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Comparative study of the deformation behavior of hexagonal magnesium–lithium alloys and a conventional magnesium AZ31 alloy
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Acta Materialia 57 (2009) 2229–2242 www.elsevier.com/locate/actamat
Comparative study of the deformation behavior of hexagonal magnesium–lithium alloys and a conventional magnesium AZ31 alloy T. Al-Samman * Institut fu¨r Metallkunde und Metallphysik, RWTH Aachen, Kopernikusstr. 14, 52056 Aachen, Germany Received 6 November 2008; received in revised form 21 January 2009; accepted 23 January 2009 Available online 28 February 2009
Abstract Specimens of a conventional magnesium AZ31 alloy and a binary a-solid solution Mg4Li alloy with similar starting textures and microstructure were subjected to plane strain deformation under various deformation temperatures ranging from 298 K to 673 K. Lithium addition to magnesium exhibited remarkable room temperature ductility improvement owing to enhanced activity of non-basal slip, particularly, hc + ai-slip mode. Furthermore, the addition of lithium to magnesium seemed to reduce the plastic anisotropy, typical for commercial magnesium alloys. This was evident in the flow curves and texture development obtained at 200 °C and 400 °C. At 400 °C prismatic slip gains strong influence in accommodating the imposed deformation. In terms of thermal stability against microstructure coarsening at elevated temperatures, the lithium containing alloy undergoes significant grain growth following recrystallization. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Texture; Slip; Twinning; Plastic deformation; Ductility
1. Introduction Below 498 K plastic deformation of polycrystalline magnesium is limited to basal (0 0 0 1)h1 1 2 0i-slip and to pyramidal f1 0 1 2gh1 0 1 1i-twinning [1]. The magnesium crystal, thus has only three geometrical and two independent slip systems compared to aluminum, for example, which has 12 {1 1 1}h1 10i geometrical and five independent slip system for arbitrary shape change [2]. Pure magnesium and conventional magnesium alloys therefore have a tendency to embrittlement at low deformation temperatures due to intergranular failure, and localized transcrystalline fracture along twinned regions, particularly upon compression [3], or along basal (0 0 0 1) planes for very coarse grains [1]. It is proposed that the room temperature ductility of magnesium can be enhanced by:
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(i) Grain refinement obtained via severe plastic deformation or conventional dynamic recrystallization, which can promote grain boundary sliding and/or grain boundary rotation allowing easier accommodation of the overlap or void displacement between two adjacent grains [4]. (ii) Altering the texture observed in conventional magnesium alloy sheet material. Typical wrought magnesium alloys usually exhibit strong textures with little variations between them. The typical sheet basal texture places the vast majority of grains, such that their basal (0 0 0 1) planes are close to the plane of the sheet. With such an orientation, i.e. c-axis compression, it is difficult to deform. In the literature, many indications can be found about the improved properties of magnesium alloys containing additions of yttrium (Y) and rare earth elements (RE), such as cerium (Ce), neodymium (Nd) and gadolinium (Gd). Ball et al. [5] reported that an Mg–Y–Nd mischmetal-based alloy can develop more random textures during extrusion compared to conventional alloys. They found that a more randomized texture strongly
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.01.031
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reduces the tension–compression asymmetry, which is an attractive benefit of controlling the texture to enhance formability. Another study by Laser et al. [6] showed that texture modification is also possible in other alloys without much additions of expensive RE elements. They studied four different alloys based on conventional AZ31 containing additions of cerium and calcium. The alloys were hot rolled and tensile tested at room temperature. It was observed that the addition of less than 1 wt.% Ce mischmetal and Ca reduced the peak intensity in the basal pole figure to only 3 multiples of a random distribution (MRD), as compared to 9 MRD for a conventional AZ31 alloy. Not only was the texture weaker, but the qualitative character of the texture was also altered. (iii) Lithium addition. The Mg–Li system (Fig. 1a) [7] is quite interesting. Not only because the addition of lithium offers a great potential to enhance the room temperature ductility of the material (brittle to ductile transition temperature is suppressed to below liquid nitrogen temperature) [8], but also because it allows to substantially lower the density of the already light magnesium even more (qLi = 0.58 g/cm3 vs. qMg = 1.74 g/cm3) (Fig. 1b) [9]. Hence, magnesium– lithium alloys are the lightest structural metals.
Fig. 1. (a) Mg–Li binary phase diagram [7]. (b) The effect of lithium addition on the alloy density [9].
The purpose of the current study is to investigate the influence of lithium addition on the various deformation modes and their contribution to deformation, particularly, on the development of crystallographic texture and microstructure during deformation at room temperature and at higher deformation temperatures of 200 °C and 400 °C. The deformation behavior, texture and microstructure development of the investigated Mg–Li alloy are compared with those of a commercial magnesium AZ31 alloy, deformed exactly under the same conditions. The findings of this study will be used in future investigations to set directions for new Mg–Li–X alloy development. 2. Background of Mg–Li alloys Lithium addition to magnesium has an important effect on the crystal structure, i.e. the lattice parameters. With increasing lithium content the c/a axial ratio decreases (from 1.624 for pure Mg to 1.607 for Mg–17 at.% Li, close to the solid-solubility limit) [10]. Between 17 and 30 at.% Li contents magnesium–lithium alloys exhibit two phase structures consisting of the a-magnesium rich (hcp) and b-lithium rich (bcc) phases. The b single phase structure exists for Li contents greater than 30 at.% (Fig. 1a). The body centered cubic b phase is soft and ductile. Hence, magnesium alloys with such high lithium content are very ductile. Nevertheless, and aside from high production costs, these alloys suffer from low temperature instability. The improved room temperature ductility of a-solid solution Mg–Li alloys is attributed to a substantial increase in non-basal slip activity, particularly, prismatic slip. While prismatic slip in pure magnesium is negligible at room temperature, and only limited to regions with stress concentrations in the vicinity of grain boundaries, extensive prismatic slip over entire grains occurs in magnesium alloys containing lithium [11]. With respect to basal slip, Quimby et al. [12] reported that its critical resolved shear stress (CRSS) increased with increasing lithium content, i.e. ds/dcLi > 0, where s is CRSS for basal slip and cLi is the concentration of lithium solutes. The alloy hardening for basal slip was attributed to interaction between dislocations and Li solute atoms [13]. It should be noted that the decrease in the c/a ratio resulting from lithium addition is not the only reason for the enhanced activity of prismatic slip. Beryllium, for example, has a c/a ratio of 1.568, less than that of zirconium and titanium (1.593 and 1.588, respectively), and yet deforms primarily by basal slip, where both Zr and Ti have prismatic slip as the primary deformation mode. Another – more important – reason for the softening in non-basal slip is the influence of lithium addition on the dislocation mechanisms and the stacking fault energy related to non-basal slip planes. The complex ds/dcLi < 0 behavior, i.e. softening for prismatic slip was attributed by Ahmadieh et al. [11] to a reduction in the Peierls stress when adding lithium. The theoretical reason for this effect
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was not reported. Urakami et al. [14] performed electron microscopic investigation of prismatic slip in Mg–10.5 at.% Li single crystals deformed in tension at liquid nitrogen and ambient temperatures. Their theory for the rate controlling mechanism for low temperature prismatic slip of Mg–Li alloys was based upon friction stresses and mobility of screw dislocations. In another study, the authors proposed that prismatic slip occurs by a double kink mechanism at interaction sites between dislocations and Li atoms. The activation energy for the nucleation of double kinks is reduced by the presence of lithium [15]. The investigations mentioned earlier [8,11–15], that attributed the extensive low temperature ductility improvement of a-solid solution Mg–Li alloys to the availability of additional deformation modes, i.e. prismatic slip, were all carried out on single crystals. However, in polycrystal Mg–Li alloys, even if both basal and prismatic slip were active, accommodating deformation along the c-axis will still be a problem (considering the unidirectional nature of mechanical twinning and hence their limited ability to accommodate deformation). The fact that polycrystal hcp Mg–Li alloys can undergo c-axis compression at room temperature up to large strains with no signs of mechanical instability or fracture (as shown in the current work) suggests a significant involvement of hc + ai-slip as well. In fact, a texture simulation study by Agnew et al. [16], supported by TEM investigation of a-solid solution Mg–Li alloys from the same authors [17] reported that hc + ai-slip on f1 1 2 2g second-order pyramidal planes was indeed active during compression at ambient temperatures. 3. Experimental procedure For the current study a-solid solution Mg–15 at.% Li (Mg4Li) (wt.%) alloy and a commercial magnesium AZ31 alloy were used. Both materials were available in the asrolled and as-extruded state. The processing history of the Mg4Li alloy comprised a thickness reduction from 12 mm to 6 mm for the rolled material (same for AZ31) and a cross section reduction from Ø 30 mm to Ø 15 mm for the extruded material. The cross section reduction for the commercially extruded AZ31 material was from Ø 180 mm to Ø 60 mm. The processing temperature for the AZ31 material (both conditions extruded and rolled) was 400 °C. The Mg– Li material was thermo-mechanically processed at 300 °C. Prior to the deformation tests both materials were annealed at 350 °C for 1 h. The microstructures are shown in Fig. 3. In order to cover a wide range of investigations, four specimen types with different starting textures were investigated (Fig. 2a and 2b). The starting textures were obtained by simply cutting the samples in such a way that the basal plane was essentially parallel to any desired direction for the given plane strain geometry (Fig. 2c). The specimen dimensions (types 1, 2 and 3, Fig. 2a) for the plane strain compression (PSC) tests were 14 mm (RD) 10 mm (TD) 6 mm (ND). The dimensions of cold rolling specimens (type 4, Fig. 2b) were 60 mm 30 mm 4 mm.
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Fig. 2. Orientation of the investigated specimens (a) types 1, 2 and 3 (extruded starting condition); (b) type 4 (rolled starting condition) and (c) selection of a new reference system represented by the rolling (RD), transverse (TD) and normal direction (ND) of the channel-die geometry with respect to the original loading direction (extruded ED or rolled RD).
