Extruded Mg–Zn–Ca–Mn alloys with low yield anisotropy

Materials Science & Engineering A 558 (2012) 356–365

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Extruded Mg–Zn–Ca–Mn alloys with low yield anisotropy S.W. Xu n, K. Oh-ishi, H. Sunohara, S. Kamado Department of Mechanical Engineering, Nagaoka University of Technology, Nagaoka 940-2188, Japan

a r t i c l e i n f o

a b s t r a c t

Article history: Received 4 March 2012 Received in revised form 28 July 2012 Accepted 5 August 2012 Available online 10 August 2012

Rare-earth-free Mg–5.99Zn–1.76Ca–0.35Mn (wt%) magnesium (Mg) alloys with low yield anisotropy and moderate performances were developed by utilizing hot extrusion. The sample extruded at 350 1C exhibited a high compression/tension yield ratio close to 1 along the extrusion direction due to its high volume fraction of the dynamically recrystallized (DRXed) grains (approximately 98%), its fine DRXed grains (average of 2.7 mm) and its extremely weak basal texture (nearly the same as that of the as-homogenized sample). The sample extruded at 300 1C also exhibited a compression/tension yield ratio of 0.89 with a tensile elongation to failure of 16% and a tensile 0.2% proof stress of 289 MPa. The high compression/tension yield ratio and the medium performances of the sample extruded at 300 1C are attributable to its refined microstructure consisting of the fine DRXed grains (average of 1.4 mm) and the elongated unrecrystallized regions of approximately 12 mm in width with high number density of fine Mg6Zn3Ca2 and a-Mn precipitates. During compression test, nearly no {10–12} extension twinning was observed in the fine DRXed grains region and the formation of {10–12} extension twinning was also effectively restrained in the unrecrystallized regions due to their refined size in width, their high density of precipitates and their relatively weaker basal texture than other basaltextured Mg alloys. Crown Copyright & 2012 Published by Elsevier B.V. All rights reserved.

Keywords: Magnesium alloy Extrusion Microstructures Yield anisotropy Mechanical properties

1. Introduction In the last decade, the growing worldwide emphasis on increasing fuel efficiency and reducing emissions provides a driving force for developing the lightweight wrought magnesium (Mg) alloys because of their potential applications for weight reduction in transportation vehicles [1,2]. However, the low ductility and the high yield anisotropy limit their applications. For example, the tensile 0.2% proof stress (sTPS) of the as-extruded commercial Mg–3Al–1Zn–1Mn (AZ31B, wt%) along the extrusion direction (ED) at room temperature is 199 MPa, while its compressive 0.2% proof stress (sCPS) along the ED is only 96 MPa, which shows a serious anisotropy in compression/tension yield [1]. Similar yield anisotropy results have also been reported for other basal-textured Mg alloys, such as the extruded or rolled commercial AZ61A, AZ80A, ZK60A [1,3] and the extruded experimental Mg–Al–Ca–Mn [4], Mg–Zn–Ag–Ca [5,6], Mg–Mn–Ca [7], Mg–Mn–(Ce, Y, Nd) [8] and Mg–Zn–Y [9] alloys. The compression/ tension yield anisotropy is found to result from the differences in the active modes of deformation twinning and slip systems. The essential reason is the hexagonal close-packed (hcp) structured Mg with low symmetry of slip systems [10–13]. Recently, the as-extruded Mg alloys with improved yield anisotropy

n

Corresponding author. Tel.: þ81 258 47 9760; fax: þ81 258 47 9770. E-mail addresses: [email protected], [email protected] (S.W. Xu).

(sCPS/sTPS 40.85) have been developed in the ultrafine-grained AZ31 [11,14] and Mg–Zn–Y [9], the extruded Mg–Zn–Ho [15], Mg– Zn–Ag–Ca–Zr [5,6], Mg–Sn–Zn–Al [16] and Mg–Gd–Y–Zn–Zr [17] alloys by grain refinement, texture modification or precipitation hardening. These results are exciting; however these Mg alloys are either produced by severe plastic deformation or containing costly alloying elements such as rare earth (RE) elements, which will cause the products to become expensive [1,2]. To facilitate the wider application of wrought Mg alloys as structural materials in the automotive industry, low-cost rare-earth-free Mg alloys with the improved formability and the minimized yield strength anisotropy are much more desirable. Thus, in this study, we applied hot extrusion to a new developed Mg–Zn–Ca–Mn alloy, in which only the common alloying elements were used, to explore the possibility of achieving high formability and low yield anisotropy in the extruded Mg alloy bars.

