Effects of heat treatments on the microstructures and mechanical properties of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy

Effects of heat treatments on the microstructures and mechanical properties of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy

Materials Science and Engineering A 486 (2008) 183–192 Effects of heat treatments on the microstructures and mechanical properties of Mg–3Nd–0.2Zn–0...

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Materials Science and Engineering A 486 (2008) 183–192

Effects of heat treatments on the microstructures and mechanical properties of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy Fu Penghuai ∗ , Peng Liming, Jiang Haiyan, Chang Jianwei, Zhai Chunquan National Engineering Research Center of Light Alloys Net Forming, Shanghai Jiaotong University, 200030 Shanghai, PR China Received 30 June 2007; received in revised form 27 August 2007; accepted 27 August 2007

Abstract Microstructure and mechanical properties of as-cast and different heat treated Mg–3Nd–0.2Zn–0.4Zr (wt.%) (NZ30K) alloys were investigated. The as-cast alloy was comprised of ␣ magnesium matrix and Mg12 Nd eutectic compounds. After solution treatment at 540 ◦ C for 6 h, the eutectic compounds dissolved into the matrix and small Zr-containing particles precipitated at grain interiors. Further aging at low temperatures led to plate-shaped metastable precipitates, which strengthened the alloy. Peak-aged at 200 ◦ C for 10–16 h, fine ␤ particles with DO19 structure was the dominant strengthening phase. The alloy had ultimate tensile strength (UTS) and elongation of 300–305 MPa and 11%, respectively. Aged at 250 ◦ C for 10 h, coarse ␤ particles with fcc structure was the dominant strengthening phase. The alloy showed UTS and elongation of 265 MPa and 20%, respectively. Yield strengths (YS) of these two aged conditions were in the same level, about 140 MPa. Precipitation strengthening was the largest contributor (about 60%) to the strength in these two aged conditions. The hardness of aged NZ30K alloy seemed to correspond to UTS not YS. © 2007 Elsevier B.V. All rights reserved. Keywords: Mg–Nd–Zn–Zr; Heat treatment; Microstructure; Mechanical properties

1. Introduction The extremely low density, high specific strength and stiffness of Mg make it attractive for engineering applications [1,2]. It has been demonstrated that rare earth metals (RE) are the most effective elements to improve the strength properties of magnesium alloys especially at elevated temperatures [3,4]. More recently, a lot of works have been focused on magnesium alloy containing heavy rare earth elements, such as Mg–Gd–Y–Zr [5–8], Mg–Y–Sm–Zr [9,10], Mg–Dy–Gd–Nd [11] and Mg–Gd–Nd–Zr [12] alloys. Though these alloys show high strength, high content of heavy rare earth elements enhances alloys’ cost, which would probably limit their engineering application. Rational use of neodynium (Nd) could make it possible to develop a high strength and low cost magnesium alloy. Nd is one of light rare earth elements, with maximum solubility in solid Mg of 3.6 wt.% at eutectic temperature 545 ◦ C. Mg–Nd binary alloys have already had significant strengthening effect [13,14]. Addition of 0.5 wt.% Zn to Mg–3wt.% Nd alloy would further increase its peak-aged hardness [15]. Despite these interesting findings, there is a lack of systematic investigation ∗

Corresponding author. E-mail address: [email protected] (F. Penghuai).

0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.08.064

of little Zn containing Mg–Nd alloy, both microstructure and mechanical properties. Early studies using transmission electron microscopy (TEM) reported that the precipitation process in binary Mg–Nd alloys involved formation of G.P. zones, ␤ , ␤ , and ␤ phases [16,17]. The G.P. zones have in fact a form of discs on {1 1¯ 0 0}␣ planes of ␣-Mg, and the ␤ phase has a DO19 structure (a = 0.64 nm, c = 0.52 nm), and ␤ phase has a face-centered cubic structure with a = 0.736 nm, or a hexagonal structure with a = 0.52 nm, c = 1.30 nm [18]. The equilibrium phase ␤ has a body centered tetragonal structure with a = 1.03 nm and c = 0.593 nm. Although the precipitation sequence in Mg–Nd alloy has been well characterized, the relationship between precipitations and mechanical properties are seldom mentioned. In present study, the effects of heat treatments on mechanical properties and microstructures of gravity cast Mg–3Nd–0.2Zn–0.4Zr (wt.%) (NZ30K) alloy have been investigated. The relationship between precipitations and mechanical properties is discussed. 2. Experimental An alloy of nominal composition Mg–3Nd–0.2Zn–0.4Zr (wt.%) was prepared by high purity Mg, Zn and Mg–25 wt.% Nd, Mg–30 wt.% Zr master alloys by melting in an electrical

