Effect of Stacking Fault Energy on Evolution of Recrystallization Textures in Drawn Wires and Rolled Sheets

Transcription

1 Materials Science Forum Vols (2005) pp online at Trans Tech Publications, Switzerland 194 Effect of Stacking Fault Energy on Evolution of Recrystallization Textures in Drawn Wires and Rolled Sheets Dong Nyung Lee Research Institute of Advanced Materials and School of Materials Science and Engineering Seoul National University, Seoul , Korea Keywords: Aluminum, copper, gold, silver, aluminum bronze, brass, wire drawing, rolling, recrystallization, texture. Abstract. The drawing textures of aluminum, copper, gold, silver, and Cu-7.3% Al bronze wires are approximated by major <111>+minor <100>, except silver wire, which can have the <100> texture at extremely high reductions. The <111> component in the drawing textures of aluminum, copper, gold, and silver transform to the <100> component after recrystallization. On the other hand, the <111> deformation texture of the Cu-7.3% Al bronze wire, which has very low stackingfault-energy, remains unchanged after recrystallization. The Brass component {110}<112> in rolling textures of high stacking-fault-energy metals such as aluminum and copper alloys changes to the Goss orientation {110}<001> after recrystallization. However, the Brass orientation in rolling textures of low stacking-fault-energy fcc metals such as brass appears to change to the {236}<385> orientation after recrystallization. These results seem to be related to the stability of dislocations during annealing. Introduction The drawing textures of aluminum (stacking fault energy (SFE): 166 mj/m 2 [1]) [2], copper (SFE: 78 [1]) [3,4], gold (SFE: 45 [1]) [4], silver (SFE: 22 [1]) [5], and Cu-7.3% Al bronze (SFE: <20) [6] wires are approximated by major <111>+minor <100>, except silver wire, which can have the <100> texture at extremely high reductions. The <111> components in drawing textures of aluminum, copper, gold, and silver transform to the <100> components in their recrystallization textures. The recrystallization texture transform to the <111> texture again after prolonged annealing at high temperatures. However, the <111> drawing texture of the Cu-7.3% Al bronze wire, which has very low stacking-fault-energy, remains unchanged after annealing. The Brass component {110}<112> in rolling textures of high stacking-fault-energy metals such as aluminum [7] and Cu-Mn alloys [8,9] changes to the Goss orientation {110}<001> after annealing. However, the Brass rolling textures of low stacking-fault-energy fcc Cu-alloys transform to the {236}<385> orientation after annealing [10]. The purpose of this paper is to discuss the development of different annealing textures from the same deformation texture. Exprimental Results Drawing and Annealing Textures. The drawing textures of aluminum [2], copper [3,4], gold [4], silver [5], and Cu-7.3% Al bronze [6] wires are approximated by major <111>+minor <100>, except silver wire, which can have the <100> texture at extremely high reductions. The <111> component in drawing textures of aluminum, copper, and silver transforms to the <100> component during recrystallization. On further annealing, the density of the <111> component increases. One example is shown in Fig. 1. Licensed to Dong Nyung (dnlee@snu.ac.kr) - Seoul National University - Korea All rights reserved. No part of the contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of the publisher: Trans Tech Publications Ltd, Switzerland, (ID: /05/05,03:10:32)

2 1244 Textures of Materials - ICOTOM 14 Fig. 1. Orientation density ratio of <100> fiber component to <111> component of 90 % drawn copper wire as a function of time during annealing at 700 C [3]. On the other hand, the <111> deformation texture of the Cu-7.3% Al bronze wire, which has very low stacking-fault-energy, remains unchanged after annealing, skipping over the stage of the <111> to <100> transformation [6]. Rolling and Annealing Textures. The Brass component {110}<112> in rolling textures of high stacking-fault-energy metals such as aluminum [7] and Cu-Mn alloys [8,9] changes to the Goss orientation {110}<001> after annealing as shown in Fig. 2. However, the Brass rolling textures of low stacking-fault-energy fcc Cu-alloys transform to the {236}<385> orientation after annealing [10]. The 40 <111> relation is satisfied between the Brass orientation and the {236}<385> orientation, whereas the relation is not established between the Brass and Goss orientations. The 40 <111> relation is often addressed as a token of the oriented growth model. Fig. 2. (left) β-fiber intensity lines of Cu-16%Mn alloys after various rolling reductions [8]. (right) {111} pole figures of Cu-16%Mn alloys after complete recrystallization (97.5% rolled, annealed for 1000s at 450 C) [9].

