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Article

Effects of Ce-Rich Misch Metal on the Microstructures and Tensile Properties of as-Cast Mg-7Al-3Sn-1Zn Alloys

by
Guo-Jun Liu
1,*,
Yan-Hua Sun
2,
Nan Xia
1 and
Xiao-Fang Guan
1
1
School of Materials Science and Engineering, Nanling Campus, Jilin University, No. 5988 Renmin Street, Changchun 130025, China
2
Kunming Metallurgical Research Institute Co., Ltd., Kunming 650031, China
*
Author to whom correspondence should be addressed.
Metals 2021, 11(10), 1648; https://doi.org/10.3390/met11101648
Submission received: 18 September 2021 / Revised: 8 October 2021 / Accepted: 11 October 2021 / Published: 18 October 2021
(This article belongs to the Section Metal Casting, Forming and Heat Treatment)
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Abstract

:
The effects of small amounts of Ce-rich misch metal (Mm: 0.5, 1.0 and 2.0 wt.%) addition on the microstructure and tensile properties of as-cast Mg-7Al-3Sn-1Zn wt.% (ATZ731) alloy have been investigated. The addition of Mm restricts the formation of the Mg17Al12 phase but greatly promotes the Al4Mm phase. The proper Mm addition enhances the strength and ductility of ATZ731 alloys at both room temperature (RT) and 175 °C. ATZ731 alloys with 1.0 wt.% Mm addition exhibit an advantageous combination strength and ductility, with the ultimate tensile strength (UTS), 0.2% yield strength (YS) and elongation to failure (Ef) at 175 °C of ~148 MPa, ~102 MPa and ~28%, improved by ~14.7%, ~24.3% and ~53.8%, respectively, compared to those of ATZ731 alloy. This enhancement is primarily owing to the refined microstructures and the high thermal stability of Al4Mm at the elevated temperature in contrast with that of the Mg17Al12 phase. The fracture modes are also discussed.

