Precipitation process and effect on mechanical properties of Mg–9Gd–3Y–0.6Zn–0.5Zr alloy

Materials Science and Engineering A 454–455 (2007) 274–280

Precipitation process and effect on mechanical properties of Mg–9Gd–3Y–0.6Zn–0.5Zr alloy Z. Yang ∗ , J.P. Li, Y.C. Guo, T. Liu, F. Xia, Z.W. Zeng, M.X. Liang Department of Materials Science and Engineering, Xi’an Technological University, Xi’an 710032, China Received 5 July 2006; received in revised form 3 November 2006; accepted 7 November 2006

Abstract Precipitation process and its effect on the Vickers hardness, room temperature tensile properties of Mg–9Gd–3Y–0.6Zn–0.5Zr (GWZ930) alloy were investigated in combination with differential scanning calorimetric (DSC), transmission electron microscopy (TEM), Vickers hardness and room temperature tensile test. A four-stage precipitation sequence of supersaturated Mg solid solution, Mg(SSSS) → ␤ (144 ◦ C) → ␤ (240 ◦ C) → ␤1 (276 ◦ C) → ␤ (stable), was analyzed. A satisfied combination of high rupture tensile strength and ductility was obtained when the alloy was aged at 200 ◦ C for 63 h post extrusion. Moreover, the effective strengthening mechanisms were explicitly considered and the model prediction strength of GWZ930 alloy was proposed. It is concluded that precipitate strengthening due to the sub-micron metastable phases of ␤ and ␤ was the largest contributor to the room temperature tensile strength of the alloy GWZ930. © 2006 Elsevier B.V. All rights reserved. Keywords: Precipitation sequence; Mg–Gd–Y–Zn alloy; Mechanical properties; Strengthening

1. Introduction Precipitation hardened magnesium–rare earth alloys Mg–RE, particularly those containing heavy rare earth system, offer attractive properties for the aerospace and racing automotive industries [1]. About two decades ago, Drits et al. [2] discovered that 6 mass% Y addition to Mg–10Gd–0.6Mn alloy can increase the room temperature tensile property from 340 to 440 MPa and 300 ◦ C tensile property from 170 to 230 MPa. Rokhlin and Nikitina [3] developed high performance Mg–Gd–Y alloys, such as Mg–10Gd–5Y–0.5Mn alloy and Mg–10Gd–3Y–0.4Zr alloy, which exhibit higher specific strength at both room and elevated temperatures and good creep resistance than conventional aluminium and magnesium alloys including WE54 alloy. The tensile properties of Mg–10Gd–2Y–0.5Zr alloy at different heat treatments were conducted by He et al. [4] recently and a high rupture tensile strength up to 403 MPa was obtained. Nie et al. [5] studied the hardening response and creep resistance of Mg–6Gd–Zn alloys.

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Corresponding author. Tel.: +86 29 8320 8080; fax: +86 29 8320 8078. E-mail addresses: [email protected] (Z. Yang), [email protected] (J.P. Li). 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.11.047

They found that the presence of 1–2 wt% Zn in two ternary Mg–6Gd–Zr alloys can considerably enhance both tensile strengths at room temperature and creep strength at 175 ◦ C. However, few works on the relationship between microstructures and tensile properties of Mg–Gd–Y–Zn and Mg–Gd–Y–Zn–Zr alloys were published at the present time. Precipitation process of Mg–RE alloys had been extensively investigated by several researchers. The peak aging microstructure of Mg–RE alloys had been reported by Karimzadeh [6] and Lorimer [7]. Dispersed precipitates metastable phase ␤ and equilibrium phase ␤, and the two phases have been described as forming as plates on {1 1¯ 0 0}␣ planes of the magnesium matrix (a) phase. These prismatic precipitate plates play an important role in strengthening by providing effective barriers to gliding dislocations [8]. It is thus of interesting to determine the structure of such precipitate plates, and to identify those factors responsible for their presence in magnesium alloys. A three-stage precipitation sequence was proposed: Mg(SSSS) → ␤ → ␤ → ␤ [9–11]. More recently, Nie and Muddle [12] observed a fourth intermediate precipitate, designated ␤1 , forming between ␤ and ␤ phases in WE54. Apps et al. [13] observed the same precipitation process in the Mg–Gd–Y alloy. Almost all research results published on the precipitation behavior of Mg–RE alloys were established by transmission electron microscopy (TEM) observation. However, character-

