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Effect of multimodal microstructure evolution on mechanical properties of Mg–Zn–Y extruded alloy
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Acta Materialia 59 (2011) 3646–3658 www.elsevier.com/locate/actamat
Effect of multimodal microstructure evolution on mechanical properties of Mg–Zn–Y extruded alloy Michiaki Yamasaki a,⇑, Kenji Hashimoto a, Koji Hagihara b, Yoshihito Kawamura a b
a Department of Materials Science, Kumamoto University, 2-39-1 Kurokami, Kumamoto 860-8555, Japan Department of Adaptive Machine Systems, Graduate School of Engineering, Osaka University, 2-1, Yamadaoka, Suita, Osaka 565-0871, Japan
Received 11 January 2011; received in revised form 22 February 2011; accepted 22 February 2011
Abstract A high strength Mg–Zn–Y alloy featuring increased ductility and a multimodal microstructure is developed. The microstructure of the extruded Mg–Zn–Y alloy consists of three regions: a dynamically recrystallized a-Mg fine-grain region with random orientation; a hotworked a-Mg coarse-grain region with strong basal texture; and a long-period stacking ordered (LPSO) phase grain region. Having found that bimodal microstructure evolution in the a-Mg matrix is influenced by the morphology of the LPSO phase in the as-cast state, the authors investigate the effect of secondary dendrite arm spacing (SDAS) in the cast state on the microstructure evolution and mechanical properties of the extruded Mg–Zn–Y alloy. Mg–Zn–Y alloy ingots with various SDAS are obtained by temperaturecontrolled solidification techniques at various cooling rates. Mg–Zn–Y ingots are extruded at 623 K and an extrusion ratio of 10. A decrease in SDAS is associated with dynamic recrystallization of the a-Mg phase region and a high dispersion of fiber-shaped LPSO phase during extrusion. An increase in dynamically recrystallized a-Mg grains with very weak texture improves ductility; the effective dispersion of the hot-worked a-Mg grains with a strong basal texture and the fiber-shaped LPSO phase grains conspire to strengthen the alloy. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloys; Extrusion; Dynamic recrystallization; Texture
1. Introduction Standing first on the list of light structural metallic materials, Mg alloys have been attracting great attention as one of the promising materials for automobile, railway and aerospace applications where weight reduction is of importance. The most important alloying system is Mg– Al alloy with additional alloying elements such as Zn, Mn, Ca, Si, Sr and rare earth (RE). Commercially available AZ- and AM-series alloys are ruling the market [1,2], but much work is being carried on in China regarding Mg– RE alloys [3–7]. RE elements have been used for decades in Europe, North America, and Japan to improve the mechanical property and creep resistance of Mg alloys [8,9]. Although commercially available ZE- and WE-series ⇑ Corresponding author. Tel.: +81 96 342 3705; fax: +81 96 342 3710.
E-mail address: [email protected] (M. Yamasaki).
alloys have so far found little real industrial application, extensive research work has been done with RE-containing Mg alloys with a view to rendering these of practical use in light of their creep behavior [10–18], precipitation hardening [19–30], etc. In the last decade, novel Mg–Zn–Y and Mg–Zn–late lanthanoid (Gd, Tb, Dy, Ho, Er, and Tm) alloys consisting of a-Mg matrix and long-period stacking ordered (LPSO) structure phase have been developed by rapidly solidified powder metallurgy [31] or conventional ingot metallurgy [32–36] to achieve significant improvement in alloy strength and ductility. Different approaches were taken in these efforts. The Mg–Zn–Y and Mg–Zn–late lanthanoid alloys are characterized by a unique microstructure called a/LPSO multi-phase structure. They differ greatly from other RE-containing Mg alloys with respect to microstructure and mechanical performance. The atomistic structure of the LPSO phase was examined using conventional and
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.02.038
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novel analyzing methods such as high-resolution transmission electron spectroscopy (TEM) [37,38], high-angle annular dark-field scanning transmission electron spectroscopy [39–41] and 3D-atom probe [42]. The LPSO phase was found to have a (0 0 0 1) basal plane which is the same as that in Mg (2H, Ramsdell notation), but its stacking periodicity was lengthened 18-fold (18R) or 14-fold (14H) along the c-axis. The 18R- and 14H-LPSO phases have chemical modulation, in which solute elements are enriched in four atomic layers on the closely packed plane at sixand seven-period intervals, respectively [41]. Plastic deformation behavior of Mg12YZn with 18R-LPSO structure was also investigated using directionally solidified crystals. The (0 0 0 1)h1 1 2 0i basal slip was found to be a dominant operative deformation mode in the 18R-LPSO phase, whose critical resolved shear stress (CRSS) was estimated to be 10–30 MPa [43]. The LPSO phase is expected to act as an alloy-strengthening factor in Mg alloys. Besides the unique microstructure of the LPSO phase, LPSO phasecontaining Mg alloys present another subject of inquiry concerning the effects of plastic working on their mechanical properties. Many previous studies found that the conventional thermo-mechanical treatments such as extrusion and rolling caused a marked enhancement in strength without impairing the ductility that is necessary for practical plastic working, though the as-cast alloy showed low strength [32–35,44–46]. Several factors such as a-Mg matrix grain refinement, the formation of suitable texture and the introduction of kink-deformation bands in the LPSO phase [47–50] are presumed to come into play in the enhancement of mechanical properties of the alloy during extrusion, but the strengthening mechanisms of the extruded alloys are yet to be unraveled in detail. Meanwhile, feasibility studies have been under way for the LPSO phase-containing Mg–Zn–RE alloys because LPSO phase-containing Mg alloys have been attracting attention as promising wrought alloys, and mass-production and scaling-up techniques for Mg–Zn–RE alloys are urgently sought in industrial circles. Scaling-up ingot casting often causes grain coarsening owing to slower cooling during solidification. The influence of casting-related grain coarsening and dendrite growth on the microstructure evolution of extruded Mg–Zn–RE alloys also remains to be clarified. These considerations show that it is important to clarify the relationship between microstructure evolution and the mechanical properties of extruded Mg–Zn–Y alloys with an LPSO phase. This study focused on the secondary dendrite arm spacing (SDAS) as a factor playing an important role in the fabrication of extruded alloys from the standpoint of both fundamental research and the feasibility of scaling-up and mass production. The morphology of the LPSO phase in cast ingots can be controlled by changes in the SDAS. Several Mg–Zn–Y alloys with varying SDAS were cast at various cooling rates ranging from 0.01 to 10 K s1, and in this study the influence of initial microstructure to the evolution of microstructure during
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extrusion and the resulting mechanical properties were examined. The ductility enhancement and strengthening mechanisms are discussed in relation to the formation of a multimodal microstructure in extruded Mg–Zn–Y alloys. 2. Experimental procedure Mg97Zn1Y2 (at.%) cast ingot was produced by highfrequency induction melting of pure Mg (99.99 wt.%), Zn (99.9 wt.%) and Y (99.9 wt.%) metals in a carbon crucible. The molten alloys kept at 1023 K were cast in a steel mold with a hole 30 mm in diameter and 100 mm in length in an argon atmosphere, or solidified in a carbon crucible with a hole 70 mm in diameter and 80 mm in length at various cooling rates. Four cooling rates of 0.06 K s1, 0.18 K s1, 5.7 K s1, and 9.6 K s1 were derived from temperature-controlled solidification. In gravity casting, the initial temperatures of the steel mold were kept at 360 K and 673 K by electric heating during solidification. The wall thickness of the mold used in this study was 20 mm. In crucible cooling, the molten alloys were heated to 1023 K and then cooled at two cooling rates, controlled by changing induction heating input. The cooling history of the alloy during the solidification was monitored with thermocouples installed at three different positions (bottom, middle and top of the solidified alloys) in the mold and the crucible. Cooling rates, Rc (K s1), in each solidification were estimated by the following equation: Rc ¼
ðT m T s Þ t
ð1Þ
where Tm is the liquidus temperature measured at middle position of solidified alloy, Ts is the solidus temperature, and t is the solidification time in seconds. These ingots were machined into round bars 29 mm in diameter and 70 mm in length. The round billets were extruded at an extrusion ratio of 10, extrusion temperature 623 K, and ram speed 2.5 mm s1. In this study, axisymmetric extrusion was adopted using a round die; long rod-like alloys with diameter 9 mm and length 600 mm were extruded from the ascast ingots. The constituent phases in the alloy were examined by X-ray diffractometry (XRD; Bruker AXS D8 Discover with Cu Ka radiation). XRD data were collected in the 2h diffraction angle range of 20–90°. The step width and counting time in each step were 0.02° and 1 s, respectively. The microstructure of the specimens was observed by confocally optical microscopy (OM; Lasertec C-130), scanning electron microscopy (SEM; JEOL JSM-7001F) and transmission electron microscopy (TEM; JEOL JEM-2000FX). The specimens were polished mechanically and then chemically with a mixture of 80 ml ethyl alcohol and 20 ml nitric acid. Finally, they were etched with a mixture of 4.2 g picric acid, 10 ml acetic acid, 80 ml ethyl alcohol and 10 ml water for OM observation of the microstructure. The grain size in the alloys was measured by the liner intercept method in OM observation. SEM images were taken in the secondary
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electron mode and the back-scattered electron (BSE) mode. Specimens were polished by paper lapping and ion-milling (JEOL Cross Section Polisher SM-09010) for SEM observation. The volume fraction of the LPSO phase was computed from binary images converted from gray-scale SEM–BSE images using the NIH image software package. TEM specimens were thinned by mechanical polishing using ion-milling with a Fischione Model 1010. The textures of the extruded alloys were analyzed by electron back-scatter diffraction pattern analysis (EBSD) with orientation imaging microscopy (OIM; TSL Solutions K.K.) and by pole figure analysis (PF) performed with an XRD with a collimator of diameter 1 mm. PF data were collected within specimen tilt angle (v) range 0–86° and specimen rotation angle (/) range 0–360°. Both angles v and / were increased stepwise in increments of 2°, and the counting time in each step was 1 s for angle /. Tensile tests were carried out using an Instron testing machine (Instron Model 5584) at room temperature with an initial strain rate of 5 104 s1. The gauge section of tensile specimens was 2.5 mm and 15 mm in diameter and length, respectively. In the ambient-temperature testing, strain was measured with an electrical strain gauge. The tensile axis was held parallel to the direction of extrusion. The 0.2% proof strength was employed as yield strength.
