Journal of Materials Science & Technology  2019 , 35 (10): 2365-2374 https://doi.org/10.1016/j.jmst.2019.05.053

Orginal Article

Development of extruded Mg-6Er-3Y-1.5Zn-0.4Mn (wt.%) alloy with high strength at elevated temperature

Mingquan Zhanga, Yan Fenga, Jinghuai Zhanga*, Shujuan Liub, Qiang Yangc, Zhuang Liud, Rongguang Lie, Jian Mengc, Ruizhi Wua

a Key Laboratory of Superlight Material and Surface Technology, Ministry of Education, College of Material Science and Chemical Engineering, Harbin Engineering University, Harbin 150001, China
b Department of Materials Physics and Chemistry, Harbin Institute of Technology, Harbin 150001, China
c State Key Laboratory of Rare Earth Resources Utilization, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Changchun 130022, China
d Faculty of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, China
e School of Mechanical Engineering, Shenyang University of Chemical Technology, Shenyang 110142, China

Corresponding authors:   *Corresponding author.E-mail address: jinghuaizhang@gmail.com (J. Zhang).

Received: 2019-03-16

Revised:  2019-05-3

Accepted:  2019-05-6

Online:  2019-10-05

Copyright:  2019 Editorial board of Journal of Materials Science & Technology Copyright reserved, Editorial board of Journal of Materials Science & Technology

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Abstract

A new Mg-6Er-3Y-1.5Zn-0.4 Mn (wt.%) alloy with high strength at high temperature was designed and extruded at 350 °C. The as-extruded alloy exhibits ultimate tensile strength of 301 MPa, yield strength (along ED) of 274 MPa and thermal conductivity of 73 W/m·K at 300 °C. Such outstanding high-temperature strength is mainly attributed to the formation of nano-spaced solute-segregated basal plane stacking faults (SFs) with a large aspect ratio throughout the entire Mg matrix, fine dynamically recrystallized (DRXed) grains of 1-2 μm and strongly textured un-DRXed grains with numerous sub-structures. Microstructural examination unveils that long period stacking ordered (LPSO) phases are formed in Mg matrix of the as-cast alloy when rational design of alloy composition was employed, i.e. (Er + Y): Zn = 3: 1 and Er: Y = 1: 1 (at.%). It is worth mentioning that it is the first report regarding the formation of nano-spaced basal plane SFs throughout both DRXed and un-DRXed grains in as-extruded alloy with well-designed compositions and processing parameters. The results provide new opportunities to the development of deformed Mg alloys with satisfactory mechanical performance for high-temperature services.

Keywords: Mg alloys ; Stacking faults ; Mechanical properties ; Thermal conductivity

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Mingquan Zhang, Yan Feng, Jinghuai Zhang, Shujuan Liu, Qiang Yang, Zhuang Liu, Rongguang Li, Jian Meng, Ruizhi Wu. Development of extruded Mg-6Er-3Y-1.5Zn-0.4Mn (wt.%) alloy with high strength at elevated temperature[J]. Journal of Materials Science & Technology, 2019, 35(10): 2365-2374 https://doi.org/10.1016/j.jmst.2019.05.053

1. Introduction

Magnesium (Mg) alloys, as the lightest metallic structural materials, offer significant potential for improving weight reduction and energy efficiency in automobile, aerospace and military fields [[1], [2], [3], [4]]. With the rapid development of various industries, higher requirements are put forward for the performances of Mg alloys such as high-temperature mechanical properties and heat-conducting properties, which can ensure the mechanical stability and uniform temperature distribution to prolong the service life of alloy components in high temperature conditions [5]. However, at present, the relevant properties of widely used commercial Mg alloys such as AZ and AM series are unsatisfactory, especially high-temperature strength [6,7].