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Fig. 1. Optical micrographs of Mg–3Nd–0.2Zn–0.4Zr (wt. %) alloy: (a) as-cast and (b) solution treated at 540 ◦ C for 6 h.

resistance furnace under protection of mixture gas of SF6 , CO2 and air and cast in permanent mould [19] at pouring temperature 740 ± 5 ◦ C and mould temperature of 200 ± 5 ◦ C. The actual chemical composition of the alloy was determined to be Mg–3.001Nd–0.2696Zn–0.3752Zr (wt.%) by an inductively coupled plasma analyzer (ICP). Specimens cut from the cast ingot were first solution treated at 540 ◦ C for 6 h and quenched into hot water at ∼70 ◦ C, then subsequently aged at 160, 200 ◦ C and 250 ◦ C in an oil-bath, respectively. Vickers hardness testing was taken using 5 kg load and holding time of 30 s. Tensile test samples were cut into rectangular tensile specimens with dimensions of 10 mm width, 2 mm thickness and 30 mm gauge length by an electric-sparking wire-cutting machine. Tensile testing was carried out on a Zwick/Roell-20 kN material test machine at a cross-head speed of 1 mm/min at room temperature. Microstructure was examined in an optical microscope (OM) and a JEOL-2010 TEM and operating at 200 kV. Phase analyses were carried out with X-ray diffractometer (XRD). Fracture surface was investigated in a scanning electron microscope (SEM).

3. Results 3.1. Microstructure Fig. 1 shows the microstructures of NZ30K alloy in ascast and solution-treated conditions. The microstructure of as-cast alloy (Fig. 1a) consists of ␣ magnesium matrix and eutectic compounds. By XRD analysis (Fig. 2a), the eutectic compounds are Mg12 Nd as indicated in Fig. 1a. The average grain size of as-cast alloy is about 50 ␮m. After solution treatment, the eutectic compounds have almost dissolved into the matrix as illustrated by XRD pattern (Fig. 2b). Detail investigation reveals small gathered precipitates at grain interior. They are ZrH2 , Zn2 Zr3 and some Zr-containing particles which are not identified now [20]. The average grain size of solution-treated alloy is also about 50 ␮m. Unlike Mg–Gd–Y–Zr [5] and Mg–Y–Sm–Zr [10] alloys, the grains nearly do not coarsen during solution treatment in NZ30K alloy.

Fig. 2. XRD pattern of investigated Mg–3Nd–0.2Zn–0.4Zr (wt.%): (a) as-cast state and (b) T4 state.

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Fig. 3. Hardness evolution as a function of aging time during isothermal aging at 160, 200, and 250 ◦ C.

3.2. Aging characteristics Fig. 3 shows hardness curves of NZ30K alloy isothermally aged at 160, 200, and 250 ◦ C. The peak hardness decreases with increasing aging temperature. Aged at 160 ◦ C, the alloy takes 600 h to get the peak hardness of 73 HV, which is too long to be acceptable for the industry application. While aged at 200 ◦ C, only about 10–16 h are needed to obtain the peak hardness of 73 HV. Alloy aged at 250 ◦ C exhibits the lowest peak hardness of 65 HV and shortest peak-aged time of only half an hour. Obviously, specimen aged at 200 ◦ C for 10–16 h has a higher hardness with shorter time and is taken as optimum aging process. Fig. 4 shows the microstructures of NZ30K alloy in different aged conditions. After peak-aged at 200 ◦ C (Fig. 4a), the alloy shows similar microstructure to that of solution-treated one (Fig. 1b). The Zr-containing particles still exist at grain interior. The micrographs and corresponding selected area electron diffraction (SAED) patterns of precipitates in peak-aged condition are shown in Fig. 5. According to [0 0 0 1]␣ image (Fig. 5a), the plate-shaped precipitate has an average size less