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Fig. 1. (a) Precipitation hardening curves of GWZ930 alloy aged at different temperatures and (b) ln(HV ) vs. ln(t − t0 ) plot for the samples aged at 200 and 225 ◦ C.

izing structure evolution of the supersaturated Mg–RE alloy or accurately determining the forming temperature of precipitation phases with TEM are difficult. Moreover, the effect of precipitation phases and process on mechanical properties, such as Vickers hardness, tensile properties of Mg–Gd–Y alloys is not very clear at the present time. Differential scanning calorimetric (DSC) have been proved to be a powerful tool in characterizing the structure evolution of supersaturated magnesium alloys, giving also useful hints into design and modifications of heat treatments [16]. As well known, a better understanding of the precipitation process and microstructural constituents in Mg–Gd–Y–Zn alloys is essential in optimizing alloy design and preparation processing. The aim of this paper is to report the results of precipitation sequence and its effect on the mechanical properties of Mg–9Gd–3Y–0.5Zn–0.5Zr (GWZ930) alloy which were identified by DSC, TEM, Vickers hardness and tensile test. The relationship between precipitation phases and strengthening mechanism of the alloy is also discussed. 2. Experimental procedure The alloy Mg–8.8Gd–3.1Y–0.6Zn–0.5Zr (wt%) (GWZ930) was prepared with high purity (99.9%) Mg, Zn, Y, Mg–30.6Gd and Mg–30.33Zr (wt%) master alloys in a mild steel crucible under the protection of a mixed atmosphere of SF6 (10 vol%) and CO2 (bal.). The chemical compositions of the alloy are determined by inductively coupled plasma atomic emission spectrum (ICP-AES) apparatus. The ingots in diameter 45 mm were homogenized at 793 K for 10 h followed by quenching into hot water at 333 K, and then were hot extruded into rods in diameter 12 mm. The extrusion ratio is 16:1 and the extrusion temperature is 523 K. The extruded samples were machined into 6 mm gauge in diameter and 30 mm gauge length for tensile test, and 12 mm in diameter and 5 mm in length for aging treatment. The aging treatment was carried out at temperature 473–573 K with different times in an electric furnace. Specimens for optical microscopy (Neophot30) observation were prepared by standard techniques and were etched in 30% HNO3 with ethanol. Thin foils for TEM were prepared by ion-beam thinning operation with an incidence angle of 15◦ and were examined in a high

resolution TEM (JSM2010) with an attached Link INCA EDX system. Tensile test was carried out on a Zwick-10 kN material test machine at a crosshead speed of 5 mm/min. DSC traces were made on all the samples with DSC823e apparatus in a protective pure argon atmosphere, using pure magnesium as a reference. Hardness tests have been performed with a Vickers micro-indenter with a load of 100 N on the same sample used for DSC measurements. The reported values are the average of five indentations. 3. Results and discussion 3.1. Age hardening behavior Fig. 1a shows the precipitation hardening curves of GWZ930 alloy aged at temperature 200, 225 and 250 ◦ C, respectively. The peak hardness and corresponded aging time decrease with increasing aging temperature. When aging time at 200 ◦ C is less than 5 h, the hardness increased slightly. The hardness increased rapidly when aging time is longer than 5 h, and the peak hardness (∼127.2 Hv) is obtained when aging is about 63 h. When aged at 225 ◦ C, the specimen exhibits the peak hardness of 110.7 Hv, but decreases peak ageing time to about 40 h. The specimen aged at 250 ◦ C has the minimum peak hardness of 100.2 Hv, corresponding the aging time being about 6 h. After that, the hardness dropped rapidly. Precipitation is a phase transformation process. Its kinetics can be described by a Johnson–Mehl–Avrami (JMA) equation [14]: f = 1 − exp(−kt m )

(1)

where f is the degree of transformation of new phase fraction, t the time duration after the start of the transformation, and k (the rate constant) and m (the Avrami index) are two parameters characterizing the transformation process. From Figs. 1 and 2, we can conclude that when the alloy is aged at 200 ◦ C for less than 6 h, the hardness increase is due to the precipitation of ␤ phase. Vostrˇıy et al. [25] pointed out that formation of the ␤ phase is reflected by a decrease of electrical resistivity and a slight hardening in the temperature interval 180–220 ◦ C. This research results agree with the present results except the for-

