The Cube Texture Revisited

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The Cube Texture Revisited ARTICLE in MATERIALS SCIENCE FORUM · DECEMBER 2011 DOI: 10.4028/www.scientific.net/MSF.702-703.3

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1 AUTHOR: Bevis Hutchinson Swerea - Swedish Research 119 PUBLICATIONS 1,854 CITATIONS SEE PROFILE

Available from: Bevis Hutchinson Retrieved on: 24 February 2016

Materials Science Forum Vols. 702-703 (2012) pp 3-10 © (2012) Trans Tech Publications, Switzerland doi:10.4028/www.scientific.net/MSF.702-703.3

The Cube Texture Revisited Bevis HUTCHINSON SwereaKIMAB (Swedish Institute for Metals and Corrosion Research) Box 55970, SE-10216 Stockholm, Sweden email: [email protected] Keywords: cube texture, nucleation, growth, recovery, solute, precipitation

Abstract. The cube texture in rolled and annealed fcc metals and alloys has long fascinated metallurgists because of its high symmetry and extreme sharpness. This paper demonstrates and analyses the texture perfection that is developed in a copper sample. Reasons that have been advanced to explain the development of the texture during recrystallisation and grain growth are discussed. Orientation selectivity is favoured during growth but more particularly in the nucleation stage. Especial attention is paid to the rapid recovery which cube oriented crystals undergo on heating after plane strain deformation and which is the basis for its uniquely preferential nucleation. Various metallurgical factors are known to affect the strength of the cube texture in practice and explanations for some of these are presented. Introduction The title of this paper should not be taken as meaning that the cube texture is no longer a subject of interest. New information is reported frequently on this subject. However, it does not retain the level of fascination that it exercised during the second half of the twentieth century. Then, it was one of the burning topics of texture research that was debated vigorously at almost every meeting on textures or recrystallisation. Now we can look upon it in a more dispassionate manner and hopefully present a balanced viewpoint. That, in any event, is the intention of the present paper. Why did the cube texture arouse such strong reactions among researchers? At least three reasons can be put forward to explain this and the present review will broadly follow these lines. Firstly, the generation of cube texture during annealing is such a remarkable occurrence. Hsun Hu stated that ‘the formation of cube texture is a great wonder of nature and the mechanism of its formation has been a scientific curiosity of the first order for many decades’ [1]. The cold rolling texture of fcc metals like copper and aluminium is complex and only moderately strong. Even a simplified description of this requires nine separate components (2xBr, 2xCu, 4xS and Goss) and yet after recrystallisation it can transform into one single extremely sharp texture component that has the highest possible symmetry. This is a fascinating thing to try to understand. Secondly, the cube texture was for decades the main test-bed for the competing theories of oriented growth and oriented nucleation. Adherents of both these viewpoints could claim support from the evolution of cube texture in recrystallisation and this stimulated continuing research and progress in the field of texture. Its influence therefore spread widely. Thirdly, cube texture is by no means of purely academic interest. Its development, prevention or strict control are central to a surprising number of industrial processes and products so research into it has been driven by commercial as well as academic demands. Table 1 gives a probably incomplete summary of practical applications where the control of cube texture is applied to different materials. Further discussion of practicalities will come later.

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Table 1 Examples of cube texture control in industrial processes and products Material Application Requirement Aluminium alloys Deep drawn boxes and cans Balanced texture with mixture of cube and R components to minimize earing in both annealed and rolled conditions Copper sheet Processing of thin sheet and Minimise cube texture as its presence foil leads to bad ductility and band breaks Pure aluminium Foil for electrolytic condensers Maximise cube texture to optimize tunnel etching to increase the surface area Nickel-iron Soft magnetic material for field Maximise cube texture to utilize the easy permalloy screening magnetization directions Nickel alloys Substrate for high temperature Sheet surface with (100)[001] orientation superconducting phases for epitaxial microstructure control Copper Photo-voltaic connection flat direction along the wire minimise wires on silicon solar cells Young’s modulus to reduce loading and fracture of the silicon wafers The cube texture phenomenon Cube texture is observed to develop after rolling and annealing in most fcc metals and alloys that have medium to high stacking fault energies. Most experimental data refer to aluminium, copper and nickel-based materials but other metals including low carbon steel after hot rolling in the austenite range are also susceptible. Heavy cold reductions, typically >80%, are required if the texture is to develop strongly after annealing. These conditions are necessary but not sufficient as the cube texture is fickle and many things can interfere with its development as will be discussed later. Nevertheless, when all the conditions are fulfilled the effect is quite remarkable. Figure 1 presents some observations on tough pitch copper that was processed to a fine grained condition (12µm) and then cold rolled 91% in the laboratory before annealing for 1 hour at 900ºC. Textures were measured using EBSD on nearly 5000 widely spaced points for both the as-rolled and annealed conditions and orientation maps were made on smaller areas of the annealed sheet.