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than 10 nm in width and 1–2 nm in thickness. Precipitates on {1 1¯ 0 0}␣ (indicated by grey triangle in Fig. 5a) and {1 1 2¯ 0}␣ (by black triangle) prism planes of the matrix are both observed, which is consistent with the investigation in Mg–3 wt.% Nd alloy [15] peak-aged at 200 ◦ C for 32 h. [0 0 0 1]␣ SAED pattern (Fig. 5b) does not reveal any additional reflections, which is also similar to that in Mg–3 wt.% Nd alloy [15]. According to [1 1 2¯ 0]␣ image (Fig. 5c), the precipitate has an average size about 50 nm in length along [0 0 0 1] direction of reciprocal space and several nanometers in thickness, which looks thicker than that along [0 0 0 1] direction of reciprocal space in [1 1¯ 0 0]␣ image (Fig. 5e). [1 1 2¯ 0]␣ SAED pattern (Fig. 5d), reveals unambiguously the existence of additional streaks at 1/2 {1 1¯ 0 0}␣ , which parallel to [1 1¯ 0 0] direction of reciprocal space, and confirms that there are precipitates on {1 1¯ 0 0}␣ prism planes. The ␤ phases with DO19 structure in Mg–RE alloy [21] usually form these kind of streaks. Hence, the precipitates here may be the ␤ phases. [1 1¯ 0 0]␣ SAED pattern (Fig. 5f), also reveals the presence of additional streaks, which parallel to [1 1 2¯ 0] direction of reciprocal space. The similar continuous streaks were once reported in Mg–0.5 at.% Nd [18] alloy, which was identified to be G.P. zones. The additional streaks here are discontinuous streaks. Therefore, the discontinuous streaks are probably corresponding to ␤ metastable phases, not G.P. zones. Hence, in 200 ◦ C peak-aging condition, the precipitates are probably mainly ␤ phase, which may have two kinds of habit planes, {1 1¯ 0 0}␣ and {1 1 2¯ 0}␣ prism planes. After aged at 250 ◦ C for 10 h, the alloy shows interesting microstructure after etching. The grains show different colors with different orientations (Fig. 4b), which may indicate different precipitates from that in 200 ◦ C peak-aged condition (Fig. 4a). Fig. 6 shows TEM images and SAED patterns of precipitates along different beam directions. [0 0 0 1]␣ image and SAED pattern suggest the existence of plate-shaped precipitates with habit plane parallel to {1 1¯ 0 0}␣ . The precipitates become larger in size and less in dense than that in 200 ◦ C peak-aged condition. [0 0 0 1]␣ SAED pattern does not reveal any additional reflections. According to [1 1 2¯ 0]␣ SAED pattern (Fig. 6d), additional reflections exist at 1/4 {1 1¯ 0 0}␣ , 2/4 {1 1¯ 0 0}␣ and 3/4 {1 1¯ 0 0}␣ and the precipitate seems to be ␤ phase. XRD pat-

Fig. 4. Optical images of Mg–3Nd–0.2Zn–0.4Zr (wt.%): (a) 540 ◦ C × 6 h + 200 ◦ C × 16 h and (b) 540 ◦ C × 6 h + 250 ◦ C × 10 h.

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Fig. 5. TEM image and corresponding diffraction pattern showing precipitates in peak-aged condition (200 ◦ C × 10 h): (a and b) along [0 0 0 1]␣ zone axis and corresponding SAED pattern; (c and d) along [1 1 2¯ 0]␣ zone axis; (e and f) along [1 1¯ 0 0]␣ zone axis.

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Fig. 6. TEM images and SAED patterns showing precipitates aged at 250 ◦ C for 10 h: (a and b) along [0 0 0 1]␣ zone axis; (c and d) along [1 1 2¯ 0]␣ zone axis.

tern (Fig. 7) confirms that the precipitates are ␤ phase with fcc structure (a = 0.742 nm) [22].

the alloy has the higher elongation than that in 200 ◦ C peakaged condition, nearly the same as the solution-treated condition, meanwhile the UTS drops to 265 MPa.

