Keywords: Slip systems, Twinning, Deformation Texture, Recrystallization Texture, Dynamic Recrystallization, Texture Simulation

Transcription

1 Materials Science Forum Online: ISSN: , Vols , pp doi: / Trans Tech Publications, Switzerland Texture Development in pure Mg and Mg Alloy AZ31 G. Gottstein, T. Al Samman Institut für Metallkunde und Metallphysik, RWTH Aachen, Aachen, Germany Keywords: Slip systems, Twinning, Deformation Texture, Recrystallization Texture, Dynamic Recrystallization, Texture Simulation Abstract Texture evolution in pure Mg and Mg alloy AZ31 during deformation and annealing was investigated. The poor low temperature ductility can be attributed to both, insufficient shear systems and unfavorable deformation geometry. Static recrystallization was shown to proceed discontinuously despite little texture change. High temperature deformation was accompanied by dynamic recrystallization with similar texture development as during static recrystallization. 1. Introduction Mg alloys have been subject of intense research during recent years, owing to their excellent properties, such as low density and high specific strength, and because of the growing interest in weight savings of the automotive and aerospace industries. However, magnesium being a hexagonal close-packed metal has a limited ductility and poor formability at room temperature due to an insufficient number of operative slip and twinning systems. At elevated temperatures the workability of magnesium substantially increases as additional slip systems become available by thermal activation and, therefore, wrought Mg alloys are hot formed [1]. Besides ductility, Mg alloys also suffer from strong mechanical anisotropy due to the development of pronounced crystallographic textures. Texture development during processing of hexagonal materials has been subject of several studies [2-11], but due to the complicated deformation geometry and the lower crystal symmetry there is still need for systematic investigations into the mechanisms of texture evolution in these materials. This paper is concerned with the origin and development of crystallographic texture in magnesium and magnesium alloy AZ31 (3%Al,1%Zn). In principle, crystallographic texture can arise or can be changed during solidification, deformation, recrystallization, and phase transformations. Since magnesium and its commercial alloys do not undergo massive phase transformations or a change of crystal structure in the solid state and because solidification textures are usually weak except for special processing conditions, we will confine our consideration here to deformation and recrystallization textures. 2. Deformation textures 2.1 Deformation geometry Since crystalline materials have to deform by simple shear in order to conserve the crystal structure and simple shear is associated with a lattice rotation, the orientation of a crystal will usually change during plastic deformation. This orientation change depends on the deformation mechanisms and their geometry. In magnesium there are two dominant deformation mechanisms, crystallographic slip and mechanical twinning (Fig. 1). In contrast to cubic crystals, in particular fcc materials, there is a large variety of slip systems and mechanical twinning systems in hexagonal crystals and their activation depends in particular on the c/a ratio of the material. Magnesium has a ratio c/a = which is close to the ideal ratio of close packed spheres c/a(ideal) = 8 / 3 = Therefore, basal slip (0001)<1120 > is most easy and strongly preferred among all slip systems. Less densely packed slip planes or even larger Burgers vectors (slip directions) are more difficult to activate, therefore, respective slip systems operate only at a higher critical resolved shear stress. There are 3 All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, (ID: , Pennsylvania State University, University Park, USA-06/03/16,19:57:06)