Addition of rare earth (RE) to forming Mg-RE or Mg-RE-Zn alloys has been recognized as an effective way to improve the strength at both room and elevated temperatures [[8], [9], [10], [11], [12]]. At present, some room temperature (RT) high-strength or even ultrahigh-strength Mg alloys have been fabricated by deformation processing or deformation processing + aging treatment. For example, the ultimate tensile strength (UTS) and yield strength (YS) of Mg-10Gd-5.7Y-1.6Zn-0.7 Zr (wt.%) alloy can reach to 542 MPa and 473 MPa respectively fabricated via hot extrusion and aging [13]. The high strength of these alloys is mainly attributed to the formation of nano-metastable phases and/or long period stacking ordered (LPSO) phases, and the morphology of basal plane LPSO phases has great influence on the improvement in alloy performance [9,14,15]. A few studies also reported the high temperature mechanical properties of these Mg-RE or Mg-RE-Zn alloys. Yang et al. [16] reported that the UTS and YS of hot-extruded Mg-8Gd-3Yb-1.2Zn-0.5 Zr (wt.%) alloy are approximately 425 MPa and 413 MPa at RT, and 204 MPa and 191 MPa at 300 °C, respectively. Meng et al. [17] reported that the UTS and YS of as-extruded Mg-7Y-4Gd-1.5Zn-0.4 Zr are about 331 MPa and 228 MPa at RT, and 230 MPa and 123 MPa at 300 °C, respectively. It can be found that compared with traditional Mg alloys, these alloys have high RT strength and high-temperature strength, but their strength still decreases significantly at elevated temperatures.

On the other hand, thermal conductivity is also very important for the high-temperature applications of Mg alloys [18]. Generally speaking, for thermal conductivity, the best is pure metal (pure Mg) with a perfect lattice structure without any defects. The more defects, the lower the thermal conductivity, and on the contrary, the current alloy strengthening is essentially defect strengthening (solid solution, secondary phase, grain boundary, etc.), that is, the more defects, the higher the strength. At present, the high thermal conductivity Mg alloys such as Mg-Zn [19], Mg-Mn [20] series are generally low-alloying (alloying elements also have relatively low solid solubility) alloys, which can reduce the degree of defects, and thus the strengthening effect is also inadequate due to the almost inversely strengthening and thermal conductivity mechanisms.

Based on the literature review, it can be found that the current research and development of new Mg alloys mostly emphasizes the single strength or thermal conductivity. In this study, we designed a new extruded Mg-6Er-3Y-1.5Zn-0.4 Mn (wt.%) alloy having excellent high temperature strength even at 300 °C and suitable thermal conductivity. The alloy design is (Er + Y): Zn = 3: 1 and Er: Y = 1: 1 (at.%) in order to form special strengthening structure, i.e. fine basal plane stacking faults (SFs) or LPSO phases, according to the principal found from the previous experimental and theoretical calculation results [21,22]. But it should be noted that LPSO plates are formed in the most previous reported Mg-RE-Zn alloys, while numerous fine basal plane SFs are formed in our alloy through reasonable heat treatment and deformation process combined with the alloy composition design. The addition of a small amount of Mn is mainly considered to be beneficial to the corrosion resistance [23]. The alloy composition can also be written as Mg-0.9Er-0.9Y-0.6Zn-0.2 Mn (at.%). This paper reports the microstructure, tensile properties and thermal conductivity of the new alloy. The results make up a basis for the development of high-temperature high performance deformed Mg alloys.

2. Experimental procedures

The cast ingot of studied alloy Mg-6Er-3Y-1.5Zn-0.4 Mn (wt.%) was prepared by direct-chill semi-continuous casting method. High purity Mg (99.98 wt.%), Zn (99.95 wt.%) and Mg-20Er (wt.%), Mg-20Y (wt.%), Mg-10 Mn (wt.%) master alloys were melted in a mild steel crucible at 750 °C under the protection of Ar atmosphere. The melt was casted into the cooling crystallizer at $\widetilde{7}$ 10 °C. The semi-continuous casting process was carried out with the casting speed of about 100 mm/min. Finally, a cylindric ingot with the length of about 600 mm and diameter of 100 mm was fabricated successfully. The chemical composition of the obtained ingot was examined by an inductivity coupled plasma atomic emission spectroscope (ICP-AES) and the result was listed in Table 1. The differential scanning calorimetry (DSC) test was done with the heating rate of 5°/min using a Netzsch STA449F3 apparatus. The ingots were solid-solution treated at 545 °C for 12 h in a vacuum heat treating furnace, and then quenched into warm water of $\widetilde{9}$ 0 °C. Before extrusion, the ingots were machined into billets with diameter of 80 mm. After preheated at 350 °C for 2 h, the billet was extruded into the sheet with width of 50 mm and thickness of 8 mm at the same temperature under a ram speed of 0.1 mm/s and an extrusion ratio of 13.