3.3. Mechanical properties 3.4. Fracture Tensile properties at room temperature of NZ30K alloy in different conditions are shown in Fig. 8. After solution treatment, the alloy has higher ultimate tensile strength (UTS) and elongation than that of as-cast alloy, however, the yield strength (YS) drops a little. After consequent aging, large improvements of UTS and YS are observed, but YS seems not sensitive to different aging treatments. The alloy in T6 condition has nearly the same YS, about 140 MPa, regardless of the aging process (Fig. 8). The highest UTS is reached when the alloy peak-aged at 200 ◦ C for 10–16 h, about 300–305 MPa. The elongation is greatly reduced from 22% in solution-treated condition to 11% in 200 ◦ C peak-aged condition. After aged at 250 ◦ C for 10 h,

Fig. 9 shows the optical microstructures of ruptured samples perpendicular to the fracture surface, which deformed at room temperature during tensile test. Secondary cracks near the fracture surface are observed. NZ30K alloy shows different secondary crack morphologies in different thermal conditions. In as-cast alloy, the secondary micro-cracks are observed along grain boundaries (Fig. 9a), mainly inside the eutectics. After solution treatment, the eutectics along grain boundaries disappear. The secondary cracks mainly locate inside the grains (Fig. 9b). Followed by low temperature aging at 200 ◦ C for 10 h, the secondary cracks are also at grain interiors (Fig. 9c).

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Fig. 7. XRD pattern of investigated Mg–3Nd–0.2Zn–0.4Zr (wt.%) after aging at 250 ◦ C for 10 h.

Aging at 250 ◦ C for 10 h, the secondary cracks mainly exist at grain boundaries (Fig. 9d). The locations of secondary cracks may reveal their initiation position. The secondary cracks at grain boundaries usually indicate they initiated there. But those at grain interiors may initiate at grain interiors or first start at grain boundaries and extend into grain interiors. Therefore, in

Fig. 8. Tensile properties of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy at room temperature in different conditions: as-cast; solution treated; T61: 540 ◦ C × 6 h + 200 ◦ C × 10 h; T62: 540 ◦ C × 6 h + 200 ◦ C × 16 h; T63: 540 ◦ C × 6 h + 250 ◦ C × 10 h.

as-cast and aged at 250 ◦ C for 10 h condition, the cracks probably mainly initiate at grain boundaries. For solution-treated and 200 ◦ C peak-aged condition, it is unclear whether the cracks form at grain interiors or grain boundaries. Fig. 10 shows SEM micrographs of fracture surfaces of NZ30K alloy in various conditions. In as-cast alloy, the failure surfaces are composed of fractured eutectics and cleavage planes

Fig. 9. Optical images of a longitudinal section of the fracture surface of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy in different conditions: (a) as-cast; (b) solution treated; (c) 540 ◦ C × 6 h + 200 ◦ C × 10 h; (d) 540 ◦ C × 6 h + 250 ◦ C × 10 h.

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Fig. 10. Fractography of tensile specimens of Mg–3Nd–0.2Zn–0.4Zr (wt.%) alloy in different conditions: (a) as-cast; (b) solution treated; (c) 540 ◦ C × 6 h + 200 ◦ C × 10 h; (d) 540 ◦ C × 6 h + 250 ◦ C × 10 h.

(Fig. 10a). It indicates that after the micro-crack forms inside the eutectics, it progresses along the grain boundaries by connecting with other micro-cracks or across the grains as cleavage fracture. Therefore, the fracture mode of as-cast NZ30K alloy is quasicleavage. When the alloy is subjected to solution treatment, the fracture surfaces are mainly composed of cleavage planes (Fig. 10b). The solution-treated NZ30K alloy shows transgranular cleavage fracture. After peak-aged at 200 ◦ C, the fracture surfaces are also characterized by cleavage planes (Fig. 10c). Hence the fracture mode of NZ30K alloy peak-aged at 200 ◦ C is also transgranular cleavage fracture, the same as that in solution-

treated condition. Aged at 250 ◦ C for 10 h, the fracture surface shows lots of tear ridges (Fig. 10d), which is corresponding to local plastic deformation at grain boundary and high elongation (Fig. 7), and cleavage planes. It probably indicates that the alloy in this condition has a mixed transgranular and intergranular fracture. 4. Discussion Rare earth metals (RE) are the most effective elements to improve the strength properties of magnesium alloys. The typi-