For room temperature deformation, cold rolling and plane strain compression using a channel-die device were chosen. Extension occurred only in the longitudinal direction of the channel-die (parallel to RD, Fig. 2c) since deformation in the transversal direction was suppressed, owing to the deformation geometry (lateral constraints). Deformation experiments at elevated temperature comprised channel-die compression only. Deformation temperatures and strain rates ranged between 200 °C to 400 °C and 102 s1 to 104 s1, respectively. Hexagonal boron nitride (h-BN) powder was used as a lubricant for minimizing friction between sample and channel-die. For cold rolling few drops of industrial oil was used for lubrication. After completion of the tests, specimens for optical microscopy were shortly ground with a 4000 grit SiC paper and subsequently mechanically polished with diamond paste down to 3 and 1 lm, respectively. Final polishing was performed using a colloidal silica solution. After polishing, specimens were etched in acetic picral to visualize grains and grain boundaries. Same preparation procedure was used for both alloys. The macrotexture was determined in the mid-plane of deformed specimens by measuring incomplete pole figures (5° 6 a 6 75°) in the back reflection mode using CoKa radiation. A set of six measured pole figures [f1 0 1 0g; f0 0 0 2g; f1 0 1 1g; f1 0 1 2g; f1120g, and {10 13}] was used to calculate the orientation distribution function (ODF) using the positivity method [18]. 4. Results 4.1. Initial condition Fig. 3 shows the initial microstructures of the alloys AZ31 and Mg–Li following annealing at 350 °C for 1 hour. For the extruded Mg4Li material, the microstructure
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Fig. 3. Initial microstructures prior to the PSC tests. (a) Mg4Li extruded condition; (b) AZ31 extruded condition; (c) Mg4Li rolled condition; (d) AZ31 rolled condition.
revealed equiaxed grains with a homogeneous grain size distribution (Fig. 3a). Interestingly, the extruded microstructure showed no signs of deformed elongated grains. Such grains are usually a common microstructural feature of commercially extruded AZ31 material (Fig. 3b).The rolled Mg4Li material exhibited a microstructure that was similar to that of the extruded material. The mean grain size (linear intercept) was larger than that of the extruded material (80 lm for the rolled material compared to 50 lm for the rolled one) (Fig. 3c). Some deformation twins from previous processing were still visible after annealing. No second phase precipitates were observed. The average grain sizes of the rolled and the extruded AZ31 alloy were 40 lm and 25 lm, respectively. The starting textures of the different test series of both alloys are shown in Fig. 4. For specimen types 1 and 2 obtained from the extruded materials (compare with Fig. 2a) the textures showed little difference in terms of texture characteristics, yet the difference was noticeable in terms of texture strength. AZ31 starting textures were much more pronounced. This is attributed to the different chemical composition, as well as to the different extrusion
ratio for both alloys. The textures of specimen type 3 (also obtained from the extruded materials) showed perceptible qualitative difference of texture components. AZ31 starting texture, i.e. RD h1 0 1 0i was also much more pronounced. Specimen type 4 obtained from the rolled Mg4Li material exhibited an unconventional rolling texture, characterized by a relatively weak intensity (5 mrd) and a large h0 0 0 1i scatter around ND. The rolled AZ31 material revealed a typical strong rolling texture (12 mrd) with a maximum intensity split. 4.2. Room temperature deformation 4.2.1. Cold rolling The cold rolling behavior of the alloys AZ31 and Mg4Li is shown in Fig. 5. As expected, the commercial AZ31 material exhibited typical brittle response and failed at a total thickness reduction of 26% despite low reduction per each rolling pass (0.1 mm/pass). By contrast, the lithium containing material remained in excellent shape and revealed no signs of cracks (Fig. 5a). During further cold rolling, the Mg4Li material conveyed remarkable ductility
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Fig. 4. Starting textures (prior to PSC) represented by means of (0 0 0 2)- and {1 0 1 0}- recalculated pole figures. Intensity levels 1:2:3:4:5:6:7:8:9: 10:12:14:16:18:20.
reaching a total thickness reduction of 86% (Fig. 5b). The surface quality was similar to the quality usually obtained for AZ31 by hot rolling at 400 °C with intermediate annealing. The corresponding microstructures (observed in the transverse plane) are shown in Fig. 6. The microstructure of the failed AZ31 material is presented in Fig. 6a, while Fig. 6b and c depict the microstructures of Mg4Li specimens at 26% and 86% thickness reduction, respectively. The rolling textures are shown in Fig. 7 in terms of basal (0 0 0 2) and prismatic {1 0 1 0 } recalculated pole figures.
4.2.2. Channel-die compression The mechanical response of AZ31 and Mg4Li materials during room temperature plane strain deformation at a constant strain rate of 104 s1 using a channel-die device is shown in Fig. 8 in terms of stress–strain curves. Stress–strain data after e = 0.6 were not taken into account due to friction between sample and channeldie, which shows in the flow curve as a false increase of stress at high strains. Furthermore, it should be noted that in case of room temperature deformation the recorded stress-strain data do not actually show true val-
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different contribution of the deformation mechanisms during compression. In the current case, i.e. room temperature deformation the fact that the presented textures were measured at two completely different failure strains should be also considered. 4.3. Elevated temperature deformation
Fig. 5. Cold rolling behavior of Mg4Li. (a) Direct comparison with AZ31 at 26% total thickness reduction (AZ31 failed). (b) Mg4Li cold rolled up to 86% final thickness reduction demonstrating remarkable room temperature formability.
ues, but rather nominal ones because of the large heterogeneity of deformation. As evident from the flow curves, the specimens containing lithium demonstrated a ductile room temperature deformation behavior reaching failure strains above 1.2 (not shown in Fig. 8). By contrast, AZ31 specimens underwent failure by shear localization at failure strains ranging from ef = 0.15 to 0.24. The cold deformation microstructures after channel-die compression are depicted in Fig. 9a and b for AZ31 (compressed to failure) and Mg4Li (compressed to e = 1), respectively. The microstructure of both transverse (CD-RD) and compression surfaces (RDTD) was examined. In case of Mg4Li no noticeable difference between the two surfaces was observed. For AZ31 the microstructure of the compression surface comprised some microcracks that were probably initiated from deformation twins. For the different types of each alloy the microstructure development was similar, and hence, only one microstructure for each alloy is presented. The microstructure of Mg4Li remained coarser than the microstructure of AZ31. The evolution of macrotextures of the different specimen types of AZ31 and Mg4Li materials during room temperature PSC tests to failure is presented in Fig. 10. A comparison of the results reveals noticeable differences between both materials in terms of texture characteristic and maximum intensity. This is mainly attributed to the
For investigating the elevated temperature deformation behavior of both materials channel-die compression tests at 200 °C and 400 °C at a constant strain rate of 104 s1 were carried out. Hot rolling experiments were not included in the current investigations because of the difficulty of controlling the rolling temperature. The flow curves of the investigated materials at elevated temperatures are shown in Fig. 8b for 200 °C and in Fig. 8c for 400 °C, respectively. As evident, the flow behavior was influenced by the deformation conditions (temperature and strain rate), the starting texture and the chemical composition of the deformed material. This is discussed in more details below. In analogy with stress-strain data presented in Fig. 8a for room temperature deformation, the curves in Fig. 8b and c show only data up to a true strain of e = 0.6. The tests were, however, carried out up till e = 1.4. The corresponding microstructure development is shown in Fig. 9c–f for PSC tests at 200 °C (c and d), and 400 °C (e and f). It should be noted, that it was clear from the beginning that the Mg4Li material would always exhibit a coarser microstructure than AZ31, and therefore, it was given no attention in Fig. 9 to show micrographs taken at the same magnification. 5. Discussion 5.1. Deformation behavior and formability Commercial magnesium sheet does not currently exhibit sufficient formability. Modifying the c/a ratio of its hexagonal close packed lattice and the crystallographic texture by lithium addition offers a great potential to improve the room temperature formability of magnesium wrought alloys. This is clearly demonstrated in Figs. 5 and 8a which show the cold deformation response of a binary Mg4Li alloy in comparison with a standard AZ31 alloy in terms
Fig. 6. Cold rolling microstructures of (a) AZ31 and (b) Mg4Li at 26% total thickness reduction. (c) Microstructure of Mg4Li at 86% final thickness reduction. The micrographs were taken from the transverse plane of the rolled specimens.
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Fig. 7. Macrotexture evolution after cold rolling of (a) AZ31 to failure (26%), (b) Mg4Li at 26% thickness reduction and (c) Mg4Li at 86% final thickness reduction. The textures are presented in terms of (0 0 0 2)- and {1 0 1 0}-recalculated pole figures. Levels: 1:2:3:4:5:6:7:8:9:10:12.
of sheet formability and stress-strain flow curves. While the standard wrought alloy AZ31 failed at typical thickness reduction and strain values during cold rolling and cold channel-die compression, the lithium containing alloy exhibited a substantial ductility enhancement with no signs of mechanical instability or fracture up to very high strains (e > 1). This remarkable ductility improvement is attributed (as stated earlier) to alloy softening for non-basal slip resulting not only from the slight decrease of the axial c/a ratio, but also from the influence of lithium solutes on the dislocation dynamics. Although the samples of both materials were tested under the same deformation conditions, i.e. the same temperature and strain rate, a markedly different mechanical response for each specimen type was observed (particularly shows in Fig. 8a). This is attributed to the different starting texture each specimen type had. Each initial texture favors specific combination of deformation mechanisms, which is reflected during deformation by a different flow behavior. Comparing the flow curves of both materials during PSC deformation at room temperature indicates much lower flow stress values for Mg4Li than the commercial alloy AZ31. This is also true in case of elevated temperature
Fig. 8. Stress–strain curves of AZ31 and Mg4Li specimens during channel-die compression tests (a) at room temperature (b) at 200 °C (c) at 400 °C. e_ = 104 s1.