2 624 Textures of Materials - ICOTOM 14 different <1120 > directions in the (0001) plane and thus, there are 3 basal slip systems, only 2 of them are independent. The <1120 > direction is also parallel to the prism planes {10 10 } and to the pyramidal planes {10 11}. Slip on the systems {10 10 }<1120 > and {10 11}<1120 > is referred to as prismatic slip and pyramidal slip, respectively. They share with the basal slip systems the slip direction in the basal plane, which is also referred to as <a> slip, but the prismatic and pyramidal slip systems have larger critical resolved shear stresses. Fig. 1: Potential shear systems in Mg: basal <a>, prismatic and pyramidal <a>, pyramidal <c+a>, and tensile twin. Since basal, prismatic, and pyramidal slip have a slip direction parallel to the basal plane, they cannot accommodate any deformation out of the basal plane, i.e. they do not provide 5 independent slip systems. For an arbitrary deformation also slip vectors with a component perpendicular to the basal plane (i.e. parallel to the c-axis) have to be activated, which is referred to as <c+a> slip, e.g. on the {1122 } pyramidal plane, i.e. the pyramidal slip systems {1122 }<1123> (Fig. 1). This requires a substantially larger slip vector and thus, a markedly higher Peierls stress, which makes itself felt as a much higher critical resolved shear stress. Crystals can also deform by mechanical twinning. While there are a variety of mechanical twinning systems in hexagonal crystals, the only twinning system that has been observed to operate in Mg is the so-called tension twin, i.e. the {10 12 }<10 11> twinning system (Fig. 1). It is called a tension twin, since it can only be operated for c/a < 3 if there is a tensile component parallel to the c-axis. By contrast, this twinning system cannot be operated if a compressive stress component acts along the c-axis, which is the case during rolling, when the basal plane is essentially parallel to the rolling plane. Therefore, twinning in Mg during rolling will only occur, if the crystal orientation is far from the basal orientation. On the other hand, twinning is accompanied by a large orientation change, namely a 86.3 <1120 > rotation. An orientation with its basal plane essentially perpendicular to the rolling plane will become reoriented by twinning into an orientation close to the basal orientation, where further mechanical twinning will cease. Plastic deformation in Mg and its alloys is, therefore, strongly affected by: (a) the orientation of the crystal, i.e. the texture of the material prior to deformation and (b) the deformation conditions, i.e. temperature and strain rate, since the operation of non-basal slip systems requires thermal activation. 2.2 Effect of initial texture The effect of crystal orientation on deformation and recrystallization can be readily studied in hexagonal materials, since thermomechanical processing (e.g. hot extrusion) usually leads to a very pronounced texture with the basal plane parallel to the deformation direction. Therefore, one can cut samples in such a way that the basal plane is essentially parallel to any desired

3 Materials Science Forum Vols direction for a given deformation geometry. We chose plane strain compression (PSC) as the deformation mode since in contrast to rolling it allows to conduct experiments at defined deformation conditions, i.e. constant temperature and strain rate and the samples can be quenched immediately after the test (Fig. 2). Compared to uniaxial compression, plane strain deformation breaks the cylindrical symmetry which may conceal orientation changes with crystal rotations about the cylinder axis in subsequent texture measurements. Fig. 2: schematic illustration of the channel-die device. We selected 3 different starting orientations (Fig. 3), one with the basal plane parallel to the compression plane ((A) in Fig. 3), and two orientations with the basal plane perpendicular to the compression plane ((C,D) in Fig. 3). The latter two orientations differed in such a way that either the c-axis (type C) or a <10 10 > axis (type D) were parallel to the extrusion (rolling direction) (Fig. 4). Old reference system: RD0, ND0, TD0 New reference system: Type A: RD//RD0, ND//ND0, TD//TD0 Type C: RD//ND0, ND//RD0, TD//TD0 Type D: RD//RD0, ND//TD0, TD//ND0 Fig. 3: Sample types and their reference system for the Channel-Die experiments. The three orientations showed a conspicuously different deformation behaviour and texture formation during PSC. In orientations with the basal plane parallel to the compression plane (type A), twinning was suppressed, and the basal slip systems experienced virtually no shear stress. At low temperatures where non-basal systems could not be sufficiently thermally activated, the specimens showed low ductility (Fig. 5) and the texture did not change during deformation, both for low and high temperature deformation (Fig. 6). AZ31, TypeA max=11.7 AZ31, TypeC max=10.56 AZ31, TypeD max=9.39 Fig. 4: Starting textures ({0002}-pole figures) used for the channel-die experiments, schematic illustration of the three different specimen types with respect to the applied deformation: compression (ND) and tension (RD).