3 Materials Science Forum Vols Discussion In order to understand the experimental results, we have to understand the evolution of rectystallization textures and grain growth textures. A major driving force for recrystallization is dislocation energy, and a major driving force for grain growth is grain boundary energy. The recrystallization textures will be discussed based on the strain-energy-release maximization (SERM) model advanced by the author [11]. Once the dislocation density decreases due to recrystallization, grain growth can be controlled by grain boundary energy and mobility, resulting in a change in texture. Thus, the recrystallization texture can differ from the grain growth texture. The SERM Model. The model is based on the presumption that the stored energy due to dislocations is a major driving force for recrystallization and has been elaborated in recent reviews [12,13]. In the SERM model, the recrystallization texture is determined by the orientation relationship between the internal stress field associated with the dislocation array in the deformed matrix and the elastic anisotropy of recrystallized new grains. The strain-energy release can be maximized when the direction of absolute maximum internal stress due to dislocations in the deformed material is parallel to the minimum Young s modulus direction (MYMD) in recrystallized grains. The absolute, maximum internal stress direction (AMSD) due to dislocations can be calculated by superposition of stress fields of individual dislocations. The dislocation array in severely deformed materials is very complicated. The dislocations generated during deformation can be of edge, screw, and mixed types. However, their Burgers vectors can be determined by the deformation state and texture, and they can be approximated by a stable or low-energy array of edge dislocations. It can be shown that the absolute maximum normal-stress direction is along the Burgers vector. For multiple slip, the AMSD is calculated by the vector sum of active slip directions, taking their activities into account. The activity of each slip direction is linearly proportional to the dislocation density on the corresponding slip system, which is roughly proportional to shear strain γ on the slip system. Fig. 3. (a) Tetrahedron showing slip systems {111}<110> of fcc metal having <111> fiber texture, and (b) slip systems leading to <100> orientation (W.A. stands for wire axis). Axisymmetrically Drawn Fcc Wires. The texture transition from <111> to <100> was discussed based on the SERM model [11,12]. It is briefly introduced here. The main components in the global deformation texture of the drawn fcc wire are <111> and <100>, with the <111> component having a higher density than the <100> component. When the <111> or <100> fiber-