deformation. While AZ31- type 1 specimen exhibited a flow stress maximum of 400 MPa at e = 0.18 before fracture, Mg4Li specimen of the same type exhibited a flow stress value of 265 MPa at the same strain. With increasing strain some strain hardening takes place, but the maximum strain values remain under 300 MPa (without considering friction at e > 0.6). This modest work hardening behavior at room and elevated temperatures (see also Fig. 8b and c) is one of the aspects binary Mg–Li alloys suffer from, when compared with other magnesium alloys (AZ, ZK and AM series). During elevated temperature deformation an interesting behavior of the Mg4Li alloy was recorded. Comparing the flow curves of both materials at 200 °C (Fig. 8b) and 400 °C (Fig. 8c) shows that the mechanical response, i.e. flow behavior of the Mg4Li spec-
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Fig. 9. Deformation microstructures of AZ31 (a, c and e) and Mg4Li (b, d and f) after channel-die compression at room temperature (a and b); at 200 °C (c and d); and at 400 °C (e and f). The strain corresponding to micrograph (a) was ef = 0.14. For micrograph (b) e = 1. For micrographs (c–f) e = 1.4. e_ = 104 s1. tt, tensile twins; ct, compression twins.
imens seemed to be not very much affected by the starting texture. By contrast, the flow curves and hardening response of the AZ31 specimens were strongly dependent on the initial orientation of grains and the evolution of texture during deformation. A comparison of the recorded yield stress data for AZ31 and Mg4Li is depicted in Fig. 13. Obviously, increasing the deformation temperature resulted for both materials a significant drop of the yield stress. Mg4Li specimens exhibited much lower yield stress values than AZ31. Interestingly, the lithium containing specimens showed more or less the same yield stress strength, irrespective of specimen type (starting texture). In case of AZ31, the yield stress revealed different values for different starting orientations. With respect to the influence of the starting texture and the deformation mechanisms on the flow behavior of the
specimens (particularly AZ31) during channel-die compression at RT and 200 °C, two kinds of plastic flow can be observed: (1) If mechanical twinning was more favorable than crystallographic slip for deformation, then the work hardening part of the flow curve would result from twinning driven deformation (characterized by a concave shape of the flow curve and significant work hardening regime). (2) If at incipient deformation mechanical twinning was inhibited (very low Schmid factor or for deformation geometrical reasons), and hence, crystallographic slip was the primary deformation mechanism, the work hardening part of the flow curve would be less evident than the previous case (slip driven deformation).
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Fig. 10. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at room temperature at a constant strain rate of 104 s1. The strains corresponding to the presented textures are ef = 0.14 for AZ31 and e = 1 for Mg4Li. Levels: 1:2:3:4:5:6:7:8:9:10:12:14:16:18:20:25.
The hardening effect of twins can be seen in the flow curves of type1 specimens (particularly for AZ31, Fig. 8a and b), which possessed a starting orientation that is very favorable for tensile twinning (c-axes perpendicular to compression axis, Fig. 4). Twin-induced hardening in magnesium has been frequently reported [16,19]. It is important to note that the conspicuous hardening effect of twins is attributed to the stress required to initiate slip in the twinned regions of the matrix rather than to the activation stress of the twinning process itself. Crystallographic slip within the twin interior is hard because (i) f1 0 1 2g-tensile twinning reorients the basal planes in the twinned parts of the matrix into unfavorable orientations for easy basal slip, and (ii) the mean free path of dislocation movement within the twin interior is too short in comparison with the matrix. The above aspects are the reason why the flow stress in grains undergoing tensile twinning strongly rises with strain compared to other specimen types, in which grains are primarily undergoing crystallographic slip. During deformation at 400 °C the flow behavior of all specimen types was driven by slip as the dominant deformation mechanism. Mechanical twinning in commercial magnesium alloys is usually limited to deformation temperatures below 250 °C.
5.2. Microstructure development During room temperature deformation the development of the microstructure of failed AZ31 specimens revealed two common features: high fraction of twinned areas and microcracks. Previous investigations on AZ31 [3] showed that deformation at ambient temperatures can activate f1 0 1 1g-compression twinning and double twinning. Both twinning types can be distinguished from the typical tensile twins by their thin morphological appearance. In Fig. 5a the rolling microstructure of the failed specimen at 26% thickness reduction show deformation twins propagating along the transverse section at approximately ±45° to the rolling direction. The microstructure also reveals several locations of voids (microcracks) with the same alignment of twins, i.e. ±45° to RD. According to earlier investigations [3], these twin-sized voids could be initiated from the preexisting twins. The same microstructural combination of high volume fraction of twins and microcracks was also found in the compression surface of the RTPSC deformed specimen (type 1) to failure (Fig. 9a). By contrast, the optical micrographs of Mg4Li specimens deformed at room temperatures show lower volume fraction of twins and no signs of voids. The fact that both twin
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types, i.e. tensile (broad) and compression (thin) twins were evident in the deformation microstructure (Fig. 9b) suggests that mechanical twinning remains an important deformation mechanism despite sufficient availability of non-basal slip modes. The microstructure of Mg4Li specimen cold rolled up to a final thickness reduction of 86% comprised large-scale deformation heterogeneities in the form of very fine shear bands aligned at approximately ±35° from the rolling direction RD. This is quite expected at such high level of deformation. Similar shear band microstructure can be found in AZ31 during large strain hot rolling (LSHR) up to high thickness reductions (usually >75%). During deformation at higher temperatures (200 °C) both investigated materials showed common microstructural features comprising partial dynamic recrystallization (DRX) and a grain microstructure that is elongated towards the expansion direction of the deformed specimen, i.e. rolling direction RD. The unique directionality of the microstructure observed in Fig. 9c and d is a creation of the strain path of plane strain deformation. Usually, a material undergoing DRX at deformation temperatures between 200 °C and 250 °C exhibits a typical stress-strain behavior characterized by an initial work hardening regime (reaching a peak stress) followed by rapid work softening. In the current investigations this was mostly evident in case of AZ31-type 1 specimen (Fig. 8b). In other specimen types (also Mg4Li specimens) the work softening was less evident, and the flow curves showed steady-state behavior straight after work hardening. This could suggest that the kinetics of DRX is influenced by the texture, as it was reported recently by del Valle [20]. The grain size recrystallized dDRX, however, was found to be insensitive to texture [20,21]. The grain size recrystallized for AZ31 and Mg4Li was 4 lm and 10 lm, respectively. With regard to nucleation of DRX, it was shown in a previous study [21] that DRX can take place at twin boundaries and within twin interiors, provided that the vast majority of grains undergo massive twinning. In case of Mg4Li twinning was not important, and hence DRX nucleated mostly at grain boundaries and triple junctions (Fig. 14). During deformation at elevated temperatures (400 °C) the deformed Mg4Li specimen revealed a coarse granular microstructure comprised of grains with intercept lengths of 100–200 lm (Fig. 9f). This was obviously an indication of grain growth. By contrast, AZ31 specimens undergoing plastic deformation at the same conditions (400 °C/104 s1) did not reveal any signs of noticeable grain growth after DRX (owing to the eutectic constituent Mg17Al12 dissolved in the matrix). The microstructure was fully recrystallized and the grain size remained comparatively very small (d 20 lm). This concludes that the binary Mg4Li alloy behaves similarly to pure magnesium deformed or annealed at such elevated temperatures, whereat the occurrence of grain growth is inevitable.