2. Experimental procedures The alloy used in this study was prepared by a permanent mold (PM) cast in-house using master alloy of Mg–1.3 wt%Mn, and pure Mg, Zn and Ca. The castings were melted under the cover gas mixture of SF6 and CO2 in a steel crucible. After the alloy was molten it was poured into a heated mold (approximately 250–300 1C) producing a cylindrical ingot of 50 mm diameter and 250 mm length. The composition (wt%) of the alloy was Mg– 5.99Zn–1.76Ca–0.35Mn, which is denoted ZXM620. After casting,

0921-5093/$ - see front matter Crown Copyright & 2012 Published by Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.08.012

S.W. Xu et al. / Materials Science & Engineering A 558 (2012) 356–365

billets of 43 mm diameter and 37 mm height were machined from the ZXM620 ingot and homogenized at 385 1C for 24 h. Then these homogenized billets were extruded at 300 1C, 350 1C and a ram rate of 0.1 mm/s. Hereinafter, the ZXM620 alloys extruded at 300 1C and 350 1C are denoted ZXM620-300 and ZXM620-350, respectively. The reduction ratio used was 20, and this yielded extrudates of 9.6 mm diameter. The microstructures of the as-homogenized, as-extruded and the deformed as-extruded alloy samples were observed with an optical microscope (OM; Olympus BX51M) and a transmission electron microscope (TEM; JEOL JEM-2100F) equipped with an energy dispersive X-ray spectroscope (EDX) operating at 200 kV. The grain size, the volume fraction of the dynamically recrystallized (DRXed) grains (DRX ratio) and the grain orientation information were obtained by an EDAX-TSL electron backscattered diffraction (EBSD) system with an orientation imaging microscope (OIM; OIM5.2). To ensure statistical rigor, in excess of 5000 grains were examined for each condition. In this study, all of the micrographs and images of the as-extruded samples were taken from the longitudinal sections. Tensile specimens were machined from the cylindrical extrudates such that the tensile testing was carried out parallel to the ED. The tensile specimens had a gauge length of 20 mm and a diameter of 4 mm. Compressive specimens were section from the same extrudates as tensile specimens and had a length of 25 mm, a diameter of 10 mm and were compressed parallel to the ED. The tensile and compressive tests were performed using a Shimadzu Autograph AG-I (50 kN) at room temperature with an initial strain

357

rate of 1  10  3 s  1. The 0.2% proof strength was employed as the yield strength. In order to examine the microstructural evolution during compression test, the compression test was interrupted after the stress–strain curve showed a yielding.

3. Results 3.1. Microstructure of the as-homogenized ZXM620 After homogenization treatment at 385 1C for 24 h, the interdentritic regions have been significantly reduced, but the as-homogenized ZXM620 still shows a microstructure consisting of dendritic a-Mg grains with the intermetallic phase, and examples are given in Fig. 1a and b. The intermetallic phase was characterized by the selected area electron diffraction (SAED) analysis; the results are shown in Fig. 1c and d. These experimentally observed diffraction patterns are consistent with the schematic and experimental ones for the Mg6Zn3Ca2 phase reported by Oh-ishi et al. [18]; thus based on the phase equilibrium [19,20], the intermetallic phase is identified to be the Mg6Zn3Ca2 ternary phase. The average grain size of the as-homogenized ZXM620 is approximately 130 mm with a random initial grain orientation (the maximal intensity is approximately 3.2, EBSD results not shown). TEM analysis also reveals the existence of precipitates with varying morphologies distributed throughout the a-Mg matrix of the as-homogenized ZXM620 (Fig. 2a and b). The size (length) of the precipitates is found to vary between 20 and 70 nm. EDX

Mg6Zn3Ca2

10µm

100µm

B = [10 10]

B = [0001]

-1.-1.0 1.-2.0 0.0.0 2.-1.0

0.-1.0

0.-1.2

0.0.0

0.0.2

Fig. 1. (a,b) Optical micrographs of the as-homogenized ZXM620 at different magnifications. (c,d) Selected area electron diffraction patterns (SAED) of the intermetallic phase taken from the beam directions of [0001] and [101¯0], respectively.

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Zn

200nm

200nm Mn

Ca

a -Mn

B=[011]

4.0.0 2.0.0 0.-2.2

200nm

200nm

Fig. 2. (a) TEM bright filed (BF) image and (b) HADDF image of precipitates in the as-homogenized ZXM620, (c–f) their corresponding EDX mapping analysis showing that the precipitates mainly contain Mn, (d) SAED pattern of the Mn precipitate taken from the beam direction of [011].

ZXM620-300

1mm ZXM620-350

1mm

10µm

10µm

Fig. 3. (a,b) General optical micrographs of the as-extruded ZXM620 on the sections parallel to the ED, (c,d) enlarged optical micrographs. (a) and (c) ZXM620-300, (b) and (d) ZXM620-350. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

analysis conducted on the precipitates (Fig. 2b–e) indicates that they consist mainly of Mn. Because such Mn-containing precipitates were also observed even in the PM-cast sample (results not shown); thus they may precipitate during cooling after solidification and remained after homogenization [21]. Further SAED analysis conducted on the precipitate (Fig. 2f) indicates that these precipitates are a-Mn, which is consistent with the prediction by the phase diagram at this composition [22] and the results reported by Grobner et al. [21] and Celikin [23]. Furthermore, due to the peristaltic nature of Mg rich portion of the Mg–Mn binary system (Mg–0.35 wt%Mn in this study), intragranular areas have more dissolved Mn than the grain boundary regions [21–23]. Consequently, a-Mn precipitation is mainly from the a-Mg matrix in the grain interiors during solidification as the alloy cools through the solvus.