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Fig.1 (a) {100} pole figures for cold rolled copper, (b) {100} pole figure after annealing, (c) grain boundaries in annealed copper for misorientations 1º, 3º, 6º and 9º.

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It is clear that the recrystallisation texture is a sharp single orientation but it is not possible to represent it accurately using the conventional plotting routines because of this extreme sharpness. A different approach has been adopted based on the cumulative distribution of misorientations relative to the exact cube orientation. Since the spread of orientations in this case is nearly isotropic, it can be expressed using a single angle, φ, and is shown in Fig.2. Half of all grains lie within 4.2º of the precise cube orientation and none deviate by more than 15º. It should be noted here that not a single annealing twin was seen in this annealed copper sheet. To quantify the result in terms of the standard multiples of random density, the function in Fig.2 was fitted to a series of overlapping 4th order polynomials which could then be differentiated and compared with the differential form of the Mackenzie expression for a random texture, 24(φ – sin(φ))/π. The result of normalisation is shown by the solid line in Fig.3. Unfortunately, this procedure also fails at very small angles where the statistics are small and also the argument φ becomes comparable with the precision of the EBSD method. For the smallest deviations, those below 2º, the fraction of measurements is simply compared with the integral Mackenzie expression, 24(1-cos(φ))/π, which produces the dashed horizontal line in the figure. Within this range all details of the texture are lost. Nevertheless, the amazing sharpness of the texture is seen from its peak strength which in the present material exceeds 1400 x random. The only other texture that can rival this in sharpness is the Goss component in grain oriented silicon steel and that requires extreme care in all stages of the complex production route whereas the cube texture can be simply produced with the most basic laboratory facilities. Such a cube texture is indeed ‘a great wonder of nature’. Although the cube texture develops first during primary recrystallisation, it continues to increase in strength during subsequent grain growth as smaller and less favoured grains are consumed by growth of the larger cube grains. There is accordingly an influence of the annealing temperature on the perfection of the texture as is shown in Fig.4 from various sources in the literature [2].

Fig.2 Cumulative distribution of deviations from cube orientation

Fig.3 Normalised texture strength for the cube texture

Fig.4 Cube texture strengthens during grain growth

Origin of the cube texture - nucleation and growth The long running dispute between oriented growth, OG, and oriented nucleation, ON, as mechanisms for recrystallisation texture formation often centred on the cube texture and this has been reviewed on a number of occasions, e.g. [1,3,4]. It seems now probable that both factors contribute. The mobility of grain boundaries varies with their misorientation and in the case of aluminium the fastest moving boundaries are reported to have 40º relationships [5,6]. In copper the maximum is sometimes reported as closer to 30º . These conditions are usually interpreted as being close to the Σ7 CSL relationship (38.2º ) that was first identified by Kronberg and Wilson [7] in copper sheets containing a strong cube texture. However, mobile grain boundaries show a significant spread with respect to both these angles and axes. The four symmetrical S-orientations in the rolled texture all have approximately 38º relationships to the cube texture, Fig.5, so a large proportion of the deformed matrix is suitably oriented for consumption by the cube texture on annealing. Transformation of experimental rolling textures using oriented growth relationships have been carried out on numerous occasions and these confirm

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that the cube orientation is predicted on this type of basis. Figure 6 is taken from the earliest of these computer models [8]. These models indicate that the predicted peak is actually closer to {100}, the ND rotated cube orientation, and this has sometimes been observed in some aluminium alloys which may support the OG mechanism [9] although alternative explanations have been given for this as described below.

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(c) Fig.5 {111} pole figures, (a) cold rolled, (b) recrystallised and (c) the 40º relationship between the cube orientation and one of the S-components of the rolling texture.

Fig.6 Predicted ODF for recrystallised aluminium based on 40º relationships according to OG theory [8].