5.3. Deformation mechanisms and texture evolution The starting texture of the Mg4Li rolled material (Fig. 4, type 4) showed a typical basal orientation (basal planes lying close to the rolling plane), yet with a higher maximum spread around the normal direction of the sheet ND than usual. There was no evident preference for any crystallographic direction to align with the rolling direction. Additionally, the texture intensity was nearly half of that of rolled AZ31 (Fig. 4, type 4) taking into account similar thickness reduction. After cold rolling at 26% thickness reduction (failure reduction for AZ31) the Mg4Li specimen developed a double peak texture at ±15° about the transverse direction of the sheet TD (Fig. 7b, (0 0 0 2) pole figure). Double peak basal textures are commonly observed in hot rolled AZ31 sheet. Cold rolled AZ31 usually develops a sharp single peak basal texture similar to the one shown in Fig. 7a. The basal poles intensity in the (0 0 0 2) pole figure of Mg4Li after cold rolling was still nearly half of that of AZ31 at the same thickness reduction. Both cold rolled materials developed a preferred crystallographic alignment of the a-directions of the crystals with the rolling direction (Fig. 7a and b, {1 0 1 0} pole figure). During further deformation of the Mg4Li material up to 86% total thickness reduction the splitting of the central basal peak remained present in the rolling texture, but the peak intensity was increased up to 11 mrd (multiples of random distribution) (Fig. 7c). The preferred h1 1 2 0i RD orientation of the a-directions ({1 0 1 0} pole figure) was still evident, however, with less spread around TD. Obviously, the observed differences in the present rolling texture of Mg4Li, i.e. the central basal pole splitting and the considerably lower peak intensity compared to AZ31 (compare Fig. 7b and 7a) are a consequence of lithium addition. The texture development of the specimen types obtained from the extruded materials during channel-die compression at room temperature is shown in Fig. 10. As expected, different starting textures affected the activation of deformation mechanisms and their contribution rate in accommodating the imposed strain during PSC. This can be understood on the basis that the activation of crystallographic slip on basal and non-basal planes, as well as the activation of mechanical twinning was easier for some orientations than for others. This yielded different deformation scenarios seen by the shape of the flow curves and the evolution of the deformation textures. The initial texture of type 1 specimens (Fig. 4) suppressed basal slip under compressive loading since the basal planes were nearly parallel to the compression axis (very low Schmid factor). Under this condition non-basal slip modes were expected to gain influence on deformation during incipient stages. Prismatic slip can be considered the next slip system – after basal slip – to be activated most easily in magnesium. However, owing to the plane strain geometry only grains oriented such that their c-axes are parallel to TD can be deformed by prismatic slip. The rest of the grains, being the vast majority, cannot be deformed by this slip
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mode since it would cause sample broadening, which is suppressed by the lateral constraints of the channel-die. In this case prismatic slip would not contribute extensively to deformation. For the same geometrical reasons (broadening) hc + ai-pyramidal slip, which is highly required to provide more satisfactory accommodation of the imposed strain was only favorable in grains with their c-axes parallel to RD. However, in case of AZ31 its high critical shear stress (CRSS) at ambient temperatures renders its activation difficult. Grains with their c-axes oriented other than parallel to TD were very much favorable for tensile twinning since they were loaded in compression parallel to the basal planes. f1 0 1 2g-tensile twins were found to comply with Schmid law [22], and its CRSS in AZ31 is much lower than the CRSS for hc + ai-pyramidal slip. As a matter of fact, the twinning signature in AZ31 specimens (type 1) was clearly evident in the texture evolution (Fig. 10) and in the deformation microstructure (Fig. 9a). Tensile twinning readily took place at incipient strains reorienting the basal planes of grains by 86° about h1 1 2 0i, so that they became closely aligned with the compression plane forming a very sharp basal orientation with a maximum intensity of 28 mrd. The above aspects strongly suggest that depending on the deformation conditions (in the present case, the temperature since the strain rate was not changed) and the chemical composition (lithium addition) either non-basal slip or twinning becomes very important once basal slip ceases to operate any further or when it is suppressed at the beginning of deformation. In analogy to type 1, the same analysis of deformation mechanisms based on the initial orientation and deformation geometry can be done for the other specimen types. Interestingly, in cases where tensile twinning was inhibited (types 2 and 3), the samples were less prone to develop a basal texture, and the textures present showed even much weaker peak intensity. Although the initial extrusion textures in Fig. 4 were somewhat similar for AZ31 and Mg4Li, a comparison between the cold PSC textures of the two materials shows obvious differences in terms of texture characteristic and peak intensity. Mg4Li type 1 texture (recorded at a much higher strain than its counterpart of AZ31) was less inclined to build a strong basal orientation. The texture development of types 2 and 3 (showing high similarity even with type 1) exhibited a sort of a RD h1 0 1 0i fiber texture with a very evident maximum intensity splitting of the basal poles. It is well known from polycrystal plasticity texture simulations that splitting of the basal poles about TD is attributed to an increased activity of the non-basal hc + ai-pyramidal slip mode during compressive deformation. Central basal pole splitting is prevalently observed during hot deformation of wrought commercial magnesium alloys such as AZ31, but in the present study it was also evident in the basal texture of the cold rolled Mg4Li specimen (Fig. 7) and in the basal pole figures of Mg4Li types 2 and 3 specimens compressed at room temperature (Fig. 10). hc + aislip is usually a thermally activated process that requires
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elevated deformation temperatures of 250 °C and higher. The current findings, however, based on the distinction of deformation texture evolution between Mg4Li and AZ31 strongly suggests the presence of hc + ai-slip in the lithium containing specimens deformed at room temperature. Such argument that lithium additions promote the activation of hc + ai-slip even at ambient temperatures is not unreasonable since there is no other explanation for the remarkable room temperature ductility of Mg–Li alloys, particularly during sheet forming, at which the majority of grains are forced to be compressed nearly parallel to their c-axes. In fact and as mentioned before, an earlier TEM experimental study by Agnew and co workers [17] investigating the dislocation structures of basal oriented a-solid solution Mg– Li specimens reported a high density and uniform distribution of hc + ai dislocations upon room temperature deformation. increasing deformation temperature With (RT ? 200 °C ? 400 °C) the texture evolution was less prone for developing a basal orientation. This behavior is consistent with the fact that with rising temperatures, twinning became less and less important for accommodating the imposed shape change and conversely, the thermally activated non-basal slip was gaining more influence on the deformation. In case of AZ31 type 1 specimens, it was obvious from Fig. 11 that twinning was still an important deformation mechanism even at 200 °C. The peak intensity, however, was much weaker than the one at room temperature. The weak prismatic texture component near to TD in the (0 0 0 2) pole figure was a remnant of the initial texture that was not able to twin. The texture characteristics and peak intensity of the Mg4Li specimens at 200 °C did not seem to differ much from one sample type to another (Fig. 11). A tendency to propagate along the transverse direction was perspicuous. The observed textures comprised a sharp RD h1 0 1 0i fiber texture component ({1 0 1 0} pole figure) with a peak intensity splitting at ±55° from RD ((0 0 0 2) pole figure). Deformation at 400 °C yielded an interesting texture development, particularly, for the lithium containing specimens (Fig. 8c). Evidently, all specimen types shared a sharp {1 1 2 0}h1 0 1 0iprism texture. This suggests that at elevated temperatures, such as 400 °C, prismatic slip becomes even more important than hc + ai-pyramidal slip. However, this conclusion cannot be generalized at the moment and has to be limited for plane strain deformation geometry only. In case of AZ31-type 1 tensile twinning was rendered completely unimportant and the deformation was driven primarily by prismatic slip indicated by the TD[0 0 0 1] orientation in the (0 0 0 2) basal pole figure. With respect to the peak intensities, the textures in AZ31 deformed at 400 °C were markedly weaker than their counterparts in Mg4Li. In case of AZ31 (where the grain size remained relatively small, Fig. 9e) this can be explained by a previous conclusion that the recrystallization texture of DRX grains plays an important role in weakening the overall texture of the specimen by counteracting the strong deformation texture of the
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Fig. 11. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at 200 °C at a constant strain rate of 104 s1 up to e = 1.4. Levels: 1:2:3:4:5:6:7:8:9:10:12.
deformed microstructure [21]. In case of Mg4Li, the strong texture evolution can be attributed to: (a) The deformation being mainly compensated by one slip mode, i.e. prismatic slip, which results in very well defined textures, as observed in Fig. 12. In other words the more deformation modes involved in accommodating the imposed strain, the more homogeneous the deformation, and the weaker the texture development and anisotropy (as in cubic metals). (b) Sharpening of the recrystallization texture during subsequent grain growth, which was evident in the coarse-grained microstructure in Fig. 9f.
6. Summary and conclusions The combined effects of lithium addition (4 wt.%), the starting texture, and the deformation temperature on the deformation behavior and the evolution of texture and microstructure in magnesium were investigated. The same investigations were carried out on a commercial AZ31
magnesium alloy and the obtained results were compared and discussed for both materials. The experimental findings of the present study are in good accordance with results reported in the literature for similar materials. The following conclusions (focusing on the Li containing alloy) can be drawn: (1) With respect to room temperature formability examined in cold rolling and channel-die compression experiments, the a-solid solution Mg4Li alloy demonstrated a remarkable ductility enhancement over a standard wrought magnesium alloy AZ31. During cold rolling a total thickness reduction of 86% was easily achieved with no signs of cracks or mechanical instabilities. (2) Alloying additions affect the balance of deformation mechanisms that contribute to deformation. This in turn affects the evolution of the crystallographic texture. Lithium addition seemed to result in a congenerous deformation behavior and texture evolution irrespective of the starting texture. This type of behavior was completely different from the one observed in AZ31, in which the stress-
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Fig. 12. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at 400 °C at a constant strain rate of 104 s1 up to e = 1.4. Levels: 1:2:3:4:5:6:7:8:9:10:12:14:16:18.
Fig. 13. Comparison of the yield strength of AZ31 and Mg4Li for different starting textures during PSC deformation at (a) 200 °C and (b) 400 °C.
strain mechanic response (flow curves) and texture evolution were very dependent on the starting texture. (3) After cold rolling Mg4Li specimens showed a basal texture with central peak intensity splitting rotated towards the rolling direction. In channel-die compression the texture developed as a consequence of the operating deformation mechanisms and their
balance, which was influenced by the deformation temperature. At room temperature, Li addition to magnesium seemed to enhance the activation of hc + ai-pyramidal slip. With rising deformation temperature the orientations seemed to spread along TD forming an RD h1 0 1 0i fiber texture component with maximum intensities at ± 55° from RD. At 400 °C the basal poles were completely
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Acknowledgments This work was funded by the Deutsche Forschungsgemeinschaft DFG under grant (GO 335/27). The authors also thank Dr. Dirk Bormann from the Institut fu¨r Werkstoffkunde at the Leibniz Universita¨t Hannover for the kind donation of the magnesium–lithium material. References
Fig. 14. DRX in Mg4Li specimen deformed by channel-die compression at 200 °C/104 s1.
rotated towards TD developing a strong {1 1 2 0}h1 0 1 0i-prism texture. Elevated deformation temperatures (400 °C) render prismatic slip in Mg4Li subjected to PSC more important than hc + ai-pyramidal slip. (4) During deformation at 200 °C dynamic recrystallization (DRX) took place in Mg4Li specimens mainly at grain boundaries and triple junctions (locations of high stress concentration). The deformed grains – being the majority of the microstructure – were elongated in the extension (rolling) direction. At 400 °C the deformed specimen revealed a coarse granular microstructure comprised of grains with an average intercept length of 150 lm. Microstructure coarsening at 400 °C was a consequence of grain growth. By contrast, AZ31 deformed under the same conditions did not reveal any signs of grain growth following recrystallization.
[1] Emley EF. Principles of magnesium technology. Oxford: Pergamon; 1966. [2] von Mises R. J Appl Math Mech (ZAMM) 1928;8:161–85. [3] Al-Samman T, Gottstein G. Mater Sci Eng A 2008;488:406–14. [4] Koike J, Ohyama R, Kobayashi T, Suzuki M, Maruyama K. Mater Trans 2003;44:445–51. [5] Ball EA, Prangnell PB. Scripta Metall et Mater 1994;31:111–6. [6] Laser T, Nu¨rnberg MR, Hartig Ch, Letzig D, Bormann R. Magnesium alloys and their applications. In: Kainer KU, editor. Wiley VCH; 2006. p. 1055–61. [7] Massalski TB. Binary alloy phase diagrams. ASM International; 1992. [8] Hauser FE, Landon PR, Dorn JE. Trans ASM 1956;48:986–1002. [9] Sahoo M, Atkinson JTN. J Mater Sci 1982;17:3564–74. [10] Hardie D, Parkins RN. Phil Mag 1959;4:815–25. [11] Ahmadieh A, Mitchell J, Dorn JE. Trans Metall Soc AIME 1965;233:1130–8. [12] Quimby RM, Mote JD, Dorn JE. Trans ASM 1962;55:149–57. [13] Yoshinaga Y, Horiuchi R. Trans Japan Inst Metals 1964;4:134–41. [14] Urakami A, Meshii M, Fine ME. Acta Metall 1970;18:87–99. [15] Urakami A, Meshii M, Fine ME. In: Proc. 2nd Int. Conf. Strength of Metals and Alloys (ICSMA). 1970; p. 272–6. [16] Agnew SR, Yoo MH, Tome´ CN. Acta Mater 2001;49:4277–89. [17] Agnew SR, Horton JA, Yoo MH. Metall Mater Trans A 2002;33:851–8. [18] Dahms M, Bunge HJ. Text Microstruct 1987;7:171–85. [19] Kelly EW, Hosford WF. Trans Metall Soc AIME 1968;242:654–61. [20] del Valle JA, Ruano OA. Mater Sci Eng A 2008;487:473–80. [21] Al-Samman T, Gottstein G. Mater Sci Eng A 2008;490:411–20. [22] Gharghouri MA, Weatherly GC, Embury JD, Root J. Phil Mag 1999;79:1671–95.
