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Article

Effects of Ag, Nd, and Yb on the Microstructures and Mechanical Properties of Mg‒Zn‒Ca Metallic Glasses

by
Zhuofan Liang
1,2,†,
Lianzan Yang
1,†,
Yongyan Li
1,
Xi Wang
1,*,
Chunling Qin
1,
Weimin Zhao
1,*,
Hui Yu
1 and
Zhifeng Wang
1,2,*
1
School of Materials Science and Engineering, Hebei University of Technology, Tianjin 300130, China
2
Key Laboratory for New Type of Functional Materials in Hebei Province, Hebei University of Technology, Tianjin 300130, China
*
Authors to whom correspondence should be addressed.
†
These authors contributed equally to this work.
Metals 2018, 8(10), 856; https://doi.org/10.3390/met8100856
Submission received: 19 September 2018 / Revised: 18 October 2018 / Accepted: 19 October 2018 / Published: 20 October 2018
(This article belongs to the Special Issue Effect of Rare Earth Additions on the Microstructure, Mechanical Properties and Corrosion of Magnesium Alloys)
Browse Figures
Versions Notes

Abstract

:
Mg‒Zn‒Ca metallic glasses are regarded as promising biodegradable materials. Previous studies on this alloy system have mostly focused on the composition regions with a large critical size (Dc) for the formation of metallic glasses, while this paper investigates the composition regions with a small Dc, which has been overlooked by researchers for a long time. The effects of the addition of Ag, Nd, and Yb elements on the microstructure and mechanical properties of Mg‒Zn‒Ca metallic glasses were studied. It was found that the Mg‒Zn‒Ca metallic glass exhibits a single and uniform amorphous structure with a compressive strength of 590 MPa. After the addition of a small amount of Ag into the alloy, the amorphous matrix is retained and new precipitate phases that lead to the decrease of the compressive strength are formed. The addition of the rare earth elements Nd and Yb changes the microstructure from a single amorphous matrix to a large number of quasicrystal phases, which results in an increase in compressive strength. The compressive strength of the Mg‒Zn‒Ca‒Yb alloy increases to 606.2 MPa due to the formation of multi-layered swirling solidified structure and a large number of small quasicrystals with high microhardness. Moreover, this study can be considered as a useful supplement to the existing studies on the Mg‒Zn‒Ca alloy system; it also provides new ideas for designing the microstructure and spatial structure of quasicrystal containing alloys with high performances.