3. Results 3.1. Microstructure of cast ingot Fig. 1 shows optical micrographs of Mg97Zn1Y2 alloys that were solidified at various cooling rates. Mg97Zn1Y2 alloy consisted of an a-Mg matrix and a LPSO phase, with dendritic primary crystals of a-Mg phase. Grain size decreased with increasing cooling rates. The morphology of dendritic primary crystals was also observed in the Mg97Zn1Y2 alloy, and SDAS decreased with increasing cooling rates. Grain size and SDAS in the as-cast state are shown in Table 1. The changes in the grain size dg and SDAS d2, in Mg97Zn1Y2 cast alloy as functions of cooling rate Rc are summarized in Fig. 2. A commonly accepted empirical relationship between SDAS and cooling rate is given by [51,52] d 2 ¼ A RBc
ð2Þ
where A and B are constants for a given alloy system. The values of the constants for Mg97Zn1Y2 alloy were determined by power approximation to be A = 37.1 and B = 0.33. The solid line in Fig. 2 represents power approximation of the experimental results. Fig. 3 shows the XRD pattern of the Mg97Zn1Y2 alloy solidified at
Fig. 1. Optical micrographs of Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1, and (d) 9.6 K s1.
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Table 1 Grain size and SDAS of the Mg97Zn1Y2 alloy solidified at various cooling rates.
crystals of the a-Mg phase, thus improving the dispersion of the LPSO phase in the cast state.
Cooling rate (K s1)
Grain size (mm)
SDAS (lm)
0.061 0.18 5.7 9.6
0.91 ± 0.25 0.86 ± 0.21 0.43 ± 0.19 0.41 ± 0.16
88 ± 8.2 68 ± 7.9 25 ± 3.3 15 ± 2.8
3.2. Microstructure and mechanical properties of extruded alloy
Fig. 2. Change in the grain size and SDAS of Mg97Zn1Y2 cast alloys as a function of cooling rate.
cooling rate 9.6 K s1. Peaks originating from the 18RLPSO phase and the a-Mg matrix phase were detected [37]. Fig. 4 shows SEM–BSE images of Mg97Zn1Y2 cast alloys at various solidification cooling rates. Bright and lamellar contrast corresponds to the LPSO phase; the average volume fraction of LPSO phase in Mg97Zn1Y2 was 24%, no matter how the cooling rate varied. Increasing the cooling rate led to refinement of the dendritic primary
Fig. 3. XRD patterns of Mg97Zn1Y2 alloys solidified at cooling rates of 9.6 K s1.
Optical micrographs of extruded Mg97Zn1Y2 alloys taken from a longitudinal section are shown in Fig. 5. The direction of extrusion is parallel to the horizontal direction in all the figures presented in this paper. Three feature areas can be seen: the dynamically recrystallized (DRX-ed) a-Mg fine-grain region; the hot-worked a-Mg coarse-grain region; and the 18R-LPSO phase region. Changes in the volume fractions of two types of matrices, the DRX-ed a-Mg region and the hot-worked a-Mg region, are shown as functions of SDAS in Fig. 6. Although the average grain size of the DRX-ed a-Mg fine-grain region in the extruded alloys was 2.5 lm regardless of different SDAS, the volume fraction of the DRX-ed a-Mg region increased with decreasing SDAS. This suggests that the SDAS is a major determinant of the volume fraction of the DRX-ed grains region in the extruded alloy as well as the LPSO phase morphology in the as-cast state. Fig. 7 shows SEM–BSE images of extruded Mg97Zn1Y2 alloys taken from the longitudinal section. SEM–BSE images revealed the changes in dispersion of the LPSO phase that were affected by a change in SDAS. Furthermore, many kink-deformation bands were seen to have been introduced in the LPSO phase, as reported previously [33,43,44,48,53]. To describe the change in the dispersion of LPSO quantitatively, the dispersion level DL (lm1) was defined for discussion in this study by the following expression [50]: P Nv DL ¼ P ð3Þ Lv where Lv (lm) is the length of a vertical line segment L drawn perpendicular to the extrusion direction in Fig. 7, and Nv is the number of LPSO phase P intersected by segment L. The total length of segments Lv is more than 5 mm. Fig. 8 shows the relationship between the dispersion of LPSO phase in the extruded Mg97Zn1Y2 alloys and the SDAS in the as-solidified state. Reduction of SDAS promoted dispersion and refinement of the fiber-shaped LPSO phase during extrusion. In order to clarify texture evolution during extrusion, PF and inverse pole figure (IPF) measurements were performed by XRD and EBSD, respectively. Measurements were made with two extruded Mg97Zn1Y2 alloys that were cast at cooling rates of 0.061 and 9.6 K s1. The former ascast ingot had SDAS of 88 lm and the latter 15 lm (Table 1). Hereinafter, extruded Mg97Zn1Y2 alloys that were solidified at cooling rates of 0.061 and 9.6 K s1 will be called 88-lm-extruded Mg97Zn1Y2 alloy and 15-lmextruded Mg97Zn1Y2 alloy, respectively. The {1 0 1 0} X-ray pole figures of the 88-lm-extruded Mg97Zn1Y2 alloy
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Fig. 4. SEM-BSE images of Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1 and (d) 9.6 K s1.
and the 15-lm-extruded Mg97Zn1Y2 alloy are shown in Fig. 9. The measured planes were transverse cross-section (perpendicular to the direction of extrusion) and longitudinal section of the extruded bars. The extruded alloys had a strong basal texture. The <1 0 1 0> axis of the hexagonal close packed (hcp) structure is parallel to the direction of extrusion. Fig. 10 presents the IPF map obtained from a longitudinal section of the 88-lm-extruded Mg97Zn1Y2 alloy. Colors designate the crystallographic orientation with respect to the sample normal, coded as shown in the color key triangle. In the IPF mapping, a confidence index (CI) was used to indicate the degree of confidence as to whether the orientation calculation is correct [54]. The shaded areas correspond to the regions where a CI value is <0.1, i.e. the OIM system cannot analyze the Kikuchi diffraction patterns. The large shaded areas most closely represent the 18RLPSO phase grains, which fact was confirmed by comparing them with the SEM images. IPF mapping also revealed that there were two kinds of matrices, that is, the DRX-ed fine-grain region (DRX-ed region) and the hot-worked coarse-grain region where recrystallization did not occur (worked region). It is noteworthy that dynamic recrystallization occurs in the a-Mg matrix next to the LPSO phase. This suggests that a strain sufficient to induce
recrystallization is concentrated in the neighborhood of the LPSO phase. Fig. 11 shows IPF maps of DRX-ed regions and worked regions in the 88-lm-extruded Mg97Zn1Y2 alloy shown in Fig. 10. In order to express clearly the relationship between extrusion direction and texture, Fig. 12 shows EBSD-derived pole figures taken from a wide-scan region (160 320 lm) of the 88-lmextruded Mg97Zn1Y2 alloy and a narrow-scan region (80 100 lm) selected from only the DRX-ed zone. Wide-scan EBSD-derived pole figures and X-ray pole figures also revealed that there was a strong basal texture in the extruded alloy. However, DRX-ed zone scanning of EBSD-derived pole figure revealed that the basal fiber texture in DRX-ed grains is much weaker, and the DRX-ed grains had almost randomized crystallographic orientation. Fig. 13 presents the IPF maps obtained from a longitudinal section of the 15-lm-extruded Mg97Zn1Y2 alloy. The area of the DRX-ed region expanded, and the area of the worked grains region decreased in comparison with the 88-lm-extruded Mg97Zn1Y2 alloy. Fig. 14 shows IPF maps of DRX-ed regions and worked regions in the 15-lmextruded Mg97Zn1Y2 alloy. DRX-ed regions were identified from grain size distribution. The DRX-ed grains ranged in size from less than a micron to 4 lm, the size peaking at 1 lm. In this study, a grain measuring 4 lm
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Fig. 5. Optical micrographs of extruded Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1 and (d) 9.6 K s1. Micrographs are of longitudinal section.
Fig. 6. Change in the volume fractions of two types of matrices (the DRXed a-Mg region and the hot-worked a-Mg region) as a function of SDAS.
or less was considered to be a DRX-ed grain. Fig. 15 shows EBSD-derived pole figures from the entire EBSD scan
region (160 320 lm) of the 15-lm-extruded Mg97Zn1Y2 alloy. The 15-lm-extruded Mg97Zn1Y2 alloy had a weaker basal texture than the 88-lm-extruded Mg97Zn1Y2 alloy, because of its reduced region, which has a strong h1 0 1 0i texture. Because the OIM system was unable to analyze the Kikuchi diffraction patterns from the LPSO phase owing to the large amount of residual strain, TEM observation was performed to confirm the morphology and texture evolution of the LPSO phase. Fig. 16 shows a TEM image of the fiber-shaped LPSO phase embedded in the DRXed a-Mg grain region of the 15-lm-extruded Mg97Zn1Y2 alloy. Almost all the LPSO phases observed in this study were elongated along the extrusion direction and had a strong basal texture in which the c-axes of the grains were perpendicular to the direction of extrusion. The basal slip was strongly suppressed in the extruded alloy, most probably because the Schmid factor was very small when the tensile direction was parallel to the extrusion direction, so that very high yield stress arose accompanied by the formation of a deformation kink. Fig. 17 shows changes in the yield tensile strength, ultimate tensile strength and elongation of the extruded alloys
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Fig. 7. SEM–BSE images of extruded Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1, and (d) 9.6 K s1. Micrographs are of longitudinal section.
4. Discussion
Fig. 8. Relationship between the extruded Mg97Zn1Y2 alloys and the SDAS with respect to the as-solidified state in LPSO phase dispersion.
as functions of SDAS in the as-cast state. The tensile axis was along the direction of extrusion. Tensile yield strength and elongation of the extruded alloys simultaneously increased with decreasing SDAS, although, in general, there is a trade-off between strength and elongation.