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Fig. 2. DSC traces of the alloy GWZ930 at 225 ◦ C for different aging times (scanning rate: 2 K/min).

mation temperature of ␤ phase. When the alloy is aged at 200 ◦ C for more than 6 h, the formation of ␤ phase causes hardening increasing rapidly, and the peak hardness is achieved at 63 h, which is coincided with the result of reference [26]. When aging time is longer than 100 h, the formation of coarse and lamellar-shaped ␤1 (about 250 nm) causes the hardening decreasing quickly. Subsequently, the ␤1 phase will transfer to the stable coarse lamellar-shaped ␤ phase (1–2 ␮m) [26] when aging time is longer than 100 h or aging temperature is higher than 200 ◦ C. In the early stage at 200 ◦ C (for 6–63 h), from the precipitation hardening theories based on the interaction of dislocations and precipitates of ␤ and ␤ , i.e., the strengthening mechanism, a simple relation between hardness or strength and aging time can be derived as: Ht − H0 = Hv = K(t − t0 )n

(2)

where H0 is the initial hardness or strength at time t0 , Ht the hardness or strength at time t, K the temperature dependent hardness rate constant, and n is a time exponent, which is slightly temperature dependent. This time exponent n relates to the Avrami index m according to Eqs. (3) and (4): n=

1 m + 4 2

ln Hv − ln K = n ln(t − t0 )

(3) (4)

The early stage hardening kinetic can be explained based on the theory presented above. The slope of the linear regression curves is the value of n in Eq. (3), and is related to the Avrami index m. A large n represents a higher dependence on time of precipitation and hardening, and vice versa, as shown in Fig. 1b. When the specimen is aged at 200 ◦ C, n = 0.95, m = 1.4, then K = 0.021. When it is aged at 225 ◦ C, n = 0.63, m = 0.76, then K = 0.23. The relations between hardness and aging time at 200 and 225 ◦ C were given by: H200

◦C

= 0.021(t − t0 )0.95

H225

◦C

= 0.23(t − t0 )0.63

Thus, the appropriate JMA equation for GWZ930 alloy aged at 200 ◦ C for 6–63 h is as: f = 1 − exp(−k1 t 1.4 )

(5)

JMA equation for GWZ930 alloy aged at 225 ◦ C for 6–40 h is as: f = 1 − exp(−k2 t 0.76 )

(6)

where f is the total fraction of ␤ and ␤ phases, k1 and k2 are the rate constants corresponded to aging temperatures. 3.2. Precipitation process and microstructure Fig. 2 shows the DSC traces of the alloy GWZ930 in different aging states. No noticeable exothermal or endothermic signal of precipitation was observed in DSC traces of GWZ930 alloy when it was aged at 225 ◦ C for 6 h. When aged at the same temperature for 10 and 40 h, two calorimetric signals which are exothermal signal at 142–145 ◦ C and endothermic signal at 239–249 ◦ C are obviously observed. When aging time reaches 100 h, only one endothermic signal at 275–278 ◦ C is observed. It was obviously observed from Fig. 2 that the exothermic signals of 142–145 ◦ C and endothermic signal of 239–249 ◦ C happened in GWZ930 alloy which was aged at 225 ◦ C for 10 or 40 h. ␤ and ␤ phases are observed in Figs. 3 and 4. Fine globular ␤ precipitates with DO19 hexagonal super-cell, formed at 144 ◦ C, were observed forming on {1 0 1¯ 0} and [0 0 0 1]Mg zone axes in GWZ930 alloy aged at 225 ◦ C for 40 h. These globular precipitates distributed homogeneously and coherented with Mg (α = 2αMg , C = CMg ), with size up to 15 nm, as shown in Fig. 3. The ␤ phase forms in the first stages of precipitation at lower temperature (150–200 ◦ C), which was proposed by majority of works on magnesium–rare earth alloys [22–24,10]. The present result of DSC just conforms to previous work. Fig. 4 shows centre dark field and bright field images of globular ␤ precipitates and the corresponding SAD pattern from the same specimen of Fig. 3. The globular ␤ precipitates, formed at 240 ◦ C, with C base centre orthorhombic (cbco) and fine size up to 20 nm, were observed in GWZ930 alloy. The intermediate ␤ phase which formed at ageing temperatures between 200 and

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277

Fig. 3. (a) Centre dark field, (b) bright field image of globular ␤ precipitates taken along [0 0 0 1]Mg and (c) the corresponding SAED pattern, recorded from an alloy sample aged for 40 h at 225 ◦ C.

