Microstructures and mechanical properties of AZ91D alloys with Y addition

Microstructures and mechanical properties of AZ91D alloys with Y addition

Materials Science and Engineering A 515 (2009) 152–161 Contents lists available at ScienceDirect Materials Science and Engineering A journal homepag...

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Materials Science and Engineering A 515 (2009) 152–161

Contents lists available at ScienceDirect

Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

Microstructures and mechanical properties of AZ91D alloys with Y addition Zude Zhao, Qiang Chen ∗ , Yanbin Wang, Dayu Shu Southwest Technique and Engineering Institute, Chongqing 400039, PR China

a r t i c l e

i n f o

Article history: Received 8 December 2008 Received in revised form 21 February 2009 Accepted 11 March 2009 Keywords: AZ91D alloy Yttrium Extrusion Tensile properties Microstructure

a b s t r a c t As-cast AZ91D + xY (x = 0, 0.5, 1, 2 mass%) magnesium alloys were prepared by a simple addition of Yrich interalloy to AZ91D alloy. Influences of Y on the microstructures and tensile properties of AZ91D alloy were investigated. Influences of extrusion temperature on the microstructure and tensile properties of AZ91D + 2Y alloy were investigated. Moreover, microstructures and the resulting tensile of extruded AZ91D + 2Y alloy were also investigated under different heat-treatment conditions. The results show that the addition of Y to AZ91D alloy refines primary ␣-Mg matrix and ␤-Mg17 Al12 phase. The addition of 2 mass%Y forms rod-shaped Al2 Y phase. After hot extrusion, the grain sizes of AZ91D and AZ91D + 2Y alloys are greatly refined through dynamic recrystallisation. At extrusion temperature below 300 ◦ C, the microstructure of AZ91D + 2Y alloy is characterized by fine and equiaxed grains, which leads to a high tensile strength but relatively low elongation. With increasing extrusion temperature up to 400 ◦ C, more equiaxed grains with larger size are obtained. A low strength and large elongation are achieved compared with those extruded at low temperature. The best combination of both high strength and large elongation takes place at 325 ◦ C. As for extruded AZ91D + 2Y alloy, solution treatment at 413 ◦ C (T4) causes parts of broken ␤-Mg17 Al12 phase to be dissolved into ␣-Mg matrix and the deterioration of tensile properties. During aging at 216 ◦ C (T5), discontinuous precipitates appear at grain boundaries. Solid solution followed by aging (T6) treatment increases ultimate tensile strength but decreases yield strength and elongation. © 2009 Elsevier B.V. All rights reserved.

1. Introduction Magnesium alloys have drawn a great deal of attention from automotive and aerospace industries, due to their low density, highspecific strength and good damping capacity [1]. The widely used magnesium alloys belong to Mg–Al–Zn series alloys, which have excellent castability and low cost [2–4]. However, the application of these alloys has been limited because of their poor mechanical properties. To meet the challenges of new application, the mechanical properties of Mg–Al–Zn series alloys need to be further improved [5]. Recently, several investigators [6–9] have reported that the addition of Ca, Sr, Bi, Sb or rare earth (RE) elements was an effective way to improve the mechanical properties of Mg–Al–Zn series alloys. Yuan et al. [6] have conducted tensile tests on as-cast AZ91 alloy treated by combined addition of Bi and Sb. Results showed that a small amount of combined addition of Bi and Sb to AZ91 alloy increased the yield strength (YS) and creep resistance significantly at elevated temperature up to 200 ◦ C. Hirai et al. [7] have reported that with the combined addition of 1.0 wt.%Ca and 0.5 wt.%Sr, the ultimate tensile strength (UTS) of as-cast AZ91 alloy increased from

∗ Corresponding author. Tel.: +86 023 68792238. E-mail address: [email protected] (Q. Chen). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.03.030

