, good structural and textural stability during deformation enables one to precisely define the orientation relationship when the ‘primary nucleus/grain’ is formed. The basic research technique employed local orientation measurements based on scanning electron microscopy (SEM) equipped with electron backscattered (EBSD) facility. Experimental The single crystal bars of single phase Al-1%wt.Mn were grown by directional solidification and the nominal (110)[1-1-2] and (110)[00-1] oriented samples were carefully cut from the bars to dimensions of 10 x 10 x 10 mm3 (height x width x length) using a wire saw. Crystal orientations were checked by X-ray diffraction. In fact some the initial crystal orientations used here were ‘non-ideal’ and samples were rotated 2o - 4o away from the ideal position. The deformation was carried out by plane strain compression (PSC) at 293 K to a true strain ε=ln(hinitial/hfinal ) of 0.51 (or 40% thickness reduction), where a clear microband structure, with relatively weak scattering of the initial orientation (i.e. without high angle boundary segments), was observed. In order to limit friction between the sample, the punch and the walls of the channel-die, each sample was wrapped in PTFE films. To investigate the interrelation between as-deformed and recrystallization textures, the deformed samples were cut perpendicular to TD into ~2.5 - 3.0 mm sheets by means of a wire saw and then annealed (in an air furnace) to the first stages of primary recrystallization, typically at 450 - 650 K for less than 240 s. Prior to annealing all samples were mechanically ground on all faces to 4000 grid SiC paper and electropolished. All specimens were water quenched post-annealing. The microstructures were analysed on the ND/ED plane in the sample mid-thickness. The microtextures were characterized by SEM using a JEOL 6500F or Zeiss Supra 55VP equipped with EBSD facility. The applied step size ranged between 0.1 µm and 2 µm. Post-processing analysis of the orientation maps was performed using HKL Technology Channel 5 software. Results and Discussion ‘Isotropic’ Distribution of Misorientations in the Deformed Microstructure The as-deformed microstructures of both orientations are composed of two sets microbands of approximately the same volumes (Figs. 1a and b). The textures show strong stability of the initial orientation, as presented many times in the past, e.g. [10-12]. Over most of the sample rectangular cells were observed as a result of microband intersections, with dimensions of 200 500 nm. At this strain level the traces of dislocation walls correspond to the traces of {111} planes. The majority of adjacent microbands displayed opposite rotation senses and the

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misorientation angles between particular pixels of the orientation maps were significantly below 8o. Statistical analysis of the frequency distributions of the misorientation angle with respect to the average orientation is presented in Fig. 1c, based on whole orientation maps. The plots clearly demonstrate again that deviations above 8o are not present in the deformed structures, as seen also by typical line scans (wrt. first measured point); the latter is presented in Fig. 2 for Goss{110}-oriented single crystal. In the both orientations slightly higher values of misorientations are observed for line scans inclined at (+/-) 45o to ED.

Fig. 1. Microstructure of brass {110} (a) and Goss{110} (b) oriented single crystals (ND/ED plane). Misorientation frequency distribution (with respect to average orientation) of Goss and brass orientated single crystals (c). Orientation maps presented as a function of ‘local misorientation’ between neighbouring pixels, where: whiter pixels showing higher misorientation (range of misorientations: 0o (white) – 10o (black)). After a strain of 0.5 the spread of misorientation axes (between pixels in the as-deformed state) is small. In the case of Goss-oriented crystal the maximum density of misorientation axes occurred in the vicinity of TD. There is, however, a spread away from TD so that the most frequent axes were near the closest and axes to TD. In the case of brass-oriented crystal, the misorientation axes display a systematic departure from TD. The maximum misorientation axis is located between TD and the intersection trace of the most active {111} planes (along direction). New Grain Orientations Despite the scatter of the recrystallized grain orientations, it can be claimed with certainty that they are not random and only a finite number of groups of recrystallized grain orientations were

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observed. Most of the orientation groups are symmetrically situated with respect to certain axes, i.e. the ‘texture image’ of the first recrystallized grains is quite symmetrical with respect to the external coordinate directions. This symmetry results from positive and negative rotations of the new grain orientations around axes grouped near selected the poles of the deformed state.

Fig. 2. Typical misorientation line scan (wrt. first measured point) showing small misorientations between particular microbands (d). A primary nucleus of homogeneous orientation is related to the adjacent deformation orientation by a 25o - 45o rotation around an axis located near the pole of the {111} planes of its deformed environment. This rotation mechanism is not accidental. Very often it was noticed that if a grain forms by a rotation around a near axis, then a neighbouring grain appears characterized by a rotation around another axis. However, a general rule is that the axes of recrystallized grains are grouped around one of the {111} plane normals of the deformed state but only rarely coincide with them. ‘Directionality’ of Grain Growth The orientation maps measured in the ND/ED section of Goss and brass oriented samples confirm that most grains are characterized by a strongly elongated shape and the straight high angle grain boundaries were observed along traces of {111} planes. This anisotropic growth of new grains in a relatively homogeneous deformed structure suggest the mechanisms responsible for rapid grain growth. The typical partly recrystallized microstructure observed in central areas of the ND/ED section of Goss-oriented single crystal after annealing at 696 K for 25 s is presented in Fig. 3a, as ‘local misorientation’ map. The orientation spreads are qualitatively described by the frequency distributions of the misorientation angle with respect to the average orientation (Fig. 3b). The plots clearly demonstrate a changing spread of misorientations during annealing, from 1-8o for the deformed sample (see Fig. 1c) to 1-20o for recovered areas (of Gossoriented sample). If the (first) recrystallized grains were taken into account the second (broader) peak between 25-45o is observed. As expected, the average misorientation angle from the average orientation is small for the as-deformed microstructure (see Fig. 1c), whereas in the case of recovered microstructure increases significantly. The orientation of deformed/recovered areas (Fig. 4a) is similar to that observed after deformation, whereas the crystal lattice of elongated grains exhibit a well-defined rotation with respect to the deformed/recovered areas (Figs. 4b and c). It is clear that the longest grain direction runs mostly along the traces of the most active planes during strain, which are symmetrically inclined to the ND/TD plane. Another rule is that

