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Development of the “brass” texture component during the hot deformation of Al–6Cu–0.4Zr
Acta Materialia 52 (2004) 4281–4289 www.actamat-journals.com
Development of the ‘‘brass’’ texture component during the hot deformation of Al–6Cu–0.4Zr P.S. Bate *, Y. Huang, F.J. Humphreys Manchester Materials Science Centre, The University of Manchester, Grosvenor Street, Manchester, M1 7HS, UK Received 18 May 2004; received in revised form 27 May 2004; accepted 28 May 2004 Available online 19 June 2004
Abstract Texture and microstructure development during hot plane strain compression of Al–6Cu–0.4Zr has been examined using EBSD. Starting with a random texture in the cast condition, the material developed a typical b-fibre texture during the early stages of plane strain compression at a temperature of 375 °C. At strains higher than about 2, the ‘‘brass’’ texture component, {0 1 1} h2 1 1i, began to dominate the texture. The development of such a texture has been observed previously in aluminium alloys, particularly when fine particles effecting Zener pinning of grain boundaries are present. In the present case, measurements indicate that grain boundary migration during deformation was responsible for the development of the strong ‘‘brass’’ texture. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Texture; Grain growth; Aluminium; Dynamic boundary migration
1. Introduction The crystallographic texture of face centred cubic (FCC) metals following plane strain compression, such as near the mid-plane in flat rolling, is typically concentrated on orientations lying along a fibre – the b-fibre – between the ‘‘brass’’ orientation, {0 1 1} h2 1 1i, and the ‘‘copper’’ orientation, {2 2 5} h5 5 4i. In high stacking fault energy metals such as aluminium, the orientation densities along that fibre are relatively uniform. When the stacking fault energy is low, the ‘‘brass’’ component dominates, especially at relatively high strain: a good example of this was given by Pospiech et al. [1] in silver. There has been a significant volume of research on this effect of stacking fault energy, and several potential mechanisms have been proposed. Most of these relate to the effect of mechanical twinning, either directly [2] or indirectly [3]; the tendency for mechanical twinning being higher in low SFE cubic metals. One significant observation is that an increase in the temperature of *
Corresponding author. Tel.: +44-161-200-8842; fax: +44-161-2003586. E-mail address: [email protected] (P.S. Bate).
deformation in low SFE materials will generally reduce the tendency to form a texture dominated by the ‘‘brass’’ component [4,5], which would be expected as mechanical twinning is generally less significant at higher temperatures. It is interesting, then, that elevated temperature deformation of some aluminium alloys can lead also to strong ‘‘brass’’ textures. This has been mainly observed in aluminium alloys intended for superplastic forming [6], but also in the Al–1Mn–1Mg alloy used for beverage can bodies [7,8]. It is clear that stacking fault energy is highly unlikely to be a factor here, and other mechanisms have been suggested, for example the idea that non-octahedral slip i.e. the operation of slip systems other than {1 1 1} h1 1 0i, could be significant at the elevated temperatures where the strong ‘‘brass’’ texture develops [9,10]. Another possibility is that migration of grain boundaries could be occurring during hot deformation, leading to a change in the texture. It was noted by Higginson and Bate [11] that broad-front strain-induced boundary migration (SIBM) could help explain a transient increase in the ‘‘brass’’ component during recrystallisation of aluminium, and that this might also offer an explanation for the occurrence of the high densities
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of ‘‘brass’’ texture in hot rolled aluminium. Knutsen et al. [12] used the idea of boundary migration during deformation to account for high levels of ‘‘cube’’ texture following hot compression in aluminium. The development of the ‘‘brass’’ texture during hot deformation is much more extreme in the superplastic aluminium alloys, for example that reported by Blackwell and Bate [13], than in the Al–1Mn–1Mg alloy. The superplastic alloys need a fine grain size after processing, and so have a content of fine particles which lead to stabilisation of the grain size via Zener pinning. This may also indicate a role for boundary migration in the evolution of deformation texture, because Zener drag could create a threshold for migration and so increase selectivity of the grain growth. In any case, it is important to investigate this aspect of texture development in one of those superplastic alloys, and Al–6Cu–0.4Zr (‘‘Supral 100’’) was used.