4 626 Textures of Materials - ICOTOM 14 Fig. 5: Deformation behaviour of Mg at RT and elevated temperatures; bottom: prior to deformation; center: after room temperature rolling; top: after rolling at 300 C. Type A max= 7.14 Type C max= 5.01 Type D max= 7.18 Type A max= 2.59 Type C max= 2.1 Type D max= Fig. 6: Textures after PSC of specimens type A (70%), C (20%), and D (50%) of AZ31 at 200 C; 2 1 strain rate s. Top row: {0002} pole figures; bottom row {10 10 }pole figures. In case of orientation C with a {1120 } plane parallel to the compression plane and the c-axis in rolling direction, basal, prism, and pyramidal slip remained inactive because they would lead to broadening which was suppressed by the channel-die geometry. Mechanical twinning was able to accommodate the strain in the beginning of deformation, but since this would reorient the crystals with the basal plane close to the compression plane (Fig. 6), twinning ceased to operate at larger degrees of deformation. Ductility improved over specimens of orientation A, but at low temperatures the strain to fracture remained small. In contrast, orientation D revealed a substantial improvement of ductility over samples with orientations A and C at temperatures as low as 100 C. In this case, prismatic slip could provide the necessary shape change during channel-die compression and rendered twinning unimportant. This caused the initial texture to remain mostly unchanged and thus, to maintain the favorable slip geometry up to large strains. These experiments were conducted on pure Mg and the major commercial wrought Mg-alloy AZ31 with similar results [12]. It is noted that in no case microstructure evolution was homogeneous. Shear banding was evident in all samples and would finally cause failure, although markedly less pronounced in samples of orientation D. Raising the deformation temperature to 200 C substantially improved the ductility although the texture changed little compared to lower deformation temperatures. This can be associated with the beginning of thermally activated <c+a> slip at elevated temperatures that increasingly accomplished the through thickness deformation. 2.3 Texture Simulation Due to the multiplicity of slip and twinning systems with different critical resolved shear stresses it is difficult to visualize deformation texture development and to associate deformation textures with deformation mechanisms. Texture simulation can reveal trends of

5 Materials Science Forum Vols texture evolution for given preferences of particular slip or twinning mechanisms. We used the LAPP code [13] to simulate plane strain compression textures for hexagonal materials including basal slip, prism slip, <c+a> pyramidal slip and {10 12 }<10 11> twinning with various ratios of critical resolved shear stresses (Fig. 7) to promote or disfavor specific systems. The starting texture was always random. If basal slip was predominant, a basal texture always did form. If basal slip was disfavored compared to other slip systems, the basal texture disappeared and the basal plane became parallel to the sheet plane normal with the c- axis essentially aligned with the rolling direction. These textures are not observed in Mg except for special geometries (see Sec. 2.1) because basal slip is generally predominant. The influence of <c+a> slip leads always to a splitting of the basal texture about TD. Max= 9.3 Max=9.36 Max=14.6 Max=44.24 τ c (b.)=1; τ c (<c+a>pyr.)=99 τ c (b.)=1; τ c (<c+a>pyr.)=3 τ c (b.) = 1; τ c (pr.) = 10 τ c (<c+a>pyr.) = 10 τ c (b.) = 1; τ c (prism.)= 0.5 τ c (<c+a>pyr.) = 0.5 Max=5.01 Max=11.3 Max=8.8 Max=10.2 τ c (b.)=1; τ c (<c+a>pyr.)=99 τ c { 10 12} tw. = 2 τ c (b.)=1; τ c (<c+a>pyr.)=3 τ c { 10 12} tw. = 2 τ c (b.) = 1; τ c (pr.) = 10 τ c (<c+a>pyr.) = 10 τ c { 10 12} tw. = 4 τ c (b.) = 1; τ c (prism.)= 0.5 τ c (<c+a>pyr.) = 0.5 τ c { 10 12} tw. = 4 Fig. 7: Results of texture simulations for different combination of slip systems and different relative CRSS values. The initial texture was random, and the total thickness reduction was set to 40%. Top: without twinning; bottom: with twinning. By contrast, addition of prism slip will lead to a scatter of the texture about RD. It is noted, however, that simulations will always enforce activation of slip systems that are necessary to accommodate the imposed strain, no matter how large the necessary stress. In reality, the propensity of a material to unstable flow at high stresses as obvious from the many shear bands, observed particularly at low deformation temperatures will cause premature cracking and fracture, i.e. a deviation from the predictions of simulation. Nevertheless, the observed tendencies are helpful for an interpretation of the observed textures with respect to potential deformation mechanisms. 3. Recrystallization Textures During annealing of cold deformed Mg the deformed microstructure changed to a distinct granular arrangement (Fig. 8). Obviously, the deformed material undergoes static recrystallization during annealing. Recrystallization is understood to proceed by nucleation of strain free grains and their subsequent growth until complete impingement. In fcc materials the recrystallization texture is usually different from the deformation texture and in many cases even more pronounced. By contrast, the recrystallization of Mg and its alloys is usually not accompanied with an obvious change of crystallographic texture, and this is frequently