3.2. Microstructure of the as-extruded ZXM620 Fig. 3 shows the optical micrographs of the as-homogenized ZXM620 after hot extrusion at 300 1C and 350 1C. Their corresponding TEM and EBSD analysis results are shown in Figs. 4–6. The ZXM620-300 does not show a fully recrystallized microstructure, but shows a typical extrusion microstructure consisting of the DRXed a-Mg fine grain region, the elongated strips of unrecrystallized (unDRXed) a-Mg region and the fragmented Mg6Zn3Ca2 intermetallic phase region (indicated by blue, red and yellow arrows in Fig. 3a and c, respectively). This microstructure was further examined using TEM, and a typical example is shown in Fig. 4a. The fragmented Mg6Zn3Ca2 intermetallic phase with a diameter varying between 0.1 and 0.5 mm are observed pinning at the DRXed grain boundaries. Furthermore,

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ZXM620-300

359

ZXM620-350 DRXed grain

UnDRXed region

DRXed grain

Mg6Zn3Ca2

0.5 µm ZXM620-300

ZXM620-300

Mg

3

4

2 5

1

200nm

200nm ZXM620-300

Zn

Ca

200nm

200nm ZXM620-300

ZXM620-300

Mn

Mg6Zn3Ca2

B = [1213]

0.-2.2

2.0.0 0.0.0

200nm Fig. 4. TEM analysis results obtained from the longitudinal sections of the as-extruded ZXM620. (a,b) TEM BF images, (c) HAADF image, (d–g) corresponding EDX elemental mappings of the region shown in (c), (h) SAED pattern of the Zn–Ca containing spherical precipitate marked as 1 shown in (c). (a, c–h) ZXM620-300, (b) ZXM620-350.

extremely fine phases are also present. These phases are composed of small particles of 10–100 nm diameters, and their number density is higher in the unDRXed region. EDX analysis

was carried out on these phases; the results are shown in Fig. 4c–g. Two types of particles are characterized: the spherical ones (indicated by no. 1 and 2 in Fig. 4c) are found to be enriched in

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b

a

50µ µm

c

50µm

d

DRXed region

ED

Max: 3.4

ED

50µm

UnDRXed region

Max: 12.2

Fig. 5. EBSD analysis results of the ZXM620-300. (a) IPF map, (b) (0001)/112¯0S basal slip Schmid factor distribution map when the tensile force is along the ED, (c) (101¯2)/101¯1¯S extension twinning Schmid factor distribution map when the compression force is along the ED, (d) (0001) and (101¯0) pole figures of the DRXed region and the unDRXed region.

Zn and Ca (Fig. 4e and f); the rectangular ones (indicated by no. 3–5 in Fig. 4c) are found to be enriched mainly in Mn (Fig. 4g), as observed in the as-homogenized condition (Fig. 2). The formation of Zn–Ca containing spherical precipitates suggests that dynamic precipitation occurred during hot extrusion. These spherical precipitates are identified to be the Mg6Zn3Ca2 phase by the SAED analysis (Fig. 4h). The formation of Mg6Zn3Ca2 precipitates was previously reported in the aged Mg–0.5 wt%Ca–4.2 wt%Zn [18] and Mg–1.6 wt%Ca–3.2 wt%Zn [24] alloys. Because the composition of the ZXM620 alloy is located near the two-phase equilibrium consisting of a-Mg and Mg6Zn3Ca2 [18,19]; thus our result is consistent with these reports. The grain orientation information of the ZXM620-300 was also examined using EBSD; the results are shown in Fig. 5. The unDRXed regions (indicated by arrows in Fig. 5a–c) have the stronger basal texture (Fig. 5a and d), the lower Schmid factor for the (0001)/112¯0S basal slip when the tensile force is along the ED (Fig. 5b) and the higher Schmid factor for the (101¯2)/101¯1¯S extension twin when the compression force is along the ED (Fig. 5c) compared to the DRXed grains region. Namely, theoretically, when tension along the ED, the ZXM620-300 shows a bimodal grain structure consisting of the ‘‘high strength’’ unDRXed regions and the relatively ‘‘ductile’’ fine DRXed grains region; on the contrary, when compression along the ED, the ZXM620-300 shows a bimodal grain structure consisting of the ‘‘easily twinned’’ unDRXed regions and the relatively ‘‘hardly twinned’’ fine DRXed grains region.

Increasing the extrusion temperature from 300 1C to 350 1C causes that (i) the unDRXed regions are minimally visible in the ZXM620-350 (compare Fig. 3b with a, and Fig. 3d with c); as a result, the DRX ratio increases from 70% 75 (for the ZXM620-300) to 98% 72 (for the ZXM620-350); and (ii) the number density of the fine precipitates decreases (compare Fig. 4b with a). Because the unDRXed regions exhibit the microstructural features as introduced above, the dramatic decreasing of the unDRXed region in the ZXM620-350 leads to that (i) the maximal basal texture intensity decreases from 11.2 (for the ZXM620-300, Figs. 6a) to 3.7 (for the ZXM620-350, Fig. 6b); (ii) the average Schmid factor for the basal slip (Fig. 6c) increases from 0.18 (for the ZXM620300) to 0.24 (for the ZXM620-350); and (iii) the average Schmid factor for the extension twin (Fig. 6d) decreases from 0.42 (for the ZXM620-300) to 0.38 (for the ZXM620-350). Furthermore, the average grain misorientation angles increases from 49 to 55 (Fig. 6e) and the average DRXed grain size also increases from 1.4 mm to 2.7 mm (Fig. 6f). 3.3. Mechanical properties of the as-extruded ZXM620 The as-extruded ZXM620 were tested in tension and compression to evaluate their yield point, ductility and tension/compression asymmetry. Fig. 7 shows the nominal stress–strain curves of the tension and compression tests. There is a progressive increase in the strength but no loss of the tensile ductility with the lowering of the extrusion temperature, as clearly evident in