This study has shown that there are three typical regions in extruded Mg–Zn–Y alloys: the DRX-ed fine-grain region; the hot-worked a-Mg coarse-grain region; and the fiber-shaped LPSO phase region. Although the LPSO phase is presumed to help strengthen the alloys, it is necessary to clarify which region does what to help improve the strength and ductility of the extruded alloy. TEM observation revealed that the LPSO phase elongated along the direction of extrusion had a strong basal texture in which the c-axes of the grains were perpendicular to the direction of extrusion, as shown in Fig. 16. This basal fiber texture is very suitable for the strengthening of the alloy, since the operation of the basal slip, which is the dominant deformation mode also in the LPSO phase, is strongly hindered by the texture. Very recently, Mayama et al. [59] demonstrated by calculation that a large CRSS ratio of a basal slip to a non-basal slip system produced a strong basal texture owing to crystal rotation during extrusion, bringing the h1 0 1 0i axis of LPSO phase into a position parallel to the direction of extrusion. The LPSO phase as well as the hcp-Mg has a large CRSS ratio of the (0 0 0 1)h1 1 2 0i basal slip to the non-basal slip. Moreover, in the LPSO
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Fig. 9. {1 0 1 0} X-ray pole figures of the extruded Mg97Zn1Y2 alloy: (a) transverse section of 88-lm-extruded alloy; (b) longitudinal section of 88-lmextruded alloy; (c) transverse section of 15-lm-extruded alloy; (d) longitudinal section of 15-lm-extruded alloy.
phase, both the twin deformation and the dynamic recrystallization hardly occur during extrusion, though they easily occur in the hcp-Mg metal. Taken altogether, the LPSO phase can be considered capable of developing a strong basal texture owing to crystal rotation during extrusion. The difference in plastic deformation behavior between the LPSO and the a-Mg phase accounts for the formation of the fiber-shaped LPSO phase-reinforced metal matrix composite during extrusion. Therefore, the dispersion of the fiber-shaped LPSO phase can be considered to be of particular significance. As shown in Fig. 8, the degree of dispersion of the LPSO phase increased with decreasing SDAS in the as-cast state, though the volume fraction of LPSO phase did not change. Fig. 18 shows the effect of the LPSO phase dispersion on alloy strengthening. An increase in LPSO phase dispersion brought about an improvement in the tensile yield strength of the alloy in proportion. It should be noted that twinning was not found in the LPSO phase-containing Mg–Zn–Y extruded alloys, while some kink-deformation bands were observed. Matsuda et al. [60] reported that a densely developed LPSO phase prevented the growth of {1 0 1 2} deformation twins in the a-Mg matrix in rapidly solidified Mg97Zn1Y2 alloy. Very recently, Hagihara et al. [43,44,48] reported that the dominant operative deformation mode in
0i the 18R-LPSO phase was identified as a (0 0 0 1)h1 1 2 basal slip and that twin deformation never occurred. Based on the findings above, the present authors are inclined to presume that the highly dispersive fiber-shaped LPSO phase grains, along with the significantly refined a-Mg matrix grains, in the extruded alloys suppressed twin deformation. Furthermore, the formation of a kink-deformation band in the LPSO phase and a-Mg phase region may strengthen the alloy. The reason is that, since kink-deformation bands are macroscopically formed perpendicularly to the primary slip direction, they effectively counteract basal slipping [61]. The fiber-shaped LPSO phase with kink-deformation bands is one of the important factors that come into play in strengthening of the Mg–Zn–Y alloys. However, in consideration of the a-Mg matrix phase, the existence of two kinds of matrices seems to greatly influence the ductility of the alloy. Fig. 19 shows the relationship between the volume fractions of two kinds of matrices and the mechanical properties of extruded Mg97Zn1Y2 alloy. It is easy to understand that an increase in the volume fraction of fine grains is conductive to alloy strengthening through the Hall–Petch relation [55,56]. Furthermore, Fig. 19 suggests that a DRX-ed a-Mg finegrain region with almost randomized crystallographic
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Fig. 11. IPF maps and IPF of (a and b) DRX-ed regions and (c and d) worked regions in the 88-lm-extruded Mg97Zn1Y2 alloy.
Fig. 10. IPF map obtained from a longitudinal section of the 88-lmextruded Mg97Zn1Y2 alloy.
orientation helps to improve elongation in the extruded alloys. Average Schmid factors are available for prediction of the response of a textured polycrystalline material to a given stress state. Average Schmid factors for the (0 0 0 1)h1 1 2 0i slip of the 88-lm- and 15-lm-extruded Mg97Zn1Y2 alloys were estimated as 0.17 and 0.23, respectively. These figures suggest that the basal slip can be so readily operative in 15-lm-extruded Mg97Zn1Y2 alloys that it helps to enhance ductility. In addition, Koike et al. [57] proposed new deformation mechanisms in the fine-grained Mg alloys. They discussed the activation of non-basal slip systems attributed to compatibility stress at grain boundaries produced during large elongation. When grain size is so small that the grain-boundary
Fig. 12. EBSD-derived pole figures of the 88-lm-extruded Mg97Zn1Y2 alloy: (a) wide-scan region (160 320 lm); (b) narrow-scan region (80 100 lm) selected solely from the DRX-ed zone.
affected region can cover the entire grain, the ductility of the Mg alloy is improved. Therefore, the fine DRX-ed grains formed near an LPSO grain are very likely to do much to improve ductility. On close examination of the hot-worked matrix region, the worked a-Mg grains were found to have a strong basal texture, as shown in Figs. 11 and 14. When the tensile direction was parallel to the direction of extrusion, the Schmid factor for the basal slip in the hot-worked a-Mg grains
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Fig. 14. IPF maps and IPF of (a and b) DRX-ed regions and (c and d) worked regions in the 15-lm-extruded Mg97Zn1Y2 alloy.
Fig. 15. EBSD-derived pole figures from the entire EBSD scan region (160 320 lm) of the 15-lm-extruded Mg97Zn1Y2 alloy. Fig. 13. IPF maps obtained from a longitudinal section of the 15-lmextruded Mg97Zn1Y2 alloy.
turned out to be a small value. In addition, since most of the basal planes in the grains of extruded Mg alloys lie parallel to the direction of extrusion, {1 0 1 2} twinning is also inhibited in tensile deformation along the direction of extrusion. This phenomenon led to achievement of high yield strength [58]. The textured a-Mg grains embedded in fine-grained a-Mg matrix works as a kind of reinforced component. As SDAS diminishes, the alteration of the worked coarse grains into a fiber shape might enhance the effect of composite strengthening, considering the possibility that the increase in dispersion may compensate the decrease in their volume fraction to some extent. Fig. 20 shows the relationship between tensile yield strength and elongation of extruded Mg97Zn1Y2 alloys. Mg96Zn2Y2 extruded alloys containing LPSO phase [49],
Fig. 16. TEM image of the fiber-shaped LPSO phase embedded in DRXed a-Mg grain region of the 15-lm-extruded Mg97Zn1Y2 alloy.
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Fig. 17. Yield tensile strength, ultimate tensile strength and elongation of the extruded alloys as a function of SDAS in as-cast state.
Fig. 18. Change in the yield strength of the extruded alloy as a function of LPSO phase dispersion.
Mg–Zn–RE extruded alloys not containing an LPSO phase, and conventional extruded alloys (AZ31B, AZ61A, AZ91D, AM60, WE54) are included in this figure for purposes of comparison. The extrusion conditions (temperature, extrusion ratio and ram speed) were the same with all Mg–Zn–RE alloys. In general, an increase in strength due to solution hardening or precipitation hardening accompanies a decrease in the ductility of the materials. However, this trade-off between strength and ductility does not occur in the extruded Mg alloy with an LPSO phase under discussion. Recently, it has been reported that bimodal nanostructured metal materials with micrometer-sized grains embedded inside a matrix of nanocrystalline and ultra-fine grains exhibit high strength and ductility at once [62,63]. In these materials, micrometer-sized grains obtained through secondary recrystallization produced strain hardening sufficient to
Fig. 19. Relation between the volume fractions of two kinds of matrices and the mechanical properties of extruded Mg97Zn1Y2 alloy.
Fig. 20. Relation between the tensile yield strength of extruded Mg97Zn1Y2 and Mg96Zn2Y2 alloys and the elongation.