2. Experiments 2.1. Material and deformation The material used in this investigation was a highpurity base Al–6 wt.% Cu–0.4 wt.% Zr alloy, supplied by Alcan International. This is the same nominal composition as the commercial superplastic alloy SupralÒ 100. The material was chill cast, followed by an ageing treatment at 360 °C for 16 h and air-cooling to precipitate fine ZrAl3 particles. The alloy was then solution treated at 500 °C for 2 h and water quenched to take most of the copper into solid solution. Specimens of dimensions 15 mm 10 mm 10 mm were machined from the heat treated material for deformation. Plane strain deformation was carried out on an Instron testing machine at a temperature of 375 °C with a constant strain rate of 103 s1 to various strains using a stainless steel channel die. The convention used for the geometry of flat rolling is adopted here, with ND being the compression direction, RD the extension direction and TD the direction constrained by the channel. During deformation, the specimen was surrounded with thin PTFE sheet to provide lubrication. A two-stage deformation procedure was used to produce large strains. In this procedure, specimens were first deformed to a reduction of about 85% (i.e. to an thickness of about 2.3 mm) and were cut into pieces of dimensions 2.3 mm 10 mm 10 mm. After surface dressing, 6 of these deformed pieces were stacked on top of each other, forming an equivalent specimen of dimensions 13.8 mm 10 mm 10 mm, and were then deformed in the channel die under the same conditions as those used in the previous stage to produce further reductions of up to 85%. Using this two-stage deformation procedure, a total reduction of 97.5% – a total strain of about 3.7 –
was obtained. Previous investigations have shown that this hot deformed alloy is resistant to recrystallisation and static coarsening, and therefore a two stage deformation such as described above is valid. All specimens were quenched in water within 5 s of the end of deformation. 2.2. Metallography The deformed specimens were sectioned normal to TD through the centre of the specimen and were then mechanically polished and electro-polished ( in a solution of 30% nitric acid in methanol at 12 V and at )30 °C for 75 s) for microstructural characterization and texture analysis. The specimens were examined using a CAMSCAN Maxim 2040 FEGSEM, by backscattered electron imaging and by electron backscattered diffraction (EBSD). EBSD orientation maps were obtained using an HKL Channel acquisition system and processed using VMAP – in-house software developed for microstructural characterization. The as-cast structure had near-equiaxed grains with a mean intercept size of 45 lm. Micrographs of deformed material are shown in Fig. 1. The development of microstructure is superficially very simple. The original grains become elongated in the RD, compressed in ND and there is the generation of substructure within the grains, manifest as low angle boundaries. Some of those substructural boundaries will accumulate high misorientations during deformation, leading to a reduction in grain size, but any gross boundary migration will have the opposite effect. There is clear evidence that spheroidisation is occurring in the high strain microstructures, indicating that boundary migration is occurring. The most convenient way to assess those types of effects is to measure mean linear intercept values between high angle boundaries in the ND (KND ). In the absence of boundary generation or gross migration, the intercepts should reduce in line with the macroscopic reduction in height of the specimens. These results are shown in Fig. 2, where the intercepts, defined by boundaries with misorientation angles >15°, are shown for the main texture components. In all components, KND decreases at a rate greater than that due to the simple geometric change at strains less than about 2. This is presumably because of the generation of new high angle boundaries by grain fragmentation. At higher strains, KND becomes greater than the geometrical value, indicating significant grain growth. This growth was much greater for ‘‘brass’’ oriented grains than for those of other texture components. 2.3. Texture The crystallographic texture was calculated from EBSD measurements made using stage scanning. The scan step size was 20 lm for the as-cast material, and
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Fig. 1. EBSD micrographs (orientation maps) of material compressed at 375 °C to the strains indicated. The grey scales are based on the orientation, with high angle (x > 15°) boundaries delineated in black and other, low angle (x > 1:5°) in pale grey.
Fig. 2. The evolution of boundary spacings in the compression direction with strain. The standard error estimates for these data are about the size of the symbols.
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was reduced with the specimen strain to 5 lm for the highest level of compression. This ensures efficient sampling, without a high level of redundancy involving multiple measurements of one grain (or sub-grain). Be-
tween 3 104 and 1.25 105 orientations were measured on each specimen. Continuous orientation distribution functions (ODFs) were derived from the discrete orientation data using harmonic series, with series truncation
Fig. 3. Sections at u2 ¼ 45° of the orientation distribution functions of material following plane strain compression to the strains indicated, at a temperature of 375 °C. These show the ‘‘copper’’, at u1 90° and U 30°, and the ‘‘brass’’, at u1 54° and U 90°, texture components.
Fig. 4. Densities along the b-fibre for material compressed at 375 °C for different strains. The orientation densities are the maxima in the vicinity on the ideal fibre orientations at the u2 values i.e. the exact location of the fibre in {u1 , U} is not pre-assumed.
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Fig. 5. Orientation densities of the principal ideal texture components as functions of strain for material compressed at 375 °C. Maximum orientation densities in the vicinity of the ideal orientations at constant u2 were used, with u2 ¼ 67:5° for the ‘‘S’’ component.
at lmax ¼ 22, and assuming orthotropic specimen symmetry. The texture of the as-cast material was essentially random. The textures of deformed material could all be described in terms of components lying along the b-fibre typical of FCC metals deformed in plane strain. Representative sections through the ODFs of compressed material are given in Fig. 3, and orientation densities along the b-fibre are shown in Fig. 4. It is clear that at strains up to about 2 there is a reasonably uniform orientation density along the b-fibre and the level of the orientation density of that fibre is increasing. At higher strain, however, the ‘‘brass’’ component increases in density and there is an absolute reduction in the ‘‘copper’’ density and, after a strain of 3.5, a reduction in the density of the intermediate b-fibre orientations ( typified by an ‘‘S’’ orientation). After a strain of 3.7, about 60% of the material lies within 15° of the ideal ‘‘brass’’ orientations and this is the only significant component of the texture. This situation is further demonstrated in Fig. 5, where the orientation densities at the three representative orientations are shown as functions of strain.