Fig. 4. TEM images of globular ␤ precipitates and the corresponding SAED pattern in GWZ930 alloy aged at 225 ◦ C for 40 h: (a) centre dark field, (b) bright field of and (c) SAED patter taken along [0 0 0 1]Mg .

250 ◦ C, were verified by Antion et al. [10]. Further ageing led to the lamellar-shaped precipitates forming on {1 0 1¯ 0}Mg in GWZ930 alloy aged at 225 ◦ C for 100 h, with size up to 20 nm in longitudinal and 200–300 nm in transverse. Micro-diffraction patterns, as shown in Fig. 5, indicated this phase has a fcc structure and a lattice parameter of 0.74 nm, This is identical to the ␤1 phase which were observed in both alloys WE54 [15] and Mg–15Gd [26]. Therefore, it is creditable that the exothermic signals of 142–145 ◦ C are due to the formation of the metastable phase ␤ , and the endothermic signal of 239–249 ◦ C is due to transformation of the metastable ␤ phase into ␤ , as shown in

Fig. 2. Finally, the endothermic signal at 275–278 ◦ C is due to the formation of the metastable ␤1 phases. The precipitation sequence ends with the formation of equilibrium ␤ phase which develops in a fcc structure (a = 2.223 nm) after ageing for a long time at 250 ◦ C and or above it [5]. Comparing the effects of aging time on Vickers hardness and DSC traces, no ␤ precipitate was detected by DSC when the alloy was aged at 225 ◦ C for 6 h. Precipitation reaction of Mg(SSSS) → ␤ → ␤ happened when aging time reaches 40 h. The specimens consists of ␤ and ␤ , corresponding to the peak hardness. After that, the Vickers hardness decreased with increasing of the aging time because of disappearance of ␤ and ␤ which coherent or semi-coherent with ␣-Mg and formation of coarse block-shaped ␤1 phase. Fig. 6 shows the effect of age treatment post extrusion on the microstructure of GWZ930 alloy. The average grain size, dynamic recrystallized, of extrusion GWZ930 alloy is 8 ␮m. The microstructure aged at 225 ◦ C for 40 h post extrusion consists of fine dynamic recrystallized ␣-Mg solid solution and precipitate phases of ␤ and ␤ , though ␤ and ␤ cannot be watched in optical micrograph. The grain size of ␣-Mg solid solution is about 5 ␮m. 3.3. Mechanical properties

Fig. 5. GWZ930 alloy aged at 225 ◦ C for 100 h viewed along [0 0 0 1]Mg .

Tensile properties of the alloy at different conditions are listed in Table 1. Compared with cast specimens, ultimate tensile strength (UTS) and tensile yield strength (TYS) at as-extruded specimens increased slightly, by about 38 and 67 MPa, respec-

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Fig. 6. Optical micrographs of GWZ930 alloy of (a) extruded and (b) aged at 225 ◦ C for 40 h post extrusion.

tively. The elongation, however, is greatly increased from 7 to 17.6%. In the condition of extrusion followed with ageing treatment, UTS of 430 MPa and TYS of 375 MPa were obtained in the treatment of extruded + 200 ◦ C/63 h. However, the elongation decreased from 17.6 to 9.5% in comparing with extruded states. Moreover, the effect of aging time on tensile properties is just as the effect of aging time on Vickers hardness. The aging time getting peak hardness corresponds to the maximum TYS and UTS. The effect of aging time on elongation is just reverse to that on UTS and TYS. 4. Predictions of strengthening contributions Based on the age hardening behavior and the result of mechanical properties of GWZ930 alloy, it is considered that the strengthening mechanism involves the following three contributors: solid solution strengthening, grain boundary strengthening and precipitation strengthening. 4.1. Solid solution strengthening The yield strength related to the solid solution effect is expressed by [17]: σy = σu +