1. Introduction

Mg alloys have attracted increasing attention for automobile, subway, and biomedical applications due to their high specific strength, ease of recycling and degradability [1,2]. Mg-Al serials alloys currently lead the market of commercial casting Mg alloys; however, their wide application as metal structure materials is still limited because of their weak strength at temperatures above 120 °C due to the softening β-Mg17Al12 phase with a eutectic temperature of 437 °C [3,4]. Therefore, further study on Mg-Al alloy is needed to enlarge its application.
Recently, Mg-Al-Sn (AT) alloys have been showing promise. The addition of Sn to Mg alloy has many advantages [5,6,7,8,9,10,11]. Sn causes solid solution strengthening (solubility of Sn in Mg is 14.85 wt.% at the eutectic temperature of 561 °C), precipitation strengthening, and creep performance at elevated temperatures in binary Mg-Sn alloys. It also provides a beneficial effect on the elongation via decreasing of stacking fault energy of pure Mg proved by first principle calculation [8,9,10,11]. However, the resultant mechanical properties of Mg-Sn binary alloys cannot satisfy the commercial requirements, stimulating the fast development of AT based Mg alloys [12]. There are many papers on AT alloys with low Al contents, but researches with Al contents above 6 wt.% are few [5,8,13,14,15,16]. Luo et al. [17] investigated AT alloys with varying additions of Al (5–9 wt.%) and Sn (1–5 wt.%) prepared by permanent mould, in which the optimised AT72 alloy had enhanced strength and ductility in contrast with the commercial AZ91 alloy. The volume fraction of Mg2Sn and Mg17Al12 are 0.33% and 6.46%, respectively. Meanwhile, the coarse eutectic phases along the grain boundaries decrease the ductility of AT alloys with the increasing addition of Al and Sn. To modify these coarse microstructures, Zn was introduced into Mg-Al-Sn alloys [18,19]. Moreover, the addition of Zn refined microstructures (grain size and secondary phase) attributed to the super strength and ductility compared to Mg-Al-Sn alloys [18,19,20]. Wang and Pan confirmed that 0.5–2.0 wt.% Zn addition refined the microstructures of AT82 alloys contributing to the enhanced ductility [18,19]. Liu et al. [20] reported that adding 3 wt.% Zn decreased the solid solubility of both Sn and Al in Mg-9Al-6Sn alloy and produced fine-scale Mg2Sn and Mg17Al12 particles. Moreover, Zn addition remarkably refined the Mg2Sn phase in as-cast Mg-Sn alloys and changed the orientation relationships between Mg2Sn and α-Mg under peak-aged conditions [21,22]. Based on the development of ATZ alloys and considering the mechanical properties of ATZ alloys at high temperatures, the content of Sn should be appropriately increased, while that of Al should be somewhat reduced in the designed alloys referring to AT82 alloy. Therefore, Mg-7Al-3Sn-1Zn wt.% (ATZ731) alloy was designed as the based alloy in this work, which is expected to have the desirable tensile property. However, the inherent properties of Mg17Al12 with a low melting point limit the strength at an elevated temperature. Hence, it is crucial to reduce the volume fraction of Mg17Al12, which is simultaneously replaced by the high thermal stability of Al-containing phases in ATZ731 alloy, to enhance their high-temperature strength [23,24,25].
To achieve this goal for ATZ731 alloys, alloying might be a feasible method to modify the microstructures and properties. Ce-rich misch metal (Mm), as one kind of rare-earth (RE) element, was often used in Mg and Al alloys and altered the corrosion resistance, formality, and mechanical properties, as well as the microstructures of these alloys [23,24,25,26,27]. Since Mm is cheaper than pure RE elements, it is preferred in many industries. Moreover, alloying element Al will react with the added Mm to form the high-melting Al-Mm phases in Mg-Al alloys [25]. Therefore, Mm may be a suitable candidate for ATZ Mg alloys to qualify the microstructures and mechanical properties. When minor Mm was introduced into Mg-Al alloys, the thermally stable AlxMmy intermetallic compounds were formed with the decreasing fraction of Mg17Al12, resulting in superior strength at high temperatures [26]. The values of x and y in AlxMmy greatly depend on the fabrication processes and the mass ratio of Al/Mm. Generally, there are three intermetallic compounds, such as Al11Mm3, Al4Mm, and Al2Mm, that usually appear in Mm-containing Mg-Al alloys [25]. Employed 1.25 wt.% Mm in AZ91 alloy refined the grain size and produced rod-like Al4(Ce, La) compound [28], whereas 0.3–0.6 wt.% Mm addition caused a decrease of the mean diameters of both α-Mg and Mg17Al12 precipitations in AZ91 alloy [27]. The addition of 0.44 wt.% Mm improved the strength and ductility of die-cast AZ91D alloy at 170 °C and attributed to the high thermal stability Al-Mm phases, which effectively hindered the grain boundaries glide and cracks propagation [29]. Another paper reported that the simultaneous additions of 0.7 wt.% Mn and Mm refined grains and produced the new need-like Al4Mm phase in wrought AZ61 alloy, which conduces to the nucleation of dynamic recrystallisation at elevated temperatures [30]. Note that, 0.9 wt.% (0.4 Y combined with 0.5 Nd) [23], 1.5 wt.% Mm [31] and 0.44 wt.% Mm [29] is required to the optimal mechanical properties for AT42, AZ81 and AZ91D alloy, respectively. However, there is a contrary opinion that a small addition of Mm (0.2–1.2 wt.%) combined with Mn led to the grain coarsening in AZ31, AZ61 and AZ91 alloy prepared by sand casting [32]. The mentioned papers mainly focused on the microstructures and tensile properties of Mm-adding Mg-Al-Zn alloys, while the studies on the Mm-adding ATZ alloys are very few. The tensile properties of as-cast ATZ alloys at high temperatures still are not reported. Most importantly, the appropriate addition of Mm is very different for alloys with various compositions to gain optimal performance.
Thus, this study investigates the effect of the small addition of cheap Mm on the microstructures and tensile properties of as-cast Mg-7Al-3Sn-1Zn wt.% alloy. The fracture mode was discussed with fracture morphologies.