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cal heat treatments of these Mg–RE alloys are high temperature solution and consequent low temperature aging. Here the effect of solution and solution-plus-aging treatment on microstructure, mechanical properties and fracture behaviors are discussed. Four contributions to the YS above that pure Mg, denoted by σ Mg , can be identified traditionally for cast NZ30K alloy. They are: secondary phase (eutectic compounds) strengthening (σ sp ), solid solution strengthening (σ ss ), grain boundary strengthening (σ gb ) and precipitation strengthening (σ ppt ). YS of pure Mg is 21 MPa and that of Mg–(0.4–1.0 wt.%) Zr, whose grain size is estimated to be 30–40 ␮m [23], is to be 55 MPa by zirconium refinement [16]. Thus, the grain boundary strengthening contribution in refined as-cast Mg is estimated to be 55 − 21 = 34 MPa [5]. The as-cast NZ30K alloy consists of ␣ magnesium matrix and eutectic Mg12 Nd compounds (Fig. 1a). The content of other elements in as-cast ␣-Mg matrix is low and can be ignored to estimate the solid solution strengthening effect. Therefore in as-cast alloy, besides pure Mg, only second phase strengthening (σ sp ) and grain boundary strengthening (σ gb ) contribute to the strength. The alloy has low YS (88 MPa) in as-cast condition (Fig. 7). As discussed above, for rough estimate, the σ gb is about 34 MPa in as-cast alloy. The second phase strengthening contribution is about 88 − 55 = 33 MPa. The eutectics are complex in shape with large aspect ratio and are prone to fracture under local stress concentration during tensile test, just as large Si particles in Al–Si alloy [24] and Mg17 Al12 compounds in Mg–Al alloys [25]. The micro-cracks form easily by the fracture of eutectic compounds (Fig. 9a). Hence, the as-cast alloy has lowest fracture strength, about 175 MPa. As the soft grain interior has little resistance to the crack propagation, after the crack forming, it travels both intergranularly by connecting with other micro-cracks and transgranularly by cleavage fracture of part of grains. Small cleavage planes are left on the fracture surface by transgranular fracture of grains (Fig. 10a). The fracture pattern of as-cast NZ30K alloy is quasi-cleavage. After solution treatment, the eutectics along the grain boundary dissolve into the matrix (Fig. 1b). Only Small Zr–H and Zr–Zn particles [20] precipitate at grain interiors. These small particles’ influence on mechanical properties is unclear. As they are small in size and little in density, their influence on YS is ignored here. No small RE enriched quadrate phases are left as in Mg–Y–Sm–Zr [9] and Mg–Gd–Y–Zr [5] alloys. In solutiontreated alloy, besides pure Mg, only solid solution strengthening (σ ss ) and grain boundary strengthening (σ gb ) contribute to the strength. The alloy has lowest YS (85 MPa) in solution-treated condition. The σ gb is also about 34 MPa as no grains coarsen during solution treatment. The solid solution strengthening contribution is about 85 − 55 = 30 MPa. In solution-treated condition, the initiation and propagation of cracks may become much harder than that of as-cast condition for two reasons: (1) as the eutectics disappear, micro-cracks would not initiate easily by the fracture of eutectics; (2) the grain interior is strengthened by solid atoms, therefore, there would be more resistance for the cracks to across the grains. Hence, solution-treated alloy shows higher elongation (Fig. 7), about