3.1εGC1/2 700

to 0.61 at% (3.82 wt%) at 200 ◦ C. Assuming the solubility of Gd + Y in solid magnesium is 3.82 wt% (0.61 at%) at 200 ◦ C for peak aging time, the concentration of solute in at% is approximately 0.61%. To isolate the effect of the solute, yield strength may be plotted against the square root of the solute concentration. A linear fit on such a plot provides a measure of ε. Yield strengths were given by Rokhlin [11] for Mg–Gd alloy rods at T6 condition (as shown in Fig. 7). These rods had yield strengths of 21, 60 and 75 MPa for Mg contents of 0, 0.5 and 1%, respectively. Analyzing these data we can get the ␧ value of 0.74 for Mg for σ y in MPa. Thus, calculated contribution of solid solution strengthening to TYS of GWZ930 alloy is approximately 64 MPa. 4.2. Grain boundary strengthening The standard Hall–Petch equation, Eq. (8), was employed to relate the yield strength σ y of the material to the average grain size d: σy = σ0 + Kd −1/2

In this equation, σ 0 is the intrinsic resistance of the lattice to dislocation motion and d is the grain size. d = 10 ␮m at peak

(7)

where ε is an experimental constant, G the shear modulus of the matrix, C the concentration of the solute in at%, and σ u is the TYS of pure magnesium. The shear modulus of magnesium is approximately 1.66 × 104 MPa. The solid solubility of Gd in Mg has been reported as 4.53 at% or 23.49 wt% at 548 ◦ C and decreases exponentially, with a decrease in temperature, Table 1 The room temperature tensile properties of GWZ930 alloy at different states Condition

σ 0.2 (MPa)

σ b (MPa)

EL (%)

As-cast As-extruded Extruded + 200 ◦ C/40 h Extruded + 200 ◦ C/63 h Extruded + 200 ◦ C/100 h Extruded + 200 ◦ C/126 h

170 208 310 375 340 320

230 297 395 430 422 407

7.0 17.6 13.7 9.5 12.9 14.3

(8)

Fig. 7. TYS of Mg–Gd alloys at 20 ◦ C and T6 condition [11].

Z. Yang et al. / Materials Science and Engineering A 454–455 (2007) 274–280

aging state (see Fig. 6b). The value of K is determined by [27]:   τc 4Gb 1/2 (9) K=M (1 − υ)π where M is the Taylor factor, τ c the critical stress for slip to break through grain boundaries, υ the Poisson’s ratio (υ = 0.35), and b is the magnitude of the Burgers vector (3.21 × 10−10 m for Mg [19]). When τ c is the constant at the condition of d > 1 ␮m or ε > 0.35, K is the constant (280–320 MPa ␮m(−1/2) for Mg [19]). The calculated contribution of grain boundary strengthening to the strength is approximately 88–125 MPa. 4.3. Precipitation strengthening The precipitation strengthening involves the following four processes with the corresponding contributions [20]: (i) the dislocation–particle interaction associated with the Orowan process, σ 0 ; (ii) the load transfer from the matrix to particles, σ T ; (iii) the generation of dislocations due to the difference between the thermal expansions of the matrix and particles, σ g ; (iv) the generation of dislocations due to the geometric requirements in the course of deformation, σ f . The strength of the extruded solid material can be represented in the form: σ = σ0 + σT + σg + σf

(10)

The contribution of the load transfer from the matrix to particles can be defined by the relationship:   σT = σS 21 fv (11) where σ s (100–150 MPa for Mg alloys [20]) is the yield strength of the matrix (determined by the supersaturated magnesium solid solution) and the volume fraction of dispersed particles at peak hardness, fv is about 10% (see Fig. 4b). The contribution of this process is approximately equal to 5–7.5 MPa. For the globular precipitate ␤ in GWZ930 alloy, dp can be calculated by Eq. (13) [21]: dp =

πdt 4

(13)

For the block-shaped ␤ precipitates in GWZ930 alloy, dp can be calculated by Eq. (14) [21]  πdt dt dp = (14) 4 For plates-shaped precipitates ␤1 in GWZ930 alloy, dp can be calculated by Eq. (15) [21]:  πdt tt dp = (15) 4 where dt is the uniform diameter and tt is thickness (dp  tt ) of prismatic precipitate. When dt = 15 nm for precipitate ␤ (as shown Fig. 3), dp = 11.75 nm. When dt = 20 nm for precipitate ␤ (as shown Fig. 4), dp = 17.72 nm and when tt = 20 nm, dt = 250 nm for precipitate ␤1 (as shown Fig. 5), dp = 62.6 nm.