6 628 Textures of Materials - ICOTOM 14 interpreted as continuous recrystallization, i.e. strong recovery. On the other hand, microstructural evolution during annealing of deformed Mg appears to proceed quite analogous to fcc metals, namely discontinuously. (a) (b) Fig. 8: OIM Map for isothermally annealed condition. (a) 200 C/15 sec. and, (b) 200 C/5 min. One problem with an investigation of the recrystallization texture is the strong basal texture of rolled Mg. No matter what the initial orientation distribution, except for very special cases, a basal texture (basal plane parallel to the rolling plane) readily develops due to crystallographic slip and twinning. However, the in-plane texture is much less pronounced although the orthorhombic specimen symmetry ought to prevent a homogeneous (0001) fiber texture. A specific difficulty in hexagonal crystals is the 6 fold rotation symmetric in the basal plane (Fig. 1). Depending on the activation of slip systems in the basal plane, either a <1120 > (single slip) or <10 10 > (double slip) direction will become parallel to the rolling direction, both of which are only 30 apart. If the deformation texture consists of mainly two components (0001) <1120 > and (0001) <10 10 > then the scatter of the components generates a pole figure of rolled Mg that is almost a (0001) fiber texture. Lücke and coworkers [14] conducted growth selection experiments on deformed Zn single crystals, and they found that such recrystallized grains would grow preferentially, which had a 30 <0001> orientation relationship to the deformed matrix. If indeed the deformation texture in Mg relates to the corresponding recrystallization texture by a 30 <0001> rotation, the texture change would be impossible to recognize in a (0001) pole figure, and because of the high symmetry in the basal plane it is even difficult to recognize it in a {1010} pole figure. We conducted a very accurate analysis of the in-plane {10 10 } pole distribution during annealing of warm deformed (150 C) pure Mg at different temperatures (Fig. 9). Evidently, after warm rolling, the maximum intensity of {10 10 } poles was rotated 30 from the rolling direction. On annealing, however, this intensity distribution changed and the maximum intensity shifted to the rolling direction [15]. This shift was even strengthened at higher temperatures (see Sec. 4). This result can be readily interpreted by a 30 <0001> preference during nucleation and growth of the recrystallized grains. It is still to be determined whether the nuclei bear already this orientation relationship or whether this preference develops by competition during nucleus growth. In any event, the results demonstrate that recrystallization proceeds discontinuously and not by recrystallization in-situ. f(g) ϕ 1 As Rolled 200 C/1h 300 C/1h 350 C/1h Fig. 9: (0001) fibre intensity as function of the rotation angle about the fibre axis for isochronally annealed samples and the as rolled condition For the sake of completion it is mentioned that Mg also undergoes substantial grain growth at temperatures above 400 C. Sometimes very large grains were observed due to abnormal grain growth, the original of which remains to be investigated. Except for the very large grains, the recrystallization texture was changed very little, but decomposed eventually into individual orientations once the grain size became excessively large.

7 Materials Science Forum Vols High Temperature Deformation and Dynamic Recrystallization At deformation temperatures above 250 C non basal and <c+a> slip are sufficiently activated to convey excellent formability to the material for sheet production. However, at the same time the material is liable to undergo recrystallization during deformation, i.e. dynamic recrystallization. It is of particular interest to elucidate the texture forming mechanisms during hot working since most commercial Mg alloys will be fabricated to semi-finished products by such processing. We investigated the combined effect of non-basal slip and dynamic recrystallization on the texture development in pure Mg and alloy AZ31. The starting material was produced by hot extrusion. The extruded pure Mg had a pronounced basal texture (Fig. 10), while the as extruded AZ31 had the appearance of an <10 10 > RD fiber texture, so that the basal plane was parallel but arbitrarily rotated about RD (Fig. 11). Upon PSC of pure Mg at 200 C, 300 C, and 400 C, the textures changed in such a way that a basal component texture (0001) <10 10 > developed at 200 C while an increasingly sharp (0001) fiber texture prevailed at higher temperatures (Fig. 10). In AZ31 the RD<10 10 > fiber texture was essentially conserved during elevated temperature deformation, although the intensity distribution along the fiber was liable to some scatter with a basal component and a 70 RD rotated component always prevailing as the main orientations in the texture (Fig. 11). The corresponding microstructures revealed a recrystallized granular structure with growing grain size for rising deformation temperatures (Fig. 12). Interestingly, the {10 10 } pole figure also revealed the development although weak of an intensity in TD direction corresponding to a 90 <0001> rotation of the (0001) <10 10 > main texture component. The grain size in AZ31 remained comparably small, even during deformation at 400 C (Fig. 13). Mg, initial max=9.52 Mg, 200 C max=12.14 Mg, 300 C max=51.3 Mg, 400 C max=62.86 Mg, initial max=3.23 Mg, 200 C max=3.39 Mg, 300 C max=5.81 Mg, 400 C max=7.38 Fig. 10: Texture of AZ31 prior to PSC and after 70% PSC at different temperatures. Top row: (0002) pole figures; bottom row: {10 10 } pole figures.