3. Discussion The evolution of texture, in terms of the relative densities of the ‘‘copper’’ and ‘‘brass’’ components, with strain is similar to that reported by Bate and Oscarsson [7] in Al–1Mn–1Mg, although the orientation density of
the ‘‘brass’’ component at high strains was much greater in the present case. This texture evolution was paralleled by changes in the spacings of high angle grain boundaries (HAGB) in the compression direction (Fig. 2) which strongly suggest that migration of those boundaries is occurring during the hot deformation. The changes in HAGB are also consistent with the relative increase in ‘‘brass’’ texture: the spacing of boundaries of grains of that texture component became significantly greater than those of other b-fibre orientations. 3.1. Crystal plasticity Although the evolution in texture of low SFE metals at low temperatures with strain proceeds in a nominally similar manner – i.e. a b-fibre forms at relatively small strains and becomes dominated by the ‘‘brass’’ component with increasing deformation – the mechanisms proposed for that low temperature, low SFE, texture can reasonably be discounted for explaining the present results. There is no evidence of deformation twinning, nor of any other intragranular shearing, and in any case there is little likelihood of the SFE of the aluminium alloy being low. While non-octahedral glide cannot be discounted in aluminium, there has been criticism of its use for explaining the evolution of deformation texture during high temperature deformation of aluminium, for example by Leffers [14]. The immediate problem is that significant activity of non-octahedral systems must only
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occur after a certain level of strain to explain the present results, and it is not easy to see why that might happen. Slip activity on a particular system could be affected by the distance between boundaries on that system, and if this were a factor it would become more significant as the boundary spacing became small. Irrespective of what classes of slip systems are present, however, most slip occurs on systems with quite similar spatial orientations and so would have similar slip lengths between the high angle boundaries. It seems unlikely that a model based on simple crystal plasticity could account for the texture changes. The metallography reported above shows that the increase in the ‘‘brass’’ texture component is associated with the growth of ‘‘brass’’ grains. The grain size, in terms of dimensions in the ND, increases relative to the deformed initial grain size at strains greater than about 2, which is the strain at which the increase in ‘‘brass’’ texture becomes noticeable. There are questions about whether this effect is due to boundary migration, and if it is, what causes this to occur. 3.2. Orientation coalescence The possibility that the relative increase in grain size in the compression direction might be caused by the formation of a strong ‘‘brass’’ texture, rather than grain growth leading to that texture, needs to be considered. In an extreme case, the formation of a very strong texture can mean that many boundaries would develop low misorientations, giving an apparent increase in grain
size. To assess this effect, misorientation distribution functions (MODFs) were calculated from textures on the basis of a random likelihood of orientations adjoining: so-called uncorrelated MODFs. These are shown in Fig. 6 for a random texture and for the textures measured after strains of 1.0 and 3.5. It is clear that while the fraction of boundaries with misorientation angles less than 15° increases with texture development, it is still actually rather small. After a strain of 1.0, about 95% of uncorrelated misorientations are high angle, and after a strain of 3.5, this only reduces to 90%. The effect is, then, unlikely to be significant. There is a noticeable increase in the fraction of boundaries near x ¼ 60° when the ‘‘brass’’ texture begins to dominate: this is simply a consequence of the two variants of ‘‘brass’’ which occur, when sheet (orthotropic) symmetry is involved, being related by a crystal rotation of 60° about h1 1 1i. 3.3. Boundary migration The growth of ‘‘brass’’ oriented grains appears, then to be a real effect. Different substructural densities in different grain orientations are fundamental to the change in texture associated with this grain growth. The ‘‘brass’’ oriented material must have a lower substructural energy density than the other main components for the texture change to occur by dynamic boundary migration. This type of orientation dependence is well known in a-Fe [15,16], where the energy density varies with the Taylor factor, which can be interpreted as the
Fig. 6. Uncorrelated misorientation distribution functions, as functions of the misorientation angle x only, for two experimental textures and a random texture. These were calculated from 5000 discrete orientations i.e. about 1.2 107 misorientations. MODFs calculated using the harmonic coefficients were very similar to these ones.
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extent of slip activity relative to the overall strain rate. It has been suggested [17] that this correlates with an increased tendency for inhomogeneous deformation in high Taylor factor (M) orientations. The Taylor factors for fully constrained plane strain deformation (MFC ) of b-fibre orientations are shown in Fig. 7, along with the ratio of those full constraint factors to ones where RD– ND shear has been relaxed, MRC . That ratio will give some measure of the tendency for shear inhomogeneity. Both MFC and the ratio decrease going from ‘‘copper’’ to ‘‘brass’’ along the fibre, and if the situation occurring in steel also applies to this aluminium alloy then there will be a small, but definite, difference in substructural energy which can drive the preferential growth of ‘‘brass’’ oriented grains. There is also metallographic evidence for that energy difference. The high angle boundaries aligned normal to the ND will migrate in an approximately planar manner. Such migration of a planar high-angle boundary between two grains, A and B, as a result of a difference in average substructural energy content and in the presence of a Zener drag is determined by: CA CB fv ð1Þ k k >ar ; A
B
where C are the mean sub-boundary energies relative to that for a high angle boundary and k are the subgrain boundary mean linear intercepts. The Zener drag term depends on mean particle radius, r, and volume frac-
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tion, fv , with a a numerical factor usually taken to be 0.75 in the present context (i.e. pressure divided by twice the high angle boundary tension). Following previous work [18], reasonable values for the radius and volume fraction of ZrAl3 dispersoids in the material are 10 nm and 0.007, giving a value for the righthand side of the inequality given above of about 0.5 lm1 . The substructural energy densities of different texture components were calculated from EBSD data, using the Read–Shockley equation for relative boundary energy. The values for ‘‘copper’’ and ‘‘S’’ were similar, and were consistently higher than those for ‘‘brass’’. That difference, however, was rather small, about 0.01–0.03 lm1 . Values given by Theyssier and Driver [19] for aluminium deformed at 400 °C translate to a slightly higher figure, about 0.06 lm1 . These are about an order of magnitude less than required to drive the type of boundary migration necessary to increase the ‘‘brass’’ texture. There will be a population of subgrain boundaries with very low misorientations, less than the cut-off used for the EBSD data, but this is unlikely to increase the overall driving pressures by a factor of more than about 2. 3.4. Vertex modelling Despite the strong indications from the metallographic measurements that differential grain growth is occurring, the mean substructural energy difference
Fig. 7. The full constraint Taylor factor, MFC , for plane strain and the ratio of full constraint to relaxed constraint Taylor factors, MFC =MRC , as functions of location on the b-fibre. The fibre orientations were taken from the texture of material deformed to e ¼ 2, and the form of M vs. u2 was very similar to that using M calculated with 10° wide Gaussian spreads about the individual orientations.