It seems unlikely that an OG mechanism based on ‘special’ grain boundaries is sufficient to account for the cube texture for a number of reasons. In the first case, several of the predicted components in Fig.6 are not found in practice which casts doubts on its generality. Secondly, there are many cases where virtually identical deformation textures lead to quite different recrystallisation textures. For example, copper in a fine grained condition prior to rolling yields, after annealing, the extreme cube texture shown in Fig.2 but in a coarse grained initial condition gives the same rolling texture but only a very weak final cube component. A simple transformation model cannot account for this. Thirdly, the condition for maximum growth rate is far from exact. Despite the assumption of a unique relationship in the models, the predicted intensity in the cube orientation may be underestimated by a factor of at least 100. Experimental evidence for an enhanced growth rate of cube grains as compared to other orientations is conflicting. Some measurements [10,11] showed a somewhat higher growth rate while others did not [12]. Thus, while it would be wrong to discount any contribution from OG arising from ‘special’ boundaries, it is evident this alone cannot account for the remarkable perfection that is seen in this ‘scientific curiosity of the first order’. An additional growth condition that favours the cube texture arises from the rather few cube nucleation sites that are typically present in the cold rolled metal. The nuclei are few and far between which means that a growing cube grain will very seldom encounter a region of nearly similar orientation. Since low angle grain boundaries are well known to have very low mobility, the growing cube grains are seldom hindered by the presence of sluggish segments on their moving boundaries. Other nuclei, arising by SIBM or shear banding, tend to have orientations near to the deformation texture and so their growth is more inhibited. This was formalised under the name of ‘orientation pinning’ by Juul Jensen. Growing cube grains are not susceptible to orientation pinning [13] except in cases where the material prior to rolling already contains a strong cube texture which is then significantly retained in the deformed structure, such that final cube texture strength is much reduced. The spacing between bands of specific deformation textures becomes an important factor governing the relative sharpness of cube and other textures as discussed by Doherty et al. [11].

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Although early theories paid attention to the role of nucleation in texture selection [3], there was for many years no experimental evidence for the existence of cube oriented regions within the deformation substructure in cold rolled metals. Major TEM studies by Hu [14] and others [3] failed to detect any cube orientation in cold rolled copper and, as a result, it was widely held that the orientation arose during annealing either by repeated grain rotation and coalescence or by a kind of un-twinning process, although both of these processes were known to be physically unreasonable. Subsequently, after about 1980, the previously so elusive cube oriented substructures became readily apparent in TEM and then later in EBSD studies. Numerous quantitative measurements [e.g. 15-18] of these structures have since been made in many kinds of fcc metals and alloys and it is clear that the origin of the final texture can be traced back to these pre-existing cube regions which are present, albeit sparsely distributed, among the major deformation texture components. The key question then becomes ‘why are just these cube elements in the substructure so much more potent at developing into recrystallised grains than others’? In most observations it is reported that the cube oriented regions within the deformed structure exist as elongated bands that comprise larger subgrains and smaller misorientations than in the other texture components. They have therefore lower stored energy than their surroundings. For this reason and because of their elongated geometry, they can readily bulge out and then continue to grow as recrystallised grains with mobile high angle boundaries. An example of such an elongated cube band at the very start of recrystallisation [15] is presented in Fig.7.

Fig.7 A cube oriented band in rolled copper at the very start of recrystallisation Dillamore and Katoh [19] proposed a model by which the cube orientation could form within transition bands during rolling and this received some support from TEM observations in copper [15]. However, in aluminium, quite different lattice rotations have been found [16] and the transition band model is no longer widely supported. In most cases it appears that the deformed cube bands are residues of grains that possessed this orientation in the initial structure [11,17]. It is most significant that the cube regions have a quite exceptional ability to develop into recrystallised grains, irrespective of how they originated in the deformed structure. For example, Kamijo et al. [20] found that heavily rolled single crystals of aluminium having the stable S-orientation recrystallised to cube texture. Although the origin of the cube nuclei is not known in this case, they probably arose from some form of local deformation heterogeneity. Several views have been presented regarding the favourable nucleation condition. Duggan et al. [21] showed that cube bands in copper generated recrystallisation preferentially when adjacent to deformed grains in the S-orientation. Since their boundaries then have a Σ7 character, this behaviour was interpreted as ‘micro growth selection’ taking place at the nucleation stage. Nes et al [17] confirmed the observation in aluminium but explained it by the Σ7 boundaries having lower energy and so fulfilling the nucleation condition more readily than for other boundaries. Experiments [22] have also shown the S-orientation to possess higher stored energy than the other major components and this is another factor that would favour its being engulfed by the bulging cube grains. In view of the current consensus that nucleation is the dominating factor in cube texture development and that the special low energy substructure inside the cube bands is largely responsible for this, the salient question becomes ‘why do cube bands have this special low energy