Table 1   Chemical composition of the investigated alloy (wt.%).

AlloyErYZnMnMg
Mg-6Er-3Y-1.5Zn-0.4 Mn5.722.931.690.43Bal.

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The microstructure, phase structure and composition were characterized by optical microscopy (OM), scanning electron microscopy (SEM) equipped with an X-ray energy-dispersive spectrometry (EDS) system and a NordlysMax3 electron backscattered diffraction (EBSD) system, X-ray diffraction (XRD) and transmission electron microscopy (TEM, FEI Tecnai G20). Specimens for SEM characterization were etched by a mixture of 1 mL acetic acid, 0.42 g picric acid, 1 mL H2O and 7 mL ethanol. Thin foils for TEM observation were prepared by argon ion thinning technique (PIPS, Gatan 691). The tensile specimens were cut from the sheet with the tensile direction parallel to extrusion direction. Tensile tests were performed using an Instron 5869 testing machine with an initial strain rate of 1 × 10-3 s-1 at RT and high temperatures. Before each high temperature test, a 15 min holding was applied to balance the testing temperature. The tensile data were an average value of at least four tensile specimens.

The samples for thermal diffusivity measurement were cut from the as-cast and as-extruded alloys in the shape of disks with a diameter of 12.7 mm and a thickness of 2 mm. For the as-extruded sheet, the test plane is the extrusion direction-transverse direction (ED-TD) plane and the thermal diffusivity test direction is the normal direction (ND). The thermal diffusivity was measured on a Netzsch LFA427 apparatus via the laser flash method in the temperature range from 20 °C to 500 °C. The room temperature density was obtained by Archimedes method. The density at elevated temperature was calculated using the following equation [24]:

ρ = ρ0-0.156(T-298)

where ρ0 (g/cm3) is the density at room temperature, and T (K) is the absolute temperature. Thermal conductivity (λ) was calculated via the following equation:

λ=α·ρ·CP

where α (m2/s) is the thermal diffusivity, ρ (g/cm3) is the density, and Cp (J/(g·K)) is the specific heat capacity. The Cp at each temperature can be measured by NetzschSTA449F3 apparatus.

3. Results and discussion

3.1. Microstructure of as-cast Mg-6Er-3Y-1.5Zn-0.4 Mn alloy

Fig. 1 shows OM and SEM micrographs of as-cast Mg alloy specimens. It reveals a relatively uniform microstructure, which is composed of α-Mg dendrites and intermetallic phases along grain boundary or interdendritic regions. The dendritic arm spacing (DAS) is about 30-60 μm measured by linear intercept method. According to the corresponding SEM image (Fig. 1(b)), the secondary phase presents lamellar morphology, which is the typical morphological feature of LPSO phase [25]. Fig. 2 presents XRD patterns of the studied alloy in as-cast, solid-solution and as-extruded states. The results suggest that the alloy in as-cast state is mainly composed of Mg matrix phase and Mg12Y1Zn1-type phase, which is known as LPSO phase [26,27].

Fig. 1.   (a) OM and (b) SEM micrographs of the as-cast alloy.

Fig. 2.   XRD patterns of (a) as-cast, (b) solid-solution treated and (c) as-extruded samples.