Acta Materialia 57 (2009) 2229–2242 www.elsevier.com/locate/actamat
Comparative study of the deformation behavior of hexagonal magnesium–lithium alloys and a conventional magnesium AZ31 alloy T. Al-Samman * Institut fu¨r Metallkunde und Metallphysik, RWTH Aachen, Kopernikusstr. 14, 52056 Aachen, Germany Received 6 November 2008; received in revised form 21 January 2009; accepted 23 January 2009 Available online 28 February 2009
Abstract Specimens of a conventional magnesium AZ31 alloy and a binary a-solid solution Mg4Li alloy with similar starting textures and microstructure were subjected to plane strain deformation under various deformation temperatures ranging from 298 K to 673 K. Lithium addition to magnesium exhibited remarkable room temperature ductility improvement owing to enhanced activity of non-basal slip, particularly, hc + ai-slip mode. Furthermore, the addition of lithium to magnesium seemed to reduce the plastic anisotropy, typical for commercial magnesium alloys. This was evident in the flow curves and texture development obtained at 200 °C and 400 °C. At 400 °C prismatic slip gains strong influence in accommodating the imposed deformation. In terms of thermal stability against microstructure coarsening at elevated temperatures, the lithium containing alloy undergoes significant grain growth following recrystallization. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Texture; Slip; Twinning; Plastic deformation; Ductility
1. Introduction Below 498 K plastic deformation of polycrystalline magnesium is limited to basal (0 0 0 1)h1 1 2 0i-slip and to pyramidal f1 0 1 2gh1 0 1 1i-twinning [1]. The magnesium crystal, thus has only three geometrical and two independent slip systems compared to aluminum, for example, which has 12 {1 1 1}h1 10i geometrical and five independent slip system for arbitrary shape change [2]. Pure magnesium and conventional magnesium alloys therefore have a tendency to embrittlement at low deformation temperatures due to intergranular failure, and localized transcrystalline fracture along twinned regions, particularly upon compression [3], or along basal (0 0 0 1) planes for very coarse grains [1]. It is proposed that the room temperature ductility of magnesium can be enhanced by:
*
Tel./fax: +49 241 80 26871. E-mail address: [email protected]
(i) Grain refinement obtained via severe plastic deformation or conventional dynamic recrystallization, which can promote grain boundary sliding and/or grain boundary rotation allowing easier accommodation of the overlap or void displacement between two adjacent grains [4]. (ii) Altering the texture observed in conventional magnesium alloy sheet material. Typical wrought magnesium alloys usually exhibit strong textures with little variations between them. The typical sheet basal texture places the vast majority of grains, such that their basal (0 0 0 1) planes are close to the plane of the sheet. With such an orientation, i.e. c-axis compression, it is difficult to deform. In the literature, many indications can be found about the improved properties of magnesium alloys containing additions of yttrium (Y) and rare earth elements (RE), such as cerium (Ce), neodymium (Nd) and gadolinium (Gd). Ball et al. [5] reported that an Mg–Y–Nd mischmetal-based alloy can develop more random textures during extrusion compared to conventional alloys. They found that a more randomized texture strongly
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.01.031
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reduces the tension–compression asymmetry, which is an attractive benefit of controlling the texture to enhance formability. Another study by Laser et al. [6] showed that texture modification is also possible in other alloys without much additions of expensive RE elements. They studied four different alloys based on conventional AZ31 containing additions of cerium and calcium. The alloys were hot rolled and tensile tested at room temperature. It was observed that the addition of less than 1 wt.% Ce mischmetal and Ca reduced the peak intensity in the basal pole figure to only 3 multiples of a random distribution (MRD), as compared to 9 MRD for a conventional AZ31 alloy. Not only was the texture weaker, but the qualitative character of the texture was also altered. (iii) Lithium addition. The Mg–Li system (Fig. 1a) [7] is quite interesting. Not only because the addition of lithium offers a great potential to enhance the room temperature ductility of the material (brittle to ductile transition temperature is suppressed to below liquid nitrogen temperature) [8], but also because it allows to substantially lower the density of the already light magnesium even more (qLi = 0.58 g/cm3 vs. qMg = 1.74 g/cm3) (Fig. 1b) [9]. Hence, magnesium– lithium alloys are the lightest structural metals.
Fig. 1. (a) Mg–Li binary phase diagram [7]. (b) The effect of lithium addition on the alloy density [9].
The purpose of the current study is to investigate the influence of lithium addition on the various deformation modes and their contribution to deformation, particularly, on the development of crystallographic texture and microstructure during deformation at room temperature and at higher deformation temperatures of 200 °C and 400 °C. The deformation behavior, texture and microstructure development of the investigated Mg–Li alloy are compared with those of a commercial magnesium AZ31 alloy, deformed exactly under the same conditions. The findings of this study will be used in future investigations to set directions for new Mg–Li–X alloy development. 2. Background of Mg–Li alloys Lithium addition to magnesium has an important effect on the crystal structure, i.e. the lattice parameters. With increasing lithium content the c/a axial ratio decreases (from 1.624 for pure Mg to 1.607 for Mg–17 at.% Li, close to the solid-solubility limit) [10]. Between 17 and 30 at.% Li contents magnesium–lithium alloys exhibit two phase structures consisting of the a-magnesium rich (hcp) and b-lithium rich (bcc) phases. The b single phase structure exists for Li contents greater than 30 at.% (Fig. 1a). The body centered cubic b phase is soft and ductile. Hence, magnesium alloys with such high lithium content are very ductile. Nevertheless, and aside from high production costs, these alloys suffer from low temperature instability. The improved room temperature ductility of a-solid solution Mg–Li alloys is attributed to a substantial increase in non-basal slip activity, particularly, prismatic slip. While prismatic slip in pure magnesium is negligible at room temperature, and only limited to regions with stress concentrations in the vicinity of grain boundaries, extensive prismatic slip over entire grains occurs in magnesium alloys containing lithium [11]. With respect to basal slip, Quimby et al. [12] reported that its critical resolved shear stress (CRSS) increased with increasing lithium content, i.e. ds/dcLi > 0, where s is CRSS for basal slip and cLi is the concentration of lithium solutes. The alloy hardening for basal slip was attributed to interaction between dislocations and Li solute atoms [13]. It should be noted that the decrease in the c/a ratio resulting from lithium addition is not the only reason for the enhanced activity of prismatic slip. Beryllium, for example, has a c/a ratio of 1.568, less than that of zirconium and titanium (1.593 and 1.588, respectively), and yet deforms primarily by basal slip, where both Zr and Ti have prismatic slip as the primary deformation mode. Another – more important – reason for the softening in non-basal slip is the influence of lithium addition on the dislocation mechanisms and the stacking fault energy related to non-basal slip planes. The complex ds/dcLi < 0 behavior, i.e. softening for prismatic slip was attributed by Ahmadieh et al. [11] to a reduction in the Peierls stress when adding lithium. The theoretical reason for this effect
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was not reported. Urakami et al. [14] performed electron microscopic investigation of prismatic slip in Mg–10.5 at.% Li single crystals deformed in tension at liquid nitrogen and ambient temperatures. Their theory for the rate controlling mechanism for low temperature prismatic slip of Mg–Li alloys was based upon friction stresses and mobility of screw dislocations. In another study, the authors proposed that prismatic slip occurs by a double kink mechanism at interaction sites between dislocations and Li atoms. The activation energy for the nucleation of double kinks is reduced by the presence of lithium [15]. The investigations mentioned earlier [8,11–15], that attributed the extensive low temperature ductility improvement of a-solid solution Mg–Li alloys to the availability of additional deformation modes, i.e. prismatic slip, were all carried out on single crystals. However, in polycrystal Mg–Li alloys, even if both basal and prismatic slip were active, accommodating deformation along the c-axis will still be a problem (considering the unidirectional nature of mechanical twinning and hence their limited ability to accommodate deformation). The fact that polycrystal hcp Mg–Li alloys can undergo c-axis compression at room temperature up to large strains with no signs of mechanical instability or fracture (as shown in the current work) suggests a significant involvement of hc + ai-slip as well. In fact, a texture simulation study by Agnew et al. [16], supported by TEM investigation of a-solid solution Mg–Li alloys from the same authors [17] reported that hc + ai-slip on f1 1 2 2g second-order pyramidal planes was indeed active during compression at ambient temperatures. 3. Experimental procedure For the current study a-solid solution Mg–15 at.% Li (Mg4Li) (wt.%) alloy and a commercial magnesium AZ31 alloy were used. Both materials were available in the asrolled and as-extruded state. The processing history of the Mg4Li alloy comprised a thickness reduction from 12 mm to 6 mm for the rolled material (same for AZ31) and a cross section reduction from Ø 30 mm to Ø 15 mm for the extruded material. The cross section reduction for the commercially extruded AZ31 material was from Ø 180 mm to Ø 60 mm. The processing temperature for the AZ31 material (both conditions extruded and rolled) was 400 °C. The Mg– Li material was thermo-mechanically processed at 300 °C. Prior to the deformation tests both materials were annealed at 350 °C for 1 h. The microstructures are shown in Fig. 3. In order to cover a wide range of investigations, four specimen types with different starting textures were investigated (Fig. 2a and 2b). The starting textures were obtained by simply cutting the samples in such a way that the basal plane was essentially parallel to any desired direction for the given plane strain geometry (Fig. 2c). The specimen dimensions (types 1, 2 and 3, Fig. 2a) for the plane strain compression (PSC) tests were 14 mm (RD) 10 mm (TD) 6 mm (ND). The dimensions of cold rolling specimens (type 4, Fig. 2b) were 60 mm 30 mm 4 mm.