228 to 251 MPa. The increase in the UTS of AZ91 alloy could be attributed to grain refinement. Wu et al. [8] studied the effect of combined addition of Ca and RE elements on the mechanical properties of AZ91D alloy. The results show that with the combined addition of 1 mass%Ca and 1 mass%RE elements to AZ91D alloy, the UTS increased by 15.9%. Zhang et al. [9] have reported the effect of Y-rich misch metal on the microstructure and mechanical properties of die cast AZ91 alloy. Results showed that small amount of Y addition to AZ91 alloy caused significant refinement of ␣ phase and eutectic ␤ phase. At room temperature, the UTS, YS and elongation of die cast AZ91 + 0.8 wt.%Y were 270 MPa, 160 MPa and 11%, respectively. Y-rich interalloy is abundant in China and its price is much lower than that of pure yttrium metal. However, to the best of our knowledge, Y has been little used in magnesium alloys except in Mg–Al–Zn [9], WE (WE43 and WE54) [10] and Mg–Zn–Zr series alloys [11] up to now. Therefore, it is necessary to develop the applications of Y-rich interalloy in magnesium alloys and study the effect of Y on microstructure and mechanical properties of magnesium alloys. The comparison in physical characteristics between Mg and Y is given in Table 1. In the present work, AZ91D alloy was used as the master alloy and different amounts of Y were added into AZ91D alloy. Influences of Y on the microstructure and mechanical properties of as-cast and extruded AZ91D alloy were studied. At the same time, the effects of extrusion temperature on the AZ91D alloy

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Table 1 Physical characteristics of magnesium and yttrium. Element

Atomic number

Atom weight

Atom radius (nm)

Melting point (K)

Special density (g/mm3 )

Crystal structure

Mg Y

12 39

24 89

0.162 0.182

922 1799

1.74 4.50

HCP HCP

Table 2 Chemical composition of interalloy (wt.%). Elements

Composition

Y La Nd Dy Ho Er Tm Yb Lu Fe Mg

20.7 0.019 0.029 0.020 1.179 2.09 0.027 0.003 0.102 0.028 Bal.

Table 3 Chemical composition of studied AZ91D with the addition of Y (mass%). Composition

AZ91D AZ91D + 0.5Y AZ91D + 1Y AZ91D + 2Y

Mg

Al

Zn

Mn

Y

Bal. Bal. Bal. Bal.

9.32 9.18 9.11 8.92

1.15 1.21 1.25 1.18

0.21 0.14 0.08 0.09

0 0.46 0.96 1.76

containing 2 mass%Y were studied. Moreover, we also reported a study of the microstructural evolution and the resulting mechanical properties of extruded AZ91D alloy containing 2 mass%Y under different heat-treatment conditions. 2. Experimental procedures In the present research, as-cast AZ91D was chosen as the master alloy. The AZ91D alloys with different amounts of Y were prepared by melting AZ91D alloy in an electric resistance furnace being flushed continuously with SF6 and CO2 gas (mixing ratio was 1:100) at 750 ◦ C. And then Y element in the form of Y-rich interalloy was added into the metal. The chemical composition of Y-rich interalloy was shown in Table 2. The metal was isothermally held at 760 ◦ C for about 20 min to ensure that the Y was completely dissolved. After that the metal was poured in a metal mold with a diameter of 70 mm. The compositions of AZ91D + Y alloys are shown in Table 3. To reveal grain boundaries of alloys, as-cast AZ91D and

AZ91D + 2Y samples were reheated to 560 ◦ C (semi-solid interval) for 10 min isothermal holding, respectively. During partial remelting, Ar was used as a protective atmosphere to prevent oxidization. When samples were reheated to 560 ◦ C for 10 min holding, samples were taken out immediately for water quenching. During extrusion, as-cast AZ91D + 2Y alloys were extruded into bars with a diameter of 13 mm at 250, 300, 325, 350 and 400 ◦ C, respectively. Moreover, as-cast AZ91D alloys were also extruded into bars with a diameter of 13 mm at 300 ◦ C. The extrusion ratio was 16:1. AZ91D + 2Y samples, produced by extrusion at 300 ◦ C, were subjected to different heat-treatment processes. The solution treatment (T4) was carried out at 413 ◦ C for up to 5 h in air with the protection of carbon power on the surface of samples. The samples were cooling in air. Those samples that had been subjected to solution treatment at 413 ◦ C for up to 5 h were then aged at 216 ◦ C for up to 5 h (T6). Moreover, some extruded samples were aged at 216 ◦ C for up to 5 h (T5). The mechanical properties were measured according to ASTM B557 on cylindrical samples with a reduced section of 6 mm and a gauge length of 50 mm using an Instron 5569 testing machine at a cross-head speed of 1 mm/min. The microstructures were examined by optical microscopy (OM), X-ray diffraction (XRD) and scanning electron microscopy (SEM) equipped with an energy dispersive X-ray spectrometer (EDS). The samples for OM and SEM were prepared by standard technique of grinding with SiC abrasive and polishing with a diamond spray (0.5 ␮m). As-cast samples were etched in 4% aqueous nitric acid. For as-extruded samples, they were etched in a solution of 100 ml ethanol, 6 g picric acid, 5 ml acetic acid and 10 ml water. Grain sizes were measured using a mean linear intercept method described in the ASTM specification E112-96. For each sample, measurements were taken from the whole sectioned area with 300–400 intercepts counted per sample. 3. Results 3.1. Microstructure 3.1.1. Microstructure of as-cast AZ91D with the addition of Y Fig. 1 shows microstructures of as-cast AZ91D and AZ91D + 2Y alloys. As shown in Fig. 1(a), as-cast AZ91D alloy consisted of primary ␣-Mg matrix and eutectic phase. The eutectic phase precipitated as discontinuous network at grain boundaries. With the