22%, and fracture strength, about 210 MPa than that of as-cast alloy. As discussed in Section 3.4, though second cracks are observed at grain interiors (Fig. 9b), it is still not sure where the cracks initiate. After the formation of crack, it progresses mainly transgranularly as the fracture surface consists of nearly all cleavage planes (Fig. 10b). The solution-treated NZ30K alloy has transgranular cleavage fracture pattern. After peak-aged at 200 ◦ C for 10–16 h, dispersive precipitates form inside of grains (Fig. 5). These precipitates are plate-shaped and identified to be ␤ . The Zr–H and Zr–Zn particles are still there (Fig. 4a). In this condition, the solid solution strengthening contribution decreases due to the fact that the matrix is depleted of solute as the precipitation process proceeds. According to Mg–Nd phase diagram [26], the solubility of Nd in solid magnesium is reduced from 3.6 wt.% at the eutectic temperature to 0.08 wt.% at 200 ◦ C. Therefore, the solid solution strengthening can be ignored in 200 ◦ C peak-aged condition. Only grain boundary strengthening (σ gb ) and precipitation strengthening (σ ppt ) contributes to the strength here. As the YS at 200 ◦ C peak-aged condition is 137–142 MPa and the σ gb is still about 34 MPa, the precipitation strengthening contribution is about (137–142) − 55 = 82–87 MPa. The plate-shaped ␤ precipitation is the main reason to strengthen NZ30K alloy in 200 ◦ C peak-aged condition. Though the grain interiors are strengthened by the ␤ precipitations, the fracture pattern is still similar to that of solutiontreated condition. Secondary cracks are also found inside the grains (Fig. 9c). After the formation of crack, it progresses mainly transgranularly (Fig. 10c). The alloy shows highest fracture strength and lowest elongation, about 300–305 MPa and 11%, respectively. The fracture pattern of 200 ◦ C peak-aged NZ30K alloy could also be regarded as transgranular cleavage pattern. After aged at 250 ◦ C for 10 h, the strengthening precipitates change to ␤ (Figs. 6 and 7), which is larger in size and less in density. Similar to the peak-aged condition at 200 ◦ C, the alloy in this condition also has YS about 140 MPa. The precipitation strengthening contribution is also about 140 − 55 = 85 MPa. The initiate of the cracks are mainly along grain boundaries (Fig. 9d). After the formation of crack, it probably travels both intergranularly and transgranularly. Lots of tear ridges, which always indicate local plastic deformation at grain boundaries, and cleavage planes can be found on the fracture surface (Fig. 10d). Different from the one in 200 ◦ C peak-aged condition, the alloy in this condition has lower work hardening rate and higher elongation, which leads to moderate fracture strength, about 265 MPa. The fracture pattern of 250 ◦ C × 10 h aged NZ30K alloy could be regarded as trans-granular and intergranular mixed fracture. The change of crack initiation, work hardening rate and UTS of different aging treatments are mainly associated with different strengthening phases: kind, shape, orientation and distribution, volume fraction and number density. NZ30K has different strengthening phases in different aging condition: in 200 ◦ C peak-aged condition, ␤ precipitates with DO19 structure are the dominant phase, while at 250 ◦ C aged for 10 h, ␤ precipitates with fcc structure are the dominant phase. Just as in

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Table 1 Strengthening contributions in Mg–3Nd–0.2Zn–0.4Zr alloy at room temperature Strengthening contribution

Pure Mg strength Second phase strengthening Solid solution strengthening Grain boundary strengthening Precipitation strengthening Experimental yield strength

As-cast

Solution treated

Peak-aged at 200 ◦ C

Aged at 250 ◦ C for 10 h

MPa

%

MPa

%

MPa

%

MPa

%

21 33 0 34 0 88

24 38 0 38 0 100

21 0 30 34 0 85

25 0 35 40 0 100

21 0 0 34 82–87 137–142

15 0 0 24 61 100

21 0 0 34 85 140

15 0 0 24 61 100

Al [27] and Ni [28] alloys, metastable ␤ and ␤ phases probably have their different deformation mechanism during tensile test. The ␤ phase, coherent with matrix, which may exhibit precipitates shearing mechanism, tends to lead to an inhomogeneous distribution of slip, with shorter elongation. The ␤ phase, semi-coherent with matrix, which may have precipitate bypass mechanism, tends to homogenize the slip distribution, with higher elongation. The ␤ precipitates under 200 ◦ C peakaged condition are denser and smaller than ␤ precipitates under 250 ◦ C × 10 h aged condition, as shown in Figs. 5 and 6, and also possess a lager volume fraction since the solubility decreases with decrease the aging temperature. The two metastable phases have similar shape, orientation and distribution. It leads to higher hardness, work hardening rate and UTS in 200 ◦ C peak-aged condition than that in 250 ◦ C × 10 h aged condition. However, the YS of 200 ◦ C peak-aged condition is nearly the same as that in 250 ◦ C × 10 h aged condition. The hardness of aged NZ30K alloy seems corresponding to UTS not YS. The hardness of 200 ◦ C peak-aged alloy is 73 HV, corresponding to UTS 300–305 MPa. The hardness of 250 ◦ C × 10 h aged alloy is 60 HV, corresponding to UTS 265 MPa. The hardness and UTS match well. The strengthening contributions are summarized in Table 1. It is obvious that precipitation strengthening is the largest contributor to the strength either in 200 ◦ C peak-aged condition or 250 ◦ C × 10 h aged condition. About 60% strength comes from precipitation strengthening. The remarkable precipitation strengthening is attributable to plate-shaped precipitates formed on prismatic planes of Mg matrix. Different aging treatments lead to different precipitates. In 200 ◦ C peak-aged condition, ␤ is the dominant phase strengthening the alloy, led to high work hardening rate, higher UTS and shorter elongation. In 250 ◦ C × 10 h aged condition, ␤ is the strengthening phase, led to lower work hardening rate, lower UTS and higher elongation. 5. Conclusion (1) Mg–3Nd–0.2Zn–0.4Zr (wt.%) (NZ30K) alloy in as-cast, solution-treated and aged condition contains ␣ magnesium matrix + Mg12 Nd eutectic compounds, supersaturated solid solution + Zr–H particles + Zr–Zn particles and solid solution + Zr–H particles + Zr–Zn particles + plate-shaped precipitates, respectively. (2) In solution-treated condition, as the eutectic compounds dissolve into the matrix, NZ30K has higher strength and