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Fig. 2. Orientation of the investigated specimens (a) types 1, 2 and 3 (extruded starting condition); (b) type 4 (rolled starting condition) and (c) selection of a new reference system represented by the rolling (RD), transverse (TD) and normal direction (ND) of the channel-die geometry with respect to the original loading direction (extruded ED or rolled RD).
For room temperature deformation, cold rolling and plane strain compression using a channel-die device were chosen. Extension occurred only in the longitudinal direction of the channel-die (parallel to RD, Fig. 2c) since deformation in the transversal direction was suppressed, owing to the deformation geometry (lateral constraints). Deformation experiments at elevated temperature comprised channel-die compression only. Deformation temperatures and strain rates ranged between 200 °C to 400 °C and 102 s1 to 104 s1, respectively. Hexagonal boron nitride (h-BN) powder was used as a lubricant for minimizing friction between sample and channel-die. For cold rolling few drops of industrial oil was used for lubrication. After completion of the tests, specimens for optical microscopy were shortly ground with a 4000 grit SiC paper and subsequently mechanically polished with diamond paste down to 3 and 1 lm, respectively. Final polishing was performed using a colloidal silica solution. After polishing, specimens were etched in acetic picral to visualize grains and grain boundaries. Same preparation procedure was used for both alloys. The macrotexture was determined in the mid-plane of deformed specimens by measuring incomplete pole figures (5° 6 a 6 75°) in the back reflection mode using CoKa radiation. A set of six measured pole figures [f1 0 1 0g; f0 0 0 2g; f1 0 1 1g; f1 0 1 2g; f1120g, and {10 13}] was used to calculate the orientation distribution function (ODF) using the positivity method [18]. 4. Results 4.1. Initial condition Fig. 3 shows the initial microstructures of the alloys AZ31 and Mg–Li following annealing at 350 °C for 1 hour. For the extruded Mg4Li material, the microstructure
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Fig. 3. Initial microstructures prior to the PSC tests. (a) Mg4Li extruded condition; (b) AZ31 extruded condition; (c) Mg4Li rolled condition; (d) AZ31 rolled condition.
revealed equiaxed grains with a homogeneous grain size distribution (Fig. 3a). Interestingly, the extruded microstructure showed no signs of deformed elongated grains. Such grains are usually a common microstructural feature of commercially extruded AZ31 material (Fig. 3b).The rolled Mg4Li material exhibited a microstructure that was similar to that of the extruded material. The mean grain size (linear intercept) was larger than that of the extruded material (80 lm for the rolled material compared to 50 lm for the rolled one) (Fig. 3c). Some deformation twins from previous processing were still visible after annealing. No second phase precipitates were observed. The average grain sizes of the rolled and the extruded AZ31 alloy were 40 lm and 25 lm, respectively. The starting textures of the different test series of both alloys are shown in Fig. 4. For specimen types 1 and 2 obtained from the extruded materials (compare with Fig. 2a) the textures showed little difference in terms of texture characteristics, yet the difference was noticeable in terms of texture strength. AZ31 starting textures were much more pronounced. This is attributed to the different chemical composition, as well as to the different extrusion
ratio for both alloys. The textures of specimen type 3 (also obtained from the extruded materials) showed perceptible qualitative difference of texture components. AZ31 starting texture, i.e. RD h1 0 1 0i was also much more pronounced. Specimen type 4 obtained from the rolled Mg4Li material exhibited an unconventional rolling texture, characterized by a relatively weak intensity (5 mrd) and a large h0 0 0 1i scatter around ND. The rolled AZ31 material revealed a typical strong rolling texture (12 mrd) with a maximum intensity split. 4.2. Room temperature deformation 4.2.1. Cold rolling The cold rolling behavior of the alloys AZ31 and Mg4Li is shown in Fig. 5. As expected, the commercial AZ31 material exhibited typical brittle response and failed at a total thickness reduction of 26% despite low reduction per each rolling pass (0.1 mm/pass). By contrast, the lithium containing material remained in excellent shape and revealed no signs of cracks (Fig. 5a). During further cold rolling, the Mg4Li material conveyed remarkable ductility
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Fig. 4. Starting textures (prior to PSC) represented by means of (0 0 0 2)- and {1 0 1 0}- recalculated pole figures. Intensity levels 1:2:3:4:5:6:7:8:9: 10:12:14:16:18:20.
reaching a total thickness reduction of 86% (Fig. 5b). The surface quality was similar to the quality usually obtained for AZ31 by hot rolling at 400 °C with intermediate annealing. The corresponding microstructures (observed in the transverse plane) are shown in Fig. 6. The microstructure of the failed AZ31 material is presented in Fig. 6a, while Fig. 6b and c depict the microstructures of Mg4Li specimens at 26% and 86% thickness reduction, respectively. The rolling textures are shown in Fig. 7 in terms of basal (0 0 0 2) and prismatic {1 0 1 0 } recalculated pole figures.
4.2.2. Channel-die compression The mechanical response of AZ31 and Mg4Li materials during room temperature plane strain deformation at a constant strain rate of 104 s1 using a channel-die device is shown in Fig. 8 in terms of stress–strain curves. Stress–strain data after e = 0.6 were not taken into account due to friction between sample and channeldie, which shows in the flow curve as a false increase of stress at high strains. Furthermore, it should be noted that in case of room temperature deformation the recorded stress-strain data do not actually show true val-
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different contribution of the deformation mechanisms during compression. In the current case, i.e. room temperature deformation the fact that the presented textures were measured at two completely different failure strains should be also considered. 4.3. Elevated temperature deformation
Fig. 5. Cold rolling behavior of Mg4Li. (a) Direct comparison with AZ31 at 26% total thickness reduction (AZ31 failed). (b) Mg4Li cold rolled up to 86% final thickness reduction demonstrating remarkable room temperature formability.
ues, but rather nominal ones because of the large heterogeneity of deformation. As evident from the flow curves, the specimens containing lithium demonstrated a ductile room temperature deformation behavior reaching failure strains above 1.2 (not shown in Fig. 8). By contrast, AZ31 specimens underwent failure by shear localization at failure strains ranging from ef = 0.15 to 0.24. The cold deformation microstructures after channel-die compression are depicted in Fig. 9a and b for AZ31 (compressed to failure) and Mg4Li (compressed to e = 1), respectively. The microstructure of both transverse (CD-RD) and compression surfaces (RDTD) was examined. In case of Mg4Li no noticeable difference between the two surfaces was observed. For AZ31 the microstructure of the compression surface comprised some microcracks that were probably initiated from deformation twins. For the different types of each alloy the microstructure development was similar, and hence, only one microstructure for each alloy is presented. The microstructure of Mg4Li remained coarser than the microstructure of AZ31. The evolution of macrotextures of the different specimen types of AZ31 and Mg4Li materials during room temperature PSC tests to failure is presented in Fig. 10. A comparison of the results reveals noticeable differences between both materials in terms of texture characteristic and maximum intensity. This is mainly attributed to the
For investigating the elevated temperature deformation behavior of both materials channel-die compression tests at 200 °C and 400 °C at a constant strain rate of 104 s1 were carried out. Hot rolling experiments were not included in the current investigations because of the difficulty of controlling the rolling temperature. The flow curves of the investigated materials at elevated temperatures are shown in Fig. 8b for 200 °C and in Fig. 8c for 400 °C, respectively. As evident, the flow behavior was influenced by the deformation conditions (temperature and strain rate), the starting texture and the chemical composition of the deformed material. This is discussed in more details below. In analogy with stress-strain data presented in Fig. 8a for room temperature deformation, the curves in Fig. 8b and c show only data up to a true strain of e = 0.6. The tests were, however, carried out up till e = 1.4. The corresponding microstructure development is shown in Fig. 9c–f for PSC tests at 200 °C (c and d), and 400 °C (e and f). It should be noted, that it was clear from the beginning that the Mg4Li material would always exhibit a coarser microstructure than AZ31, and therefore, it was given no attention in Fig. 9 to show micrographs taken at the same magnification. 5. Discussion 5.1. Deformation behavior and formability Commercial magnesium sheet does not currently exhibit sufficient formability. Modifying the c/a ratio of its hexagonal close packed lattice and the crystallographic texture by lithium addition offers a great potential to improve the room temperature formability of magnesium wrought alloys. This is clearly demonstrated in Figs. 5 and 8a which show the cold deformation response of a binary Mg4Li alloy in comparison with a standard AZ31 alloy in terms
Fig. 6. Cold rolling microstructures of (a) AZ31 and (b) Mg4Li at 26% total thickness reduction. (c) Microstructure of Mg4Li at 86% final thickness reduction. The micrographs were taken from the transverse plane of the rolled specimens.