sustain a reasonably uniform deformation to the extent of large elongation [64,65]. In the extruded Mg–Zn–Y alloys under discussion, bimodal grain size and texture distribution may warrant both high strength and ductility of the alloys. The LPSO phase plays such an important role in multimodal microstructure formation during normal extrusion in this alloy system. The results suggest that the DRX-ed a-Mg fine grains do much to improve the ductility of the alloys. The micrometer-sized grains that have a random crystallographic orientation are extremely effective in enhancing ductility, because Mg metal shows strong plastic anisotropy derived from the hcp crystal structure [57]. The findings that fine grains of hcp-Mg enhance the ductility of the alloys and that coarse a-Mg and LPSO phase grains strengthen the alloys are of interest. Both the fine grains and the coarse grains act in a manner opposite
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to the bimodally microstructured face-centered cubic (fcc)Cu [62,63] or fcc-Al [64,65]. Heterogeneity of the microstructure in Mg alloys has often been discussed. Ion et al. [66] reported that the centers of old grains (cores) and equiaxed subgrains near the old grain boundaries (mantle region) were developed in an Mg–0.8 wt.%Al extruded alloy deformed 16% at 533 K. The former region had a strong basal texture and kink deformation banding, while the latter region had a random crystallographic orientation. They reported that in deformation <600 K, lattice rotation occurred in the regions around the grain boundaries by the basal slip accompanied by kinking due to the relative operating difficulty of the non-basal slip, while the cores of the grains remained relatively undeformed. This leads to a dynamic recrystallization involving the dynamic polygonization of rotated lattice (mantle) regions adjacent to grain boundaries, resulting in the formation of a bimodal microstructure. The existence of the LPSO phase is considered to accelerate the formation of refined recrystallized grains around them by a similar mechanism. In addition, Stanford et al. reported that micro-alloying additions of RE weakened the texture in extruded binary Mg-based alloys and that their extruded alloys showed good ductility [67– 69]. These results suggest that optimum plastic deformation and addition of optimum alloying element can produce a controlled heterogeneous microstructure in Mg alloys. In this study, SDAS in an a/LPSO two-phase Mg alloy was found to be an important factor for controlling the heterogeneous microstructure. A decrease in SDAS involves the dispersion of the LPSO phase in the as-cast state, resulting in the expansion of the DRX-ed region. In particular, basal-textured fiber-shaped LPSO phase grains embedded in a non-textured fine-grained a-Mg matrix can produce an improvement in both the strength and the ductility of the alloy simultaneously. 5. Conclusions Multimodally microstructured Mg–Zn–Y alloys were prepared by casting and extrusion processes. Mg–Zn–Y alloy ingots with various SDAS were extruded, and then tensile-tested at room temperature. The alloys exhibited high strength and reasonable ductility for practical purposes of use. OM and SEM observations revealed that the extruded alloy had three typical regions: a DRX-ed a-Mg fine-grain region; a hot-worked a-Mg coarse-grain region; and a fiber-shaped LPSO phase region. EBSD analysis characterized two matrices: the DRX-ed fine grains had random crystallographic orientation, and the hotworked coarse grains had a strong texture with basal planes parallel to the direction of extrusion. The former grain region seems to account for the ductility of the alloys and the latter seems to strengthen them. In addition, the LPSO phase acts as a fiber-like reinforcement in effectively increasing the strength of the alloy due to its strong basal fiber texture and containing deformation kink bands. As
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to alloy strengthening, one should not fail to note that there is another factor at work when the alloys are strengthened. That is an increase in LPSO phase dispersion. SDAS in the cast state had a considerable effect on bimodal a-Mg matrix microstructure evolution and LPSO phase dispersion in the extruded Mg–Zn–Y alloys. Acknowledgements This work is supported by the Kumamoto Prefecture CREATE Project from JST and the Grants-in-aid for Young Scientists (A) No. 2086050 and Scientific Research (A) No. 19206075 from MEXT, Japan. References [1] Froes FH, Eliezer D, Aghion E. JOM 1998;50(9):30. [2] Luo AA. Int Mater Rev 2004;49:13. [3] Yang Z, Li JP, Zhang JX, Lorimer GW, Robson J. Acta Metall Sin (Engl Lett) 2008;21:313. [4] Ma C, Liu M, Wu G, Ding WJ, Zhu YP. Mater Sci Eng A 2003;349:207. [5] Xu DK, Liu L, Zu YB, Han EH. Mater Sci Eng A 2006;420:322. [6] He SM, Peng LM, Zeng XQ, Ding WJ, Zhu YP. Mater Sci Eng A 2006;433:175. [7] Peng Q, Wang J, Wu Y, Wang L. Mater Sci Eng A 2006;433:133. [8] Mordike BL, Ebert T. Mater Sci Eng A 2001;302:37. [9] Pekguleryuz MO, Kaya AA. Adv Eng Mater 2003;5:886. [10] Suzuki M, Sato H, Maruyama K, Oikawa H. Mater Sci Eng A 1998;A252:248. [11] Anyanwu IA, Kamado S, Kojima Y. Mater Trans 2001;42:1212. [12] Suzuki M, Sato H, Maruyama K, Oikawa H. Mater Sci Eng A 2001;A319–321:751. [13] Wang JG, Hsiung LM, Nieh TG, Mabuchi M. Mater Sci Eng A 2001;A315:81. [14] Mordike BL. Mater Sci Eng A 2002;A324:103. [15] Maruyama K, Suzuki M, Sato H. Metall Mater Trans A 2002;33:875. [16] Suzuki M, Kimura T, Koike J, Maruyama K. Scripta Mater 2003;48:997. [17] Smola B, Stulikova I, Pelcova´ J, Mordike BL. J Alloy Compd 2004;378:196. [18] Nie JF, Gao X, Zhu SM. Scripta Mater 2005;53:1049. [19] Nie JF, Muddle BC. Scripta Mater 1999;40:1089. [20] Vostry P, Smola B, Stulikova I, von Buch F, Mordike BL. Phys Stat Sol (a) 1999;175:491. [21] Nie JF, Muddle BC. Acta Mater 2000;48:1691. [22] Anyanwu IA, Kamado S, Kojima Y. Mater Trans 2001;42:1206. [23] Smola B, Stulikova I, von Buch F, Mordike BL. Mater Sci Eng A 2002;A324:113. [24] Apps PJ, Karimzadeh H, King JF, Lorimer GW. Scripta Mater 2003;48:1023. [25] Honma T, Ohkubo T, Hono K, Kamado S. Mater Sci Eng A 2005;A395:301. [26] Gao X, He SM, Zeng XQ, Peng LM, Ding WJ, Nie JF. Mater Sci Eng A 2006;A431:322. [27] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2006;47:2109. [28] Nishijima M, Yubuta K, Hiraga K. Mater Trans 2007;48:84. [29] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2008;49:227. [30] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2009;50:1747. [31] Kawamura Y, Hayashi K, Inoue A, Masumoto T. Mater Trans 2001;42:1172.
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M. Yamasaki et al. / Acta Materialia 59 (2011) 3646–3658
[32] Yamasaki M, Anan T, Yoshimoto S, Kawamura Y. Scripta Mater 2005;53:799. [33] Yoshimoto S, Yamasaki M, Kawamura Y. Mater Trans 2006;47:959. [34] Kawamura Y, Kasahara T, Izumi S, Yamasaki M. Scripta Mater 2006;55:453. [35] Kawamura Y, Yamasaki M. Mater Trans 2007;48:2986. [36] Yamasaki M, Sasaki M, Nishijima M, Hiraga K, Kawamura Y. Acta Mater 2007;55:6798. [37] Luo ZP, Zhang SQ. J Mater Sci Lett 2000;19:813–5. [38] Itoi T, Seimiya T, Kawamura Y, Hirohashi M. Scripta Mater 2004;51:107. [39] Abe E, Kawamura Y, Hayashi K, Inoue A. Acta Mater 2002;50:3845. [40] Zhu YM, Morton AJ, Nie JF. Acta Mater 2010;58:2936. [41] Yokobayashi H, Kishida K, Inui H, Yamasaki M, Kawamura Y. MRS Symp Proc 2011;1295, doi:10.1557/opl.2011.36. [42] Ping DH, Hono K, Kawamura Y, Inoue A. Philos Mag Lett 2002;82:543. [43] Hagihara K, Yokotani N, Umakoshi Y. Intermetallics 2010;18:267. [44] Hagihara K, Kinoshita A, Sugino Y, Yamasaki M, Kawamura Y, Yasuda HY, et al. Intermetallics 2010;18:1079. [45] Itoi T, Takahashi K, Moriyama H, Hirohashi M. Scripta Mater 2008;59:1155. [46] Itoi T, Inazawa T, Kuroda Y, Yamasaki M, Kawamura Y, Hirohashi M. Mater Lett 2010;64:2277. [47] Morikawa T, Kaneko K, Higashida K, Kinoshita D, Takenaka M, Kawamura Y. Mater Trans 2008;49:1294. [48] Hagihara K, Kinoshita A, Sugino Y, Yamasaki M, Kawamura Y, Yasuda HY, et al. Acta Mater 2010;58:6282. [49] Yamasaki M, Hashimoto K, Hagihara K, Kawamura Y. Mater Sci Forum 2010;654–656:615. [50] Hirano M, Yamasaki M, Hagihara K, Higashida K, Kawamura Y. Mater Trans 2010;51:1640.
[51] Mehrabian R. In: Mehrabian R, Kear BH, Cohen M, editors. Rapid solidification processing: principles and technologies, proceedings of the international conference on rapid solidification processing. Baton Rouge, LA: Claitor’s Pub. Division; 1978. p. 9–27. [52] Khedkar P, Nash P. J Mater Sci Lett 1986;5:1273. [53] Shao XH, Yang ZQ, Ma XL. Acta Mater 2010;58:4760. [54] Wright SI. Fundamentals of automated EBSD. In: Schwartz AJ, Kumar M, Adams BL, editors. Electron backscatter diffraction in materials science. New York: Kluwer Academic/Plenum Publishers; 2000. p. 58. [55] Armstrong RW, Codd I, Douthwaite RM, Petch NJ. Philos Mag 1962;7:45. [56] Sambasiva Rao G, Prasad YVRK. Metal Trans 1982;13A:2219. [57] Koike J, Kobayashi T, Mukai T, Watanabe H, Suzuki M, Maruyama K, et al. Acta Mater 2003;51:2055. [58] Kleiner S, Uggowitzer PJ. Mater Sci Eng A 2004;A379:258. [59] Mayama T, Noda M, Chiba R, Kuroda M. Int J Plasticity, in press, doi:10.1016/j.ijplas.2011.02.007. [60] Matsuda M, Ii S, Kawamura Y, Ikuhara Y, Nishida M. Mater Sci Eng A 2004;A386:447. [61] Higashida K, Takamura J, Narita N. Mater Sci Eng 1986;81:239. [62] Wang Y, Chen M, Zhou F, Ma E. Nature 2002;419:912. [63] Wang Y, Ma E. Acta Mater 2004;52:1699. [64] Witkin D, Lee Z, Rodriguez R, Nutt S, Lavernia E. Scripta Mater 2003;49:297. [65] Han BQ, Lee Z, Witkin D, Nutt S, Lavernia EJ. Metal Mater Trans A 2005;36A:957. [66] Ion SE, Humphreys FJ, White SH. Acta Metall 1982;30:1909. [67] Cottam R, Robson J, Lorimer G, Davis B. Mater Sci Eng A 2008;A485:375. [68] Stanford N, Barnett MR. Mater Sci Eng A 2008;A496:399. [69] Stanford N. Mater Sci Eng A 2010;A527:2669.