279

The strengthening due to the difference between the thermal expansions of the matrix and particles is given by [20]:   12TCfv 1/2 (16) σg = αGb bdp where α is a constant, T the temperature increment, C (CMg − C␤ ) the difference between the thermal expansion coefficients of the matrix and particles under investigation, and dp is the mean planar diameter of the obstacles. When α = 1.25 and fυ = 12.6%, CMg = 2.61 × 10−5 K, and C␤ = 7.5 × 10−6 K, dp = 11.75 nm for precipitate ␤ and dp = 17.72 nm for precipitate ␤ . Assuming the value fraction ratio of ␤ /␤ is 1:10 at the peak aging state, the fv of ␤ is 1 and 9% for ␤ precipitate, then the contribution is estimated as 217 MPa. At the over-aging state, the volume fraction, fv , of ␤ and ␤ will change to ␤1 phase completely, the contribution to strength is then estimated as 86 MPa. The contribution associated with the generation of dislocations due to the geometric requirements during deformation can be written as [20]:   8fv γ 1/2 σf = αGb (17) bdp where γ is the shear strain calculated with the use of the Taylor factor. According to [20], this contribution is estimated as 217 MPa. However, at the over-aging state, the contribution of precipitation strengthen of ␤1 is estimated as 86 MPa. For globular and block-shaped precipitates of uniform diameter dp that formed on {1 0 1¯ 0} or {2 1¯ 1¯ 0} prismatic planes of a-Mg, the appropriate version of the Orowan equation for magnesium alloys containing prismatic precipitate plates is given as Eq. (15) [18,21]: σ0 =

Gb 0.785dt ln √ b 2π 1 − υ((0.779/ fv ) − 0.785)dt √

(18)

When υ = 0.35, dt = 15 nm, fv = 1% for precipitate ␤ and dt = 20 nm, fv = 9% for precipitate ␤ . This contribution is estimated as 149 MPa. However, at the over-aging state, lamellar-shaped ␤1 , (dt  tt ) forms on {1 0 1¯ 0} or {2 1¯ 1¯ 0} prismatic planes of a-Mg. The appropriate version of the Orowan equation for magnesium alloys containing prismatic precipitate plates is given as Eq. (16) [18,21] σ0 =

√

2π 1 − υ(0.825  0.886 dp tt × b



Gb ln dp tt /fv − 0.393dp − 0.886tt ) (19)

When tt = 20 nm, dt = 250 nm, and dp = 62.6 nm for precipitate ␤1 . The contribution of precipitation strengthen of ␤1 is estimated as 80.8 MPa. Therefore, the total contribution of precipitation strengthening of ␤ and ␤ is about 588 MPa at the peak aging state. However, the total contribution of precipitation strengthening of ␤1 is about 260 MPa at the over-aging state.

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Based on the estimations to all strengthening mechanisms, it is suggested that precipitation strengthening is the largest contributor to the strength of the alloy GWZ930. The sum of all contributions should approximately equal 0.2% proof stress or the tensile yield strength [20]. The calculated total contribution should reach 740–770 MPa when the alloy was strengthened fully with precipitation of ␤ and ␤ , which is considerably larger than experimental strength. The difference between the value of calculation and experiment comes from the error of the parameters chosen in calculation and the defects which come from the preparation process of the alloy. On the other hand, the calculated total contribution should reach 412–449 MPa when the alloy was strengthened with precipitate of ␤1 , which is close well to the value of experimental strength. 5. Conclusions (1) A four-stage precipitation sequence of GWZ930 (Mg–9Gd– 3Y–0.5Zn–0.5Zr) alloy, Mg(SSSS) → ␤ (144 ◦ C) → ␤ (240 ◦ C) → ␤1(276 ◦ C) → ␤, was verified by the method of combination of DSC and TEM. The precipitation phases at peak aging state consists of ␤ and ␤ metastable phases. (2) Aging time has the same influence on the room temperature tensile properties and Vickers hardness of GWZ930 alloy. The peak hardness corresponds to the maximum ultimate tensile strength and the minimum ductility. (3) Precipitation strengthening is the largest contributor to GWZ930 alloy, while grain boundary strengthening also contributes significantly. (4) The calculated tensile yield strength of GWZ930 alloy based on the prediction model implied that the strengths of the alloy aged at 200 ◦ C for 63 h post extrusion are much higher than that aged at 200 ◦ C for 100 h. Acknowledgements Financial support by the Chinese Foundation Research Projection under Grant No. 5133001A4 and TEM analysis rendered by Mr. Han Zhiling and Ms. Liu Chunxia are gratefully acknowledged.

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