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Fig. 7. Macrotexture evolution after cold rolling of (a) AZ31 to failure (26%), (b) Mg4Li at 26% thickness reduction and (c) Mg4Li at 86% final thickness reduction. The textures are presented in terms of (0 0 0 2)- and {1 0 1 0}-recalculated pole figures. Levels: 1:2:3:4:5:6:7:8:9:10:12.
of sheet formability and stress-strain flow curves. While the standard wrought alloy AZ31 failed at typical thickness reduction and strain values during cold rolling and cold channel-die compression, the lithium containing alloy exhibited a substantial ductility enhancement with no signs of mechanical instability or fracture up to very high strains (e > 1). This remarkable ductility improvement is attributed (as stated earlier) to alloy softening for non-basal slip resulting not only from the slight decrease of the axial c/a ratio, but also from the influence of lithium solutes on the dislocation dynamics. Although the samples of both materials were tested under the same deformation conditions, i.e. the same temperature and strain rate, a markedly different mechanical response for each specimen type was observed (particularly shows in Fig. 8a). This is attributed to the different starting texture each specimen type had. Each initial texture favors specific combination of deformation mechanisms, which is reflected during deformation by a different flow behavior. Comparing the flow curves of both materials during PSC deformation at room temperature indicates much lower flow stress values for Mg4Li than the commercial alloy AZ31. This is also true in case of elevated temperature
Fig. 8. Stress–strain curves of AZ31 and Mg4Li specimens during channel-die compression tests (a) at room temperature (b) at 200 °C (c) at 400 °C. e_ = 104 s1.
deformation. While AZ31- type 1 specimen exhibited a flow stress maximum of 400 MPa at e = 0.18 before fracture, Mg4Li specimen of the same type exhibited a flow stress value of 265 MPa at the same strain. With increasing strain some strain hardening takes place, but the maximum strain values remain under 300 MPa (without considering friction at e > 0.6). This modest work hardening behavior at room and elevated temperatures (see also Fig. 8b and c) is one of the aspects binary Mg–Li alloys suffer from, when compared with other magnesium alloys (AZ, ZK and AM series). During elevated temperature deformation an interesting behavior of the Mg4Li alloy was recorded. Comparing the flow curves of both materials at 200 °C (Fig. 8b) and 400 °C (Fig. 8c) shows that the mechanical response, i.e. flow behavior of the Mg4Li spec-
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Fig. 9. Deformation microstructures of AZ31 (a, c and e) and Mg4Li (b, d and f) after channel-die compression at room temperature (a and b); at 200 °C (c and d); and at 400 °C (e and f). The strain corresponding to micrograph (a) was ef = 0.14. For micrograph (b) e = 1. For micrographs (c–f) e = 1.4. e_ = 104 s1. tt, tensile twins; ct, compression twins.
imens seemed to be not very much affected by the starting texture. By contrast, the flow curves and hardening response of the AZ31 specimens were strongly dependent on the initial orientation of grains and the evolution of texture during deformation. A comparison of the recorded yield stress data for AZ31 and Mg4Li is depicted in Fig. 13. Obviously, increasing the deformation temperature resulted for both materials a significant drop of the yield stress. Mg4Li specimens exhibited much lower yield stress values than AZ31. Interestingly, the lithium containing specimens showed more or less the same yield stress strength, irrespective of specimen type (starting texture). In case of AZ31, the yield stress revealed different values for different starting orientations. With respect to the influence of the starting texture and the deformation mechanisms on the flow behavior of the
specimens (particularly AZ31) during channel-die compression at RT and 200 °C, two kinds of plastic flow can be observed: (1) If mechanical twinning was more favorable than crystallographic slip for deformation, then the work hardening part of the flow curve would result from twinning driven deformation (characterized by a concave shape of the flow curve and significant work hardening regime). (2) If at incipient deformation mechanical twinning was inhibited (very low Schmid factor or for deformation geometrical reasons), and hence, crystallographic slip was the primary deformation mechanism, the work hardening part of the flow curve would be less evident than the previous case (slip driven deformation).
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Fig. 10. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at room temperature at a constant strain rate of 104 s1. The strains corresponding to the presented textures are ef = 0.14 for AZ31 and e = 1 for Mg4Li. Levels: 1:2:3:4:5:6:7:8:9:10:12:14:16:18:20:25.
The hardening effect of twins can be seen in the flow curves of type1 specimens (particularly for AZ31, Fig. 8a and b), which possessed a starting orientation that is very favorable for tensile twinning (c-axes perpendicular to compression axis, Fig. 4). Twin-induced hardening in magnesium has been frequently reported [16,19]. It is important to note that the conspicuous hardening effect of twins is attributed to the stress required to initiate slip in the twinned regions of the matrix rather than to the activation stress of the twinning process itself. Crystallographic slip within the twin interior is hard because (i) f1 0 1 2g-tensile twinning reorients the basal planes in the twinned parts of the matrix into unfavorable orientations for easy basal slip, and (ii) the mean free path of dislocation movement within the twin interior is too short in comparison with the matrix. The above aspects are the reason why the flow stress in grains undergoing tensile twinning strongly rises with strain compared to other specimen types, in which grains are primarily undergoing crystallographic slip. During deformation at 400 °C the flow behavior of all specimen types was driven by slip as the dominant deformation mechanism. Mechanical twinning in commercial magnesium alloys is usually limited to deformation temperatures below 250 °C.
5.2. Microstructure development During room temperature deformation the development of the microstructure of failed AZ31 specimens revealed two common features: high fraction of twinned areas and microcracks. Previous investigations on AZ31 [3] showed that deformation at ambient temperatures can activate f1 0 1 1g-compression twinning and double twinning. Both twinning types can be distinguished from the typical tensile twins by their thin morphological appearance. In Fig. 5a the rolling microstructure of the failed specimen at 26% thickness reduction show deformation twins propagating along the transverse section at approximately ±45° to the rolling direction. The microstructure also reveals several locations of voids (microcracks) with the same alignment of twins, i.e. ±45° to RD. According to earlier investigations [3], these twin-sized voids could be initiated from the preexisting twins. The same microstructural combination of high volume fraction of twins and microcracks was also found in the compression surface of the RTPSC deformed specimen (type 1) to failure (Fig. 9a). By contrast, the optical micrographs of Mg4Li specimens deformed at room temperatures show lower volume fraction of twins and no signs of voids. The fact that both twin
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types, i.e. tensile (broad) and compression (thin) twins were evident in the deformation microstructure (Fig. 9b) suggests that mechanical twinning remains an important deformation mechanism despite sufficient availability of non-basal slip modes. The microstructure of Mg4Li specimen cold rolled up to a final thickness reduction of 86% comprised large-scale deformation heterogeneities in the form of very fine shear bands aligned at approximately ±35° from the rolling direction RD. This is quite expected at such high level of deformation. Similar shear band microstructure can be found in AZ31 during large strain hot rolling (LSHR) up to high thickness reductions (usually >75%). During deformation at higher temperatures (200 °C) both investigated materials showed common microstructural features comprising partial dynamic recrystallization (DRX) and a grain microstructure that is elongated towards the expansion direction of the deformed specimen, i.e. rolling direction RD. The unique directionality of the microstructure observed in Fig. 9c and d is a creation of the strain path of plane strain deformation. Usually, a material undergoing DRX at deformation temperatures between 200 °C and 250 °C exhibits a typical stress-strain behavior characterized by an initial work hardening regime (reaching a peak stress) followed by rapid work softening. In the current investigations this was mostly evident in case of AZ31-type 1 specimen (Fig. 8b). In other specimen types (also Mg4Li specimens) the work softening was less evident, and the flow curves showed steady-state behavior straight after work hardening. This could suggest that the kinetics of DRX is influenced by the texture, as it was reported recently by del Valle [20]. The grain size recrystallized dDRX, however, was found to be insensitive to texture [20,21]. The grain size recrystallized for AZ31 and Mg4Li was 4 lm and 10 lm, respectively. With regard to nucleation of DRX, it was shown in a previous study [21] that DRX can take place at twin boundaries and within twin interiors, provided that the vast majority of grains undergo massive twinning. In case of Mg4Li twinning was not important, and hence DRX nucleated mostly at grain boundaries and triple junctions (Fig. 14). During deformation at elevated temperatures (400 °C) the deformed Mg4Li specimen revealed a coarse granular microstructure comprised of grains with intercept lengths of 100–200 lm (Fig. 9f). This was obviously an indication of grain growth. By contrast, AZ31 specimens undergoing plastic deformation at the same conditions (400 °C/104 s1) did not reveal any signs of noticeable grain growth after DRX (owing to the eutectic constituent Mg17Al12 dissolved in the matrix). The microstructure was fully recrystallized and the grain size remained comparatively very small (d 20 lm). This concludes that the binary Mg4Li alloy behaves similarly to pure magnesium deformed or annealed at such elevated temperatures, whereat the occurrence of grain growth is inevitable.