S.W. Xu et al. / Materials Science & Engineering A 558 (2012) 356–365

ZXM620 -300

ED

ZXM620-300 ZXM620-350

Number Fraction

Number Fraction

Max: 3.7

0.4

0.18

0.12

0.06

0.0

0.1

0.2 0.3 0.4 Schmid Factor

0.2 0.1

0.5

0.08

0.04

30 60 90 Misorientation Angle (degree)

0.2

0.3 0.4 Schmid Factor

0.24

ZXM620-300 ZXM620-350 Area Fraction

Number Fraction

0.12

ZXM620-300 ZXM620-350

0.3

0.0

0.00

0.00

ZXM620 -350

ED

Max: 11.2

361

0.5

ZXM620-300 ZXM620-350

0.16

0.08

0.00

2

4

6

8

10

Diameter [microns]

Fig. 6. Comparison of the EBSD analysis results of the ZXM620-300 and ZXM620-350, showing the effect of extrusion temperature. (a), (b) (0001) and (101¯0) pole figures, (c) (0001)/112¯0S basal slip Schmid factor distribution histogram when the tensile fore is along the ED, (d) (101¯2)/101¯1¯S extension twin Schmid factor distribution histogram when the compression fore is along the ED, (e) grain misorientation angles distribution histogram, (f) DRXed grains distribution histogram.

600

Compression Stress (MPa)

500 400 300

P1

200

Tension

P2

ZXM620-300 ZXM620-350

100 0

plateau occur in the as-extruded ZXM620, as expected in wrought Mg alloys [1–17,23], the ZXM620-350 and the ZXM620-300 still show high sCPS of 213 MPa and 256 MPa, respectively. Thus the sCPS/sTPS ratios of the ZXM620-300 and the ZXM620-350 are 0.89 and 0.97, respectively.

0

4

8

4. Discussion 4.1. High sCPS/sTPS ratio

12

16

Strain (%) Fig. 7. Nominal stress–strain curves obtained from the as-extruded ZXM620 under tension and compression.

Fig. 7. The tensile sTPS of the ZXM620-350 and the ZXM620-300 are 219 MPa and 289 MPa, respectively, and the corresponding tensile elongations to failure (eTF) are 15% and 16%. There is not much strain hardening for the as-extruded ZXM620. The ultimate tensile strengths (sUTS) of the ZXM620-350 and the ZXM620-300 are 267 MPa and 310 MPa respectively. The ultimate strengths in compression also increase with a decreasing extrusion temperature, and ranging from approximately 464 MPa (for ZXM620-350) to 509 MPa (for ZXM620-300). The compressive elongations to failure are from 13% (for ZXM620-350) to 11% (for ZXM620-300). Although, a yield drop just after yielding and a subsequent

This study has shown that the sCPS/sTPS ratios of the asextruded ZXM620 are very high, approximately 0.89 and 0.97 for the ZXM620-300 and ZXM620-350, respectively. Fig. 8 compares the reported sCPS/sTPS ratios and the eTF of the as-extruded commercial and experimental Mg alloys with those of the asextruded ZXM620 in this study. The sCPS/sTPS ratios of the as-extruded ZXM620 are higher than those reported for the as-extruded commercial AZ [1,3] and ZK60A [1] Mg alloys, the as-extruded experimental Mg–Al–Ca–Mn [4], Mg–Zn–Ag–Ca [5,6], Mg–Zn–Ag–Ca–Zr [5,6], Mg–Mn–Ca [7] and Mg–Mn–(Y, Ce) [8] alloys. Recently, Sasaki et al. [16] reported a sCPS/sTPS ratio of 0.90 for the as-extruded Mg–Sn–Zn–Al alloy, Singh et al. [9] reported a sCPS/sTPS ratio of 0.91–0.94 for the as-extruded Mg– Zn–Y alloys and Homma et al. [17] reported a sCPS/sTPS ratio of 0.95 for the as-extruded Mg–Gd–Y–Zn–Zr alloy, which are higher than that of the ZXM620-300 but still lower than that of the ZXM620-350. Note that Sasaki et al. [16] achieved the low yield asymmetry in an alloy extruded at 250 1C, in a system where the

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Comp / Tension yield ratio

1.0

Mg-Gd-Y-Zn-Zr [17]

AZ80 [3] Mg-Sn-Zn-Al [16] Mg-Al-Ca-Mn [4]

ZXM2412-350

Mg-Zn-Y [9]

Mg-Zn-Y [9]

ZXM2412-300 Mg-Zn-Ag-Ca-Zr [5, 6] ZK60A [1]

AZ80A [1]

0.8

Mg-Mn-Y [8] Mg-Mn-Y [8]

Mg-Mn-Ca [7] Mg-Mn-Ce [8] Mg-Mn [8]

Mg-Mn-Y [8]

Mg-Zn-Ag-Ca [5, 6]

Mg-Mn-Ce [8]

0.6

AZ61A [1] Mg-Mn-Ca [7]