Acta Materialia 59 (2011) 3646–3658 www.elsevier.com/locate/actamat
Effect of multimodal microstructure evolution on mechanical properties of Mg–Zn–Y extruded alloy Michiaki Yamasaki a,⇑, Kenji Hashimoto a, Koji Hagihara b, Yoshihito Kawamura a b
a Department of Materials Science, Kumamoto University, 2-39-1 Kurokami, Kumamoto 860-8555, Japan Department of Adaptive Machine Systems, Graduate School of Engineering, Osaka University, 2-1, Yamadaoka, Suita, Osaka 565-0871, Japan
Received 11 January 2011; received in revised form 22 February 2011; accepted 22 February 2011
Abstract A high strength Mg–Zn–Y alloy featuring increased ductility and a multimodal microstructure is developed. The microstructure of the extruded Mg–Zn–Y alloy consists of three regions: a dynamically recrystallized a-Mg fine-grain region with random orientation; a hotworked a-Mg coarse-grain region with strong basal texture; and a long-period stacking ordered (LPSO) phase grain region. Having found that bimodal microstructure evolution in the a-Mg matrix is influenced by the morphology of the LPSO phase in the as-cast state, the authors investigate the effect of secondary dendrite arm spacing (SDAS) in the cast state on the microstructure evolution and mechanical properties of the extruded Mg–Zn–Y alloy. Mg–Zn–Y alloy ingots with various SDAS are obtained by temperaturecontrolled solidification techniques at various cooling rates. Mg–Zn–Y ingots are extruded at 623 K and an extrusion ratio of 10. A decrease in SDAS is associated with dynamic recrystallization of the a-Mg phase region and a high dispersion of fiber-shaped LPSO phase during extrusion. An increase in dynamically recrystallized a-Mg grains with very weak texture improves ductility; the effective dispersion of the hot-worked a-Mg grains with a strong basal texture and the fiber-shaped LPSO phase grains conspire to strengthen the alloy. Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Magnesium alloys; Extrusion; Dynamic recrystallization; Texture
1. Introduction Standing first on the list of light structural metallic materials, Mg alloys have been attracting great attention as one of the promising materials for automobile, railway and aerospace applications where weight reduction is of importance. The most important alloying system is Mg– Al alloy with additional alloying elements such as Zn, Mn, Ca, Si, Sr and rare earth (RE). Commercially available AZ- and AM-series alloys are ruling the market [1,2], but much work is being carried on in China regarding Mg– RE alloys [3–7]. RE elements have been used for decades in Europe, North America, and Japan to improve the mechanical property and creep resistance of Mg alloys [8,9]. Although commercially available ZE- and WE-series ⇑ Corresponding author. Tel.: +81 96 342 3705; fax: +81 96 342 3710.
E-mail address: [email protected] (M. Yamasaki).
alloys have so far found little real industrial application, extensive research work has been done with RE-containing Mg alloys with a view to rendering these of practical use in light of their creep behavior [10–18], precipitation hardening [19–30], etc. In the last decade, novel Mg–Zn–Y and Mg–Zn–late lanthanoid (Gd, Tb, Dy, Ho, Er, and Tm) alloys consisting of a-Mg matrix and long-period stacking ordered (LPSO) structure phase have been developed by rapidly solidified powder metallurgy [31] or conventional ingot metallurgy [32–36] to achieve significant improvement in alloy strength and ductility. Different approaches were taken in these efforts. The Mg–Zn–Y and Mg–Zn–late lanthanoid alloys are characterized by a unique microstructure called a/LPSO multi-phase structure. They differ greatly from other RE-containing Mg alloys with respect to microstructure and mechanical performance. The atomistic structure of the LPSO phase was examined using conventional and
1359-6454/$36.00 Ó 2011 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2011.02.038
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novel analyzing methods such as high-resolution transmission electron spectroscopy (TEM) [37,38], high-angle annular dark-field scanning transmission electron spectroscopy [39–41] and 3D-atom probe [42]. The LPSO phase was found to have a (0 0 0 1) basal plane which is the same as that in Mg (2H, Ramsdell notation), but its stacking periodicity was lengthened 18-fold (18R) or 14-fold (14H) along the c-axis. The 18R- and 14H-LPSO phases have chemical modulation, in which solute elements are enriched in four atomic layers on the closely packed plane at sixand seven-period intervals, respectively [41]. Plastic deformation behavior of Mg12YZn with 18R-LPSO structure was also investigated using directionally solidified crystals. The (0 0 0 1)h1 1 2 0i basal slip was found to be a dominant operative deformation mode in the 18R-LPSO phase, whose critical resolved shear stress (CRSS) was estimated to be 10–30 MPa [43]. The LPSO phase is expected to act as an alloy-strengthening factor in Mg alloys. Besides the unique microstructure of the LPSO phase, LPSO phasecontaining Mg alloys present another subject of inquiry concerning the effects of plastic working on their mechanical properties. Many previous studies found that the conventional thermo-mechanical treatments such as extrusion and rolling caused a marked enhancement in strength without impairing the ductility that is necessary for practical plastic working, though the as-cast alloy showed low strength [32–35,44–46]. Several factors such as a-Mg matrix grain refinement, the formation of suitable texture and the introduction of kink-deformation bands in the LPSO phase [47–50] are presumed to come into play in the enhancement of mechanical properties of the alloy during extrusion, but the strengthening mechanisms of the extruded alloys are yet to be unraveled in detail. Meanwhile, feasibility studies have been under way for the LPSO phase-containing Mg–Zn–RE alloys because LPSO phase-containing Mg alloys have been attracting attention as promising wrought alloys, and mass-production and scaling-up techniques for Mg–Zn–RE alloys are urgently sought in industrial circles. Scaling-up ingot casting often causes grain coarsening owing to slower cooling during solidification. The influence of casting-related grain coarsening and dendrite growth on the microstructure evolution of extruded Mg–Zn–RE alloys also remains to be clarified. These considerations show that it is important to clarify the relationship between microstructure evolution and the mechanical properties of extruded Mg–Zn–Y alloys with an LPSO phase. This study focused on the secondary dendrite arm spacing (SDAS) as a factor playing an important role in the fabrication of extruded alloys from the standpoint of both fundamental research and the feasibility of scaling-up and mass production. The morphology of the LPSO phase in cast ingots can be controlled by changes in the SDAS. Several Mg–Zn–Y alloys with varying SDAS were cast at various cooling rates ranging from 0.01 to 10 K s1, and in this study the influence of initial microstructure to the evolution of microstructure during
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extrusion and the resulting mechanical properties were examined. The ductility enhancement and strengthening mechanisms are discussed in relation to the formation of a multimodal microstructure in extruded Mg–Zn–Y alloys. 2. Experimental procedure Mg97Zn1Y2 (at.%) cast ingot was produced by highfrequency induction melting of pure Mg (99.99 wt.%), Zn (99.9 wt.%) and Y (99.9 wt.%) metals in a carbon crucible. The molten alloys kept at 1023 K were cast in a steel mold with a hole 30 mm in diameter and 100 mm in length in an argon atmosphere, or solidified in a carbon crucible with a hole 70 mm in diameter and 80 mm in length at various cooling rates. Four cooling rates of 0.06 K s1, 0.18 K s1, 5.7 K s1, and 9.6 K s1 were derived from temperature-controlled solidification. In gravity casting, the initial temperatures of the steel mold were kept at 360 K and 673 K by electric heating during solidification. The wall thickness of the mold used in this study was 20 mm. In crucible cooling, the molten alloys were heated to 1023 K and then cooled at two cooling rates, controlled by changing induction heating input. The cooling history of the alloy during the solidification was monitored with thermocouples installed at three different positions (bottom, middle and top of the solidified alloys) in the mold and the crucible. Cooling rates, Rc (K s1), in each solidification were estimated by the following equation: Rc ¼
ðT m T s Þ t
ð1Þ
where Tm is the liquidus temperature measured at middle position of solidified alloy, Ts is the solidus temperature, and t is the solidification time in seconds. These ingots were machined into round bars 29 mm in diameter and 70 mm in length. The round billets were extruded at an extrusion ratio of 10, extrusion temperature 623 K, and ram speed 2.5 mm s1. In this study, axisymmetric extrusion was adopted using a round die; long rod-like alloys with diameter 9 mm and length 600 mm were extruded from the ascast ingots. The constituent phases in the alloy were examined by X-ray diffractometry (XRD; Bruker AXS D8 Discover with Cu Ka radiation). XRD data were collected in the 2h diffraction angle range of 20–90°. The step width and counting time in each step were 0.02° and 1 s, respectively. The microstructure of the specimens was observed by confocally optical microscopy (OM; Lasertec C-130), scanning electron microscopy (SEM; JEOL JSM-7001F) and transmission electron microscopy (TEM; JEOL JEM-2000FX). The specimens were polished mechanically and then chemically with a mixture of 80 ml ethyl alcohol and 20 ml nitric acid. Finally, they were etched with a mixture of 4.2 g picric acid, 10 ml acetic acid, 80 ml ethyl alcohol and 10 ml water for OM observation of the microstructure. The grain size in the alloys was measured by the liner intercept method in OM observation. SEM images were taken in the secondary
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electron mode and the back-scattered electron (BSE) mode. Specimens were polished by paper lapping and ion-milling (JEOL Cross Section Polisher SM-09010) for SEM observation. The volume fraction of the LPSO phase was computed from binary images converted from gray-scale SEM–BSE images using the NIH image software package. TEM specimens were thinned by mechanical polishing using ion-milling with a Fischione Model 1010. The textures of the extruded alloys were analyzed by electron back-scatter diffraction pattern analysis (EBSD) with orientation imaging microscopy (OIM; TSL Solutions K.K.) and by pole figure analysis (PF) performed with an XRD with a collimator of diameter 1 mm. PF data were collected within specimen tilt angle (v) range 0–86° and specimen rotation angle (/) range 0–360°. Both angles v and / were increased stepwise in increments of 2°, and the counting time in each step was 1 s for angle /. Tensile tests were carried out using an Instron testing machine (Instron Model 5584) at room temperature with an initial strain rate of 5 104 s1. The gauge section of tensile specimens was 2.5 mm and 15 mm in diameter and length, respectively. In the ambient-temperature testing, strain was measured with an electrical strain gauge. The tensile axis was held parallel to the direction of extrusion. The 0.2% proof strength was employed as yield strength.
