Effects of heat treatment on microstructure and mechanical properties of Mg–2·6Sm–1·3Gd–0·6Zn

Abstract

The microstructure and mechanical properties of as cast and different heat treated Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr (wt-) alloy were investigated. The alloy in the as cast state is mainly composed of α-Mg matrix and (Mg,Zn)₃(Sm,Gd)₁ eutectic compounds. After solution treatment at 510°C for 4 h, the eutectic compound dissolved into the matrix completely. A good combination of tensile strength and ductility was obtained when the alloy was aged at 200°C for 32 h; the ultimate tensile strength, yield strength and elongation are 282 MPa, 185 MPa and 6·1 respectively. Precipitation strengthening is the largest contributor to strength (∼67) in 200°C peak aged alloy, which is related to plate shaped precipitate phases formed on prismatic planes of the α-Mg matrix.

Keywords

Mg–Sm–Gd–Zn–Zr alloy Microstructure Heat treatment Mechanical property

Introduction

Magnesium alloy is the lightest of all metals used as the basis for constructional alloys.¹ Among the applied Mg alloys, alloys containing rare earth (RE) elements are known to offer excellent mechanical properties by solid solution strengthening and precipitation strengthening, such as WE (Mg–Y–Nd–Zr) and QE (Mg–Ag–Nd–Zr) type alloy systems.^2–4 The precipitation sequence in WE series alloys has been reported to be Mg_ssss→β″(D019)→β′(bct)→β₁(fcc)→β (fcc) during aging at 250°C.^{5

5,6}

Rokhlin reported that the strengthening in Mg alloys can be promoted effectively by alloying two RE elements belonging to different subgroups (cerium and yttrium ones).⁷ Recently, much of the investigations have been performed on Mg–Gd based alloys with other RE elements (i.e. Y, Nd, Dy and Sc).^8–15 Samarium is one of the light RE elements belonging to the cerium group, with maximum solubility in solid Mg of 5·8 wt- at eutectic temperature 803 K. There were only a few researches on Sm containing Mg alloy. Zheng^{16

16,17} researched the microstructure, mechanical properties and creep behaviour of Mg–3Sm–0·5Zn–0·4Zr (wt-) alloy; the ultimate tensile strength (UTS), yield strength (YS) and elongation ϵ are 189 MPa, 94 MPa and 11·7 respectively. Li et al.¹⁸ researched the effects of heat treatments on the microstructure and mechanical properties of Mg–Y–Sm–Zr alloy. Zhang et al.¹⁹ researched phases in as cast and solution treated Mg–6Gd–4Sm–0·4Zr (wt-) alloy using transmission electron microscopy (TEM). The intermetallic phase in the as cast microstructure had a face centred cubic crystal structure (a = 2·2879 nm) with a composition of Mg_6·2(Sm_0·56Gd_0·44) and was dissolved after solution treatment. The UTS of the T6 Mg–6Gd–4Sm–0·4Zr (wt-) alloy is up to 305·05 MPa, and the ϵ is 5·11.²⁰ Furthermore, Zn is one of the most important elements that can further improve the strengthening response of Mg–RE alloys.^21–25 Until now, there are no systemic research on Mg–Sm–Gd–Zn–Zr alloy. In the present work, the effects of heat treatment on the microstructure and mechanical properties of Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr alloy have been systematically investigated.

Experimental

The experimental Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr (wt-) alloy was prepared from 99·9 mass-Mg, Mg–33Gd, Mg–31Sm and Mg–30Zr (wt-) master alloys in an electrical resistance furnace in steel crucible under protective gas consisting of SF₆ (1 vol.-) and CO₂ (bal.) in order to prevent burning of the melts. The melting alloys were homogenised at ∼780°C for 15 min and then cast into a preheated steel mould at ∼720°C, and the size of the ingot is 200×150×20 mm. The chemical composition of the obtained alloy ingot was analysed using inductively coupled plasma atomic emission spectrometry. Solution treatment was carried out at 510 or 530°C under CO₂ atmosphere for different times and then quenched in water at 70–80°C. Aging heat treatment was performed at 175, 200 and 225°C in an air electric resistance furnace for different times (1, 2, 4, 8, 12, 16, 32, 64, 128 and 256 h).