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in cold rolled fcc metals’? It has been pointed out that cube oriented crystals have a small Taylor factor in rolling [23] and so require less dislocation activity in deformation. This should reasonably result in a lower density of dislocations being retained. However, there are other orientations that actually have lower Taylor factors and, in fact, all crystal orientations for which lies along ND or RD have the identical Taylor factor. Sometimes these appear as recrystallisation textures such as the {100} component mentioned earlier in connection with OG but they never achieve the sharpness of the true cube texture and frequently they are not present at all. There is evidence that the low stored energy in deformed cube bands results from an enhanced recovery rate. In aluminium and also in copper this probably includes dynamic recovery that occurs simultaneously with the deformation. In iron-nickel alloys, Zaefferer [18] found little recovery during cold rolling and that cube bands in the as-rolled condition had substructures that were of similar density to those in other orientations. However, during heating, they recovered much more than the other structures and so achieved the low energy condition that allowed them to bulge out as cube nuclei at commencement of recrystallisation. The only model that can explain the specific advantage of cube bands to undergo recovery and rapid nucleation is that proposed by Ridha and Hutchinson [15]. They noted that during the plane strain deformation in rolling, four sets of slip systems are active but with only two types of Burgers vectors. Furthermore, these two slip directions are orthogonal to one another, Fig.8. This means that the dislocations have minimal elastic interactions and do not form stable junctions or networks. They should accordingly be much freer to undergo thermally activated recovery than the dislocations in crystallites of other orientations. The condition of strict orthogonality is absolutely unique to the cube orientation in rolling deformation. No other texture component at all experiences this same favourable situation.

Fig.8 Stereogram with slip planes and directions (solid symbols) for a cube oriented crystal in plane strain rolling

Fig.9(a) Stress strain curves for cube textured sheet tested in simple tension with interruption and recovery anneal

Fig.9(b) Stress strain curves for cube textured sheet tested in plane strain tension with interruption and recovery anneal

The dislocation-based recovery model has received qualitative support since it was put forward some 25 years ago but here we present independent and quantitative evidence in support of it. Specimens of the very strongly cube oriented sheet shown in Fig.1 were tested in tension along the cube axis to a strain of 10%. They were then recovery annealed for 1 hour at 300ºC and re-tested. In one case, the sample was a typical long specimen in simple tension. Alternatively, the sample had supporting plates of brass brazed on to it so that the free gauge length was much shorter than the width of the strip. This forces the same kind of plane strain deformation to exist as in flat sheet rolling. The stress-strain curves before and after interrupting the tests with the anneal are shown in Fig.9 (a) and (b). In simple tension as many as 8 slip systems are equally stressed but there is virtually no reduction in flow stress on re-loading showing that almost no recovery took place. In contrast, the plane strain tensile tests, where only 4 parallel/orthogonal slip systems are active, showed very extensive recovery. The initial hardening was almost identical in the two cases but reduction in substructure on heating is much more advanced when the dislocations only interact weakly. Exactly the same condition may be assumed to occur in cube bands in cold rolled sheet.

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Why is the cube texture so vulnerable? There are many examples of the intensely sharp cube textures shown earlier but they are by no means universal in rolled and annealed fcc metals. Many metallurgical factors can conspire to weaken or even eliminate this texture completely. In order to understand this it must be recognised that although the cube nuclei are very effective, they are also few in number. To establish a strong final cube texture necessitates that the grains grow extensively until they impinge on one another. Since they are few in number, a sharp texture, therefore, necessitates a large final grain size as is invariably observed. The natural advantage of the cube grains can be upset either by retarding their growth so that other, less favoured, nuclei have time to develop or by introducing alternative viable nucleation sites. In either case the final grain size is reduced and the texture becomes ‘diluted’ by other orientations. In the case of copper, many elements in solid solution create solute drag that inhibits the growth range of the cube grains. This effect can be very drastic as shown in Fig.10. For example, a pure 100% cube texture typically is associated with a grain size of about 100µm. If the nuclei only grow to half this distance due to competition from others, then the volume of cube orientation is reduced by a factor of (½)3, or to about 12%. If the growth distance is restricted to one quarter, then the volume fraction of cube becomes less than 1% and the texture would not be recognised at all. Note that these considerations imply no change in the deformation texture or the dislocation substructure that is present in the deformed cube bands. The content of impurity necessary to achieve this can be very small indeed. Concentrations given in Table 2 are approximate but they are sufficient to raise the recrystallisation temperature substantially (∆Trecryst) and markedly reduce the final grain size. Table 2 Approximate impurity concentrations (wt%) that eliminate cube texture in annealed copper sheet as well as the rise in recrystallisation temperature caused by these impurity contents Impurity P As Sb Sn Ge Al Mn Zn Concentration 100º >120º ~100º

(b) (a) Fig.10 Recrystallisation textures after cold rolling 95% and annealing for (a) pure Cu and (b) Cu-0.1% P