To further confirm the LPSO phase, phase morphology and structure were characterized by TEM, as shown in Fig. 3. In terms of morphology, they can be roughly divided into two types. One presents the fine lamellar morphology (Fig. 3(a)), and the corresponding selected area electron diffraction (SAED) pattern shows the extra diffraction spots at ±n/7(0001)α (n is an integer), which indicates it has a 14H LPSO structure (an ordered hexagonal structure, a =1.112 nm, c =3.647 nm) [26,28]. The other presents relatively coarse plate-like shape (Fig. 3(c)), and the corresponding SAED pattern shows the extra diffraction spots at ±n/3(0001)α, which suggests it has a 18R LPSO structure according to the previous studies (an ordered base-centered monoclinic structure, a =1.112 nm, b =1.926 nm, c =4.689 nm, and β = 83.25°) [26,28]. All the characterization results indicate that based on the reasonable design of alloy composition, i.e. (Er + Y): Zn = 3: 1 and Er: Y = 1: 1 (at.%), only secondary phases with LPSO structure are formed in Mg matrix of the new designed Mg-0.9Er-0.9Y-0.6Zn-0.2 Mn (at.%) alloy in the as-cast state. The result is consistent with the reported experimental and theoretical calculation law for formation of LPSO phase in Mg-RE-Zn alloys [21,22].

Fig. 3.   TEM images and corresponding SAED patterns of the LPSO phases: (a) and (b) 14H LPSO phase; (c) and (d) 18R LPSO phase.

3.2. Microstructure of solid-solution treated Mg-6Er-3Y-1.5Zn-0.4 Mn alloy

According to the experimental design, in order to obtain numerous nano-spaced basal plane SFs through extrusion, as-cast alloy should be solid-solution treated. Therefore, DSC test was carried out to select the appropriate solid-solution temperature, and then numerous careful solid-solution experiments were conducted. The final solid-solution temperature and time were determined through microscopic observation. Fig. 4(a) shows the DSC heating curve of as-cast alloy. It is evident that its eutectic temperature is about 552 °C, and the starting temperature is about 547 °C. Based on the DSC result, solid-solution treatments at 530 °C, 540 °C and 545 °C with different time were designed as the pre-experiment. Finally, the optimum solid-solution condition is found to be 545 °C and 12 h. XRD pattern (Fig. 2) shows that the peaks of LPSO phase almost disappear after the solid-solution treatment. Fig. 4b shows the microstructure of the alloy solid-solution treated at 545 °C for 12 h. It can be seen that except for very few of them almost all the coarse LPSO plates are dissolved into Mg matrix as compared to the microstructure of as-cast alloy (Fig. 1(a) and (b)), and the residual LPSO compounds are already changed to the relatively fine particles.

Fig. 4.   (a) DSC heating curve of the as-cast alloy and (b) SEM micrograph of the solid-solution treated alloy.

3.3. Microstructure of solid solution treated Mg-6Er-3Y-1.5Zn-0.4 Mn alloy after extrusion at 350 °C

Fig. 5 shows SEM images of solid solution treated Mg-6Er-3Y-1.5Zn-0.4 Mn alloy after extrusion at 350 °C. The as-extruded sample exhibits the typical bimodal microstructure composed of fine equiaxed dynamic recrystallized (DRXed) grains and coarse deformed un-DRXed grains elongated along ED, mainly due to the insufficient deformation. In addition, almost no other secondary phases are found by SEM observation. This typical bimodal microstructure has been reported in deformed Mg alloys, including Mg-RE or Mg-RE-Zn based alloys [[29], [30], [31], [32], [33]]. However, it is worth noting that the most key features of the alloy designed in this study are not the bimodal microstructure, but numerous fine lamellar structures in both fine DRXed grains and deformed un-DRXed grains. These features were first time reported in extruded Mg alloys. As shown in Fig. 5, these lamellae are parallel in a grain and have different orientations in different DRXed grains, and they are parallel and roughly along ED in un-DRXed grain. XRD pattern cannot identify these lamellae in the as-extruded alloy (Fig. 2). Anyway, the microstructure, especially the fine lamellae structure in grains, requires further examination.

Fig. 5.   (a) SEM image of the as-extruded alloy and (b) high-magnified SEM image showing the typical microstructure of DRXed region.