Mg-Mn [1]

0.4

0

5

10

AZ31B [1]

15

20

25

30

Tensile elongation to failure (%) Fig. 8. Relation between the sCPS/sTPS ratio and tensile elongation to failure of the as-extruded commercial and experimental Mg alloys.

a

grain size had been stabilized; Singh et al. [9] and Homma et al. [17] achieved the low yield asymmetry in an alloy containing a large amount of RE elements. In contrast, the high sCPS/sTPS ratios of the as-extruded ZXM620 are achieved after the extrusion at higher temperatures in an alloy without any RE elements. This is an exciting result and prompts us to clarify the reason by examining the microstructure development during compression test. Figs. 9 and 10 show the EBSD results and the TEM images obtained from the interrupted as-extruded ZXM620 samples just after yielding under the compression test (corresponding to the positions P1 and P2 indicated in Fig. 7 (compressive strain of 270.5%)). In both the interrupted ZXM620-300 and ZXM620-350, the {101¯2} extension twins are formed (indicated by arrows in Fig. 9a–d), leading to the rotation of some (0001) poles to the ED (Fig. 9g and h) and a typical concave-down shape in the compressive stress–strain curves (Fig. 7) [10–13,25,26]. However, the twinning volume fractions are only approximately 18% in the interrupted ZXM620-300 (mainly in the unDRXed regions, Fig. 9a

b

ZXM620-300

ZXM620-350

Twin 10µ µm

10µm

c

d

10µm

10µm

e

f

10µm

10µm

g

h ED

ED

Max: 2.8 Max: 12.8 Fig. 9. Microstructures of the interrupted as-extruded ZXM620 samples just after yielding under the compression test. (a,b) optical micrographs, (c,d) IPF maps, (e,f) boundaries distribution maps for the regions shown in (c) and (d). (g,h) (0001) and (101¯0) pole figures. (a, c, e and g) ZXM620-300, (b, d, f and h) ZXM620-350.

S.W. Xu et al. / Materials Science & Engineering A 558 (2012) 356–365

a

363

b

Twin Matrix

Twin boundary

0.5 µm

c

Matrix 100 nm

Dislocation

100 nm Fig. 10. TEM analysis of the twinning area in the interrupted ZXM620-300 shown in Fig. 9a, exhibiting the effect of the precipitates when the twinning is formed. (a) TEM image, (b) HADDF image and (c) TEM image for the untwined matrix near twin.

and c) and approximately 13% in the interrupted ZXM620-350 (mainly in the coarse DRXed grains with preferential orientations, Fig. 9b and d), which are much lower than that reported for the AZ31 alloy (approximately  30% under a compressive strain of 2% [25]). This result (shown in Fig. 9) indicates that for the present as-extruded ZXM620 the yield in compression is also ascribed to the activation of {101¯2} extension twinning, as observed in other as-extruded Mg alloys [25–27]; while the volume fraction of twinning is much lower, namely, the formation of extension twin was restrained in the present as-extruded ZXM620. Several factors are considered to result in this restraint: (i) Fine grain size (microstructure). Wang et al. [12] reported that the number fraction of the twinned grains and the compression/tension yield asymmetry of the basal-textured AZ31 depended strongly on the grain size. When the average grain size is smaller than 10 mm, the number fraction of twinned grains dramatically decreases to below 5%; when the average grain size is smaller than 2 mm, the sCPS/sTPS ratio becomes 1. Singh et al. [9] also reported that a reversal in the yield asymmetry was observed at about 2 mm grain size for the extruded Mg–Zn–Y alloy. Meyers et al. [28] further investigated the essential reason for the improved sCPS/sTPS ratio with decreasing grain size and found that the twinning stress increased more rapidly with decreasing grain size than the stress required for slip activation [15]. In this study, the average width of the elongated unDRXed regions and the average DRXed grain size are approximately 12 mm and 1.4 mm for the ZXM620-300, respectively. Thus, for the ZXM620-300, nearly no extension twinning is observed in the fine DRXed grains regions (Fig. 9c and d). The refined unDRXed regions, in which the {101¯2} extension twins are usually preferentially activated (Fig. 9a and c), also make an important contribution to the restraint of extension twinning, leading to the high sCPS/sTPS ratio. On the other hand,

for the ZXM620-350, the almost fully recrystallized microstructure with an average DRXed grain size of 2.7 mm would further restrain the formation of extension twinning (Fig. 6d), leading to the higher sCPS/sTPS ratio close to 1. (ii) Weak basal texture. As we introduced above, the yield anisotropy in the extruded Mg alloy can be ascribed to the activation of {101¯2} extension twinning in compression along ED but not in tension along ED, something which arises from the combination of sharp basal textures and the polarity of deformation twinning. Therefore, weakening of the basal texture and rotating grains away from the ideal alignment serve for extension twinning may serve to increase the stress required to initiate twinning (i.e., it increases the compression yield point) and thus contributes to a decrease in the yield anisotropy. Singh et al. [15] has reported that the sCPS becomes higher than the sTPS (yield asymmetry ratio 41) when the grain size of the weak-basaltextured Mg–Zn–Ho alloy is below 2 mm. In this study, the ZXM620-350 has an extremely weak texture, the maximal intensity of which is only 3.7, nearly the same as that of the as-homogenized ZXM620 (maximal intensity is 3.2) and are comparable with those of usual as-cast Mg alloys [7,15,29]. For the ZXM620-300, the fine DRXed grains region also shows an extremely weak texture (Fig. 5d) as the ZXM620350, this is another important reason for the minimally visible extension twinning in this region. Furthermore, because the {101¯2} extension twins are preferentially activated in the unDRXed regions, thus we focus on the basal texture intensity of these regions. As shown in Fig. 5d, the maximal intensity of the unDRXed regions is 12.2, which is approximately 4 times stronger than the as-homogenized ZXM620. Although this maximal intensity is not weaker than other as-extruded Mg–Ca–(Al,Zn) system alloys [7,15,30], it is much weaker than those of other basal-textured Mg alloys without Ca addition [25,31]. Recently, Standford [7] and