Fig. 1. Microstructures of as-cast: (a) AZ91D and (b) AZ91D + 2Y alloys.

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Fig. 2. Microstructure of as-cast: (a) AZ91D and (b) AZ91D + 2Y alloys were reheated to 560 ◦ C for 10 min holding.

addition of Y, coarse eutectic phase was refined, as shown in Fig. 1(b). To reveal grain boundaries of alloys in short times, as-cast AZ91D and AZ91D + 2Y alloys were reheated to 560 ◦ C for 10 min isothermal holding (Fig. 2). As shown in Fig. 2, the mean grain size of as-cast AZ91D alloy was 172 ␮m (Fig. 2(a)). However, the mean grain size decreased significantly to 60 ␮m with the addition of Y (Fig. 2(b)). XRD analysis indicates that AZ91D alloy consisted of ␣Mg matrix and ␤-Mg17 Al12 phase (Fig. 3(a)). With the addition of Y, the high-temperature stable Al2 Y phase was found in AZ91D + 2Y alloy (Fig. 3(b)). To determine the formation of new phases, SEM analysis and EDS spectrum of AZ91D and AZ91D + 2Y alloys were carried out, as shown in Figs. 4 and 5, respectively. As shown in Fig. 4, EDS analysis results indicate the existences of Al and Zn elements in the ␣-Mg matrix. The Mg/Al ratio (69.04/28.25) of the ␤-Mg17 Al12 phase was found to be a higher than the stoichiometric ratio, 1.42 (17/12). The excessive Mg content in the spectrum could be attributed to the contribution of the excess Mg by the ␣-Mg solid solution matrix [12,13]. As shown in Fig. 5, with the addition of Y, the rod-shaped phase (A point) appeared in the microstructure. EDS analysis indicates that the rod-shaped phase (A point) mainly consisted of Al and Y. For the rod-shaped phase, the Al/Y ratio (64.40/27.53) was close to 2:1 and therefore it could be identified as Al2 Y phase. The length of the rod-shaped Al2 Y precipitate measured by Image Tool software was about 24 ␮m, the width about 4 ␮m and the aspect ratio was hence about 6. The formation of Al2 Y phase was further confirmed by XRD analysis, as shown in Fig. 3(b). Note that some polygonal particles (B point) also appeared in the microstructure. For these polygonal particles, the Al/Y ratio (61.70/23.09) was about 2.64. The addition of Y may bring about the formation of Al3 Y phase. However, further TEM work was needed to verify this.