2. Materials and Methods

Pure Mg, Al, Sn, and Zn (>99.9 wt.%), and Mg-23.2 wt.% Ce-rich misch metal (Mm: Ce 65.8, La 26.52, Nd 5.82, and Pr 1.78 wt.%) master alloy were used to fabricate Mg-7Al-3Sn-1Zn-xMm (x = 0, 0.5, 1.0, and 2.0 wt.%) alloys. These alloys were marked as ATZ731, ATZE7310, ATZE7311, and ATZE7312 (Table 1), respectively, to narrate conveniently in the following. The alloys were melted by an electric-resistance furnace protected under CO2 and SF6 followed by pouring into a preheated (~200 °C) grey iron mould at ~700 °C.
The actual components of as-cast alloys were checked by an X-ray fluorescence spectrometer (XRF-1800, Shimadzu Sequential, Kyoto, Japan) with the analysis results listed in Table 1. The phase composition was analysed by X-ray diffraction (XRD; DX-2700B Cu Kα radiation at 40 kV and 30 mA) (HaoYuan, Dandong, China). Microstructures were observed by optical microscopy (OM; Carl Zeiss Axio Imager.A2m, Göttingen, Germany) and scanning electron microscopy (SEM; ZEISS VOA-18, Göttingen, Germany) with an energy dispersive spectrometer (EDS) to investigate the distribution of alloying elements. Tensile properties were evaluated by a material testing system (Instron5869, Instron, Norwood, USA) at a strain rate of 1.0 × 10−3 s−1 at RT and 175 °C for at least three samples. To keep good repeatability, at least three tensile samples were tested, and average mechanical properties were recorded. The samples were heated at 175 °C for 8 min prior to testing at this temperature. Phase evolution of ~20 mg samples was recorded by a differential scanning calorimeter (DSC, SDT-Q600, Lindon, UT, USA) with Pt crucible and the reference crucible at heating rates of 10 °C/min ranging from 675 to 300 °C under the protection of high-purity Ar atmosphere. The grain size of the solution-treated samples was measured by Nanomeaure 3.0 software [33]. The average volume fraction of secondary phases in as-cast alloys was calculated via the examination for at least ten SEM images by the software of Axio Vision version 4.8 (ZEISS VOA-18, Göttingen, Germany).
For metallographic observation, the as-cast specimens cut from the same place of the ingots were etched with a mixed solution of nitric acid (2 mL) and absolute ethanol (48 mL) for 10~15 s after grinding with 600, 1000, and 2000 mesh SiC sandpaper in turn and polishing by 0.5 μm diamond abrasion paste followed by cleaning in absolute ethanol. The tensile samples, as dogbone-like, are nominal size of the width 4 mm, gage length 30 mm and thickness 2.0 mm. The solid solution treatment of samples was heat-treated first at 420 °C for 20 h, then at 1 °C/min increased up to 480 °C holding for 2 h followed by water (~70 °C) quenching. The DSC samples cut from the ingots were discs with ~0.5 mm thickness and 3 mm diameter.