3. Results 3.1. Microstructure of cast ingot Fig. 1 shows optical micrographs of Mg97Zn1Y2 alloys that were solidified at various cooling rates. Mg97Zn1Y2 alloy consisted of an a-Mg matrix and a LPSO phase, with dendritic primary crystals of a-Mg phase. Grain size decreased with increasing cooling rates. The morphology of dendritic primary crystals was also observed in the Mg97Zn1Y2 alloy, and SDAS decreased with increasing cooling rates. Grain size and SDAS in the as-cast state are shown in Table 1. The changes in the grain size dg and SDAS d2, in Mg97Zn1Y2 cast alloy as functions of cooling rate Rc are summarized in Fig. 2. A commonly accepted empirical relationship between SDAS and cooling rate is given by [51,52] d 2 ¼ A RBc
ð2Þ
where A and B are constants for a given alloy system. The values of the constants for Mg97Zn1Y2 alloy were determined by power approximation to be A = 37.1 and B = 0.33. The solid line in Fig. 2 represents power approximation of the experimental results. Fig. 3 shows the XRD pattern of the Mg97Zn1Y2 alloy solidified at
Fig. 1. Optical micrographs of Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1, and (d) 9.6 K s1.
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Table 1 Grain size and SDAS of the Mg97Zn1Y2 alloy solidified at various cooling rates.
crystals of the a-Mg phase, thus improving the dispersion of the LPSO phase in the cast state.
Cooling rate (K s1)
Grain size (mm)
SDAS (lm)
0.061 0.18 5.7 9.6
0.91 ± 0.25 0.86 ± 0.21 0.43 ± 0.19 0.41 ± 0.16
88 ± 8.2 68 ± 7.9 25 ± 3.3 15 ± 2.8
3.2. Microstructure and mechanical properties of extruded alloy
Fig. 2. Change in the grain size and SDAS of Mg97Zn1Y2 cast alloys as a function of cooling rate.
cooling rate 9.6 K s1. Peaks originating from the 18RLPSO phase and the a-Mg matrix phase were detected [37]. Fig. 4 shows SEM–BSE images of Mg97Zn1Y2 cast alloys at various solidification cooling rates. Bright and lamellar contrast corresponds to the LPSO phase; the average volume fraction of LPSO phase in Mg97Zn1Y2 was 24%, no matter how the cooling rate varied. Increasing the cooling rate led to refinement of the dendritic primary
Fig. 3. XRD patterns of Mg97Zn1Y2 alloys solidified at cooling rates of 9.6 K s1.
Optical micrographs of extruded Mg97Zn1Y2 alloys taken from a longitudinal section are shown in Fig. 5. The direction of extrusion is parallel to the horizontal direction in all the figures presented in this paper. Three feature areas can be seen: the dynamically recrystallized (DRX-ed) a-Mg fine-grain region; the hot-worked a-Mg coarse-grain region; and the 18R-LPSO phase region. Changes in the volume fractions of two types of matrices, the DRX-ed a-Mg region and the hot-worked a-Mg region, are shown as functions of SDAS in Fig. 6. Although the average grain size of the DRX-ed a-Mg fine-grain region in the extruded alloys was 2.5 lm regardless of different SDAS, the volume fraction of the DRX-ed a-Mg region increased with decreasing SDAS. This suggests that the SDAS is a major determinant of the volume fraction of the DRX-ed grains region in the extruded alloy as well as the LPSO phase morphology in the as-cast state. Fig. 7 shows SEM–BSE images of extruded Mg97Zn1Y2 alloys taken from the longitudinal section. SEM–BSE images revealed the changes in dispersion of the LPSO phase that were affected by a change in SDAS. Furthermore, many kink-deformation bands were seen to have been introduced in the LPSO phase, as reported previously [33,43,44,48,53]. To describe the change in the dispersion of LPSO quantitatively, the dispersion level DL (lm1) was defined for discussion in this study by the following expression [50]: P Nv DL ¼ P ð3Þ Lv where Lv (lm) is the length of a vertical line segment L drawn perpendicular to the extrusion direction in Fig. 7, and Nv is the number of LPSO phase P intersected by segment L. The total length of segments Lv is more than 5 mm. Fig. 8 shows the relationship between the dispersion of LPSO phase in the extruded Mg97Zn1Y2 alloys and the SDAS in the as-solidified state. Reduction of SDAS promoted dispersion and refinement of the fiber-shaped LPSO phase during extrusion. In order to clarify texture evolution during extrusion, PF and inverse pole figure (IPF) measurements were performed by XRD and EBSD, respectively. Measurements were made with two extruded Mg97Zn1Y2 alloys that were cast at cooling rates of 0.061 and 9.6 K s1. The former ascast ingot had SDAS of 88 lm and the latter 15 lm (Table 1). Hereinafter, extruded Mg97Zn1Y2 alloys that were solidified at cooling rates of 0.061 and 9.6 K s1 will be called 88-lm-extruded Mg97Zn1Y2 alloy and 15-lmextruded Mg97Zn1Y2 alloy, respectively. The {1 0 1 0} X-ray pole figures of the 88-lm-extruded Mg97Zn1Y2 alloy
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Fig. 4. SEM-BSE images of Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1 and (d) 9.6 K s1.
and the 15-lm-extruded Mg97Zn1Y2 alloy are shown in Fig. 9. The measured planes were transverse cross-section (perpendicular to the direction of extrusion) and longitudinal section of the extruded bars. The extruded alloys had a strong basal texture. The <1 0 1 0> axis of the hexagonal close packed (hcp) structure is parallel to the direction of extrusion. Fig. 10 presents the IPF map obtained from a longitudinal section of the 88-lm-extruded Mg97Zn1Y2 alloy. Colors designate the crystallographic orientation with respect to the sample normal, coded as shown in the color key triangle. In the IPF mapping, a confidence index (CI) was used to indicate the degree of confidence as to whether the orientation calculation is correct [54]. The shaded areas correspond to the regions where a CI value is <0.1, i.e. the OIM system cannot analyze the Kikuchi diffraction patterns. The large shaded areas most closely represent the 18RLPSO phase grains, which fact was confirmed by comparing them with the SEM images. IPF mapping also revealed that there were two kinds of matrices, that is, the DRX-ed fine-grain region (DRX-ed region) and the hot-worked coarse-grain region where recrystallization did not occur (worked region). It is noteworthy that dynamic recrystallization occurs in the a-Mg matrix next to the LPSO phase. This suggests that a strain sufficient to induce
recrystallization is concentrated in the neighborhood of the LPSO phase. Fig. 11 shows IPF maps of DRX-ed regions and worked regions in the 88-lm-extruded Mg97Zn1Y2 alloy shown in Fig. 10. In order to express clearly the relationship between extrusion direction and texture, Fig. 12 shows EBSD-derived pole figures taken from a wide-scan region (160 320 lm) of the 88-lmextruded Mg97Zn1Y2 alloy and a narrow-scan region (80 100 lm) selected from only the DRX-ed zone. Wide-scan EBSD-derived pole figures and X-ray pole figures also revealed that there was a strong basal texture in the extruded alloy. However, DRX-ed zone scanning of EBSD-derived pole figure revealed that the basal fiber texture in DRX-ed grains is much weaker, and the DRX-ed grains had almost randomized crystallographic orientation. Fig. 13 presents the IPF maps obtained from a longitudinal section of the 15-lm-extruded Mg97Zn1Y2 alloy. The area of the DRX-ed region expanded, and the area of the worked grains region decreased in comparison with the 88-lm-extruded Mg97Zn1Y2 alloy. Fig. 14 shows IPF maps of DRX-ed regions and worked regions in the 15-lmextruded Mg97Zn1Y2 alloy. DRX-ed regions were identified from grain size distribution. The DRX-ed grains ranged in size from less than a micron to 4 lm, the size peaking at 1 lm. In this study, a grain measuring 4 lm
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Fig. 5. Optical micrographs of extruded Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1 and (d) 9.6 K s1. Micrographs are of longitudinal section.
Fig. 6. Change in the volume fractions of two types of matrices (the DRXed a-Mg region and the hot-worked a-Mg region) as a function of SDAS.
or less was considered to be a DRX-ed grain. Fig. 15 shows EBSD-derived pole figures from the entire EBSD scan
region (160 320 lm) of the 15-lm-extruded Mg97Zn1Y2 alloy. The 15-lm-extruded Mg97Zn1Y2 alloy had a weaker basal texture than the 88-lm-extruded Mg97Zn1Y2 alloy, because of its reduced region, which has a strong h1 0 1 0i texture. Because the OIM system was unable to analyze the Kikuchi diffraction patterns from the LPSO phase owing to the large amount of residual strain, TEM observation was performed to confirm the morphology and texture evolution of the LPSO phase. Fig. 16 shows a TEM image of the fiber-shaped LPSO phase embedded in the DRXed a-Mg grain region of the 15-lm-extruded Mg97Zn1Y2 alloy. Almost all the LPSO phases observed in this study were elongated along the extrusion direction and had a strong basal texture in which the c-axes of the grains were perpendicular to the direction of extrusion. The basal slip was strongly suppressed in the extruded alloy, most probably because the Schmid factor was very small when the tensile direction was parallel to the extrusion direction, so that very high yield stress arose accompanied by the formation of a deformation kink. Fig. 17 shows changes in the yield tensile strength, ultimate tensile strength and elongation of the extruded alloys
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Fig. 7. SEM–BSE images of extruded Mg97Zn1Y2 alloys solidified at cooling rates of (a) 0.06 K s1, (b) 0.18 K s1, (c) 5.7 K s1, and (d) 9.6 K s1. Micrographs are of longitudinal section.
4. Discussion
Fig. 8. Relationship between the extruded Mg97Zn1Y2 alloys and the SDAS with respect to the as-solidified state in LPSO phase dispersion.
as functions of SDAS in the as-cast state. The tensile axis was along the direction of extrusion. Tensile yield strength and elongation of the extruded alloys simultaneously increased with decreasing SDAS, although, in general, there is a trade-off between strength and elongation.