Fig. 6(a) shows the inverse pole figure (IPF) map of as-extruded sample, and Fig. 7(a) and (b) presents the IPF maps of DRXed grains and un-DRXed grains, respectively, further indicating the typical microstructure consisting of both the fine DRXed grains with relatively random orientations and the coarse deformed un-DRXed grains with certain orientation roughly (most of them are blue grains with <01-10> ||ED i.e. <0001 > ⊥ED). In addition, the color gradient in one deformed un-DRXed grain suggests that there still exists low angle misorientation in the grain. The average size of fine DRXed grains is just about 1.88 μm (Fig. 6(d)). The volume fractions of DRXed grains and un-DRXed grains are about 51% and 49%, respectively. Fig. 6(b) shows the grain boundary (GB) map of as-extruded sample. It can be seen that numerous low angle grain boundaries (LAGBs) (green lines) form in the deformed un-DRXed grains and subdivide them into smaller sub-structures, which agrees with color gradient of deformed grains in IPF map (Figs. 6(a) and 7 (b)). In addition, a few DRXed grains can be observed near LAGBs within deformed grains, as marked by arrow in Fig. 6(b). Thus it can be deduced that during hot extrusion process in this work, dislocations introduced by deformation in original grains gradually develop into the LAGBs, and the rapid transformation of LAGBs to HAGBs lead to the formation of fine DRXed grains, which is called as continuous DRX (C-DRX) mechanism [34,35], while mainly due to the insufficient deformation, some original grains with low energy storage merely develop into the deformed grains with sub-structures in them. Fig. 6(e) shows the quantitative analysis of the misorientation angle distributions. The result also reveals that there are relatively high number fraction ($\widetilde{2}$5%) of LAGBs. Fig. 6(c) shows the basal slip Schmid factor distribution map of as-extruded sample. On the whole, it can be seen that the coarse un-DRXed grains present the negligible basal slip Schmid factors, while the fine DRXed grains exhibit various basal slip Schmid factors, and the values of most DRXed grains are higher than those of un-DRXed grains. Fig. 6(f) shows the basal slip Schmid factor distributions, which have the trend to distribute towards the lower values, mainly due to the relatively high volume fraction of un-DRXed grains.

Fig. 6.   EBSD analysis of the as-extruded sample: (a) IPF map with the reference direction parallel to ED. The horizontal direction is parallel to ED; (b) grain boundary map; (c) basal slip Schmid factor distribution map along ED; (d) grain size distributions; (e) misorientation angle distributions; (f) basal slip Schmid factor distributions.

Fig. 7.   IPF maps of (a) DRXed grains and (b) un-DRXed grains with the reference direction parallel to ED. The horizontal direction is parallel to ED. (0001) pole figures of (c) whole region, (d) un-DRXed region and (e) DRXed region of the as-extruded sample.

Fig. 7(c-e) shows the (0001) pole figures (PFs) of whole region, un-DRXed region and DRXed region of the as-extruded sample, respectively. The (0001) PF of whole region reveals that the as-extruded alloy has a strong basal texture with quite high texture intensity of 24.3, and the basal pole is inclined at 30°-41° toward the TD. The (0001) PF of un-DRXed grains shows the very same texture component but higher texture intensity of 33.6, while the (0001) PF of DRXed grains shows very weak basal texture (texture intensity: 5.0) with large spread, indicating that the strong basal texture of this as-extruded alloy is mainly from the deformed un-DRXed grains. On the other hand, it confirms that the DRXed Mg-RE extruded alloy generally has weak basal texture.

Mainly to further study the lamellae formed in both DRXed and un-DRXed grains, the micro morphology and structure were characterized by TEM. Fig. 8 mainly presents the TEM analysis of un-DRXed region with electron beam (EB) parallel to [11 $\bar{2}$0]Mg, showing some features which cannot be found by SEM. Obviously, the dense and fine lamellae with the space between each other from just $\widetilde{2}$0 nm to $\widetilde{1}$00 nm are formed in the whole un-DRXed grain (Fig. 8(a-c)). Furthermore, it is noted that these lamellae are not straight, and microstrain and even kinking are formed on the lamellae (Fig. 8(a) and (b)). It can be inferred from careful TEM observation combined with EBSD analysis (Fig. 6(a) and (b)) that the low angle kinking (such as the position indicated by the arrow in Fig. 8(a)) of lamellae corresponds to the formation of sub-structure (i.e. LAGB) within the deformed un-DRXed grains, and the high angle kinking of lamellae are formed following with the DRXed grains (corresponding to HAGB). The deep relationships and interactions between formation of lamellae and dynamic recrystallization need further study. The corresponding selected area election diffraction (SAED) pattern having clear streaks between the diffraction spots along c-axis in Fig. 8(d) confirms the formation of basal plane stacking faults (SFs) [36,37]. Careful SAED analysis indicates almost all of these fine lamellae are basal plane SFs and very few of them have 18R-LPSO structure.