3.1.2. Microstructure evolution of the hot extruded alloys Fig. 6 shows the microstructures of the longitudinal sections of AZ91D and AZ91D + 2Y alloys extruded at 300 ◦ C with an extrusion ratio of 16:1. As shown in Fig. 6, dynamic recrystallisation (DRX) during the extrusion process has transformed the dendritic microstructures of as-cast alloys into fine and equiaxed grain microstructures. The mean grain sizes of extruded AZ91D and AZ91D + 2Y alloys were about 23 and 12 ␮m, respectively. The dark areas in Fig. 6 consisted of ␣-Mg and lamellar precipitates of ␤-Mg17 Al12 . Remarkable differences in the local Al-content have been found by EDS analysis using area acquisition with a window size of 20 ␮m × 20 ␮m. For example, in the microstructure of extruded AZ91D alloy, the bands featuring lamellar Mg17 Al12 contained about 12% Al whereas the bands free of precipitates contained only approximately 7% Al. Fig. 7 shows the microstructures of the longitudinal sections of AZ91D + 2Y alloys extruded at (a) 250, (b) 300, (c) 325, (d) 350 and (e) 400 ◦ C, respectively. As shown in Fig. 7, the microstructures of extruded alloy consisted of equiaxed grains, which indicated that the alloy has occurred DRX. During the hot extrusion process, parts of broken ␤-Mg17 Al12 phase dissolved into the ␣-Mg matrix. Comparison of Fig. 7(a)–(e) shows that DRX grain size increased with raising extrusion temperature. 3.1.3. Heat-treated microstructure Fig. 8 shows microstructures of the extruded AZ91D + 2Y alloy (300 ◦ C) heat-treated at (a) 413 ◦ C for 5 h, (b) 216 ◦ C for 5 h and (c) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h. As shown in Fig. 8(a), solid solution treatment of alloy at 413 ◦ C for 5 h (T4) led a majority of ␤-Mg17 Al12 precipitates to be dissolved into the matrix. Fig. 8(a) also shows that slight grain coarsening had occurred during T4. For the extruded AZ91D + 2Y alloy, aging at 216 ◦ C for 5 h (T5)

Fig. 3. XRD patterns of as-cast: (a) AZ91D and (b) AZ91D + 2Y alloys.

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Fig. 4. SEM images of: (a) as-cast AZ91D alloy and EDS patterns of (b) A point and (c) B point in (a).

Fig. 5. SEM images of: (a) as-cast AZ91D + 2Y alloy and EDS patterns of (b) A point and (c) B point in (a).

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Fig. 6. Microstructures in longitudinal sections of extruded: (a) AZ91D and (b) AZ91D + 2Y at 300 ◦ C.

caused the formation of Mg17 Al12 precipitates (Fig. 8(b)). In comparison with T5, more Mg17 Al12 precipitated from the ␣-Mg matrix in the T6 condition (Fig. 8(c)). Fig. 9 shows SEM microstructures of the extruded AZ91D + 2Y alloy heat-treated at (a) 413 ◦ C for 5 h, (b)

216 ◦ C for 5 h and (c) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h. It can be seen from Fig. 9(a) that the Al2 Y phase was observed and it was without any change, which indicated the stability of Al2 Y phase at higher temperature. When the alloy was aged at 216 ◦ C for 5 h (T5

Fig. 7. Microstructures in longitudinal sections of AZ91D + 2Y alloys extruded at: (a) 250 ◦ C, (b) 300 ◦ C, (c) 325 ◦ C, (d) 350 ◦ C and (e) 400 ◦ C.

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Fig. 8. The microstructures of the extruded AZ91D + 2Y alloy (300 ◦ C) heat-treated at (a) 413 ◦ C for 5 h, (b) 216 ◦ C for 5 h and (c) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h.

Fig. 9. SEM micrographs showing the microstructure of the extruded AZ91D + 2Y alloy heat-treated at: (a) 413 ◦ C for 5 h, (b) 216 ◦ C for 5 h and (c) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h.

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Fig. 10. SEM micrographs showing the microstructure of AZ91D + 2Y alloy extruded at 300 ◦ C followed by: (a) aging for 5 h and (b) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h.