This study has shown that there are three typical regions in extruded Mg–Zn–Y alloys: the DRX-ed fine-grain region; the hot-worked a-Mg coarse-grain region; and the fiber-shaped LPSO phase region. Although the LPSO phase is presumed to help strengthen the alloys, it is necessary to clarify which region does what to help improve the strength and ductility of the extruded alloy. TEM observation revealed that the LPSO phase elongated along the direction of extrusion had a strong basal texture in which the c-axes of the grains were perpendicular to the direction of extrusion, as shown in Fig. 16. This basal fiber texture is very suitable for the strengthening of the alloy, since the operation of the basal slip, which is the dominant deformation mode also in the LPSO phase, is strongly hindered by the texture. Very recently, Mayama et al. [59] demonstrated by calculation that a large CRSS ratio of a basal slip to a non-basal slip system produced a strong basal texture owing to crystal rotation during extrusion, bringing the h1 0 1 0i axis of LPSO phase into a position parallel to the direction of extrusion. The LPSO phase as well as the hcp-Mg has a large CRSS ratio of the (0 0 0 1)h1 1 2 0i basal slip to the non-basal slip. Moreover, in the LPSO
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Fig. 9. {1 0 1 0} X-ray pole figures of the extruded Mg97Zn1Y2 alloy: (a) transverse section of 88-lm-extruded alloy; (b) longitudinal section of 88-lmextruded alloy; (c) transverse section of 15-lm-extruded alloy; (d) longitudinal section of 15-lm-extruded alloy.
phase, both the twin deformation and the dynamic recrystallization hardly occur during extrusion, though they easily occur in the hcp-Mg metal. Taken altogether, the LPSO phase can be considered capable of developing a strong basal texture owing to crystal rotation during extrusion. The difference in plastic deformation behavior between the LPSO and the a-Mg phase accounts for the formation of the fiber-shaped LPSO phase-reinforced metal matrix composite during extrusion. Therefore, the dispersion of the fiber-shaped LPSO phase can be considered to be of particular significance. As shown in Fig. 8, the degree of dispersion of the LPSO phase increased with decreasing SDAS in the as-cast state, though the volume fraction of LPSO phase did not change. Fig. 18 shows the effect of the LPSO phase dispersion on alloy strengthening. An increase in LPSO phase dispersion brought about an improvement in the tensile yield strength of the alloy in proportion. It should be noted that twinning was not found in the LPSO phase-containing Mg–Zn–Y extruded alloys, while some kink-deformation bands were observed. Matsuda et al. [60] reported that a densely developed LPSO phase prevented the growth of {1 0 1 2} deformation twins in the a-Mg matrix in rapidly solidified Mg97Zn1Y2 alloy. Very recently, Hagihara et al. [43,44,48] reported that the dominant operative deformation mode in
0i the 18R-LPSO phase was identified as a (0 0 0 1)h1 1 2 basal slip and that twin deformation never occurred. Based on the findings above, the present authors are inclined to presume that the highly dispersive fiber-shaped LPSO phase grains, along with the significantly refined a-Mg matrix grains, in the extruded alloys suppressed twin deformation. Furthermore, the formation of a kink-deformation band in the LPSO phase and a-Mg phase region may strengthen the alloy. The reason is that, since kink-deformation bands are macroscopically formed perpendicularly to the primary slip direction, they effectively counteract basal slipping [61]. The fiber-shaped LPSO phase with kink-deformation bands is one of the important factors that come into play in strengthening of the Mg–Zn–Y alloys. However, in consideration of the a-Mg matrix phase, the existence of two kinds of matrices seems to greatly influence the ductility of the alloy. Fig. 19 shows the relationship between the volume fractions of two kinds of matrices and the mechanical properties of extruded Mg97Zn1Y2 alloy. It is easy to understand that an increase in the volume fraction of fine grains is conductive to alloy strengthening through the Hall–Petch relation [55,56]. Furthermore, Fig. 19 suggests that a DRX-ed a-Mg finegrain region with almost randomized crystallographic
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Fig. 11. IPF maps and IPF of (a and b) DRX-ed regions and (c and d) worked regions in the 88-lm-extruded Mg97Zn1Y2 alloy.
Fig. 10. IPF map obtained from a longitudinal section of the 88-lmextruded Mg97Zn1Y2 alloy.
orientation helps to improve elongation in the extruded alloys. Average Schmid factors are available for prediction of the response of a textured polycrystalline material to a given stress state. Average Schmid factors for the (0 0 0 1)h1 1 2 0i slip of the 88-lm- and 15-lm-extruded Mg97Zn1Y2 alloys were estimated as 0.17 and 0.23, respectively. These figures suggest that the basal slip can be so readily operative in 15-lm-extruded Mg97Zn1Y2 alloys that it helps to enhance ductility. In addition, Koike et al. [57] proposed new deformation mechanisms in the fine-grained Mg alloys. They discussed the activation of non-basal slip systems attributed to compatibility stress at grain boundaries produced during large elongation. When grain size is so small that the grain-boundary
Fig. 12. EBSD-derived pole figures of the 88-lm-extruded Mg97Zn1Y2 alloy: (a) wide-scan region (160 320 lm); (b) narrow-scan region (80 100 lm) selected solely from the DRX-ed zone.
affected region can cover the entire grain, the ductility of the Mg alloy is improved. Therefore, the fine DRX-ed grains formed near an LPSO grain are very likely to do much to improve ductility. On close examination of the hot-worked matrix region, the worked a-Mg grains were found to have a strong basal texture, as shown in Figs. 11 and 14. When the tensile direction was parallel to the direction of extrusion, the Schmid factor for the basal slip in the hot-worked a-Mg grains
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Fig. 14. IPF maps and IPF of (a and b) DRX-ed regions and (c and d) worked regions in the 15-lm-extruded Mg97Zn1Y2 alloy.
Fig. 15. EBSD-derived pole figures from the entire EBSD scan region (160 320 lm) of the 15-lm-extruded Mg97Zn1Y2 alloy. Fig. 13. IPF maps obtained from a longitudinal section of the 15-lmextruded Mg97Zn1Y2 alloy.
turned out to be a small value. In addition, since most of the basal planes in the grains of extruded Mg alloys lie parallel to the direction of extrusion, {1 0 1 2} twinning is also inhibited in tensile deformation along the direction of extrusion. This phenomenon led to achievement of high yield strength [58]. The textured a-Mg grains embedded in fine-grained a-Mg matrix works as a kind of reinforced component. As SDAS diminishes, the alteration of the worked coarse grains into a fiber shape might enhance the effect of composite strengthening, considering the possibility that the increase in dispersion may compensate the decrease in their volume fraction to some extent. Fig. 20 shows the relationship between tensile yield strength and elongation of extruded Mg97Zn1Y2 alloys. Mg96Zn2Y2 extruded alloys containing LPSO phase [49],
Fig. 16. TEM image of the fiber-shaped LPSO phase embedded in DRXed a-Mg grain region of the 15-lm-extruded Mg97Zn1Y2 alloy.
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Fig. 17. Yield tensile strength, ultimate tensile strength and elongation of the extruded alloys as a function of SDAS in as-cast state.
Fig. 18. Change in the yield strength of the extruded alloy as a function of LPSO phase dispersion.
Mg–Zn–RE extruded alloys not containing an LPSO phase, and conventional extruded alloys (AZ31B, AZ61A, AZ91D, AM60, WE54) are included in this figure for purposes of comparison. The extrusion conditions (temperature, extrusion ratio and ram speed) were the same with all Mg–Zn–RE alloys. In general, an increase in strength due to solution hardening or precipitation hardening accompanies a decrease in the ductility of the materials. However, this trade-off between strength and ductility does not occur in the extruded Mg alloy with an LPSO phase under discussion. Recently, it has been reported that bimodal nanostructured metal materials with micrometer-sized grains embedded inside a matrix of nanocrystalline and ultra-fine grains exhibit high strength and ductility at once [62,63]. In these materials, micrometer-sized grains obtained through secondary recrystallization produced strain hardening sufficient to
Fig. 19. Relation between the volume fractions of two kinds of matrices and the mechanical properties of extruded Mg97Zn1Y2 alloy.
Fig. 20. Relation between the tensile yield strength of extruded Mg97Zn1Y2 and Mg96Zn2Y2 alloys and the elongation.