4 1246 Textures of Materials - ICOTOM 14 texture component develops by axisymmetric plastic extension, the AMSD is parallel to the drawing direction. Fig. 3 shows a tetrahedron consisting of slip planes and directions for an fcc metal whose axial orientation is <111> and an octahedron consisting of slip planes and directions of an fcc metal whose axial orientation is <100>. For the <111> oriented crystal (Fig. 3 a), three m<110> directions and the planes made by the slip directions are calculated to be active slip systems, which are equally active. The AMSD, which is calculated by the vector sum of the slip directions, is parallel to the axial direction of the crystal. Therefore, the <111> axis direction of the <111> oriented metal will become parallel to the MYMD of a recrystallized grain. The MYMD of fcc metals is the <100> direction. Therefore, the deformed <111> grains will be replaced by recrystallized <100> grains. The evolution of the <100> recrystallization texture from the <100> deformation texture is also explained by SERM model. Operating slip systems of an fcc crystal elongated along the [100] direction are calculated to be ( 111)[101], ( 111)[1 10], ( 1 11)[101], ( 111)[110], ( 111)[101], ( 111)[110], ( 111)[101], and ( 111)[1 10]. The AMSD obtained from the vector sum of the slip directions on the operating slip systems [400] is again parallel to the axial orientation of [100] (Fig. 3 b). Therefore, the <100> deformation texture will remain unchanged after recrystallization, because atoms are likely to remain at the same positions unless otherwise required. According to the SERM model calculations, the <111> and <100> components in the deformation texture are going to transform to the <100> component after recrystallization. It is worth mentioning that the SERM model does not describe the mechanism, but address energetics of the texture turnover. However, the texture change during annealing might take place by the following process. The <100> oriented grains might retain their deformation texture during annealing by continuous recrystallization by recovery-controlled processes without long-range high-angle boundary migration. The <100> grains could grow at the expense of their neighboring <111> grains, which tend to assume the <100> orientation. The texture transition is related to recrystallization to reduce the stored energy caused by dislocations generated during deformation. The texture transition from the duplex texture of <100>+<111> to the <111> orientation is associated with grain growth. The grain boundaries in a material with a duplex fiber texture of {111} and {100} can be described approximately as tilt boundaries of {111} and {100} grains. The {111}-{100} boundaries will not determine which grains grow favorably. If the tilt boundary energies of the two differently textured grains are different, the grains with lower boundary energy will grow at the expense of the grains with the higher boundary energies, for the same mobility of the grain boundaries. According to many studies, the grain-boundary energy may change by 20 pct or less, but the mobility may change by orders of magnitude [14]. It has been also known that, for fcc metals, tilt boundaries migrate much faster than the twist boundaries [15]. The mobility data of tilt grain boundaries in aluminum [16] show that the average mobility of a <111> tilt boundary is increasingly higher than that of a <100> boundary with increasing temperature. If the mobility of <111> tilt boundaries is higher than that of <100> tilt boundaries, neighboring <111> grains will grow faster than neighboring <100> grains, resulting in larger <111> grains, which, in turn, will grow at the expense of <100> grains, because they now have a size advantage. The preferred growth of <111> grains is also observed in nanocrystalline deposits of elecroless Ni alloys [17,18], and electrodeposited Fe-Ni alloys [19, 20], which are devoid of dislocations. We are in position to discuss the lack of the <111> to <100> transition during annealing of drawn Cu-7.3% Al alloy wire, which has a very low stacking fault energy. The dislocation structure in low stacking-fault-energy metals is planar and homogeneous, so dislocations can be easily annihilated during annealing. If the dislocation density decreases so rapidly during annealing that the driving force for recrystallization due to dislocations decreases, the <111> to <100> transition is likely to be skipped over.

5 Materials Science Forum Vols Rolled Fcc Sheets. The evolution of the Goss orientation in the Brass oriented matrix during recrystallization was explained using the SERM model [21].When fcc crystals with the ( 110)[112] orientation are plane strain compressed along the [110] direction and elongated along the [ 1 12] direction, the relation between the strain ε 11 of specimen and shear strain rates dγ/dε 11 on active slip systems was calculated by the ε 12 and ε 23 relaxed strain rate sensitive model with the subscripts 1, 2, and 3 indicating the rolling direction (RD), the transverse direction (TD), and the normal direction (ND), respectively. The calculated results indicated that active slip systems are ( 111)[011] and ( 1 1 1)[101], on which shear strain rates do not vary with strain of specimen, indicating that the Brass orientation is stable with respect to the strain. It is noted that the active slip directions were determined so that they are at acute with the [ 1 12] elongation direction. The vector sum of the slip directions, [ 011] + [101] = [ 1 12], is AMSD, which is the same as the elongation direction or RD. The recrystallization texture is determined such that the AMSD in matrix is parallel to MYMD of recrystallized grain, the <100> directions in fcc metals. Therefore, the recrystallized grains will have the (hk0)[001] orientation. However, the minimum atomic shuffling condition gives rise to the (110)[001] orientation, because the (110) plane is shared by the deformed and recrystallized grains. This means that the ( 110)[112] deformation texture becomes the (110)[001] recrystallization texture. Similarly, the (011) [ 2 11] deformation texture is calculated to be the (011)[100] recrystallization texture. Therefore, it is concluded that the Goss recrystallization texture is linked with the brass deformation texture. It is experimentally shown that the Goss orientation in aluminum single crystal is stable with respect to plane strain compression and thermally stable [22]. This phenomenon could be explained by SERM [23]. Therefore, grains with the Goss orientation that survived during rolling are likely to act as nuclei during subsequent recrystallization and will grow at the expense of surrounding brass oriented grains which are destined to change to have the Goss orientation. However, the Brass rolling textures of low stacking-fault-energy fcc Cu-alloys transform to the {236}<385> orientation after annealing [10]. The 40 <111> relation is satisfied between the Brass orientation and the {236}<385> orientation, whereas the relation is not established between the Brass and Goss orientations. The 40 <111> relation is often addressed as a token of the oriented growth model. That is, the recrystallization texture is controlled by grain growth with a high mobility of the <111> grain boundaries. As discussed in the section of drawn wires, the dislocation density in low stacking fault materials is likely to decrease so rapidly during annealing that the driving force for recrystallization due to dislocations decreases. In this case, grain boundaries can control the annealing texture, resulting in the {236}<385> annealing texture. In this sense, the {236}<385> annealing texture may be a grain growth texture rather than a recrystallization texture. Conclusion The differences in annealing texture between high stacking-fault-energy metals and low stackingfault-energy metals seems to be caused by differences in stability of dislocations during annealing between high and low stacking-fault-energy metals. The dislocation structure in low stacking-faultenergy metals is planar and homogeneous, so dislocations are easy to annihilate during annealing. References [1] L.E. Murr: Interface Phenomena in Metals and Alloys (Addison-Wesley 1975) p.131. [2] H. Park and D.N. Lee: Mechanical Properties of Advanced Engineering Materials, edited by B. Xu, M. Tokuda, G. Sun, Tsinghua University Press, Beijing, (2003), p. 13.