elongation than that of as-cast alloy. Solid solution and grain boundary strengthening are the main contributors (about 75%) to the strength. (3) Peak-aged at 200 ◦ C for 10–16 h, fine plate-shaped ␤ phases precipitate and strengthen the alloy. In this condition, the alloy has highest ultimate tensile strength (UTS) and shortest elongation, 300–305 MPa and 11%, respectively. Aged at 250 ◦ C for 10 h, coarse plate-shaped ␤ phases precipitate and strengthen the alloy. In this condition, the alloy shows higher elongation and lower UTS, 20% and 265 MPa, respectively, than that of 200 ◦ C peak-aged condition. Yield strengths (YS) of these two-aged conditions are in the same level, about 140 MPa. Precipitation strengthening is the largest contributor (about 60%) to the strength in these two aged conditions. The hardness in aged NZ30K alloy seems corresponding to UTS not YS. (4) NZ30K alloy shows different fracture behaviors: in as-cast alloy, cracks form by the fracture of eutectics along the grain boundaries and propagate transgranularly; after solution treatment and peak-aged at 200 ◦ C, the alloy exhibits a transgranular cleavage fracture; after aged at 250 ◦ C for 10 h, the alloy shows probably a transgranular and intergranular mixed fracture. References [1] B.L. Mordike, T. Ebert, Mater. Sci. Eng. A 302 (2001) 37–45. [2] I.M. Baghni, Y.-S. Wu, J.-Q. Li, et al., Trans. Nonferrous Met. Soc. China 13 (2003) 1253–1259. [3] L.L. Rokhlin, T.V. Dobatkina, N.I. Nikitina, Mater. Sci. Forum 419 (2003) 291–296. [4] T. Mohri, M. Mabuchi, N. Satio, M. Nakamura, Mater. Sci. Eng. A 257 (1998) 287–294. [5] S.M. He, X.Q. Zeng, L.M. Peng, X. Gao, J.F. Nie, W.J. Ding, J. Alloy Compd. 427 (2007) 316–323. [6] S.M. He, X.Q. Zeng, L.M. Peng, X. Gao, J.F. Nie, W.J. Ding, J. Alloy Compd. 421 (2006) 309–313. [7] X. Gao, S.M. He, X.Q. Zeng, L.M. Peng, W.J. Ding, J.F. Nie, Mater. Sci. Eng. A 431 (2006) 322–327. [8] Q. Peng, J. Wang, Y. Wu, L. Wang, Mater. Sci. Eng. A 433 (2006) 133–138. [9] D. Li, Q. Wang, W. Ding, Mater. Sci. Eng. A 448 (2007) 165–170. [10] D. Li, Q. Wang, W. Ding, Mater. Sci. Eng. A 428 (2006) 295–300. [11] D. Li, J. Dong, X. Zeng, C. Lu, W. Ding, J. Alloy Compd. 439 (2006) 254–257. [12] K.Y. Zheng, J. Dong, X.Q. Zeng, W.J. Ding, Mater. Sci. Eng. A 454/455 (2007) 314–321. [13] L.L. Rokhlin, Magnesium Alloys Containing Rare Earth Metals, Taylor and Francis, London, 2003. [14] Z.-P. Luo, S.-Q. Zhang, L.-Q. Lu, G. Wei, J. Rare Earths 12 (1994) 296–298.

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