8 630 Textures of Materials - ICOTOM 14 AZ31, initial max=4.7 AZ31, 200 C max=5.2 AZ31, 300 C max=4.6 AZ31, 400 C max=5 AZ31, initial max=10 AZ31, 200 C AZ31, 300 C max=2.6 AZ31, 400 C max=5.2 max=8 Fig. 11: Texture of AZ31 prior to PSC and after 70% reduction at various temperatures. Top row: (0002) pole figures; bottom row: { } pole figures. Mg, 200 C Mg, 300 C Mg, 400 C Fig. 12: Effect of deformation temperature on the microstructure development in pure Mg. AZ31, 200 C AZ31, 300 C AZ31, 400 C Fig. 13: Microstructure of AZ31 after PSC at different temperatures. An evaluation of the strength of the texture components revealed that with increasing temperature both the basal texture component and its 30 <0001> rotated component grew. Of particular interest is the (angular) misorientation distribution at different temperatures (Fig. 14). Apparently, the distributions were far from random but two distinct maxima stand out, at low angles and around 30. The peak at low angles actually decreased with rising

9 Materials Science Forum Vols temperature from 200 C to 300 C. The low angle boundaries were obviously deformation induced subboundaries, while the peak at 30 indicated that both the (0001) <1120 > and the (0001)<10 10 > components were major elements of the texture. This is also obvious from an orientation image of the microstructure of partially recrystallized AZ31 (Fig. 15). 200 C 300 C Fig. 14: Misorientation angle distribution for AZ31 at 200 C and 300 C. The curve represents a random distribution. Fig. 15: The micrograph on the left gives an OIM map of crystal directions parallel to RD. The elongated (deformed) grains have a <1210 >, the small globular (recrystallized) grain have mostly a <10 10 > direction parallel to RD. The deformed grains had mainly an orientation < > RD while the recrystallized grains overwhelmingly developed a <12 10 > RD texture component. These observations can be readily reconciled with a concept that dynamic recrystallization comprises a superposition of deformation and static recrystallization as discussed previously. The stability of the RD fiber texture for AZ31 even at high temperatures is surprising (Fig. 11). This may also be due to dynamic recrystallization. If recrystallization promotes 30 <0001> rotations, the distribution of basal planes will not change and by this the recrystallization may actually stabilize or at least delay decomposition of the RD fiber texture. 6. Conclusions We investigated the texture formation of pure Mg and Mg alloy AZ31 during plane strain compression, static recrystallization, and dynamic recrystallization. The following results were obtained a) Low temperature deformation is difficult due to insufficient activation of non-basal slip systems. Twinning can only partially contribute to deformation, since it ceases when a basal texture is established. However, for very special geometries also substantial ductility can be achieved at low temperatures if mainly prism a slip is activated.