Tensile tests were carried out using an MTS (858) tensile testing machine at a crosshead speed of 1 mm min⁻¹ at room temperature according to ASTM standard E8-2004. The longitudinal extensometer MTS with gage length of 10 mm was used in the tests. Samples for tensile test had a rectangular cross-section (2 mm thick and 3·5 mm wide) and a gauge length of 15 mm. The constituent phases were identified by X-ray diffraction (XRD) (Rigaku D/max 2500 XRD) with Cu K_α radiation at 36 kV. Vickers hardness (HV) tests were taken using a 200 g load and holding time of 15 s. Microstructure examination was performed using a Leica optical microscope and a Sirion200 field emission scanning electron microscope (SEM) equipped with an X-ray energy dispersive spectrometer (EDS) to determine the local chemical compositions. Specimens for optical microscopy were prepared by a conventional mechanical polishing technique. The specimens were etched in a solution of 6 g picric acid, 40 mL acetic acid, 40 mL water and 100 mL ethanol. The grain size was determined on the photographs using lineal intercept method outlined in accordance with ASTM standard E112-96. The TEM observation was carried out using a Tecnai G2 at an accelerating voltage of 200 kV. The samples for TEM observation were 3 mm discs punched directly from the heat treated strips, ground to a thickness of 0·1 mm and twinjet electropolished in a solution of 5·3 g lithium chloride, 11·1 g magnesium perchlorate, 500 mL methanol and 100 mL 2-butoxy-ehanol at −40°C and 30 mA.

Results

Microstructure

The microstructure of the as cast alloy is shown in Fig. 1. It reveals that the as cast state alloy is mainly composed of α-Mg solid solution and eutectic compounds at grain boundaries. The average grain size is ∼45 μm. Typical TEM bright field (BF) image and corresponding selected area electron diffraction (SAED) pattern for eutectic phase are shown in Fig. 2. The SAED pattern of the eutectic phase (B = [0 0 1]) indicates that it has a face centred cubic (fcc) structure with a = 0·727 nm. Chemical analyses by EDS in TEM mode reveal that the average composition of the eutectic compound is Mg–(21±1·0)Sm–(3·67±1·0)Gd–(6·24±1·0)Zn (at-); therefore, the eutectic compound can be described as (Mg,Zn)₃(Sm,Gd)₁. A similar eutectic phase structure was also found in Mg–Gd–Y–Zn–Zr alloy; it is an (Mg,Zn)₃RE type compound (DO3, a = 0·7283 nm).^{26

26,27}

Figure 3 shows the differential thermal analysis result of the as cast Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr (wt-) alloy during the heating process. The first endothermic peak appears at ∼565°C, and this peak may correspond to the melting temperature of the eutectic phase. The second peak appears at ∼639°C, and this peak may be related to the melting point of the alloy.

Solution treatment temperatures are determined based on differential thermal analysis trace of the as cast alloy, with 510 and 530°C selected. Figure 4 indicates the influence of solution heat treatments on the microstructure of the alloy. The average grain size of the alloy solutionised at 510°C for 2, 4 and 8 h is 50, 55 and 82 μm respectively. The grain size of the alloy solutionised at 530°C for 4 h is 93 μm. There are still some eutectic compounds remaining at the grain boundaries after being solutionised at 510°C for 2 h, which is shown in Fig. 4a. The eutectic compounds almost dissolve into the matrix, and the average grain size is just a little larger than that of the as cast alloy after being solutionised at 510°C for 4 h. After 510°C for 4 h solution treatment, the eutectic compounds have almost dissolved into the matrix, as illustrated by XRD pattern (Fig. 5). The grain size became extremely large after solutionised at 530°C for 4 h.

Age hardening response

Figure 6 compares the hardness curves of the alloy during aging at 175, 200 and 225°C. It is found that specimens aged at different temperatures exhibit a similar age hardening process. At the initial stage of aging, the hardness increases with increasing aging time before it reaches the peak hardness. After reaching the peak hardness, the hardness gradually decreases as a result of overaging. The specimen aged at 175°C exhibits the highest hardness of 104 HV after aging for 256 h, which is too long to be acceptable for industry application. The specimen aged at 200°C takes 64 h to reach the highest hardness of 94 HV. Furthermore, the alloy shows a relatively stable hardness from 32 to 128 h. However, comparing with the specimens aged at 175 and 200°C, the specimen aged at 225°C shows the lowest peak hardness of 87 HV for 32 h and the most rapid decrease in hardness. Obviously, the specimen aged at 200°C for 32–64 h has a higher hardness with shorter time and then is taken as the optimum aging process.

Figure 7 shows a TEM BF image and the corresponding SAED pattern from the alloy aged at 175°C for 256 h, with the incident electron beam approximately parallel to _α . In this peak aged alloy, plate-like precipitates have an average size of ∼20 nm in length along [0001] _α of reciprocal space and several nanometres in thickness. In the corresponding SAED pattern (Fig. 7b), there are weak diffraction spots at 1/2 _α in the _α , which suggest that the diffraction feature is a typical D019⁵ structure. Therefore, the precipitates are identified to be β″ phase, with lattice parameters a = 2a_Mg = 0·64 nm and c = c_Mg = 0·52 nm and an orientation relationship of _β _″|| _α , _β _″|| _α .