3. Results and Discussion

3.1. Phases Composition and Microstructures

Figure 1 shows the XRD patterns of as-cast samples. ATZ731 alloys clearly consist of α-Mg, Mg17Al12, and Mg2Sn phases (Figure 1a) similar to those in Mg-Al-Sn-(Zn) alloys [34]. Small amounts of Mm addition caused the new diffraction peaks (Figure 1b–e), which are identified to be Al4Mm (most of Al4Ce) [35]. Moreover, the strength of Al4Mm intensity increased with Mm increasing. For the ATZE7312 alloy, the intensity of Al4Mm became strong, while that of Mg17Al12 at ~36° changed to weak compared with ATZ731 alloy, indicating that the increasing Mm promotes the formation of Al4Mm but suppresses the Mg17Al12 phase. From the evolution of Mg2Sn diffraction peaks, it seems that the Mm addition has a little effect on Mg2Sn may be related to the detection limit of XRD. The formation of the Al4Mm phase was explained by the different electro-negativity between Mm and Mg, Al, Sn, Zn, which suggests that Mm easily reacts with Al [23,35]. This new phase is Al4Mm, but not the Al11Mm3 phase, which appeared in Mg-Al-Mm alloy where the Mm/Al mass ratio is above 1.4 as recommended by Pettersen [36]. Su et al. [35] investigated the Al-Ce intermetallic compounds in as-cast Mg-Al-Zn-Ce alloys by thermodynamic calculation and the experiments in which the alloys were prepared by gravity casting. It was found that Al4Ce formed because of the low atomic ratio between Ce and Al as well as the dominant kinetic conditions, which is more likely to satisfy the nucleation and growth requirements of the Al4Ce compound. None of the Zn-containing phase was detected, possibly due to the low content of Zn in samples and the relative high solid solution of Zn in Mg (6.2 wt.% in eutectic temperature) according to binary Mg-Zn phase diagram associated with the other papers [19,20]. Besides, the melting temperatures of Al4Mm are higher than that of Mg17Al12 (i.e., Al4Ce of 1276 °C, Mg17Al12 of 437 °C) [35]. Thus, the thermally stable Al4Mm phase can strongly affect the tensile properties of alloys at a high temperature.
Figure 2 presents the typical SEM morphologies of as-cast samples, in which Figure 2a–d are the low magnify of ATZ731, ATZE7310, ATZE7311, and ATZE7312, and Figure 2e–h represents the local high magnify of (Figure 2a–d), respectively. Furthermore, EDS results are shown in Figure 3 and listed in Table 2. For ATZ731 alloy, masses of grey semi-continuous compounds correspond to eutectic β-Mg17Al12 phase distributed along the grain boundaries as pointed by blue arrows, while the white strip-like phases adhered to β phases belong to divorced Mg2Sn as pointed by green arrows (Figure 2a,b) based on XRD results (Figure 1) and EDS results (Figure 3). This indicates that the phase composition of ATZ731 alloy is very similar to those in Mg-7Al-2Sn alloy and Mg-8Al-2Sn-1Zn alloys [18,19]. When 0.5 wt.% Mm was added in ATZ731 alloy, the new phase of γ-Al4Mm, small size of β and Mg2Sn, and few lamella secondary precipitates of β (β′) phase can be observed compared to ATZ731 alloy. The γ phase has three different morphologies as rod-, particle- and feather-like, while the volume fraction of the feather-like is very few. Moreover, both the quantity and size of β phase were reduced in contrast with the based alloy. With further increase of Mm, the volume fraction and size of γ phase were obviously increased (Figure 2c,d), especially for feather-like γ phase segregating within grains and at the grain boundaries. While the volume fraction of β was reduced, and its morphology was changed from the coarse semi-continuous-like to block- or particle-like, indicating that the Mm addition promotes the output of γ and restricts the formation of β phase. The size of Mg2Sn clearly was reduced with the Mm increasing, possibly related to the refined β phase. It is interesting that ATZE7311 alloy showed the maximum volume fraction β’ phase pointed by white arrows among these samples. In addition, the EDS analysis of point A showed that the atomic ratio values of Al to Mm is 3.86 very near four. This further proved these phases were Al4Mm well agreed with XRD results [37]. The concentration of divorced eutectic Mg2Sn phase seems relatively stable because of the non-appearance of Sn-containing ternary phase (Figure 3), unlike that in the Mg-Sn-RE (Ce, Y, Nd) systems, where MgSnCe/Y/Nd formed decreased the fraction of Mg2Sn [17,38]. Thus, the increasing addition of Mm not only promotes the crystallisation and growth of Al4Mm but suppresses the Mg17Al12 phase. In addition, the volume fraction of secondary phases was carefully calculated with the average values are 7.1 ± 0.2, 8.6 ± 0.26, 10.8 ± 0.32, and 12.2 ± 0.37% for ATZ731, ATZ7310, ATZE7311, and ATZE7312 alloy, respectively. It indicates that the volume fraction of secondary phases increased with Mm increasing.