condition), Al was precipitated out as Mg17 Al12 at grain boundary in the form of discontinuous lamellar (Fig. 9(b) and Fig. 10(a)). When the alloy was treated at 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h, the microstructure was similar to that of sample treated by T5 condition, as shown in Fig. 9(c). Close examination of the microstructure (Fig. 10(b)) revealed that the size of discontinuous lamellar in certain areas was coarser than that of sample treated by T5 condition. 3.2. Mechanical properties and tensile fracture analysis 3.2.1. Tensile properties Fig. 11 shows the tensile test results of as-cast AZ91D, AZ91D + 0.5Y, AZ91D + 1Y and AZ91D + 2Y alloys. Compared to those of AZ91D alloy, the UTS, YS and elongation of AZ91D + 2Y alloy increased to 145.9 ± 6.1 MPa, 81.87 ± 3.6 MPa and 3.18 ± 0.62%, respectively. Fig. 12 shows the tensile test results of AZ91D and AZ91D + Y alloys extruded at 300 ◦ C. As shown in Fig. 12, the UTS, YS and elongation of extruded alloys increased from 293.3 ± 1.9 MPa, 184.3 ± 9.7 MPa and 10.44 ± 0.9% to 323.15 ± 1.6 MPa, 216.9 ± 4.3 MPa and 14.31 ± 0.4%, respectively, with the increase of Y from 0 to 2 mass%. Fig. 13 shows that the tensile test results of AZ91D + 2Y alloys extruded at 250, 300, 325, 350 and 400 ◦ C. As shown in Fig. 13, with the extrusion temperature increased from 250 and 400 ◦ C, the UTS and YS of extruded AZ91D + 2Y alloys decreased from 330.7 ± 0.7 and 225.9 ± 7.0 MPa to 308.7 ± 1.1 and 182.5 ± 1.6 MPa, respectively. However, the elongation increased from 13.4 ± 0.4 to 20.1 ± 1.6%. These results indicated that the raising extrusion temperature had a negative

Fig. 11. The tensile test results of as-cast AZ91D, AZ91D + 0.5Y, AZ91D + 1Y and AZ91D + 2Y alloys.

Fig. 12. The tensile test results of AZ91D and AZ91D + Y alloys extruded at 300 ◦ C.

effect on UTS and YS but a favorable effect on the elongation of alloys. Fig. 14 shows the tensile tests of the extruded AZ91D + 2Y alloy at different heat-treatment conditions. T4 brought about tremendous deterioration in UTS, YS and elongation. T5 improved the UTS but with a considerable reduction in YS and elongation. Moreover, T6 slightly increased ultimate tensile strength but with a considerable reduction in YS and elongation. Note that all heat-treatment conditions in the present study exerted a negative influence on YS and elongation.

Fig. 13. The tensile test results of AZ91D + 2Y alloys extruded at 250, 300, 325, 350 and 400 ◦ C.

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Fig. 14. The tensile test results of extruded AZ91D + 2Y alloy heat-treated under different conditions: (A) the extruded AZ91D + 2Y alloy; (B) 413 ◦ C for 5 h (T4); (C) 216 ◦ C for 5 h (T5); (D) 413 ◦ C for 5 h followed by aging at 216 ◦ C for 5 h (T6).

3.2.2. Tensile fracture Fig. 15 shows the SEM images of tensile fracture surface of (a) AZ91D and (b) AZ91D + 2Y extruded at 300 ◦ C. The tensile fracture surfaces of these alloys mainly consisted of tear ridges, dimples and second-phase particles. Therefore, it can be concluded that the fracture behaviors of as-extruded AZ91D with and without Y addition were similar. Fig. 16 shows the SEM images of tensile fracture surfaces of AZ91D + 2Y alloys extruded at (a) 250 and (b) 400 ◦ C. The tensile fracture surfaces of these alloys also consisted of tear ridges, dimples and second-phase particles. Note that the failure surface in Fig. 16(b) had a large number of dimples, which confirmed closely to the good ductility of AZ91D + 2Y extruded at 400 ◦ C. The dominant feature of fractures of AZ91D + Y alloys extruded at different temperatures was dimples pattern, which indicated these alloys had undergone plastic deformation before fracture. 4. Discussion XRD analysis and microstructure observations by SEM revealed that as-cast AZ91D alloy consisted of ␣-Mg matrix and ␤-Mg17 Al12 phase. The ␤-Mg17 Al12 phase in the grain boundaries was formed by the eutectic reaction L → ␣ + ␤. The as-cast AZ91D alloy exhibited