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Zhang et al. [30] have clearly proved that the addition of Ca to Mg–Mn or Mg–Zn alloys weakened the extrusion texture in a similar manner to RE, due to its large atomic radius. Therefore, except for the fine DRXed microstructure, the weak basal texture in the as-extruded ZXM620 can be also attributable to their high sCPS/sTPS ratios. (iii) Fine precipitates. This study has also shown that there is a large number of fine a-Mn and Mg6Zn3Ca2 precipitates in the matrix of the as-extruded ZXM620, and their number density is higher in the unDRXed region of ZXM620-300. Fig. 10 shows the effect of precipitates when the twinning is formed during compression test. The precipitates in the unDRXed regions of the ZXM620-300 pin at the twin boundaries (indicated by circles in Fig. 10b) or stand in the way of dislocations in the untwined matrix (Fig. 10c). Sasaki et al. [16] and Homma et al. [17] reported that these dense precipitates or stacking faults could suppress the occurrence of twinning by preventing the motion of dislocations and increasing the back stress, which the shear stress has to overcome to activate the deformation twinning [17,32]. Thus, the dense fine precipitates observed in the as-extruded ZXM620 may also attribute to their high sCPS/sTPS ratios.

4.2. Moderate mechanical properties It is also interesting to note that the as-extruded ZXM620 show good mechanical properties. Table 1 compares the reported values of the as-extruded commercial Mg alloys such as AZ and ZK60A [1], experimental Mg–6.1Zn–0.4Ag–0.2Ca [5,6], Mg–6.1Zn–0.4Ag–0.2Ca–0.6Zr (wt%) alloys [5,6], Mg–1.0Zn– 0.2Ca [30], Mg–5.3Zn–0.6Ca–0.3Mn [33] and Mg–4.7–0.5Ca [34] (all in wt%) with those of the as-extruded ZXM620. The eTF of the as-extruded ZXM620 are higher than those of the as-extruded commercial ZK60A and AZ80A, and are comparable to those of the as-extruded AZ31B and AZ61A. However, the strengths reported for the as-extruded commercial AZ series and ZK60A alloys are significantly lower than that of the ZXM620-300. The mechanical properties of the ZXM620-300 are, in fact, superior to those of the medium-performance T6-treated heat-resistant 6061 aluminum alloy (Al–1.2Mg–0.8Si (at%), sTPS and eTF are 275 MPa and 12%, respectively [35]). Furthermore, the sTPS and eTF of the ZXM620300 are even comparable to those of the as-extruded RE-containing Mg–6.1Zn–0.4Ag–0.2Ca–0.6Zr (wt%) alloy [5,6], as shown in Table 1. However, the sCPS of the as-extruded Mg–6.1Zn–0.4Ag– 0.2Ca–0.6Zr (wt%) alloy is lower than that of the ZXM620-300.

Table 1 Comparison of the reported values of the as-extruded commercial and experimental Mg alloys. Alloys (wt%)

ZXM620-300 ZXM620-350 Mg–3Al–1Zn–1Mn (AZ31B) [1] Mg–6Al–1Zn–1Mn (AZ61A) [1] Mg–8Al–Zn (AZ80A) [1] Mg–6Zn–Zr (ZK60A) [1] Mg–6.1Zn–0.4Ag–0.2Ca [5,6] Mg–6.1Zn–0.4Ag–0.2Ca–0.6Zr [5,6] Mg–1.0Zn–0.2Ca [30] Mg–5.3Zn–0.6Ca [33] Mg–5.3Zn–0.6Ca–0.3Mn [33] Mg–4.7–0.5Ca [34]

rTPS

rCPS

eTF

(MPa)

(MPa)

(%)

Ext. T (oC)