6 1248 Textures of Materials - ICOTOM 14 [3] H. Park and D.N. Lee: Metall. Mater. Trans. Vol. 34A (2003), p [4] J-H Cho, Ph.D. Thesis, Seoul National University, Seoul, Korea (2003). [5] H.-J. Shin, H.-T. Jeong and D. N. Lee: Mater. Sci. Eng. Vol. A279 (2000), p [6] S. Kim, C.-H. Choi and D.N. Lee: Mater. Sci. Forum Vol (2002), p [7] J. Hjelen, R. Ørsund and E. Nes: Acta metall. mater. Vol. 39 (1991), p [8] O. Engler: Acta mater. Vol. 48 (2000), p [9] O. Engler: Acta mater. Vol. 49 (2001), p [10] P.A. Beck and H. Hu: Trans AIME Vol. 194(1962), p. 83. [11] D.N. Lee: Scripta Metall. Mater. Vol. 32 (1995), p [12] D.N. Lee: Int. J. mech. Sci. Vol. 42 (2000), p [13] D.N. Lee: Phil. Mag. In press. [14] D. A. Molodov, in: G. Gottstein and D. A. Molodov (Eds), The First Joint International Conference on Recrystallization and Grain Growth (Springer-Verlag 2001), p. 21. [15] F. J. Humphreys and M. Hatherly, Recrystallization and Related Annealing Phenomena (Pergamon Press, Oxford, 1995). [16] A.N. Aleshin, V.Y. Aristov, B.S. Bokstein and L.S. Shvindlerman: Phys. stat. sol. Vol. A45 (1978), p [17] D.N. Lee and K.-H. Hur: Scripta Mater. Vol. 40 (1999), p [18] D.N. Lee and K.-H. Hur: Textur. Microstr. Vol. 34 (2000), p [19] F. Czerwinski, H. Li, M. Megret and J.A. Szpunar: Scripta Metall. Mater. Vol. 37 (1997), p [20] Y.-B. Park, J. Park, C.S. Ha and T.H. Yim: Mater. Sci. Forum. Vol (2002), p [21] D.N. Lee, H.-J. Shin and S.-H. Hong: Key Eng. Mater. Vol (2003), p.515. [22] M. Ferry and F.J. Humphreys: Acta mater. Vol. 44 (1996), p [23] D.N. Lee: J. Mater. Process. Tech. Vol. 117 (2001), p. 307.