5.3. Deformation mechanisms and texture evolution The starting texture of the Mg4Li rolled material (Fig. 4, type 4) showed a typical basal orientation (basal planes lying close to the rolling plane), yet with a higher maximum spread around the normal direction of the sheet ND than usual. There was no evident preference for any crystallographic direction to align with the rolling direction. Additionally, the texture intensity was nearly half of that of rolled AZ31 (Fig. 4, type 4) taking into account similar thickness reduction. After cold rolling at 26% thickness reduction (failure reduction for AZ31) the Mg4Li specimen developed a double peak texture at ±15° about the transverse direction of the sheet TD (Fig. 7b, (0 0 0 2) pole figure). Double peak basal textures are commonly observed in hot rolled AZ31 sheet. Cold rolled AZ31 usually develops a sharp single peak basal texture similar to the one shown in Fig. 7a. The basal poles intensity in the (0 0 0 2) pole figure of Mg4Li after cold rolling was still nearly half of that of AZ31 at the same thickness reduction. Both cold rolled materials developed a preferred crystallographic alignment of the a-directions of the crystals with the rolling direction (Fig. 7a and b, {1 0 1 0} pole figure). During further deformation of the Mg4Li material up to 86% total thickness reduction the splitting of the central basal peak remained present in the rolling texture, but the peak intensity was increased up to 11 mrd (multiples of random distribution) (Fig. 7c). The preferred h1 1 2 0i RD orientation of the a-directions ({1 0 1 0} pole figure) was still evident, however, with less spread around TD. Obviously, the observed differences in the present rolling texture of Mg4Li, i.e. the central basal pole splitting and the considerably lower peak intensity compared to AZ31 (compare Fig. 7b and 7a) are a consequence of lithium addition. The texture development of the specimen types obtained from the extruded materials during channel-die compression at room temperature is shown in Fig. 10. As expected, different starting textures affected the activation of deformation mechanisms and their contribution rate in accommodating the imposed strain during PSC. This can be understood on the basis that the activation of crystallographic slip on basal and non-basal planes, as well as the activation of mechanical twinning was easier for some orientations than for others. This yielded different deformation scenarios seen by the shape of the flow curves and the evolution of the deformation textures. The initial texture of type 1 specimens (Fig. 4) suppressed basal slip under compressive loading since the basal planes were nearly parallel to the compression axis (very low Schmid factor). Under this condition non-basal slip modes were expected to gain influence on deformation during incipient stages. Prismatic slip can be considered the next slip system – after basal slip – to be activated most easily in magnesium. However, owing to the plane strain geometry only grains oriented such that their c-axes are parallel to TD can be deformed by prismatic slip. The rest of the grains, being the vast majority, cannot be deformed by this slip
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mode since it would cause sample broadening, which is suppressed by the lateral constraints of the channel-die. In this case prismatic slip would not contribute extensively to deformation. For the same geometrical reasons (broadening) hc + ai-pyramidal slip, which is highly required to provide more satisfactory accommodation of the imposed strain was only favorable in grains with their c-axes parallel to RD. However, in case of AZ31 its high critical shear stress (CRSS) at ambient temperatures renders its activation difficult. Grains with their c-axes oriented other than parallel to TD were very much favorable for tensile twinning since they were loaded in compression parallel to the basal planes. f1 0 1 2g-tensile twins were found to comply with Schmid law [22], and its CRSS in AZ31 is much lower than the CRSS for hc + ai-pyramidal slip. As a matter of fact, the twinning signature in AZ31 specimens (type 1) was clearly evident in the texture evolution (Fig. 10) and in the deformation microstructure (Fig. 9a). Tensile twinning readily took place at incipient strains reorienting the basal planes of grains by 86° about h1 1 2 0i, so that they became closely aligned with the compression plane forming a very sharp basal orientation with a maximum intensity of 28 mrd. The above aspects strongly suggest that depending on the deformation conditions (in the present case, the temperature since the strain rate was not changed) and the chemical composition (lithium addition) either non-basal slip or twinning becomes very important once basal slip ceases to operate any further or when it is suppressed at the beginning of deformation. In analogy to type 1, the same analysis of deformation mechanisms based on the initial orientation and deformation geometry can be done for the other specimen types. Interestingly, in cases where tensile twinning was inhibited (types 2 and 3), the samples were less prone to develop a basal texture, and the textures present showed even much weaker peak intensity. Although the initial extrusion textures in Fig. 4 were somewhat similar for AZ31 and Mg4Li, a comparison between the cold PSC textures of the two materials shows obvious differences in terms of texture characteristic and peak intensity. Mg4Li type 1 texture (recorded at a much higher strain than its counterpart of AZ31) was less inclined to build a strong basal orientation. The texture development of types 2 and 3 (showing high similarity even with type 1) exhibited a sort of a RD h1 0 1 0i fiber texture with a very evident maximum intensity splitting of the basal poles. It is well known from polycrystal plasticity texture simulations that splitting of the basal poles about TD is attributed to an increased activity of the non-basal hc + ai-pyramidal slip mode during compressive deformation. Central basal pole splitting is prevalently observed during hot deformation of wrought commercial magnesium alloys such as AZ31, but in the present study it was also evident in the basal texture of the cold rolled Mg4Li specimen (Fig. 7) and in the basal pole figures of Mg4Li types 2 and 3 specimens compressed at room temperature (Fig. 10). hc + aislip is usually a thermally activated process that requires
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elevated deformation temperatures of 250 °C and higher. The current findings, however, based on the distinction of deformation texture evolution between Mg4Li and AZ31 strongly suggests the presence of hc + ai-slip in the lithium containing specimens deformed at room temperature. Such argument that lithium additions promote the activation of hc + ai-slip even at ambient temperatures is not unreasonable since there is no other explanation for the remarkable room temperature ductility of Mg–Li alloys, particularly during sheet forming, at which the majority of grains are forced to be compressed nearly parallel to their c-axes. In fact and as mentioned before, an earlier TEM experimental study by Agnew and co workers [17] investigating the dislocation structures of basal oriented a-solid solution Mg– Li specimens reported a high density and uniform distribution of hc + ai dislocations upon room temperature deformation. increasing deformation temperature With (RT ? 200 °C ? 400 °C) the texture evolution was less prone for developing a basal orientation. This behavior is consistent with the fact that with rising temperatures, twinning became less and less important for accommodating the imposed shape change and conversely, the thermally activated non-basal slip was gaining more influence on the deformation. In case of AZ31 type 1 specimens, it was obvious from Fig. 11 that twinning was still an important deformation mechanism even at 200 °C. The peak intensity, however, was much weaker than the one at room temperature. The weak prismatic texture component near to TD in the (0 0 0 2) pole figure was a remnant of the initial texture that was not able to twin. The texture characteristics and peak intensity of the Mg4Li specimens at 200 °C did not seem to differ much from one sample type to another (Fig. 11). A tendency to propagate along the transverse direction was perspicuous. The observed textures comprised a sharp RD h1 0 1 0i fiber texture component ({1 0 1 0} pole figure) with a peak intensity splitting at ±55° from RD ((0 0 0 2) pole figure). Deformation at 400 °C yielded an interesting texture development, particularly, for the lithium containing specimens (Fig. 8c). Evidently, all specimen types shared a sharp {1 1 2 0}h1 0 1 0iprism texture. This suggests that at elevated temperatures, such as 400 °C, prismatic slip becomes even more important than hc + ai-pyramidal slip. However, this conclusion cannot be generalized at the moment and has to be limited for plane strain deformation geometry only. In case of AZ31-type 1 tensile twinning was rendered completely unimportant and the deformation was driven primarily by prismatic slip indicated by the TD[0 0 0 1] orientation in the (0 0 0 2) basal pole figure. With respect to the peak intensities, the textures in AZ31 deformed at 400 °C were markedly weaker than their counterparts in Mg4Li. In case of AZ31 (where the grain size remained relatively small, Fig. 9e) this can be explained by a previous conclusion that the recrystallization texture of DRX grains plays an important role in weakening the overall texture of the specimen by counteracting the strong deformation texture of the
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Fig. 11. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at 200 °C at a constant strain rate of 104 s1 up to e = 1.4. Levels: 1:2:3:4:5:6:7:8:9:10:12.
deformed microstructure [21]. In case of Mg4Li, the strong texture evolution can be attributed to: (a) The deformation being mainly compensated by one slip mode, i.e. prismatic slip, which results in very well defined textures, as observed in Fig. 12. In other words the more deformation modes involved in accommodating the imposed strain, the more homogeneous the deformation, and the weaker the texture development and anisotropy (as in cubic metals). (b) Sharpening of the recrystallization texture during subsequent grain growth, which was evident in the coarse-grained microstructure in Fig. 9f.
6. Summary and conclusions The combined effects of lithium addition (4 wt.%), the starting texture, and the deformation temperature on the deformation behavior and the evolution of texture and microstructure in magnesium were investigated. The same investigations were carried out on a commercial AZ31
magnesium alloy and the obtained results were compared and discussed for both materials. The experimental findings of the present study are in good accordance with results reported in the literature for similar materials. The following conclusions (focusing on the Li containing alloy) can be drawn: (1) With respect to room temperature formability examined in cold rolling and channel-die compression experiments, the a-solid solution Mg4Li alloy demonstrated a remarkable ductility enhancement over a standard wrought magnesium alloy AZ31. During cold rolling a total thickness reduction of 86% was easily achieved with no signs of cracks or mechanical instabilities. (2) Alloying additions affect the balance of deformation mechanisms that contribute to deformation. This in turn affects the evolution of the crystallographic texture. Lithium addition seemed to result in a congenerous deformation behavior and texture evolution irrespective of the starting texture. This type of behavior was completely different from the one observed in AZ31, in which the stress-
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Fig. 12. Macrotexture evolution of AZ31 and Mg4Li specimens during channel-die compression tests at 400 °C at a constant strain rate of 104 s1 up to e = 1.4. Levels: 1:2:3:4:5:6:7:8:9:10:12:14:16:18.
Fig. 13. Comparison of the yield strength of AZ31 and Mg4Li for different starting textures during PSC deformation at (a) 200 °C and (b) 400 °C.
strain mechanic response (flow curves) and texture evolution were very dependent on the starting texture. (3) After cold rolling Mg4Li specimens showed a basal texture with central peak intensity splitting rotated towards the rolling direction. In channel-die compression the texture developed as a consequence of the operating deformation mechanisms and their
balance, which was influenced by the deformation temperature. At room temperature, Li addition to magnesium seemed to enhance the activation of hc + ai-pyramidal slip. With rising deformation temperature the orientations seemed to spread along TD forming an RD h1 0 1 0i fiber texture component with maximum intensities at ± 55° from RD. At 400 °C the basal poles were completely
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Acknowledgments This work was funded by the Deutsche Forschungsgemeinschaft DFG under grant (GO 335/27). The authors also thank Dr. Dirk Bormann from the Institut fu¨r Werkstoffkunde at the Leibniz Universita¨t Hannover for the kind donation of the magnesium–lithium material. References
Fig. 14. DRX in Mg4Li specimen deformed by channel-die compression at 200 °C/104 s1.
rotated towards TD developing a strong {1 1 2 0}h1 0 1 0i-prism texture. Elevated deformation temperatures (400 °C) render prismatic slip in Mg4Li subjected to PSC more important than hc + ai-pyramidal slip. (4) During deformation at 200 °C dynamic recrystallization (DRX) took place in Mg4Li specimens mainly at grain boundaries and triple junctions (locations of high stress concentration). The deformed grains – being the majority of the microstructure – were elongated in the extension (rolling) direction. At 400 °C the deformed specimen revealed a coarse granular microstructure comprised of grains with an average intercept length of 150 lm. Microstructure coarsening at 400 °C was a consequence of grain growth. By contrast, AZ31 deformed under the same conditions did not reveal any signs of grain growth following recrystallization.
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