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derived from EBSD data does not appear to be sufficient to drive the required boundary migration. That simple analysis ignores some important issues, however. The first of these is that it assumes uniformity of the substructure in the bands, effectively ignoring the discrete nature of that substructure. The significance of this can be demonstrated using a simple two-dimensional ’vertex’ model of the grain growth. The model used here has been shown to give similar results to three-dimensional Monte Carlo–Potts modelling of the evolution of lamellar microstructures resulting from severe deformation [20]. A uniform Zener drag can be included [21], which is reasonable with the fine pinning particles in the present material. A structure was set up comprising alternating bands of two distinct orientations, and as well as the orientation difference these had different densities of transverse low-angle boundaries, which together with the subgrain orientation distributions involved – and using the Read–Shockley formula for boundary energy – gave mean values of C=k of 9.9 and 12.5 (in terms of the nominal length unit of the model). The mean driving pressure was, then, 2.6. The evolution of that model structure with a Zener pressure of 10 is shown in Fig. 8. There has been significant HAGB migration, and a corresponding increase in the fraction containing the lower substructural energy, despite the mean driving pressure being much less than the Zener pressure. This is due to the discrete nature and the inhomogeneity of the substructure, which was perturbed at random by a maximum of 0.2 of its mean geometrical values in the start structure. Significant band migration occurred with this type of model structure up to a Zener pressure
about five times the mean driving pressure. Greater inhomogeneity of the substructure will lead to growth at even greater ratios of Zener pressure to mean driving pressure, though it is difficult to estimate the degree of that inhomogeneity in the real structure. 3.5. Dynamic effects Another important factor in the structural development is that it occurs during deformation. Very little change of microstructure or texture occurred during the static annealing, at 375 °C for 20 h, of the deformed material. The simplest effect of deformation is that it will change the geometry of the microstructure, and this can have an effect on grain growth especially in Zener pinned systems [21]. An example of this effect is shown in Fig. 9, which shows the result of a vertex model simulation similar to that shown in Fig. 8 but with overall geometry changes in line with a uniform plane strain deformation of the domain. There is noticeable grain growth, but the Zener pressure here was 15 and the fact that no new substructure was introduced means that the driving pressure has been effectively reduced. Growth is occurring in this simulation, then, with a mean driving pressure that is only about one tenth of the Zener pressure. This is in line with the experimental estimates. There are other possible effects of deformation on the grain growth. Dynamic particle coarsening has been observed in some superplastic materials and would reduce the Zener drag, but there is no evidence that this is significant with ZrAl3 dispersoids at 375 °C. Of more interest are the observations of Jazaeri and Humphreys
Fig. 8. Vertex model microstructures from a simulation of the static annealing of a banded microstructure. This is driven by a substructural pressure difference of about one quarter of the Zener pinning pressure. Both the grains and subgrains and the boundaries have grey scales corresponding to misorientation: the mean subgrain misorientation was slightly less than 2°. The light grey material has the lower substructural energy.
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Fig. 9. Vertex model microstructure from a simulation of the annealing of a banded microstructure with the geometry incrementally modified to represent the effect of a homogeneous deformation. The starting structure was the same as that shown in Fig. 8, and the total strain here is about 0.6.
[22], who found what is effectively dynamic grain growth in aluminium alloys deformed at room temperature i.e. under conditions where static annealing would be expected to be negligible. By rolling alloys with a variety of initial grain sizes, they found that there appears to be a minimum KND at about the level of the subgrain size characteristic of deformation in those alloys under the conditions of temperature and strain rate involved – typically about 0.5 lm. The rate of conventional boundary migration is insufficient to explain that result, and they suggest that the boundaries are involved in the dynamic recovery process, with the local stresses occurring during deformation perhaps aiding the grain growth. There has been work indicating that stress can influence the migration of low angle boundaries [23,24], and relatively low levels of stress can even influence the migration of high angle boundaries [25], although the general situation is far from clear. This type of effect could be contributing to the dynamic grain growth occurring in the high temperature deformation reported here, where KND is also close to the typical subgrain size expected at the temperature and strain rate used. However, boundary mobilities at 375 °C and a relatively low strain rate are unlikely to be a limiting factor in the present case.
4. Conclusions There is a transition from a texture which has an almost uniformly populated b-fibre at moderate strains to one dominated by the brass component at high strains in the hot plane strain compression of Al–6Cu– 0.4Zr. This is associated with differential dynamic grain growth, which is principally driven by different substructural energy densities in different grain orientations. The presence of Zener pinning will make this process more selective. Simple geometric effects of deformation, together with a possible influence of dynamic recovery involving the high angle boundaries when their spacing becomes small, enhance this growth.
Acknowledgements This work was partly supported by EPSRC via Grants GR/R52695 and GR/R69952. We are most grateful to Alcan International for the supply of material.