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poor mechanical properties, which could be mainly attributed to the presence of the coarse eutectic ␤-Mg17 Al12 network at grain boundaries. The ␤-Mg17 Al12 phase with a body-centered cubic (BCC) crystal structure was incoherent with the magnesium matrix with hexagonal close-packed (HCP) structure, which resulted in the fragility of the Mg/Mg17 Al12 interface [14,15]. The rod-shaped precipitates of Al2 Y synthesized in the as-cast AZ91D with Y addition. The electronegativities of Y, Mg, Mn and Al were 1.22, 1.31, 1.55 and 1.61, respectively, and the electronegativity difference was between Y and Al (Table 4). Therefore, the affinity between Y and Al was strong and synthesis of Al2 Y occurred preferentially. In contrast with those of as-cast AZ91D alloy, the tensile properties of as-cast AZ91D + Y alloy exhibited better tensile properties, which could be mainly attributed to the refinement of both ␣-Mg matrix and ␤Mg17 Al12 phase, as shown in Figs. 1 and 2. Y was a surface-active element to Mg. During the solidification, the solution atoms of Al and Y were pushed and then enriched at the front of solid–liquid interface. The aggregation of Y on the interface of ␣-Mg could bring ingredient over-cooling, effectively hindering ␣-Mg growth. In this case, the nucleation rate of ␣-Mg was increased. Moreover, Y and Mg have HCP structure (Table 1) and the lattice constant of Y is very close to that of Mg. The lattice constants of Y and Mg are a = 0.36500 and c = 0.57410 nm, and a = 0.32094 and c = 0.52104 nm, respectively [16]. The atomic radius of Y is 0.182 nm, which is also close to that of Mg (0.162 nm). According to the theory of lattice matching, Y can act as hetero-nucleating center for the ␣-Mg matrix, which promotes the nucleating ratio and refines grains. During the process of further cooling, the growth of ␤-Mg17 Al12 phase was suppressed greatly due to the enrichment of Y in solid–liquid interface. Therefore, ␤-Mg17 Al12 phase was modified from discontinuous network to fine particles [9]. The decrease in the volume fraction of ␤-Mg17 Al12 phase also led to the increase in tensile properties. Because the formation of Al2 Y phase led to the depletion of Al atoms, the decrease of amount of ␤-Mg17 Al12 phase would occur with the addition of Y (Figs. 4 and 5). Note that the presence of rod-shaped precipitates in the as-cast AZ91D + Y alloy was harmful to tensile properties. This may counteract the effect of the Y addition because of stress concentration occurring at the ends of the precipitates. During hot extrusion, banded structures (black areas) were found in the microstructures of extruded AZ91D and AZ91D + 2Y alloys (Fig. 6). Banded structures were typical for extruded products for Mg–Al alloys. Extrusion temperature was within the single-

Fig. 15. SEM images of tensile fracture surfaces of: (a) AZ91D and (b) AZ91D + 2Y alloys extruded at 300 ◦ C.

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Fig. 16. SEM images of tensile fracture surfaces of AZ91D + 2Y alloys extruded at: (a) 250 ◦ C and (b) 400 ◦ C.

phase field ␣-Mg for the investigated AZ91D and AZ91D + 2Y alloys and the intermetallic Mg17 Al12 was partially dissolved. However, time was too short to allow complete homogenization during extrusion, i.e. segregation was only reduced and elongated but not eliminated [17]. On cooling from the extrusion temperature, precipitation of lamellar Mg17 Al12 occurred in the areas of higher Al-content [17]. Comparison of Fig. 6(a)–(b) shows that the average spaces of segregation were found to be decreased with the addition of Y. The reason could be explained as follows: the formation of new phase resulted in the depletion of Al atoms, which further led to decrease of amount of Mg17 Al12 phase. Compared with that of extruded AZ91D alloy, the tensile properties of extruded AZ91D alloys with addition of Y were improved significantly. The excellent tensile properties of extruded AZ91D with the addition of Y could be attributed to grain refinement caused by the addition of Y and DRX during hot extrusion. Firstly, the recrystallised grain size depended mainly on the original grain size of the material. The finer the original grain size, the finer the recrystallised grain size after hot extrusion. Secondly, thermally stable Y-containing intermetallic compounds pinned the grain boundaries and restricted the grain growth during DRX. Some researchers [11,18] suggested that intermetallic compounds were broken into small particles and then distributed in the microstructures after hot extrusion. These small particles were responsible for the increase in the strength of materials due to Orowan mechanism. However, intermetallic compounds were observed to exist as irregular particles and distributed unevenly in the microstructure of extruded AZ91D + Y treated by different heat-treatment conditions, as shown in Fig. 9. Because the volume fraction of these unbroken intermetallic compounds with irregular shape was less than 5%, strengthening by these particles was expected to be very small. The effect of extrusion temperature on the evolution of tensile properties can be explained by examining the microstructures of