Fig. 8.   TEM analysis of the un-DRXed region: (a)-(c) bright field (BF)-TEM images; (d) SAED pattern.

Fig. 9(a) shows BF-TEM image of DRXed region. The fine DRXed grains with size about 1-2 μm and the fine lamellae within them with space similar to that of SFs ($\widetilde{2}$0-100 nm) in un-DRXed grains are formed during hot extrusion. In addition, very few 18R-LPSO laths with width of about 0.2 μm and length of about 0.7 μm are also found in DRXed region by careful TEM observation (Fig. 9(b) and (c)). It is considered that they are originated from the residual LPSO phase in the solid-solution treated alloy and broken into fine particles after hot extrusion. Due to the low content, they have little effect on the properties of alloy. Fig. 10 shows TEM analysis of a representative DRXed grain with EB // [11 $\bar{2}$ 0]Mg. Nano-spaced fine lamellae are almost uniformly distributed throughout the whole DRXed grain. Careful SAED and TEM-Fast Fourier Transform (FFT) analyses indicate that these fine lamellae are basal plane SFs.

Fig. 9.   TEM analysis of the DRX-ed region: (a) and (b) BF-TEM image of the DRXed grains; (c) SAED pattern of the lath-shaped phase.

Fig. 10.   TEM analysis of a typical DRXed grain: (a) and (b) BF-TEM image; (c) SAED pattern; (d) high-resolution-TEM image with FFT pattern.

On the whole, we have fabricated a new extruded Mg-RE-Zn alloy with special microstructure for the first time, which is obviously different from previous reported deformed Mg alloys. This alloy exhibits a bimodal microstructure, that is, about one half is the fine DRXed region with average grain size below 2 μm and the other half is the deformed un-DRXed region with numerous sub-structures. Most importantly, the nano-spaced fine basal plane SFs are almost uniformly distributed throughout both the whole DRXed and un-DRXed grains, which is the most significant difference between our alloy and those reported by other researchers. The formation of SFs and their solute-segregation have been confirmed and reported in Mg-RE-Zn alloys [38]. However, the SFs usually co-exist with other reinforcing phases (such as plate-shaped LPSO phase) in the reported Mg-RE-Zn alloys [30,38], that is, such single high number density of SFs strengthened bimodal-microstructure has not yet been reported. In this study, the formation of this novel microstructure is related to the combination effect of appropriate alloy composition design and preparation processing parameters (i.e. solid solution treatment and extrusion process). Selecting RE (Er and Y) and Zn can lower stacking fault energy and stabilize SFs or LPSO structure [36,39,40]. Controlling the ratio of RE to Zn and the appropriate processing parameters ensure the formation of profuse fine SFs rather than plate-shaped LPSO phase.

3.4. Mechanical properties

Fig. 11 shows the typical tensile stress-strain curves of the as-extruded alloy tested at RT and elevated temperatures. The average tensile properties including UTS, YS and elongation to failure (ε) are listed in Table 2. This alloy exhibits excellent tensile properties from RT to 300 °C. Both the UTS (354 MPa) and YS (316 MPa) at RT are above 300 MPa and the ε at RT can be above 8%. More importantly, the strength decreases only slightly with the increased temperature from RT to 300 °C, and the UTS can still remain at about 301 MPa and the YS keeps above 274 MPa when the temperature reaches to 300 °C. Such excellent high temperature performance is very rare in the reported deformed Mg alloys.

Fig. 11.   Typical tensile stress-strain curves of the studied alloys tested at different temperatures.

Table 2   Average tensile properties including UTS, YS and ε of the investigated alloy.