289 219 199 226 247 261 153 289

256 213 96 130 211 226 103 246

16 15 15 16 11 12 25 17

300 350 – – – – – –

135 220 272 291

– – – –

37 21 19 16

350 300 300 250

Generally, the strength of an Mg alloy is determined by the combined contributions of the following microstructural aspects [1–2,4–6,12,13,36,37]: (i) solid solution strengthening; (ii) dispersion strengthening of the intermetallic phases; (iii) precipitation strengthening; (iv) grain refinement or grain boundary strengthening; and (v) texture strengthening. As introduced above, the ZXM620-300 exhibits a typical extruded microstructure that consists of the fine DRXed regions, the elongated unDRXed regions and the fragmented Mg6Zn3Ca2 intermetallic phase. Furthermore, the ZXM620-300 also exhibits the finer DRXed grain size, the relatively stronger basal texture and the higher number density of finer precipitates than the ZXM620-350 (Figs. 4 and 6). This basal texture orientation can hinder the basal slip, which is the dominant deformation mode in hcp-Mg alloys along the ED, leading to a low Schmid factor for basal slip (Fig. 6a). Furthermore, Oh-ishi et al. [18], Bamberger et al. [24], Nie et al. [38], Oh et al. [39] and Somekawa et al. [34] have clearly demonstrated that the Mg6Zn3Ca2 precipitates are responsible for the improved hardness and strengths of Mg–Zn–Ca alloys. Thus, the greater strengths of the ZXM620-300 than those of the ZXM620-350 can be attributable to its stronger basal texture or lower Schmid factor for basal slip, finer DRXed grain (through the Hall–Petch relation [37]) and finer precipitates with higher number density (through the precipitation strengthening [40]) [1,2]. Furthermore, dispersion strengthening of the fragmented Mg6Zn3Ca2 intermetallic phase, precipitation strengthening of the a-Mn and the Mg6Zn3Ca2 particles and the fine DRXed grains strengthening are also the important reasons for the ZXM620-300 to show the better mechanical performances than other as-extruded commercial and experimental Mg alloys introduced above. As shown in Table 1 that, for the Mg–Zn–Ca-(Mn) system, the sTPS decreases and the eTF increases as the addition of alloying element decreases, mainly due to the decreasing of dispersion strengthening and precipitation strengthening effects in the as-extruded alloy samples. On the contrary, the sTPS increases as the extrusion temperature decreases, due to the finer DRXed grains and stronger basal texture. Finally, because there is a trade-off between the strengthening effect of the unDRXed regions during tension O ED and their softening effect during compression O ED, and generally it is difficult or the cost is relatively high to produce an as-extruded Mg alloy bar with a high DRX ratio close to 100%; thus it is better to refine the unDRXed regions of the as-extruded sample and use the precipitation-hardenable Mg alloys to obtain a high strength [2] with a low yield anisotropy. Our recent result [41] shows that the unDRXed regions of the as-extruded ZXM620 alloy mainly result from the {101¯2} extension twinning region formed at the beginning of the hot extrusion. Similar to the DRX, twinning is also a process of nucleation and growth, which is considered as closely related to dislocation activities. Decreasing the extrusion temperature of the ZXM620 from 350 1C to 300 1C retards the DRX process in the twinned regions; however, the twinning nucleation is promoted. Meanwhile, the growth of extension twinning is retarded and the twinned regions become the refined unDRXed regions with relatively stronger basal texture and denser precipitates (microstructural reinforcing regions). Thus, a very long region of uniform strain with almost none work hardening in the tensile stress–strain curve and the greater strengths with no loss of elongation to failure are also obtained in the ZXM620-300. Because the extrusion parameters (such as the extrusion temperatures) exert a great influence on the DRX ratio, the DRXed grain size, the precipitates and the texture of the extruded Mg alloy, as introduced above; thus it is desirable to develop the low-cost rare-earth-free ZXM620 alloys with lower yield anisotropy and much higher strengths by combining the contributions of lower-temperatures extrusion and subsequent heat treatment, the detailed analysis of which is in progress.

S.W. Xu et al. / Materials Science & Engineering A 558 (2012) 356–365

5. Summary In summary, the medium-performance rare-earth-free Mg alloys with low yield anisotropy were developed by extruding a PM cast Mg–5.99Zn–1.76Ca–0.35Mn (wt%) ingot. The sample extruded at 350 1C exhibited a high compression/tension yield ratio close to 1 along the ED. The sample extruded at 300 1C also exhibited a high compression/tension yield ratio of 0.89 with a eTF of 16% and a sTPS of 289 MPa, which are in fact, superior to those of the medium-performance T6-treated heat-resistant 6061 aluminum alloy [35]. Thus, the as-extruded ZXM620 exhibit the lowcost but medium performances with high sCPS/sTPS, which can become the potential light-weight structural materials for the automotive industry. The high compression/tension yield ratio of the sample extruded at 350 1C is attributable to its high DRX ratio of 98%, its fine DRXed grains of 2.7 mm and its extremely weak texture. The high compression/tension yield ratio and the medium performances of the sample extruded at 300 oC are attributable to its fine DRXed grains of 1.4 mm, its refined unDRXed regions of approximately 12 mm in width with high number density of fine Mg6Zn3Ca2 and a-Mn precipitates, and the relatively weak basal texture due to the addition of Ca [7,30].

Acknowledgments This work was partially supported by Grant-in-Aid for Young Scientists (B) (24760594), Grant-in-Aid for Scientific Research (A) (22246094), Grant-in-Aid for Scientific Research (B) (21360348), and for Research Activity Start-up (22860028) from JSPS, Japan. References [1] S. Kamado, H. Ohara, Y. Kojiam, Advanced Manufacturing Technologies of Magnesium Alloys, CMCbooks, Japan, 2005. [2] K. Hono, C.L. Mendis, T.T. Sasaki, K. Oh-ishi, Scr. Mater. 63 (2010) 710. [3] M. Shahzad, L. Wanger, Mater. Sci. Eng. A 506 (2009) 141. [4] S.W. Xu, K. Oh-ishi, S. Kamado, S. Uchida, T. Homma, K. Hono, Scr. Mater. 65 (2011) 269. [5] K. Oh-ishi, C.L. Mendis, T. Homma, S. Kamado, T. Ohkubo, K. Hono, Acta Mater. 57 (2009) 5593.