(c) Fig.11 Effect of grain size before cold rolling on cube texture after annealing

In copper, it has been found that when shear bands appear in the deformed structure these provide alternative viable nuclei for recrystallisation that have different orientations which are usually scattered but with some tendency towards retained rolling texture components. Shear bands also break up the cube oriented bands in the deformed structure, making these less favoured as nucleation sites and the net result is a severe weakening of the cube texture after recrystallisation [15]. The tendency for shear banding is very dependent on the initial grain size of the material and so cube texture strength can be controlled by this parameter as shown in Fig.11. Fine initial grain size also refines the spacing of orientated bands, providing more and closer spaced cube bands. This benefits the development of cube grains but also enhances the ‘orientation pinning’ effect for other competing orientations, as demonstrated in aluminium and discussed by Doherty et al. [11].

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Shear banding may also play a role in weakening the cube texture in aluminium alloys by encouraging other competing nuclei to develop [24]. However, in many practical situations, the alternative nuclei arise by particle stimulated nucleation (PSN) with nearly random character. Considerable efforts have been made to control texture in canstock alloy sheets in order to optimise their earing. Good correlations are found between the 0º/90º earing resulting from cube texture and the frequency of particles that are large enough to cause PSN as shown by the results in Fig. 12. Particles larger than the critical size for generating nuclei, dcrit, lead to weaker cube texture and smaller grain sizes after recrystallisation following hot rolling [25]. Fig.12 Increasing densities of second phase particles that are large enough to cause PSN result in weaker cube textures after recrystallisation and accordingly to less pronounced 0º/90º earing in AA3004 canstock sheet alloy.

References [1] H. Hu, Proc. 7th Risö Int. Conf. Annealing Processes-Recovery, Recrystallisation and Grain Growth, Ed. N.Hansen et al. 1986, pp. 75-92. [2] W.B. Hutchinson, E. Nes, Mater. Sci. Forum, 94-96 (1992), 385-390. [3] P.A. Beck, H. Hu, Recrystallisation, Grain Growth and Textures, ASM, Ohio 1965, pp.393-431. [4] E. Nes, W.B. Hutchinson, Proc. 10th Riso;/ Int. Symp. Ed. N. Hansen, 1989, pp. 233-250. [5] B. Liebman, K. Lücke, G. Masing, Z. Metallkde., 47 (1956) 57-61. [6] Y. Huang, F.J. Humphreys, Acta Mater., 47 (1999) 2259-2268. [7] M.L. Kronberg, F.H. Wilson, Trans. Met. Soc. AIME., 185 (1949) 501- 514. [8] T. Öztürk, Ph.D. thesis, Cambridge University, UK, 1978. [9] K. Lücke, Proc.ICOTOM-7, Holland, Netherland Soc. for Materials Science, 1984, pp.195-210. [10] D. Juul-Jensen, N. Hansen, J. Humphreys, Acta Met. 33 (1985) 2155-2162. [11] R.D. Doherty, L-C Chen, I. Samijdar, Mater. Sci. Eng., A257 (1998), 18-36. [12] E. Nes, J.K. Solberg, Mat. Sci. Tech., 2 (1986) 19-21. [13] D. Juul Jensen, Acta Mater., 43 (1995) 4117- 4129. [14] H. Hu, Textures in Research and Practice, Clausthal, Springer-Verlag, 1968, pp. 200-226. [15] A.A. Ridha, W.B. Hutchinson, Acta Met., 30 (1982) 1929-1939. [16] A.L. Dons and E. Nes, Mat. Sci. Tech. 8 (1986) 8-18. [17] O. Daaland, E. Nes, Acta Mater., 44 (1996) 1389-1411. [18] S. Zaefferer, Dr. Habil. thesis, University of Aachen, 2009. [19] I.L. Dillamore, H. Katoh, [20] T. Kamijo, A. Fujiwara, Y. Yoneda, H. Fukutomi, Acta Mater., 39 (1991) 1947-1952. [21] B.J. Duggan,K. Lücke, G. Köhlhoff, C.S. Lee, Acta Met., 41 (1993) 1921-1993. [22] A. Borbély, J.H. Driver, Proc. Rex.G.G., Aachen, Springer-Verlag , 2001, pp.635-644. [23] I. Samajdhar, R. Doherty, Scripta Met., 32 (1995) 845-850. [24] O. Engler, Scripta Mater., 44 (2001) 229 234. [25] W.B. Hutchinson, A. Oscarsson, Å. Karlsson, Mat. Sci. and Techn., 5 (1989) 1118-1127.