References [1] Pospiech J, Truszkowski W, Jura J, Krol J. In: Davies GJ, Dillamore IL, Hudd RC, Kallend JS, editors. Texture and the properties of materials (ICOTOM-4). 1976. p. 23. [2] Wassermann G. Z Metallk 1963;54:61. [3] Hutchinson WB, Duggan BJ, Hatherly M. Met Technol 1979;9:398. [4] Hu H, Cline RS. J Appl Phys 1961;32:760. [5] Goodman SR, Hu H. Trans AIME 1964;230:1413. [6] Bowen AW. Mater Sci Technol 1990;6:1058. [7] Bate P, Oscarsson A. Mater Sci Technol 1990;6:520. [8] Maurice C, Driver JH. Mater Sci Forum 1994;157–162:807. [9] Maurice C, Driver JH. Acta Mater 1997;45:4627. [10] Maurice C, Driver JH. Acta Mater 1997;45:4639. [11] Higginson R, Bate P. Acta Mater 1999;47:1079. [12] Knutsen RD, Humphreys FJ, Bate PS, Court SA. In: Szpunar JA, editor. Proceedings of the 12th international conference on textures of materials (ICOTOM-12). Ottawa, Canada: NRC Research Press; 1999. p. 617. [13] Blackwell PL, Bate PS. Met Trans 1993;24A:1085. [14] Leffers T. In: Andersen SI, Bilde-Sørensen JB, Lorentzen T, Pedersen OB, Sørensen NJ, editors. 15th Risø international symposium on materials science. Roskilde, Denmark: Risø National Laboratory; 1994. p. 387. [15] Dillamore IL, Smith CJE, Watson TW. Met Sci 1967;1:49. [16] Every RL, Hatherly M. Texture 1974;4:183. [17] Hutchinson B. Philos Trans R Soc Lond A 1999;357:1471. [18] Ridley N, Cullen EH, Humphreys FJ. Mater Sci Technol 2000;16:117. [19] Theyssier MC, Driver JH. Mat Sci Eng A 1999;272:73. [20] Prangnell PB, Hayes JS, Bowen JR, Apps PJ, Bate PS. Acta Mater 2004;52:3193. [21] Bate P. Acta Mater 2001;49:1453. [22] Jazaeri H, Humphreys FJ. Acta Mater 2004;52:3239. [23] Exell SF, Warrington DH. Philos Mag 1972;26:1121. [24] Biburger M, Blum W. Philos Mag A 1992;66:27. [25] Winning M, Gottstein G, Shvindlerman LS. Acta Mater 2002;50:353.
Development of the ‘‘brass’’ texture component during the hot deformation of Al–6Cu–0.4Zr P.S. Bate *, Y. Huang, F.J. Humphreys Manchester Materials Science Centre, The University of Manchester, Grosvenor Street, Manchester, M1 7HS, UK Received 18 May 2004; received in revised form 27 May 2004; accepted 28 May 2004 Available online 19 June 2004
Abstract Texture and microstructure development during hot plane strain compression of Al–6Cu–0.4Zr has been examined using EBSD. Starting with a random texture in the cast condition, the material developed a typical b-fibre texture during the early stages of plane strain compression at a temperature of 375 °C. At strains higher than about 2, the ‘‘brass’’ texture component, {0 1 1} h2 1 1i, began to dominate the texture. The development of such a texture has been observed previously in aluminium alloys, particularly when fine particles effecting Zener pinning of grain boundaries are present. In the present case, measurements indicate that grain boundary migration during deformation was responsible for the development of the strong ‘‘brass’’ texture. Ó 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Texture; Grain growth; Aluminium; Dynamic boundary migration
1. Introduction The crystallographic texture of face centred cubic (FCC) metals following plane strain compression, such as near the mid-plane in flat rolling, is typically concentrated on orientations lying along a fibre – the b-fibre – between the ‘‘brass’’ orientation, {0 1 1} h2 1 1i, and the ‘‘copper’’ orientation, {2 2 5} h5 5 4i. In high stacking fault energy metals such as aluminium, the orientation densities along that fibre are relatively uniform. When the stacking fault energy is low, the ‘‘brass’’ component dominates, especially at relatively high strain: a good example of this was given by Pospiech et al. [1] in silver. There has been a significant volume of research on this effect of stacking fault energy, and several potential mechanisms have been proposed. Most of these relate to the effect of mechanical twinning, either directly [2] or indirectly [3]; the tendency for mechanical twinning being higher in low SFE cubic metals. One significant observation is that an increase in the temperature of *
Corresponding author. Tel.: +44-161-200-8842; fax: +44-161-2003586. E-mail address: [email protected] (P.S. Bate).
deformation in low SFE materials will generally reduce the tendency to form a texture dominated by the ‘‘brass’’ component [4,5], which would be expected as mechanical twinning is generally less significant at higher temperatures. It is interesting, then, that elevated temperature deformation of some aluminium alloys can lead also to strong ‘‘brass’’ textures. This has been mainly observed in aluminium alloys intended for superplastic forming [6], but also in the Al–1Mn–1Mg alloy used for beverage can bodies [7,8]. It is clear that stacking fault energy is highly unlikely to be a factor here, and other mechanisms have been suggested, for example the idea that non-octahedral slip i.e. the operation of slip systems other than {1 1 1} h1 1 0i, could be significant at the elevated temperatures where the strong ‘‘brass’’ texture develops [9,10]. Another possibility is that migration of grain boundaries could be occurring during hot deformation, leading to a change in the texture. It was noted by Higginson and Bate [11] that broad-front strain-induced boundary migration (SIBM) could help explain a transient increase in the ‘‘brass’’ component during recrystallisation of aluminium, and that this might also offer an explanation for the occurrence of the high densities
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of ‘‘brass’’ texture in hot rolled aluminium. Knutsen et al. [12] used the idea of boundary migration during deformation to account for high levels of ‘‘cube’’ texture following hot compression in aluminium. The development of the ‘‘brass’’ texture during hot deformation is much more extreme in the superplastic aluminium alloys, for example that reported by Blackwell and Bate [13], than in the Al–1Mn–1Mg alloy. The superplastic alloys need a fine grain size after processing, and so have a content of fine particles which lead to stabilisation of the grain size via Zener pinning. This may also indicate a role for boundary migration in the evolution of deformation texture, because Zener drag could create a threshold for migration and so increase selectivity of the grain growth. In any case, it is important to investigate this aspect of texture development in one of those superplastic alloys, and Al–6Cu–0.4Zr (‘‘Supral 100’’) was used.