To evaluate the grain size of α-Mg for as-cast alloys, Figure 4 represents the images of the solution-treated samples. The grain size dramatically decreased from 276 ± 15 μm of ATZ731 to 190 ± 11 μm of ATZE7310, 124 ± 5 μm of ATZE7311 and 156 ± 7 μm of ATZE7312, indicating that Mm effectively refined the grain size of ATZ731 alloy (Figure 4a–d). However, the grain size of ATZE7312 further increased to ~156 μm (Figure 4d), revealing that these refinement effects became weak in ATZ731 alloy with the excessive Mm addition. This phenomenon was often observed in Mg alloys with excessive RE addition related to RE-containing phases with the high melting temperature and the possible segregation of RE [24,39]. For example, the addition of Y above 0.9 wt.% in AZ91 alloy caused the large grain size of α-Mg [27]. Besides, a large number of the undissolved Al4Mm phases can be observed on the Mm-containing samples.
The refined microstructures of ATZ731 alloy after minor Mm addition can be explained in the following. According to the solidification theory of metals, the grain size of α-Mg is mainly dependent on the degree of heterogeneous nucleation and the level of undercooling during the solidification process [40]. γ-Al4Mm phase was formed prior to α-Mg, Mg2Sn, and Mg17Al12 in Mg-Al-Sn-Zn-Mm alloys because of its high eutectic temperature during solidification. Xiao [24] suggested that the preferentially formed Al4Mm particles cannot act as the nucleation sites of α-Mg, meaning that the heterogeneous nucleation is a determinant of the refined α-Mg in Mm-added alloys [24]. Consequently, the preferential partition behaviour of ATZ731 alloys caused by the added Mm during solidification has to be considered [29,34]. Since the dissolved atoms of alloying elements were pushed forward at the α-Mg growth front leading to a great constitutional supercooling near the growth front during solidification, which contributed to the refined α-Mg (Figure 4). In the case of the β-Mg17Al12 phase in Mm-containing alloys (Figure 2 and Figure 3), its volume fraction, as well as its size, was reduced owing to the formation of Al4Mm particles as previously mentioned. As for the divorced eutectic Mg2Sn particle, it nucleated and grew to coexist with the refined β phases, and then the Mg2Sn phase was also refined, as observed in Figure 2 and Figure 3. With respect to the γ phase, its volume fraction and size rose upon the addition of Mm. The increased volume fraction of γ phase lowered not only the concentration of Mm in front of the solid/liquid interface but also the constitutional supercooling degree in the growth front of the alloy melt, causing the weak refinement effect on the microstructure as observed in Figure 2, Figure 3 and Figure 4. Additionally, supposing that 2.0 wt.% Mm completely reacted with Al in the alloys, Al still was excessive compared to Mm by calculating the atom ratio between Al and Mm, indicating that β phase should be formed during solidification even considering the relatively high solubility of Al in Mg at the eutectic temperature of binary Mg-Al alloys.
To investigate the effect of Mm contents on the solidification behaviour of ATZ731, Figure 5 shows the typical DSC curves of ATZE7310 and ATZE7311 alloys compared to that of ATZ731 alloy. All the curves exhibited two exothermic peaks, which corresponded to the transformations of α-Mg and β phase, respectively. The addition of Mm greatly decreased the initial temperatures of β phase transformation but raised the end temperature of α-Mg transformation from 587 °C of ATZ731 to 594 °C of ATZE7311. This suggests that the solidification rate of ATZE7310 and ATZE7311 alloys were faster than the based alloy, contributing to the refined α-Mg.