1. Introduction

As an environmentally friendly material, magnesium alloys have the characteristics of light weight and high reserve content in the earth crust. Therefore, they have been widely used in electronic products and medical equipment [1,2]. Compared to traditional crystalline magnesium alloys, Mg-based amorphous alloys exhibit a higher compressive strength and corrosion resistance as they have a uniform structure with almost no defects such as grain boundaries and dislocations [3,4,5,6].
Mg‒Zn‒Ca metallic glass is a typical Mg-based amorphous alloy. It is safe and non-toxic as all of its constituent elements are essential elements in the human body. The alloy is considered to be a promising biodegradable material and has attracted significant attention worldwide [6,7,8]. The existing literature on this alloy system focuses on the following aspects. Firstly, some studies attempted to obtain Mg‒Zn‒Ca alloys with a large critical size (Dc) for the formation of metallic glass. In this way, the application range of the alloys can be extended [9,10,11]. It is reported that the Mg60Zn35Ca5 alloy presents a Dc of only 1 mm and the Mg67Zn28Ca5 alloy has Dc = 4 mm [12,13]. The largest glass forming sizes are located in Mg66Zn30Ca4 and Mg67Zn29Ca4, with Dc = 5 mm [9,13,14]. Secondly, Mg‒Zn‒Ca metallic glasses may generate hydrogen gas when they are implanted into an animal’s body. Therefore, avoiding the flow-out of hydrogen gas is indispensable for promoting the application of such alloys. Löffler et al. [8] reported that when the Zn concentration is more than 28 at %, no hydrogen generation can be detected during material degradation. That means the hydrogen gas flow-out problem can be effectively solved by modifying the original composition of the alloy. Thirdly, the corrosion resistance of Mg‒Zn‒Ca alloys is relatively poor. As a result, the implanted material may lose its effect before the completion of the service period due to its excessive degradation. Improving the corrosion resistance or reducing its degradation rate is therefore an important research topic. Previous studies indicate that the corrosion resistance of Mg‒Zn‒Ca metallic glasses can be effectively enhanced by adding Sr, Ag, and a small amount of rare earth element Y [15,16,17], or synthesizing a corrosion-resistant coating on the alloys [18,19,20,21,22]. Fourthly, Mg‒Zn‒Ca metallic glasses are hard and brittle, with several studies focused on improving their plasticity, which is of great significance for their practical application. Wang et al. [23] introduced a Yb element into the Mg‒Zn‒Ca alloy and the as-formed Mg‒Zn‒Ca‒Yb amorphous ribbon showed excellent plasticity. Qin et al. [24] added Ag into the Mg‒Zn‒Ca alloy and found that the Dc is largest when the content of Ag is around 1 at %. The as-obtained Mg‒Zn‒Ca‒Ag alloy exhibited excellent mechanical properties. Wang et al. [25] prepared Mg‒Zn‒Ca‒Y rod by introducing a Y element into the Mg‒Zn‒Ca amorphous matrix. The material exhibits improved compression plasticity for the formation of endogenous dendrites as the toughness-enhancing phase.
Most of the existing works have been carried out by selecting the composition that can form metallic glass with a large Dc [26,27,28,29,30,31,32,33,34], while research on metallic glass with a small Dc is still rarely reported. If a fourth component is added into the aforementioned compositions, the changes in the structure of amorphous matrix and mechanical properties of the alloys are worth noting. If a rare earth element is added, whether a quasicrystal phase can be formed needs confirming. The studies on such a series of problems are still lacking. In this paper, the effects of the addition of Ag, Nd, and Yb elements on the microstructure and mechanical properties of the Mg‒Zn‒Ca amorphous alloys were investigated. The study would be a useful supplement to previous works and will provide important findings for further research and applications on Mg‒Zn‒Ca alloys.

2. Materials and Methods

2.1. Composition Design of the Experimental Alloys

Figure 1 shows the composition range for the metallic glass formation in a ternary Mg‒Zn‒Ca alloy system [9]. Metallic glass rods with a Dc of at least 1 mm can be formed in the yellow area. The position of the brown rhombus corresponds to the Mg‒Zn‒Ca metallic glass with Dc = 2 mm. The position of the green triangle relates to the Mg‒Zn‒Ca metallic glass with a Dc of 3 mm. The gray dot presents the composition of the alloy designed by this work, with a Dc between 2 mm and 3 mm. Owing to the relatively small Dc, when a bit of the fourth component is added to the alloy, the amorphous matrix may or may not be retained. A new precipitated phase may be formed in this situation. The composition selected in this work aims at the exploration of these changes in microstructures and mechanical properties before and after the addition of the fourth component. The compositions of the experimental alloys are shown in Table 1.

2.2. Material Preparation

High-purity Mg ingot (99.99 wt %), Zn ingot (99.99 wt %), Mg-30.5 wt % Ca master alloys, Ag ingot (99.99 wt %), Nd ingot (99.99 wt %), and Yb ingot (99.99 wt %) were melted in a crucible electric resistance furnace (SG2-5-10A) under a protection of SF6/CO2 mixed atmosphere. The melting temperatures of Mg‒Zn‒Ca, Mg‒Zn‒Ca‒Ag, Mg‒Zn‒Ca‒Nd, and Mg‒Zn‒Ca‒Yb alloys are 760 °C, 1000 °C, 1050 °C, and 850 °C, respectively, depending on the melting point of the fourth component. After stabilizing for 2 min, the melt was subsequently poured into a copper mold, followed by water cooling. The master alloys were remelted twice to ensure a homogenous composition. Finally, the alloys were induction melted and spray casted into the copper mold to obtain rods with a diameter of 2 mm and a length of 4 cm. The samples were further machined to a length of 4 mm (some samples were produced with a diameter of 3 mm and a length of 6 mm) for microstructural observation and compression tests.