2. Experiments 2.1. Material and deformation The material used in this investigation was a highpurity base Al–6 wt.% Cu–0.4 wt.% Zr alloy, supplied by Alcan International. This is the same nominal composition as the commercial superplastic alloy SupralÒ 100. The material was chill cast, followed by an ageing treatment at 360 °C for 16 h and air-cooling to precipitate fine ZrAl3 particles. The alloy was then solution treated at 500 °C for 2 h and water quenched to take most of the copper into solid solution. Specimens of dimensions 15 mm 10 mm 10 mm were machined from the heat treated material for deformation. Plane strain deformation was carried out on an Instron testing machine at a temperature of 375 °C with a constant strain rate of 103 s1 to various strains using a stainless steel channel die. The convention used for the geometry of flat rolling is adopted here, with ND being the compression direction, RD the extension direction and TD the direction constrained by the channel. During deformation, the specimen was surrounded with thin PTFE sheet to provide lubrication. A two-stage deformation procedure was used to produce large strains. In this procedure, specimens were first deformed to a reduction of about 85% (i.e. to an thickness of about 2.3 mm) and were cut into pieces of dimensions 2.3 mm 10 mm 10 mm. After surface dressing, 6 of these deformed pieces were stacked on top of each other, forming an equivalent specimen of dimensions 13.8 mm 10 mm 10 mm, and were then deformed in the channel die under the same conditions as those used in the previous stage to produce further reductions of up to 85%. Using this two-stage deformation procedure, a total reduction of 97.5% – a total strain of about 3.7 –
was obtained. Previous investigations have shown that this hot deformed alloy is resistant to recrystallisation and static coarsening, and therefore a two stage deformation such as described above is valid. All specimens were quenched in water within 5 s of the end of deformation. 2.2. Metallography The deformed specimens were sectioned normal to TD through the centre of the specimen and were then mechanically polished and electro-polished ( in a solution of 30% nitric acid in methanol at 12 V and at )30 °C for 75 s) for microstructural characterization and texture analysis. The specimens were examined using a CAMSCAN Maxim 2040 FEGSEM, by backscattered electron imaging and by electron backscattered diffraction (EBSD). EBSD orientation maps were obtained using an HKL Channel acquisition system and processed using VMAP – in-house software developed for microstructural characterization. The as-cast structure had near-equiaxed grains with a mean intercept size of 45 lm. Micrographs of deformed material are shown in Fig. 1. The development of microstructure is superficially very simple. The original grains become elongated in the RD, compressed in ND and there is the generation of substructure within the grains, manifest as low angle boundaries. Some of those substructural boundaries will accumulate high misorientations during deformation, leading to a reduction in grain size, but any gross boundary migration will have the opposite effect. There is clear evidence that spheroidisation is occurring in the high strain microstructures, indicating that boundary migration is occurring. The most convenient way to assess those types of effects is to measure mean linear intercept values between high angle boundaries in the ND (KND ). In the absence of boundary generation or gross migration, the intercepts should reduce in line with the macroscopic reduction in height of the specimens. These results are shown in Fig. 2, where the intercepts, defined by boundaries with misorientation angles >15°, are shown for the main texture components. In all components, KND decreases at a rate greater than that due to the simple geometric change at strains less than about 2. This is presumably because of the generation of new high angle boundaries by grain fragmentation. At higher strains, KND becomes greater than the geometrical value, indicating significant grain growth. This growth was much greater for ‘‘brass’’ oriented grains than for those of other texture components. 2.3. Texture The crystallographic texture was calculated from EBSD measurements made using stage scanning. The scan step size was 20 lm for the as-cast material, and
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Fig. 1. EBSD micrographs (orientation maps) of material compressed at 375 °C to the strains indicated. The grey scales are based on the orientation, with high angle (x > 15°) boundaries delineated in black and other, low angle (x > 1:5°) in pale grey.
Fig. 2. The evolution of boundary spacings in the compression direction with strain. The standard error estimates for these data are about the size of the symbols.
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was reduced with the specimen strain to 5 lm for the highest level of compression. This ensures efficient sampling, without a high level of redundancy involving multiple measurements of one grain (or sub-grain). Be-
tween 3 104 and 1.25 105 orientations were measured on each specimen. Continuous orientation distribution functions (ODFs) were derived from the discrete orientation data using harmonic series, with series truncation
Fig. 3. Sections at u2 ¼ 45° of the orientation distribution functions of material following plane strain compression to the strains indicated, at a temperature of 375 °C. These show the ‘‘copper’’, at u1 90° and U 30°, and the ‘‘brass’’, at u1 54° and U 90°, texture components.
Fig. 4. Densities along the b-fibre for material compressed at 375 °C for different strains. The orientation densities are the maxima in the vicinity on the ideal fibre orientations at the u2 values i.e. the exact location of the fibre in {u1 , U} is not pre-assumed.
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Fig. 5. Orientation densities of the principal ideal texture components as functions of strain for material compressed at 375 °C. Maximum orientation densities in the vicinity of the ideal orientations at constant u2 were used, with u2 ¼ 67:5° for the ‘‘S’’ component.
at lmax ¼ 22, and assuming orthotropic specimen symmetry. The texture of the as-cast material was essentially random. The textures of deformed material could all be described in terms of components lying along the b-fibre typical of FCC metals deformed in plane strain. Representative sections through the ODFs of compressed material are given in Fig. 3, and orientation densities along the b-fibre are shown in Fig. 4. It is clear that at strains up to about 2 there is a reasonably uniform orientation density along the b-fibre and the level of the orientation density of that fibre is increasing. At higher strain, however, the ‘‘brass’’ component increases in density and there is an absolute reduction in the ‘‘copper’’ density and, after a strain of 3.5, a reduction in the density of the intermediate b-fibre orientations ( typified by an ‘‘S’’ orientation). After a strain of 3.7, about 60% of the material lies within 15° of the ideal ‘‘brass’’ orientations and this is the only significant component of the texture. This situation is further demonstrated in Fig. 5, where the orientation densities at the three representative orientations are shown as functions of strain.