The BF TEM image and SEAD pattern of the alloy aged at 200°C for 32 h observed in the direction parallel to _α are shown in Fig. 8. A large number of ellipsoidal precipitates within the a-Mg matrix in this peak aged alloy are observed in Fig. 8a. In the diffractions at 1/4 , 2/4 and 3/4 in the _α SAED pattern (Fig. 8b), such spots can be indexed as base centred orthorhombic structure of β″ with unit cell a = 0·64 nm, b = 2·223 nm and c = 0·521 nm.^{28

28,29} An orientation relationship with the matrix in the present paper is || , || .

Figure 9 shows the TEM image and corresponding SAED pattern of the matrix in the alloy after aging at 225°C for 32 h; plate-like precipitates of β phase (Mg₅RE, fcc, a = 2·27 nm) are observed in this peak aged condition. The other precipitate is β′ phase, which become larger in size and less dense than that in the 200°C peak aged condition. In Fig. 9b, there are invisible diffraction spots along ; these diffractions are attributed to the β′ phase. The β phase's diffraction marked in Fig. 9b showed that the orientation relationship is _α || _β , _α || _β , _α || _β .

Mechanical properties

The tensile properties including UTS, YS and ϵ of different states are shown in Fig. 10. It shows that this alloy exhibits superior tensile strength after solution and aging treatment. After solution treatment, the alloy has higher YS, UTS and ϵ than that of the as cast alloy. However, ϵ decreases after aging treatment. Obviously, heat treatment has a positive influence on the mechanical properties. The alloy shows the highest tensile strength after peak aged at 200°C for 32 h among the different peak aged conditions at 175, 200 and 225°C; the UTS, YS and ϵ are 282 MPa, 185 MPa and 6·1 respectively.

Fracture

The representative fracture features of specimens are shown in Fig. 11. In the solution treated specimen, the failure surface is composed of cleavage planes and tear ridges (Fig. 11a), and the abundant tear ridges are consistent with the good ϵ. Therefore, it indicates that the fracture mode of the solution treated specimen is quasi-cleavage fracture. After peak aged at 175°C, the amount of tear ridges greatly reduced and cleavage facets became smooth (Fig. 11b), which correspond to the decrease in ductility, and it exhibited a transgranular cleavage fracture. When peak aged at 200°C, the fracture mode of the specimen partly changed to intergranular fracture (Fig. 11c). Therefore, the fracture mode of the specimen is mixed transgranular and intergranular cleavage. In 225°C peak aged specimen (Fig. 11d) , the fracture surface shows lots of tear ridges, which correspond to local plastic deformation and contribute to ductility improvement.

Discussion

In the Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr alloy, the microstructural changes have a strong influence on the mechanical properties. Based on the analysis of the microstructure and the results of the mechanical properties, the strengthening mechanism involves five factors, i.e. pure Mg σ_Mg, secondary phase (eutectic compounds) strengthening σ_sp, solid solution strengthening σ_ss, grain boundary strengthening σ_gb and precipitation strengthening σ_ppt.

The as cast alloy consists of α magnesium matrix and eutectic (Mg,Zn)₃RE compounds. The eutectic compounds distributed along the grain boundary result in the solid solution atom in the α-Mg matrix in a very low content, and so the solid solution strengthening effect can be ignored. It is obvious to find that no precipitation strengthening exists in the as cast alloy. Therefore, in the as cast alloy, only pure Mg, second phase strengthening σ_sp and grain boundary strengthening σ_gb contribute to the strength. The YS of pure Mg is 21 MPa, and the alloy with the addition of Zr (∼0·7 wt-) is elevated to be 55 MPa by zirconium refinement.³⁰ The sum of all the contributions should be equal to the YS. Thus, the grain boundary strengthening contribution can be concluded to be 55–21 = 34 MPa. The alloy has YS (140 MPa) in the as cast state. As discussed above, the second phase strengthening contribution is about 140–55 = 85 MPa, reaching 61 of the total YS.

After solution treatment, the eutectic compounds dissolve into the matrix completely. In the solution treated alloy, besides pure Mg, only solid solution strengthening σ_ss and grain boundary strengthening σ_gb contribute to the strength. The alloy has YS (145 MPa) in the solution treated state. The mechanical properties of the alloy vary with grain size, and the relationship usually follows the Hall–Petch equation; thus, we can estimate that σ_gb is still ∼34 MPa because the grains do not coarsen during solution treatment. We calculate the solid solution strengthening contribution to 145–55 = 90 MPa, which is ∼62 of the total YS.