2.3. Microstructure and Mechanical Properties

The phases were characterized via X-ray diffraction (XRD, Bruker D8) with Cu Kα radiation. The microstructure of the master alloys and compositions of precipitated phase were analyzed using a scanning electron microscope (SEM, Hitachi S4800) equipped with energy-dispersive spectroscopy (EDS). The compressive tests (five samples were tested for each composition) were conducted with an UTM5105X testing machine at room temperature. The fracture morphologies of the local area were determined by SEM.

3. Results and Discussion

Figure 2 shows the XRD results of the experimental alloys. It can be seen from Figure 2a that the XRD curve of Mg62.9Zn32.3Ca4.8 alloy exhibits typical broad (halo) peak [24], demonstrating that the alloy obtained by spray casting is indeed an amorphous alloy. When a small amount of Ag is incorporated, the broad peaks of the amorphous alloy are still visible while a few sharp crystal peaks appear, which are attributed to the formation of Mg102.08Zn39.6 phase and Ag3Mg phase. Compared to the sample with a diameter of 2 mm, the positions of diffraction peaks slightly shifted to a larger angle (to the right) for the sample with a higher diameter of 3 mm, indicating that more Ag3Mg phase precipitated in the matrix.
Figure 2b shows the XRD patterns of the Mg62.9Zn32.3Ca4.8 alloy after introducing Nd and Yb rare earth elements. It can be seen that when rare earth elements are added, the broad peaks disappear, indicating that the amorphous matrix has changed to a crystalline structure. It means a new phase has been created. It is found that the samples with diameters of 2 mm and 3 mm were mainly composed of two phases, quasicrystal I-phase [35,36,37] and α-Mg phase. Compared to the sample with a diameter of 2 mm, the diffraction peaks of the sample with Dc = 3 mm slightly shift to a small angle (to the left). Furthermore, due to the difference in composition, the position of the diffraction peaks of the Mg59.3Zn32.4Ca4.8Yb3.5 alloy shifts to a smaller angle (to the left) compared to the Mg59.8Zn33.1Ca4.7Nd2.4 alloy.
Figure 3 shows SEM images of Mg62.9Zn32.3Ca4.8 (Figure 3a) and Mg63Zn30.2Ca4.5Ag2.3 (Figure 3b,c) alloys. It can be seen from Figure 3a that the original Mg‒Zn‒Ca alloy shows a uniform microstructure with no precipitate phase. It is a typical structure of an amorphous alloy and it is consistent with the XRD data. When a small amount of Ag (2.3 at %) is added to the alloy, some precipitated phases are observed. Table 2 shows the EDS analysis of areas A, B, and C in Figure 3b. It can be seen that the composition in light gray area A is close to the Mg63Zn30.2Ca4.5Ag2.3 alloy, indicating an inheritance of the amorphous matrix. The Mg/Zn ratio in area B is 2.61, which is approximate to the composition of Mg102.08Zn39.6 phase. The composition of dark gray area C presents a relatively high Ag ratio, demonstrating the formation of an Ag3Mg phase. These results are in accord with the XRD data. In this situation, a composite material containing amorphous matrix and some precipitated phases can be confirmed. Further observations show that when the sample’s diameter reduces from 3 mm (Figure 3b) to 2 mm (Figure 3c), the microstructure does not change significantly, but the size of the spherical precipitated phases decreases significantly from 2~4 μm to 500~800 nm. The above result shows that after adding a small amount of Ag element to the alloy, the amorphous matrix is retained and finally an amorphous composite material containing spherical precipitated phases is formed.
Figure 4 shows the microstructures of the Mg‒Zn‒Ca alloy after the addition of Nd and Yb rare earth elements. It can be seen that, when the rare earth elements are added, the microstructures change significantly. After adding Nd, a typical quasicrystal petal phase [36] is formed in the alloy. When the sample’s diameter reduces from 3 mm to 2 mm, the quasicrystal petals changes from multi-lobe to nearly spherical shape [35,37], exhibiting a rapidly decreasing petal length ranging from 5~9 μm to 1~3 μm. In addition to the quasicrystal petal phase, the primary phase of the alloy matrix is α-Mg phase and a lamellar eutectic composed of quasicrystal I phase and α-Mg phase.