3. Discussion The evolution of texture, in terms of the relative densities of the ‘‘copper’’ and ‘‘brass’’ components, with strain is similar to that reported by Bate and Oscarsson [7] in Al–1Mn–1Mg, although the orientation density of
the ‘‘brass’’ component at high strains was much greater in the present case. This texture evolution was paralleled by changes in the spacings of high angle grain boundaries (HAGB) in the compression direction (Fig. 2) which strongly suggest that migration of those boundaries is occurring during the hot deformation. The changes in HAGB are also consistent with the relative increase in ‘‘brass’’ texture: the spacing of boundaries of grains of that texture component became significantly greater than those of other b-fibre orientations. 3.1. Crystal plasticity Although the evolution in texture of low SFE metals at low temperatures with strain proceeds in a nominally similar manner – i.e. a b-fibre forms at relatively small strains and becomes dominated by the ‘‘brass’’ component with increasing deformation – the mechanisms proposed for that low temperature, low SFE, texture can reasonably be discounted for explaining the present results. There is no evidence of deformation twinning, nor of any other intragranular shearing, and in any case there is little likelihood of the SFE of the aluminium alloy being low. While non-octahedral glide cannot be discounted in aluminium, there has been criticism of its use for explaining the evolution of deformation texture during high temperature deformation of aluminium, for example by Leffers [14]. The immediate problem is that significant activity of non-octahedral systems must only
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occur after a certain level of strain to explain the present results, and it is not easy to see why that might happen. Slip activity on a particular system could be affected by the distance between boundaries on that system, and if this were a factor it would become more significant as the boundary spacing became small. Irrespective of what classes of slip systems are present, however, most slip occurs on systems with quite similar spatial orientations and so would have similar slip lengths between the high angle boundaries. It seems unlikely that a model based on simple crystal plasticity could account for the texture changes. The metallography reported above shows that the increase in the ‘‘brass’’ texture component is associated with the growth of ‘‘brass’’ grains. The grain size, in terms of dimensions in the ND, increases relative to the deformed initial grain size at strains greater than about 2, which is the strain at which the increase in ‘‘brass’’ texture becomes noticeable. There are questions about whether this effect is due to boundary migration, and if it is, what causes this to occur. 3.2. Orientation coalescence The possibility that the relative increase in grain size in the compression direction might be caused by the formation of a strong ‘‘brass’’ texture, rather than grain growth leading to that texture, needs to be considered. In an extreme case, the formation of a very strong texture can mean that many boundaries would develop low misorientations, giving an apparent increase in grain
size. To assess this effect, misorientation distribution functions (MODFs) were calculated from textures on the basis of a random likelihood of orientations adjoining: so-called uncorrelated MODFs. These are shown in Fig. 6 for a random texture and for the textures measured after strains of 1.0 and 3.5. It is clear that while the fraction of boundaries with misorientation angles less than 15° increases with texture development, it is still actually rather small. After a strain of 1.0, about 95% of uncorrelated misorientations are high angle, and after a strain of 3.5, this only reduces to 90%. The effect is, then, unlikely to be significant. There is a noticeable increase in the fraction of boundaries near x ¼ 60° when the ‘‘brass’’ texture begins to dominate: this is simply a consequence of the two variants of ‘‘brass’’ which occur, when sheet (orthotropic) symmetry is involved, being related by a crystal rotation of 60° about h1 1 1i. 3.3. Boundary migration The growth of ‘‘brass’’ oriented grains appears, then to be a real effect. Different substructural densities in different grain orientations are fundamental to the change in texture associated with this grain growth. The ‘‘brass’’ oriented material must have a lower substructural energy density than the other main components for the texture change to occur by dynamic boundary migration. This type of orientation dependence is well known in a-Fe [15,16], where the energy density varies with the Taylor factor, which can be interpreted as the
Fig. 6. Uncorrelated misorientation distribution functions, as functions of the misorientation angle x only, for two experimental textures and a random texture. These were calculated from 5000 discrete orientations i.e. about 1.2 107 misorientations. MODFs calculated using the harmonic coefficients were very similar to these ones.
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extent of slip activity relative to the overall strain rate. It has been suggested [17] that this correlates with an increased tendency for inhomogeneous deformation in high Taylor factor (M) orientations. The Taylor factors for fully constrained plane strain deformation (MFC ) of b-fibre orientations are shown in Fig. 7, along with the ratio of those full constraint factors to ones where RD– ND shear has been relaxed, MRC . That ratio will give some measure of the tendency for shear inhomogeneity. Both MFC and the ratio decrease going from ‘‘copper’’ to ‘‘brass’’ along the fibre, and if the situation occurring in steel also applies to this aluminium alloy then there will be a small, but definite, difference in substructural energy which can drive the preferential growth of ‘‘brass’’ oriented grains. There is also metallographic evidence for that energy difference. The high angle boundaries aligned normal to the ND will migrate in an approximately planar manner. Such migration of a planar high-angle boundary between two grains, A and B, as a result of a difference in average substructural energy content and in the presence of a Zener drag is determined by: CA CB fv ð1Þ k k >ar ; A
B
where C are the mean sub-boundary energies relative to that for a high angle boundary and k are the subgrain boundary mean linear intercepts. The Zener drag term depends on mean particle radius, r, and volume frac-
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tion, fv , with a a numerical factor usually taken to be 0.75 in the present context (i.e. pressure divided by twice the high angle boundary tension). Following previous work [18], reasonable values for the radius and volume fraction of ZrAl3 dispersoids in the material are 10 nm and 0.007, giving a value for the righthand side of the inequality given above of about 0.5 lm1 . The substructural energy densities of different texture components were calculated from EBSD data, using the Read–Shockley equation for relative boundary energy. The values for ‘‘copper’’ and ‘‘S’’ were similar, and were consistently higher than those for ‘‘brass’’. That difference, however, was rather small, about 0.01–0.03 lm1 . Values given by Theyssier and Driver [19] for aluminium deformed at 400 °C translate to a slightly higher figure, about 0.06 lm1 . These are about an order of magnitude less than required to drive the type of boundary migration necessary to increase the ‘‘brass’’ texture. There will be a population of subgrain boundaries with very low misorientations, less than the cut-off used for the EBSD data, but this is unlikely to increase the overall driving pressures by a factor of more than about 2. 3.4. Vertex modelling Despite the strong indications from the metallographic measurements that differential grain growth is occurring, the mean substructural energy difference
Fig. 7. The full constraint Taylor factor, MFC , for plane strain and the ratio of full constraint to relaxed constraint Taylor factors, MFC =MRC , as functions of location on the b-fibre. The fibre orientations were taken from the texture of material deformed to e ¼ 2, and the form of M vs. u2 was very similar to that using M calculated with 10° wide Gaussian spreads about the individual orientations.