Temperature (°C)UTS (MPa)YS (MPa)ε (%)
25354 ± 5.8316 ± 4.28.1 ± 1.1
200336 ± 3.7298 ± 2.911.7 ± 1.6
250326 ± 2.1285 ± 3.314.5 ± 1.3
300301 ± 2.6274 ± 2.817.3 ± 2.5

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Here we consider the microstructure-tensile strength at 300 °C relationship. Since the tensile temperature used in this work is high (300 °C), the deformation for the Mg alloy should be related not only to the basal dislocation slip, but also to the non-basal dislocation slips such as prismatic slip and pyramidal slip due to the reducing critical resolved shear stress (CRSS) with increasing temperature, as well as grain boundary sliding and lattice rotation [41,42]. In addition, it has been reported that alloying with RE (such as Y) reduces the c/a ratio along with the increase in a, and raises the CRSS for basal slip but lowers the CRSS for non-basal slip [43,44]. Based on the detailed microstructure observation, it is considered that the superior high temperature strength at 300 °C is mainly related to the facts as follows. Firstly, the fine DRXed grains of 1-2 μm in size contribute to the high YS via grain boundary strengthening based on the Hall-Petch relationship [45]. Secondly, the coarse deformed un-DRXed grains present very low basal slip Schmid factors due to their strong basal texture and thus they are sufficiently improve the YS by providing a hard orientation of the basal dislocation slip during tensile test along the ED [25,46]. In addition, the sub-structures (numerous LAGBs and local strains) within the un-DRXed regions can also serve as barriers to the movement of dislocations, leading to the strengthening during tensile test. Lastly and also most importantly, the formation of high number density basal plane SFs with large aspect ratio in both DRXed and un-DRXed grains plays an important role in strengthening the alloy, especially at high temperature of 300 °C. It is considered that in Mg-RE based alloys β′ or β′′ precipitates with certain aspect ratio formed on prismatic planes of α-Mg matrix during aging are vertical to basal plane of α-Mg, providing the most effective obstacles to basal dislocation slip [47], while in the studied alloy, since the solute-segregated SFs are formed on the basal plane of α-Mg matrix, they provide relatively weak barriers to basal dislocation slip as compared to β′ or β′′ prismatic precipitates. However, the nano-spaced SFs formed on the basal plane of α-Mg matrix can give more sufficient impediment to non-basal dislocation slips compared with the prismatic precipitates and further improve the strength. Furthermore, the nano-spaced solute-segregated SFs with extremely large aspect ratio have the coherent interface with α-Mg matrix and high thermal stability [36,38,39], and thus they not only have an effective pinning effect on the dislocation movement in grains at high temperature, but also can behave as a skeleton for both the DRXed and un-DRXed grains to stabilize grains and inhibit the grain boundary sliding and lattice rotation, so as to contribute to the superior elevated temperature tensile strength at 300 °C.

To intuitively compare the strength of extruded Mg-6Er-3Y-1.5Zn-0.4 Mn alloy with those of typical high-strength heat-resistant extruded Mg alloys reported previously, their UTS and YS at 300 °C are shown in Fig. 12. By comparison, it is found that the high temperature strength of our newly developed alloy is superior to those of the alloys reported previously. In addition, its YS at 300 °C is much higher than that (below 150 MPa) of the commercial high-temperature WE43 alloy [48]. All these reported Mg alloys are Mg-RE based alloys and most of them are Mg-RE-Zn alloys, that is, the reported heat-resistant Mg alloys and our new alloy belong to the same alloy system. These reported alloys are mainly strengthened by plate-shaped basal plane LPSO phases or/and nano-scale prismatic metastable precipitates. Therefore, the superior high temperature strength of the new alloy should be related to the formation of special microstructure. It is considered that the nano-spaced fine basal plane SFs contribute much more than the plate-shaped basal plane LPSO phases to the YS due to the more effective pinning [29], and although the dense basal plane SFs with extremely large aspect ratio are weaker than nano-scale prismatic metastable precipitates in inhibiting basal dislocation slip, they play a more effective role in impeding non-basal dislocation slip and stabilizing the grain structure at high temperatures.