3.2. Tensile Properties and Fracture Morphologies

Figure 6 presents the engineering stress-strain curves of as-cast samples at RT and 175 °C. The average values of UTS, YS, and Ef are listed in Table 3. It can be seen that small amounts of Mm remarkably improved the strength and elongation to failure of ATZ731 alloy at RT and 175 °C. With Mm increasing, the average UTS, YS and Ef of samples first increased and then slightly decreased in contrast with the based alloys. Among these specimens, ATZE7311 alloy exhibited the desirable combination of strength and ductility under tested conditions with the UTS, YS and Ef of ~239 MPa, ~108 Mpa, and ~19.6% at RT, and they are ~148 MPa, ~102 Mpa, and ~28% at 175 °C, improved by ~9.6%, ~24.3%, and ~53.8%, respectively, compared to ATZ731 alloy. It needs to be mentioned that the excessive Mm addition dropped the strength and ductility compared to ATZE7311 alloy, which was similar to the RE-containing Mg alloys, such as Y-adding AZ91 and Mm-adding ZK60 alloys [38,41,42].
Generally, the fracture models of Mg alloys are a mixture of cleavage and quasi-cleavage characters due to the limited plastic manners of {001}<110> dislocation slip and {102}<101> twin for Mg alloys at RT [43]. Figure 7 shows the typical fracture morphologies of specimens. The fractures of ATZ731 alloy showed few tear ridges and cleavage facets (Figure 7a), indicating that the major fractures model at RT is cleavage fracture. For ATZ731 alloys with Mm addition, few shallow dimples, tear ridges, and cleavage facets were observed on the fractures (Figure 7b,c,e), exhibiting the cleavage and quasi-cleavage characters, which usually occurred on Mg alloys. In the case of ATZE7312 alloys tested at 175 °C, the cleavage facets, tear ridges, and the relative deep dimples (Figure 7d,f) showed that quasi-cleavage dominated the fracture model. Moreover, large amounts of fine particles are at the bottom of the dimples (Figure 7f), as expected to be beneficial to the tensile properties in contrast with that of ATZ731. Since the increasing deformation temperature of 175°C promoted the basal slip and activated the non-basal slip of Mg alloys, however, the ductility of ATZE7312 was slightly weak compared to ATZ731, possibly related to the non-uniform deformation caused by the progressive accumulation of Al4Mm phases during tensile tests [43].
The appropriate Mm addition enhanced the strength and ductility of as-cast ATZ731 alloys at RT and 175 °C ascribed to the refined microstructures. The solid solution strengthening effects of ATZ731 alloy induced by Mm addition is limited because the solid solubility of Mm is very low in α-Mg (~0.74 wt.% Ce in α-Mg) compared to those of the high solid solubility of Al, Zn, and Sn [29]. It is well known that the refined α-Mg can greatly improve strength and ductility, especially for the YS. According to the formula of Hall-Petch, YS is proportional to a square root of grain size and Taylor’s coefficient [44]. As a result, the refined α-Mg ranging from 276 μm of ATZ731 to 124 μm of ATZE7311 progressively enhanced the YS regarding the large value of Taylor’s coefficient of Mg.
For those as-cast Mg alloys, the morphology, size, and distribution of eutectic phases also greatly affected the strength and ductility by restricting the dislocation movement within grains and grain boundaries glide [44]. During the tensile deformation of samples, the dislocations were easily piled up near the grain boundaries, where plenty of hard and brittle eutectic phases existed at the interface between α-Mg and eutectic phases. Once the density of dislocation is beyond a certain extent, it will cause the local stress concentration inducing the micro-cracks near these brittle phases and finally leading to the tensile failure (Figure 7). Due to this, for as-cast ATZ731 alloy during tensile deformation, cracks easily occur and develop along the grain boundaries in which mass of semi-continuous Mg17Al12 phases are distributed (Figure 2a and Figure 7a), resulting in low strength and poor ductility. As for ATZ731 alloy modified by Mm addition, there are two main factors that determine the tensile properties. One is the reduced grain size of α-Mg, which will enhance the tensile properties; the other is the eutectic phase, including volume fraction, morphology and distribution. The latter has a great effect on strength and ductility. As for ATZE7311 alloy, the homogeneously distributed refined eutectic phases (Figure 2c and Figure 3) and the finest α-Mg play enhanced roles in high strength and good ductility. With respect to ATZ731 alloy with 2.0 wt.% Mm addition, the refined α-Mg is beneficial to the strength and ductility (Figure 4d). On the contrary, the increased volume fraction of eutectic phases associated with more Al4Mm phase highly enriched at local regions (Figure 2h and Figure 7f), such as grain boundaries, can improve the UTS but may lower the YS and Ef [34]. Once the weakening effects induced by the locally enriched eutectic phases is far stronger than the enhancing roles afforded by the refined grains, a slight decrease of average YS at RT and Ef at 175 °C possibly occur. Additionally, it has been suggested that the concentration of defects in the solidified microstructures also decrease the strength and ductility in as-cast Mg alloys. Small amounts of Mm addition raised the transformation temperature of α-Mg, causing the relatively high cooling rates and reduction of defects in ATZ731 alloy during solidification. Fewer defects in the Mm-adding alloys were also helpful to improve their strength and ductility. Thus, the appropriate Mm addition could effectively improve strength and ductility owing to the solid solution strengthening, refinement strengthening, and the secondary phase strengthening [45,46,47].