Table 4 Electronegativity of elements in AZ91D + Y alloys. Element

Electronegativity

Mg Al Mn Y

1.31 1.61 1.55 1.22

AZ91D + 2Y alloys extruded at different temperatures. As shown in Fig. 7(a), the microstructure in the alloy extruded at 250 ◦ C is mainly characterized by fine and equiaxed grains. Due to relatively low temperature, it was difficult for dislocation to rearrange via movement, which provided a strong obstacle to dislocation slip. Therefore, a high strength but low ductility was expected. In contrast, the recrystallised microstructures, more equiaxed grains with larger grain size formed at higher extrusion temperature, could supply relatively low strength but high ductility. Due to increase in extrusion temperature, it provided a great opportunity for the internal stress and dislocation density to decrease, which was beneficial for increase in elongation. The increase of extrusion temperature provided an effect similar to annealing as widely observed in AM30 [19], ultrafine-grained Cu [20], Al alloys [21,22], low-carbon steel [23] and stainless steel [24], which obviously decreased the dislocation density and internal stress, and then decreased the strength but increased the ductility. During the solution treatment (T4), parts of broken ␤-Mg17 Al12 phase were dissolved into the ␣-Mg matrix (Fig. 8(a)). However, the solution treatment at 413 ◦ C did not result in dissolution of Al2 Y phase (Fig. 9(a)). This indicates that Al2 Y phase has a high thermal stability at elevated temperature. The effect of solution treatment (T4) also led to the coarsening of grains. The final grains were dependent on the solution temperature and the time at temperature for subsequent grain growth. Under the present experimental conditions, it seems that T4 procedure did not improve the mechanical properties at all, which can be mainly attributed to grain coarsening. The decrease of strength in the procedure of T4 was in agreement with a previous observation made by Wang et al. [12]. However, contrary to their results T4 in the present study also deteriorated the ductility of the alloy. It was observed from the aged microstructures (Fig. 9(b) and 10(a)) that Al was precipitated out as Mg17 Al12 in the form of discontinuous lamellar. As well as documented in the literature, two forms of precipitates, discontinuous and continuous Mg17 Al12 of same composition originated from the supersaturated ␣-Mg solid solution [12,25]. Discontinuous precipitate occurred preferentially along the grain boundaries in the early stage of aging and later, as the solute supersaturation went down, continuous precipitates formed throughout the grain [15]. In the present study, these discontinuous precipitates played a role of hindering dislocation slip, which led to increase in UTS but decrease in both YS and elongation.

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5. Conclusions

References

1. With the addition of 2 mass%Y, coarse grain and ␤-Mg17 Al12 phase are refined and rod-shaped Al2 Y precipitate phase is observed in the microstructure. The ultimate tensile strength, yield strength and elongation of as-cast AZ91D alloys increase with the addition of Y. 2. The addition of Y has a favorable effect on reducing the grain size of the extruded AZ91D alloys. Tensile tests confirm that the ultimate tensile strength, yield strength and elongation of extruded AZ91D alloys increase with increasing content of Y in the range investigated. 3. At the extrusion temperature below 300 ◦ C, dynamic recrystallisation takes place and the microstructures are characterized by fine and equiaxed grain microstructures. These microstructures exhibit high tensile strength but low elongation. With increasing the extrusion temperature, more equiaxed grains with larger size are obtained. A low strength and large elongation are achieved compared with those extruded at low temperature. The best combination of both high strength and large elongation takes place at 325 ◦ C. The tensile fracture surfaces of extruded AZ91D and AZ91D + 2Y alloys are mainly composed of tear ridges, dimples and second-phase particles. 4. As for extruded AZ91D + 2Y alloys, T4 treatment (solution) causes parts of ␤-Mg17 Al12 phase to be dissolved into ␣-Mg matrix and slight grain coarsening. T5 (artificial aging) and T6 (solution plus artificial aging) treatments cause the precipitation of the ␤Mg17 Al12 phase. T4 treatment does not improve the mechanical properties at all. T5 and T6 treatments increase ultimate tensile strength but decreased yield strength and elongation.

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