Fig. 12.   Histogram of the high temperature strength (300 °C) comparison in the studied alloy and previously reported Mg-RE based alloys with comparative compositions [[52], [53], [54], [55], [56]].

3.5. Thermal conductivity

Fig. 13 shows the temperature dependences of thermal diffusivity and thermal conductivity of the as-cast and as-extruded alloys. It can be found that both the thermal diffusivity and the thermal conductivity of as-extruded alloy is lower than those of corresponding as-cast alloy. On the whole, the lower thermal conductivity of as-extruded alloy is due to the more defects induced after extrusion, such as the grain boundaries and SFs, which can behave as scattering sources of electrons as well as phonons. In addition, the thermal diffusivity and thermal conductivity for as-cast and as-extruded alloys have the same changing trend with the increasing temperature. The thermal conductivity increases remarkably with temperature increased from RT to 500 °C whether the alloy is in as-cast state or as-extruded state. The thermal conductivity for as-extruded alloy at RT is 49 W/m·K, while it increases to 73 W/m·K at 300 °C, which is a relatively high value for the high-strength Mg-RE alloy [49,50].

Fig. 13.   Thermal diffusivity and thermal conductivity of the as-cast and as-extruded alloys.

For Mg alloys with free electrons as the main heat carrying entities, the electrons would be impeded by atomic thermal motion and various lattice defects, thus forming the thermal resistances, i.e. phonon thermal resistance and defect thermal resistance. With the temperature increasing, the vibration of lattice intensifies which means phonon thermal resistance increases, while the defect scattering of electrons is generally elastic which means the defect thermal resistance is inversely proportional to the temperature [21,51]. As for pure Mg or Mg alloys with low alloying, defects are relatively few, and phonon thermal resistance plays a major role, thus thermal conductivity decreases with the increase of temperature [24]. However, with regard to the as-extruded Mg-6Er-3Y-1.5Zn-0.4 Mn alloy, its alloying degree is relatively high. Numerous lattice defects are caused mainly by the solute-segregated SFs, and the defect thermal resistance would take the dominant position from RT to 500 °C. Therefore, the thermal conductivity of the studied alloy increases with the increase of temperature. There have been no other reports about thermal conductivity of Mg alloys with profuse SFs, and the detailed law and mechanism remain to be further investigated.

4. Conclusions

In the present work, a new extruded Mg-6Er-3Y-1.5Zn-0.4 Mn (wt.%) alloy was fabricated, and its microstructure, mechanical properties and thermal conductivity were investigated. The principal conclusions are summarized as follows.

(1) By reasonable design of alloy composition, i.e. (Er + Y): Zn = 3: 1 and Er: Y = 1: 1 (at.%), only long period stacking ordered phase can be formed in the Mg matrix of the as-cast Mg-0.9Er-0.9Y-0.6Zn-0.2 Mn (at.%) alloy.

(2) Based on the combination effect of appropriate alloy composition design and preparation processing parameters (i.e. mainly the thorough solid solution treatment and proper extrusion temperature), only nano-spaced basal plane SFs formed in both DRXed and un-DRXed grains of the as-extruded alloy can be realized.

(3) This as-extruded alloy exhibits superior tensile strength along ED at 300 °C, and the average UTS is 301 MPa and YS is 274 MPa. The excellent high-temperature strength is mainly attributed to the formation of (a) the nano-spaced solute-segregated basal plane SFs with large aspect ratio in the whole Mg matrix, (b) the fine DRXed grains with sizes of 1-2 μm and (c) the strongly textured un-DRXed with numerous sub-structures.

(4) The thermal conductivity of the studied alloy increases remarkably with temperature increased from RT to 500 °C, and the as-extruded alloy with high strength at high temperature of 300 °C exhibits decent thermal conductivity (73 W/m·K) at 300 °C.

Acknowledgements

This work was supported by National Natural Science Foundation of China (No. 51871069), Natural Science Foundation of Heilongjiang Province of China (E2017030), Fundamental Research Funds for the Central Universities (3072019CF1004), and Foundation of State Key Laboratory of Rare Earth Resources Utilization (No. RERU2018017).


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