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the orientations of the elongated grains are related to those of the deformed crystals. The crystal lattice of each of the grain groups is rotated (positively or negatively) around diverse groups of misorientation axes. Generally, the crystal lattice rotation of elongated grains takes place around axes near the pole, corresponding to the {111} plane along which preferred grain growth occurs (Fig. 4d). Although the rotation axes approach the normal vector of the preferred slip planes they only rarely coincide with the exact location of the directions. Moreover, a small number of grains are elongated along traces of two other {111} planes. In ND/ED section they run vertically, i.e. along ND direction (Fig. 3a). None of the grains grows horizontally, i.e. parallel to TD, in this section.

Fig. 3. Directionality of grain growth along (111) plane at early stages of recrystallization (696 K for 25 s) observed in Goss{110} orientation. (a) Orientation map presented as a function of ‘local misorientation’ made in longitudinal section and corresponding (b) misorientation frequency distribution with respect to average orientation of the deformed/recovered state. The present results confirm that a significant number of recrystallized grains show growth anisotropy and the directions of fast growth coincide with traces of the {111} planes. Since the deformed microstructure is free of significant orientation gradients, this confirms that the boundary velocity is not reduced by orientation effects alone. For this reason orientation pinning is not thought to be a major contributor to the production of flat plate-like grains and the velocity of grain growth may be affected directly by the nature of the dislocation arrangements. Rapidly growing plate-like grains, along specific directions, are observed here in both orientations. In all these cases the recrystallized grains were related to the matrix by a rotation of 25-45o about axes lying near the poles. The morphology of these flat grains was in all cases such that broad faces of the plates were perpendicular to the rotation axis. Thus the fast growing rims of the plates at the head of a growing grain were high angle tilt boundaries and the slower growing broad faces were high angle twist boundaries. The boundary type across the recrystallization front can be used to elucidate the dislocation mechanisms responsible for grain growth. In an idealized form two cases are possible [9]. On the one hand, the movement of pure screw dislocations of two families (with orthogonal Burgers vectors) leads to a pure twist boundary with a rotation axis perpendicular to the boundary plane, as originally suggested by Ridha and Hutchinson [5]. In this case the misorientation axis distribution closely corresponds the

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distribution of the poles of the deformed state. On the other hand, if the dislocations stored along the plane of growth are pure edge their movement leads to a pure tilt boundary with a rotation axis (of the -type) parallel to the boundary plane, e.g. [9]. It is assumed here that this mechanism is responsible for widening the dispersion of misorientation axes, as clearly visible in Fig. 4d.

Fig. 4. The {111} pole figure (corresponding orientation map of Fig. 3a) showing orientations of deformed/recovered matrix (a). Orientations of new grains, (c) and (d), growing along 1 (111) and 2 - (11-1) plane, respectively, and corresponding (e) misorientation axes distribution. Summary The present experiments provide detailed information about the development of deformation and recrystallization microstructures together with new grain orientations in Al-1%wt.Mn alloy of a stable Goss{110} and brass{110} orientations. Based on SEM/EBSD orientation measurements the orientation relations which appear at the initial stages of recrystallization between recrystallized grain and the as-deformed areas have been analysed and possible mechanism operating during grain growth discussed.  A particular role in the formation of new grains is attributed to the mobility of low angle boundaries and to their ability to form high angle boundaries. The formation of new nuclei is shown to occur by the formation of high angle boundary segments during extended recovery as a consequence of low angle boundary migration.  A large fraction of the new grains, regardless of the crystallite orientation, develop a strongly elongated shape, despite the isotropic distribution of misorientations in the as-deformed state. This anisotropic growth coincides with the situation of the {111} planes in the initial stages of recrystallization. For both initial orientations, the most active planes during strain, located symmetrically to the ND/TD plane, play a decisive role in this process. New grains were strongly elongated with planar facets close to those of dislocation slip planes and the crystal lattice of the grains elongated along a given {111} plane undergoes a general rotation around an axis near to their normal. Acknowledgments This work was partially supported by the National Centre of Science - project no: 2014/13/B/ST8/04291, the Polish Ministry of Science and Higher Education - project no: ‘Iuventus Plus’ IP2011 036471 and the European Union under the European Social Fund project no: POKL.04.01.01-00-004/10.

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