4. Conclusions

The microstructures and tensile properties of Mg-7Al-3Sn-1Zn (ATZ731) alloys with 0–2.0 wt.% Ce-rich misch metal (Mm) were studied. The results were as follows:
(1) The as-cast ATZ731 alloy consisted of α-Mg, Mg17Al12 and Mg2Sn phases. Minor addition of Mm promoted Al4Mm but restricted Mg17Al12 phase.
(2) Minor addition of Mm refined grain size and eutectic phases of ATZ731. With the increasing addition of Mm, the Mg17Al12 phase became from semi-continuous to block- and particle-like along the grain boundaries, while Al4Mm phases changed from rod-like to feather-like distributed at grain boundaries and within grains.
(3) The appropriate Mm addition enhanced the strength and ductility of ATZ731 alloy at RT and 175 °C. ATZ731 alloy with 1.0 wt.% Mm addition exhibited the high strength and good ductility with the UTS, YS and Ef are ~239 MPa, ~108 MPa and ~19.6% at RT and ~148 MPa, ~102 MPa and ~28% at 175°C, respectively. Compared to ATZ731, they are improved by ~14.7%, ~24.3% and ~53.8% at 175 °C, mainly owing to the refined microstructures. The fractures of alloys exhibited the mixture characters of cleavage and quasi-cleavage, indicating that Mm addition did not change the fracture mode.

Author Contributions

Conceptualisation, G.-J.L. and Y.-H.S.; methodology, G.-J.L., N.X. and X.-F.G.; validation, G.-J.L., Y.-H.S., N.X. and X.-F.G.; investigation, G.-J.L., N.X. and X.-F.G.; writing—original draft preparation, G.-J.L.; writing—review and editing, G.-J.L., and N.X.; project administration, G.-J.L.; funding acquisition, G.-J.L. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by the National Natural Science Foundation of China (No. 51301074) and the Research Foundation of Education Bureau of Jilin Province, China (No.JJKH20211082KJ).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Conflicts of Interest

The authors declare no conflict of interest.