10 632 Textures of Materials - ICOTOM 14 b) Taylor simulations of plane strain compression revealed the effect of the individual deformation mechanisms on texture development, which was very different for dominant basal, prismatic a and pyramidal <c+a> slip, besides twinning. c) The annealing textures of cold deformed Mg strongly resembled the deformation textures, in particular when viewed in terms of (0001) pole figures. A more detailed analysis revealed a 30 <0001> orientation relationship between deformation and recrystallization texture. Therefore, recrystallization is a discontinuous nucleation and growth process, not a recovery process like recrystallization in-situ or continuous recrystallization. d) Hot forming (T>200 C) caused dynamic recrystallization besides thermal activation of <c+a> slip. The textures became very pronounced at high temperatures. This can be understood from the coarser dynamically recrystallized grain structure which promotes basal slip and conservation of the basal orientation during recrystallization. Acknowledgement Financial support by the Deutsche Forschungsgemeinschaft (DFG) through grant Go 335/27 is gratefully acknowledged. The authors like to thank Dr. Gehrmann and Dr. Nadella for valuable discussions. References [1] K.U. Kainer (Ed.): Magnesium-Eigenschaften, Anwendungen, Potenziale, Wiley-VCH, Weinheim 2000 [2] S. R. Agnew, O. Duygulu: Mat. Sci. Forum Vols (2003), p [3] M. H. Yoo, S. R. Agnew, J. R. Morris, K. M. Ho: Mat. Sci. Eng. A Vols (2001) p [4] S. R. Agnew, M. H. Yoo, C. N. Tomé: Acta mater. Vol 49 (2001), p [5] J. J. Fundenberger, M. J. Philippe, F. Wagner, C. Esling: Acta mater. Vol 45 (1997), p [6] M. J. Philippe, F. Wagner, F. E. Mellab, C. Esling, J. Wegria: Acta metall. et mater. Vol 42 (1994), p [7] M. J. Philippe, M. Serghat, P. Van Houtte, C. Esling: Acta metall. et mater. Vol 43 (1995), p [8] M.R. Barnett, M.D. Nave, C.J. Bettles: Mat. Sci. Eng. A Vol 386 (2004), p [9] M.D. Nave, M.R. Barnett: Scr. Mat. Vol. 51 (2004), p [10] A. Styczynski, Ch. Hartig, J. Bohlen, D. Letzig: Scr. Mat. Vol. 50 (2004), p [11] W. J. Kim, S. I. Hong, Y. S. Kim, S. H. Min, H. T. Jeong, J. D. Lee: Acta mater. Vol 51 (2003), p [12] R. Gehrmann, M. Frommert, G. Gottstein: Mat. Sci. Eng. A (2005) in press [13] U.F. Kocks, J.S. Kallend, H-R Wenk, A.D. Rollet, S.I. Wright: Preferred Orientation Package Los Alamos (1994) [14] R. Klar, K. Lücke: Z. Metallkde. Bd. 59 (1968), p [15] R.K. Nadella, I. Samajdar, G. Gottstein, in: K.U. Kainer (Ed.), Proceedings of the 6 th International Conference on Magnesium Alloys and their Applications, September 2003, Wiley-VCH, Wolfsburg, Germany, p.1052

11 Textures of Materials - ICOTOM / Texture Development in Pure Mg and Mg Alloy AZ / DOI References [1] K.U. Kainer (Ed.): Magnesium-Eigenschaften, Anwendungen, Potenziale, Wiley-VCH, einheim / ch94 [2] S. R. Agnew, O. Duygulu: Mat. Sci. Forum Vols (2003), p / [3] M. H. Yoo, S. R. Agnew, J. R. Morris, K. M. Ho: Mat. Sci. Eng. A Vols ) p /S (01) [4] S. R. Agnew, M. H. Yoo, C. N. Tomé: Acta mater. Vol 49 (2001), p /S (01)00297-X [5] J. J. Fundenberger, M. J. Philippe, F. Wagner, C. Esling: Acta mater. Vol 45 (1997), p /S (97) [6] M. J. Philippe, F. Wagner, F. E. Mellab, C. Esling, J. Wegria: Acta metall. et mater. Vol 2 (1994), p / (94) [7] M. J. Philippe, M. Serghat, P. Van Houtte, C. Esling: Acta metall. et mater. Vol ), p / (94)00329-G [8] M.R. Barnett, M.D. Nave, C.J. Bettles: Mat. Sci. Eng. A Vol 386 (2004), p /j.msea [9] M.D. Nave, M.R. Barnett: Scr. Mat. Vol. 51 (2004), p /j.scriptamat [12] R. Gehrmann, M. Frommert, G. Gottstein: Mat. Sci. Eng. A (2005) in press /j.msea [15] R.K. Nadella, I. Samajdar, G. Gottstein, in: K.U. Kainer (Ed.), Proceedings of the 6th nternational Conference on Magnesium Alloys and their Applications, September 2003, iley-vch, Wolfsburg, Germany, p / ch163