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derived from EBSD data does not appear to be sufficient to drive the required boundary migration. That simple analysis ignores some important issues, however. The first of these is that it assumes uniformity of the substructure in the bands, effectively ignoring the discrete nature of that substructure. The significance of this can be demonstrated using a simple two-dimensional ’vertex’ model of the grain growth. The model used here has been shown to give similar results to three-dimensional Monte Carlo–Potts modelling of the evolution of lamellar microstructures resulting from severe deformation [20]. A uniform Zener drag can be included [21], which is reasonable with the fine pinning particles in the present material. A structure was set up comprising alternating bands of two distinct orientations, and as well as the orientation difference these had different densities of transverse low-angle boundaries, which together with the subgrain orientation distributions involved – and using the Read–Shockley formula for boundary energy – gave mean values of C=k of 9.9 and 12.5 (in terms of the nominal length unit of the model). The mean driving pressure was, then, 2.6. The evolution of that model structure with a Zener pressure of 10 is shown in Fig. 8. There has been significant HAGB migration, and a corresponding increase in the fraction containing the lower substructural energy, despite the mean driving pressure being much less than the Zener pressure. This is due to the discrete nature and the inhomogeneity of the substructure, which was perturbed at random by a maximum of 0.2 of its mean geometrical values in the start structure. Significant band migration occurred with this type of model structure up to a Zener pressure
about five times the mean driving pressure. Greater inhomogeneity of the substructure will lead to growth at even greater ratios of Zener pressure to mean driving pressure, though it is difficult to estimate the degree of that inhomogeneity in the real structure. 3.5. Dynamic effects Another important factor in the structural development is that it occurs during deformation. Very little change of microstructure or texture occurred during the static annealing, at 375 °C for 20 h, of the deformed material. The simplest effect of deformation is that it will change the geometry of the microstructure, and this can have an effect on grain growth especially in Zener pinned systems [21]. An example of this effect is shown in Fig. 9, which shows the result of a vertex model simulation similar to that shown in Fig. 8 but with overall geometry changes in line with a uniform plane strain deformation of the domain. There is noticeable grain growth, but the Zener pressure here was 15 and the fact that no new substructure was introduced means that the driving pressure has been effectively reduced. Growth is occurring in this simulation, then, with a mean driving pressure that is only about one tenth of the Zener pressure. This is in line with the experimental estimates. There are other possible effects of deformation on the grain growth. Dynamic particle coarsening has been observed in some superplastic materials and would reduce the Zener drag, but there is no evidence that this is significant with ZrAl3 dispersoids at 375 °C. Of more interest are the observations of Jazaeri and Humphreys
Fig. 8. Vertex model microstructures from a simulation of the static annealing of a banded microstructure. This is driven by a substructural pressure difference of about one quarter of the Zener pinning pressure. Both the grains and subgrains and the boundaries have grey scales corresponding to misorientation: the mean subgrain misorientation was slightly less than 2°. The light grey material has the lower substructural energy.
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Fig. 9. Vertex model microstructure from a simulation of the annealing of a banded microstructure with the geometry incrementally modified to represent the effect of a homogeneous deformation. The starting structure was the same as that shown in Fig. 8, and the total strain here is about 0.6.
[22], who found what is effectively dynamic grain growth in aluminium alloys deformed at room temperature i.e. under conditions where static annealing would be expected to be negligible. By rolling alloys with a variety of initial grain sizes, they found that there appears to be a minimum KND at about the level of the subgrain size characteristic of deformation in those alloys under the conditions of temperature and strain rate involved – typically about 0.5 lm. The rate of conventional boundary migration is insufficient to explain that result, and they suggest that the boundaries are involved in the dynamic recovery process, with the local stresses occurring during deformation perhaps aiding the grain growth. There has been work indicating that stress can influence the migration of low angle boundaries [23,24], and relatively low levels of stress can even influence the migration of high angle boundaries [25], although the general situation is far from clear. This type of effect could be contributing to the dynamic grain growth occurring in the high temperature deformation reported here, where KND is also close to the typical subgrain size expected at the temperature and strain rate used. However, boundary mobilities at 375 °C and a relatively low strain rate are unlikely to be a limiting factor in the present case.
4. Conclusions There is a transition from a texture which has an almost uniformly populated b-fibre at moderate strains to one dominated by the brass component at high strains in the hot plane strain compression of Al–6Cu– 0.4Zr. This is associated with differential dynamic grain growth, which is principally driven by different substructural energy densities in different grain orientations. The presence of Zener pinning will make this process more selective. Simple geometric effects of deformation, together with a possible influence of dynamic recovery involving the high angle boundaries when their spacing becomes small, enhance this growth.
Acknowledgements This work was partly supported by EPSRC via Grants GR/R52695 and GR/R69952. We are most grateful to Alcan International for the supply of material.
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