sustain a reasonably uniform deformation to the extent of large elongation [64,65]. In the extruded Mg–Zn–Y alloys under discussion, bimodal grain size and texture distribution may warrant both high strength and ductility of the alloys. The LPSO phase plays such an important role in multimodal microstructure formation during normal extrusion in this alloy system. The results suggest that the DRX-ed a-Mg fine grains do much to improve the ductility of the alloys. The micrometer-sized grains that have a random crystallographic orientation are extremely effective in enhancing ductility, because Mg metal shows strong plastic anisotropy derived from the hcp crystal structure [57]. The findings that fine grains of hcp-Mg enhance the ductility of the alloys and that coarse a-Mg and LPSO phase grains strengthen the alloys are of interest. Both the fine grains and the coarse grains act in a manner opposite
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to the bimodally microstructured face-centered cubic (fcc)Cu [62,63] or fcc-Al [64,65]. Heterogeneity of the microstructure in Mg alloys has often been discussed. Ion et al. [66] reported that the centers of old grains (cores) and equiaxed subgrains near the old grain boundaries (mantle region) were developed in an Mg–0.8 wt.%Al extruded alloy deformed 16% at 533 K. The former region had a strong basal texture and kink deformation banding, while the latter region had a random crystallographic orientation. They reported that in deformation <600 K, lattice rotation occurred in the regions around the grain boundaries by the basal slip accompanied by kinking due to the relative operating difficulty of the non-basal slip, while the cores of the grains remained relatively undeformed. This leads to a dynamic recrystallization involving the dynamic polygonization of rotated lattice (mantle) regions adjacent to grain boundaries, resulting in the formation of a bimodal microstructure. The existence of the LPSO phase is considered to accelerate the formation of refined recrystallized grains around them by a similar mechanism. In addition, Stanford et al. reported that micro-alloying additions of RE weakened the texture in extruded binary Mg-based alloys and that their extruded alloys showed good ductility [67– 69]. These results suggest that optimum plastic deformation and addition of optimum alloying element can produce a controlled heterogeneous microstructure in Mg alloys. In this study, SDAS in an a/LPSO two-phase Mg alloy was found to be an important factor for controlling the heterogeneous microstructure. A decrease in SDAS involves the dispersion of the LPSO phase in the as-cast state, resulting in the expansion of the DRX-ed region. In particular, basal-textured fiber-shaped LPSO phase grains embedded in a non-textured fine-grained a-Mg matrix can produce an improvement in both the strength and the ductility of the alloy simultaneously. 5. Conclusions Multimodally microstructured Mg–Zn–Y alloys were prepared by casting and extrusion processes. Mg–Zn–Y alloy ingots with various SDAS were extruded, and then tensile-tested at room temperature. The alloys exhibited high strength and reasonable ductility for practical purposes of use. OM and SEM observations revealed that the extruded alloy had three typical regions: a DRX-ed a-Mg fine-grain region; a hot-worked a-Mg coarse-grain region; and a fiber-shaped LPSO phase region. EBSD analysis characterized two matrices: the DRX-ed fine grains had random crystallographic orientation, and the hotworked coarse grains had a strong texture with basal planes parallel to the direction of extrusion. The former grain region seems to account for the ductility of the alloys and the latter seems to strengthen them. In addition, the LPSO phase acts as a fiber-like reinforcement in effectively increasing the strength of the alloy due to its strong basal fiber texture and containing deformation kink bands. As
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to alloy strengthening, one should not fail to note that there is another factor at work when the alloys are strengthened. That is an increase in LPSO phase dispersion. SDAS in the cast state had a considerable effect on bimodal a-Mg matrix microstructure evolution and LPSO phase dispersion in the extruded Mg–Zn–Y alloys. Acknowledgements This work is supported by the Kumamoto Prefecture CREATE Project from JST and the Grants-in-aid for Young Scientists (A) No. 2086050 and Scientific Research (A) No. 19206075 from MEXT, Japan. References [1] Froes FH, Eliezer D, Aghion E. JOM 1998;50(9):30. [2] Luo AA. Int Mater Rev 2004;49:13. [3] Yang Z, Li JP, Zhang JX, Lorimer GW, Robson J. Acta Metall Sin (Engl Lett) 2008;21:313. [4] Ma C, Liu M, Wu G, Ding WJ, Zhu YP. Mater Sci Eng A 2003;349:207. [5] Xu DK, Liu L, Zu YB, Han EH. Mater Sci Eng A 2006;420:322. [6] He SM, Peng LM, Zeng XQ, Ding WJ, Zhu YP. Mater Sci Eng A 2006;433:175. [7] Peng Q, Wang J, Wu Y, Wang L. Mater Sci Eng A 2006;433:133. [8] Mordike BL, Ebert T. Mater Sci Eng A 2001;302:37. [9] Pekguleryuz MO, Kaya AA. Adv Eng Mater 2003;5:886. [10] Suzuki M, Sato H, Maruyama K, Oikawa H. Mater Sci Eng A 1998;A252:248. [11] Anyanwu IA, Kamado S, Kojima Y. Mater Trans 2001;42:1212. [12] Suzuki M, Sato H, Maruyama K, Oikawa H. Mater Sci Eng A 2001;A319–321:751. [13] Wang JG, Hsiung LM, Nieh TG, Mabuchi M. Mater Sci Eng A 2001;A315:81. [14] Mordike BL. Mater Sci Eng A 2002;A324:103. [15] Maruyama K, Suzuki M, Sato H. Metall Mater Trans A 2002;33:875. [16] Suzuki M, Kimura T, Koike J, Maruyama K. Scripta Mater 2003;48:997. [17] Smola B, Stulikova I, Pelcova´ J, Mordike BL. J Alloy Compd 2004;378:196. [18] Nie JF, Gao X, Zhu SM. Scripta Mater 2005;53:1049. [19] Nie JF, Muddle BC. Scripta Mater 1999;40:1089. [20] Vostry P, Smola B, Stulikova I, von Buch F, Mordike BL. Phys Stat Sol (a) 1999;175:491. [21] Nie JF, Muddle BC. Acta Mater 2000;48:1691. [22] Anyanwu IA, Kamado S, Kojima Y. Mater Trans 2001;42:1206. [23] Smola B, Stulikova I, von Buch F, Mordike BL. Mater Sci Eng A 2002;A324:113. [24] Apps PJ, Karimzadeh H, King JF, Lorimer GW. Scripta Mater 2003;48:1023. [25] Honma T, Ohkubo T, Hono K, Kamado S. Mater Sci Eng A 2005;A395:301. [26] Gao X, He SM, Zeng XQ, Peng LM, Ding WJ, Nie JF. Mater Sci Eng A 2006;A431:322. [27] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2006;47:2109. [28] Nishijima M, Yubuta K, Hiraga K. Mater Trans 2007;48:84. [29] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2008;49:227. [30] Nishijima M, Hiraga K, Yamasaki M, Kawamura Y. Mater Trans 2009;50:1747. [31] Kawamura Y, Hayashi K, Inoue A, Masumoto T. Mater Trans 2001;42:1172.
3658
M. Yamasaki et al. / Acta Materialia 59 (2011) 3646–3658
[32] Yamasaki M, Anan T, Yoshimoto S, Kawamura Y. Scripta Mater 2005;53:799. [33] Yoshimoto S, Yamasaki M, Kawamura Y. Mater Trans 2006;47:959. [34] Kawamura Y, Kasahara T, Izumi S, Yamasaki M. Scripta Mater 2006;55:453. [35] Kawamura Y, Yamasaki M. Mater Trans 2007;48:2986. [36] Yamasaki M, Sasaki M, Nishijima M, Hiraga K, Kawamura Y. Acta Mater 2007;55:6798. [37] Luo ZP, Zhang SQ. J Mater Sci Lett 2000;19:813–5. [38] Itoi T, Seimiya T, Kawamura Y, Hirohashi M. Scripta Mater 2004;51:107. [39] Abe E, Kawamura Y, Hayashi K, Inoue A. Acta Mater 2002;50:3845. [40] Zhu YM, Morton AJ, Nie JF. Acta Mater 2010;58:2936. [41] Yokobayashi H, Kishida K, Inui H, Yamasaki M, Kawamura Y. MRS Symp Proc 2011;1295, doi:10.1557/opl.2011.36. [42] Ping DH, Hono K, Kawamura Y, Inoue A. Philos Mag Lett 2002;82:543. [43] Hagihara K, Yokotani N, Umakoshi Y. Intermetallics 2010;18:267. [44] Hagihara K, Kinoshita A, Sugino Y, Yamasaki M, Kawamura Y, Yasuda HY, et al. Intermetallics 2010;18:1079. [45] Itoi T, Takahashi K, Moriyama H, Hirohashi M. Scripta Mater 2008;59:1155. [46] Itoi T, Inazawa T, Kuroda Y, Yamasaki M, Kawamura Y, Hirohashi M. Mater Lett 2010;64:2277. [47] Morikawa T, Kaneko K, Higashida K, Kinoshita D, Takenaka M, Kawamura Y. Mater Trans 2008;49:1294. [48] Hagihara K, Kinoshita A, Sugino Y, Yamasaki M, Kawamura Y, Yasuda HY, et al. Acta Mater 2010;58:6282. [49] Yamasaki M, Hashimoto K, Hagihara K, Kawamura Y. Mater Sci Forum 2010;654–656:615. [50] Hirano M, Yamasaki M, Hagihara K, Higashida K, Kawamura Y. Mater Trans 2010;51:1640.
[51] Mehrabian R. In: Mehrabian R, Kear BH, Cohen M, editors. Rapid solidification processing: principles and technologies, proceedings of the international conference on rapid solidification processing. Baton Rouge, LA: Claitor’s Pub. Division; 1978. p. 9–27. [52] Khedkar P, Nash P. J Mater Sci Lett 1986;5:1273. [53] Shao XH, Yang ZQ, Ma XL. Acta Mater 2010;58:4760. [54] Wright SI. Fundamentals of automated EBSD. In: Schwartz AJ, Kumar M, Adams BL, editors. Electron backscatter diffraction in materials science. New York: Kluwer Academic/Plenum Publishers; 2000. p. 58. [55] Armstrong RW, Codd I, Douthwaite RM, Petch NJ. Philos Mag 1962;7:45. [56] Sambasiva Rao G, Prasad YVRK. Metal Trans 1982;13A:2219. [57] Koike J, Kobayashi T, Mukai T, Watanabe H, Suzuki M, Maruyama K, et al. Acta Mater 2003;51:2055. [58] Kleiner S, Uggowitzer PJ. Mater Sci Eng A 2004;A379:258. [59] Mayama T, Noda M, Chiba R, Kuroda M. Int J Plasticity, in press, doi:10.1016/j.ijplas.2011.02.007. [60] Matsuda M, Ii S, Kawamura Y, Ikuhara Y, Nishida M. Mater Sci Eng A 2004;A386:447. [61] Higashida K, Takamura J, Narita N. Mater Sci Eng 1986;81:239. [62] Wang Y, Chen M, Zhou F, Ma E. Nature 2002;419:912. [63] Wang Y, Ma E. Acta Mater 2004;52:1699. [64] Witkin D, Lee Z, Rodriguez R, Nutt S, Lavernia E. Scripta Mater 2003;49:297. [65] Han BQ, Lee Z, Witkin D, Nutt S, Lavernia EJ. Metal Mater Trans A 2005;36A:957. [66] Ion SE, Humphreys FJ, White SH. Acta Metall 1982;30:1909. [67] Cottam R, Robson J, Lorimer G, Davis B. Mater Sci Eng A 2008;A485:375. [68] Stanford N, Barnett MR. Mater Sci Eng A 2008;A496:399. [69] Stanford N. Mater Sci Eng A 2010;A527:2669.