365

[6] C.L. Mendis, K. Oh-ishi, Y. Kawamura, T. Honma, S. Kamado, K. Hono, Acta Mater. 57 (2009) 749. [7] N. Stanford, Mater. Sci. Eng. A 528 (2010) 314. [8] J. Bohlen, S. Yi, D. Letzig, K.U. Kainer, Mater. Sci. Eng. A 527 (2010) 7092. [9] A. Singh, Y. Osawa, H. Somekawa, T. Mukai, Scr. Mater. 64 (2011) 661. [10] M. Knezevic, A. Levinson, R. Harris, R.K. Mishra, R.D. Doherty, S.R. Kalidindi, Acta Mater. 58 (2010) 6230. [11] J.T. Wang, D.L. Yin, J.Q. Liu, J. Tao, Y.L. Su, X. Zhao, Scr. Mater. 59 (2008) 63. [12] J. Wang, R.G. Hoagland, J.P. Hirth, L. Capolungo, I.J. Beyerleinb, C.N. Tome, Scr. Mater. 61 (2009) 903. [13] S.-H. Choi, E.J. Shin, B.S. Seong, Acta Mater. 55 (2007) 4181. [14] Q. Yang, A.K. Ghosh, Acta Mater. 51 (2006) 5159. [15] A. Singh, H. Somekawa, T. Mukai., Scr. Mater. 56 (2007) 935. [16] T.T. Sasaki, K. Yamamoto, T. Honma, S. Kamado, K. Hono, Scr. Mater 59 (2008) 1111. [17] T. Homma, N. Kunito, S. Kamado, Scr. Mater. 61 (2009) 644. [18] K. Oh-ishi, R. Watanabe, C.L. Mendis, K. Hono, Mater. Sci. Eng. A 526 (2009) 177. [19] G. Effenberg, et al., in: G. Effenberg, et al., (Eds.), Ternary Alloys, Vol. 18, H–K–Mg to Mg–Zn–Zr, MSI, Stuttgart, 2001, p. 68. [20] G. Levi, S. Avraham, A. Zilberov, M. Bamberger, Acta Mater. 54 (2006) 523. [21] J. Grobner, D. Mirkovic, M. Ohno, R. Schmid-Fetzer, J. Phase Equilib. Diffus. 26 (2005) 234. [22] A.A. Nayeb-Hashemi, J.B. Clark, ASM Handbook Volume 3 Alloy Phase Diagrams, ASM International, 1992. [23] M. Celikina, A.A. Kaya, M. Pekguleryuz, Mater. Sci. Eng. A 534 (2012) 129. [24] M. Bamberger, G. Levi, J.B. Vander Sande, Metall. Mater. Trans. A 37A (2006) 481. [25] Y.C. Xin, M.Y. Wang, Z. Zeng, M.G. Nie, Q. Liu, Scr. Mater. 66 (2012) 25. [26] S.G. Hong, S.H. Park, C.S. Lee, Acta Mater. 58 (2010) 5873. [27] M.R. Barnett, Z. Keshavarz, A.G. Beer, D. Atwell, Acta Mater. 52 (2004) 5093. [28] M.A. Meyers, O. Vohringer, V.A. Lubarda, Acta Mater. 49 (2001) 4025. [29] S.W. Xu, K. Oh-ishi, S. Kamado, H. Takahashi, T. Homma, Mater. Sci. Eng. A 542 (2012) 71. [30] B. Zhang, Y. Wang, L. Geng, C. Lu, Mater. Sci. Eng. A 539 (2012) 56. [31] S.W. Xu, N. Matsumoto, S. Kamado, T. Honma, Y. Kojima, Mater. Sci. Eng. A 517 (2009) 354. [32] A. Jain, S.R. Agnew, Magnesium Technol. (2006) 219. [33] L.B. Tong, M.Y. Zheng, S.W. Xu, S. Kamado, Y.Z. Du, X.S. Hu, K. Wu, W.M. Gan, H.G. Brokmeier, G.J. Wang, X.Y. Lv, Mater. Sci. Eng. A 528 (2011) 3741. [34] H. Somekawa, T. Mukai, Mater. Sci. Eng. A 459 (2007) 366. [35] I.J. Polmear, Light Alloys: Metallurgy of Light Metals, third ed., Edward Arnold, London, 1995. [36] R.G. Sambasiva, Y.V.R.K. Prasa, Metal. Trans. 13A (1982) 2219. [37] B. Reppich, Mater. Sci. Technol. 6 (1993) 311. [38] J.F. Nie, B.C. Muddle, Scr. Mater. 37 (1997) 1475. [39] J.C. Oh, T. Ohkubo, T. Mukai, K. Hono, Scr. Mater. 53 (2005) 675–679. [40] J.F. Nie., Scr. Mater. 48 (2003) 1009. [41] S.W. Xu, K. Oh-ishi, S. Kamado, T. Homma, Scr. Mater. 65 (2011) 875.