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Metals 11 01648 g001 550
Figure 1. XRD patterns of as-cast samples for (a) ATZ731, (b) ATZE7310, (c) ATZE7311 and (d) ATZE7312 alloys.
Figure 1. XRD patterns of as-cast samples for (a) ATZ731, (b) ATZE7310, (c) ATZE7311 and (d) ATZE7312 alloys.
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Metals 11 01648 g002 550
Figure 2. The typical SEM morphologies of as-cast samples for (a,e) ATZ731, (b,f) ATZE7310, (b,c,f,g) ATZE7311 and (d,h) ATZE7312. Note that (e–h) are the high magnification BSD of local (a–d) as pointed by the dotted line square respectively. β represents Mg17Al12 phase, while γ is Al4Mm phase.
Figure 2. The typical SEM morphologies of as-cast samples for (a,e) ATZ731, (b,f) ATZE7310, (b,c,f,g) ATZE7311 and (d,h) ATZE7312. Note that (e–h) are the high magnification BSD of local (a–d) as pointed by the dotted line square respectively. β represents Mg17Al12 phase, while γ is Al4Mm phase.
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Figure 3. SEM morphology of as-cast ATZE7311 alloy and the EDS results. (a) surface morphology, and (b–i) are the corresponding distribution of Mg, Al, Sn, Zn, Ce, La, Nd and Pr on (a), respectively. (j,k) are the point EDS results of the points of A and B in (a).
Figure 3. SEM morphology of as-cast ATZE7311 alloy and the EDS results. (a) surface morphology, and (b–i) are the corresponding distribution of Mg, Al, Sn, Zn, Ce, La, Nd and Pr on (a), respectively. (j,k) are the point EDS results of the points of A and B in (a).
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Figure 4. Optical images of solution-treated samples for (a) ATZ731, (b) ATZE7310, (c) ATZE7311 and (d) ATZE7312. The insets are the corresponding statistics results of grain size with dave represents the average values of grain size.
Figure 4. Optical images of solution-treated samples for (a) ATZ731, (b) ATZE7310, (c) ATZE7311 and (d) ATZE7312. The insets are the corresponding statistics results of grain size with dave represents the average values of grain size.
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Figure 5. The typical DSC curves of as-cast samples for ATZ731, ATZE7310 and ATZE7311.
Figure 5. The typical DSC curves of as-cast samples for ATZ731, ATZE7310 and ATZE7311.
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Figure 6. Engineering stress-strain curves of as-cast samples at (a) RT and (b) 175 °C.
Figure 6. Engineering stress-strain curves of as-cast samples at (a) RT and (b) 175 °C.
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Figure 7. The typical tensile fracture morphologies of as-cast samples tested at RT and 175 °C. (a) ATZ731 at RT; (b) ATZE7311-RT; (c) ATZE7312-RT; (d) ATZE7312-175 °C; (e) and (f) are the high magnify BSD of local (c) and (d), respectively, as pointed by the dotted-line square.
Figure 7. The typical tensile fracture morphologies of as-cast samples tested at RT and 175 °C. (a) ATZ731 at RT; (b) ATZE7311-RT; (c) ATZE7312-RT; (d) ATZE7312-175 °C; (e) and (f) are the high magnify BSD of local (c) and (d), respectively, as pointed by the dotted-line square.
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Table
Table 1. Chemical composition of the experiment alloys.
Table 1. Chemical composition of the experiment alloys.
Alloy Composition (wt.%)
AlSnZnCeLaNdPrTotal (Mm)Mg
ATZ7317.4532.9671.105-----Bal.
ATZE73107.3783.1921.0550.2260.0910.0200.0060.343Bal.
ATZE73117.2173.1881.1900.5180.2090.0460.0140.786Bal.
ATZE73127.2183.1171.0901.1880.4780.1050.0321.803Bal.
Table
Table 2. Elementary compositions of the points showed in Figure 3 and the compositions of the phases.
Table 2. Elementary compositions of the points showed in Figure 3 and the compositions of the phases.
NumberComposition (at.%)Possible Phases
MgAlSnZnCeLaNdPr
A35.3550.640.862.805.443.60.370.94Mg17Al12, Al4Mm
B62.2611.3823.932.43----Mg2Sn
C69.3426.441.352.87----Mg17Al12
D43.9646.730.932.91.962.660.290.58Al4Mm
Table
Table 3. The average values of tensile properties of as-cast samples tested at RT and 175 °C.
Table 3. The average values of tensile properties of as-cast samples tested at RT and 175 °C.
AlloyRT175 °C
UTS (MPa)YS (MPa)Ef (%)UTS (MPa)YS (MPa)Ef (%)
ATZ731192 ± 395 ± 310.2 ± 0.4129 ± 282 ± 118.2 ± 1.1
ATZE7310207 ± 491 ± 113.5 ± 0.5145 ± 287 ± 123.6 ± 1.1
ATZE7311207 ± 4108 ± 219.6 ± 0.2148 ± 3102 ± 128.0 ± 0.3
ATZE7312207 ± 494 ± 313.8 ± 0.4139 ± 690 ± 317.2 ± 0.5
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Liu, G.-J.; Sun, Y.-H.; Xia, N.; Guan, X.-F. Effects of Ce-Rich Misch Metal on the Microstructures and Tensile Properties of as-Cast Mg-7Al-3Sn-1Zn Alloys. Metals 2021, 11, 1648. https://doi.org/10.3390/met11101648

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Liu G-J, Sun Y-H, Xia N, Guan X-F. Effects of Ce-Rich Misch Metal on the Microstructures and Tensile Properties of as-Cast Mg-7Al-3Sn-1Zn Alloys. Metals. 2021; 11(10):1648. https://doi.org/10.3390/met11101648

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Liu, Guo-Jun, Yan-Hua Sun, Nan Xia, and Xiao-Fang Guan. 2021. "Effects of Ce-Rich Misch Metal on the Microstructures and Tensile Properties of as-Cast Mg-7Al-3Sn-1Zn Alloys" Metals 11, no. 10: 1648. https://doi.org/10.3390/met11101648

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