In the peak aged specimen, the secondary phase has almost dissolved into the matrix, and dispersive precipitates form inside the grains. According to the Mg–Sm phase diagram, the solubility of Sm in solid magnesium is reduced from 5·8 wt- at the eutectic temperature to 0·4 wt- at 200°C. Therefore, the solid solution strengthening can be ignored in 200°C peak aged condition, and only the precipitation and grain boundary should be considered as the main factors of preventing dislocation motion. The grain boundary strengthening contribution in peak aged alloy is the same as that in the as cast alloy since the grain size does not coarsen during aging. As the YS in the peak aged state is 182 MPa, then the precipitation strengthening contribution is about 182–55 = 127 MPa (70 of the total YS).

Based on the estimations to all strengthening contributions, the results are summarised in Table 1. In the solution treated state alloy, ϵ is improved because the compounds along the grain boundary dissolve into the matrix; these compounds may become the microcrack sources during the tensile test. Therefore, it is necessary to optimise the solution treatment parameters to dissolve the eutectic compounds completely, maximise the solution strengthening and improve the ductility.

Table 1

Strengthening contributions in Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr alloy at room temperature

Strengthening contribution	As cast		Solution treated		Peak aged at 200°C
Strengthening contribution	MPa		MPa		MPa
Pure Mg strength	21	15	21	14·5	21	11·4
Second phase strengthening	85	60·7	0	…	0	…
Solid solution strengthening	0	…	90	62·1	0	…
Grain boundary strengthening	34	24·3	34	23·4	34	18·4
Precipitation strengthening	0	…	0	…	127	70·2
Experimental yield strength	140	…	145	…	185	…

From Table 1, it is obvious that precipitation strengthening is the largest contributor to the strength in the peak aged alloy. The precipitation sequence of the alloy can be described as Mg_ssss→β″(D019)→β′(bct)→β (fcc). Different aging treatments have different influences on the microstructure of alloy, which can affect the mechanical properties of the alloy because of the different number density, distribution, size and kinds of precipitate.³¹ The most important factor to influence the mechanical properties of the alloy is aging temperature. It is known that the solid solubility of the RE element in Mg decreases along with the temperature dropping. At lower temperature aging, the alloy could precipitate more volume fraction of the second phase β″, which has D019 crystal structure and is coherent with the matrix. Such precipitates could increase the obstruction stress of dislocation motion and the maximum stress of cleavage fracture, which increase the YS and UTS of the alloy. When the aging temperature gets higher, the β″ phase will transform to β′ phase, and the β′ phase get coarsen with the prolonging of aging time. As the main strengthening phase in 200°C peak aged alloy, the β′ phase, which forms on the prismatic planes of matrix in a dense triangular arrangement, is vertical to the basal plane of α-Mg and provides the most effective obstacle to basal dislocation slip in the Mg alloy.¹⁰ Thus, the alloy in this condition has better mechanical properties than the 175°C peak aged alloy. It is indicated that the alloy has inferior mechanical properties aged at 225°C because the volume fraction and density number of precipitates are reduced.

The precipitation strengthening is attributable to the plate shaped precipitate phases formed on prismatic planes of the α-Mg matrix, which can prevent basal slipping of dislocations in hexagonal close packed Mg metals. The mechanical properties of the alloy greatly increased after the solution plus aging heat treatment. In this investigated alloy, we believe that the improved mechanical property is mainly ascribed to precipitation strengthening.

Conclusions

The microstructure and mechanical properties of Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr (wt-) alloy have been investigated. The Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr alloy in the as cast state is mainly composed of α-Mg matrix and (Mg,Zn)₃(Sm,Gd)₁ eutectic compounds. The optimal solution treatment is 510°C for 4 h, the eutectic compounds dissolve into the matrix completely and ϵ is greatly increased.

Solution and aging treatments have a positive influence on the mechanical properties of this alloy. The UTS, YS and ϵ in 200°C peak aged alloy are 282 MPa, 185 MPa and 6·1 respectively. Precipitation strengthening is the largest contributor to strength (∼67) in the peak aged alloy, which is related to the plate shaped precipitate phases formed on prismatic planes of the α-Mg matrix.

Footnotes

Acknowledgement

The authors gratefully acknowledge the financial support of the National key Basic Research and Development Program of China (grant no. 2005CB623705) for funding the work reported.

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Effects of heat treatment on microstructure and mechanical properties of Mg–2·6Sm–1·3Gd–0·6Zn–0·5Zr alloy

Abstract

Keywords

Introduction

Experimental

Results

Microstructure

Age hardening response

Mechanical properties

Fracture

Discussion

Conclusions

Footnotes

Acknowledgement

References