As seen from Figure 4c–f, when a certain amount of rare earth element Yb is added into the Mg‒Zn‒Ca ternary alloy, a composite material composed of α-Mg and quasicrystal petal phase is formed. Unlike the alloy with Nd additions, no eutectic is formed when Yb is added. When the sample’s diameter is 3 mm, the alloy’s primary structure is α-Mg matrix and the quasicrystal petals appear as the precipitated phase. When the sample’s diameter reduces to 2 mm, the quasicrystal petals are very densely distributed while the α-Mg region is greatly reduced. In this situation, the size range of the quasicrystal petal rapidly reduces from 5~7 μm (sample with Dc = 3 mm) to 2~3.5 μm (sample with Dc = 2 mm). It is further observed that the sample with Dc = 2 mm is rapidly cooled and forms structures with a liquid-flow swirling morphology since the molten metal rapidly fills up the mold. At the edge of the swirling region (region A), a multi-layer structure composed of small-size quasicrystal phases and α-Mg matrix is formed. In the center of the swirling region (region B), the solidification conditions are relatively uniform and the quasicrystal particles are uniformly distributed. The macroscopic morphologies and the compositions of regions A and B are different. The compositions of regions A and B based on EDS tests are shown in Table 3, in which the main difference is that the Mg content in region A is lower than that in region B. Such a multi-layer swirling structure provides us with a new idea for designing high-performance quasicrystal alloys.
Based on the testing results relative to Figure 4, it is found that when rare earth elements are introduced into the Mg‒Zn‒Ca alloy under the current experimental conditions, the amorphous structures cannot be retained and a large amount of Mg‒Zn‒RE quasicrystal phases forms instead. On the one hand, the rapid copper mold cooling conditions are more suitable for the nucleation process of the Mg‒Zn‒RE quasicrystal phases. Meanwhile, it can limit the growth of the quasicrystal phases to a certain extent, so that only moderate growth occurs. On the other hand, the rare earth element has a higher melting point. Therefore, Mg is excessively volatilized and consumed during a high-temperature smelting process. This makes the alloy composition deviate from the design point of the original Mg‒Zn‒Ca ternary alloy and the compositions areas with large Dc. The above reasons led to the formation of a large amount of quasicrystal phases, which replace the amorphous region as a main product.
Figure 5 and Table 4 show the measured mechanical properties of the alloys. It can be seen that the original Mg62.9Zn32.3Ca4.8 amorphous alloy shows excellent compressive strength (590 MPa). When the fourth component was added to the alloy, the variations in mechanical properties can be clearly seen. After the addition of Ag element, the compressive strength reduces to 506.5 MPa. It has been reported that when 1 at % Ag is added into the Mg65Zn30Ca5 alloy, a quaternary Mg65Zn30Ca4Ag1 amorphous alloy can be formed and its mechanical properties are better than those of the ternary amorphous alloy [24]. In this study, the Ag content is relatively high. Although the amorphous matrix is retained, the precipitated phases of Mg102.08Zn39.6 [38] and Ag3Mg [39] are formed. The hardness of the two phases is lower than that of the amorphous matrix, so the alloy’s compressive strength decreases. It is worth noting that the Ag3Mg phase possesses a long period superstructure [40], which is in favor of the plasticity maintenance of the alloy. Moreover, the majority of the alloy keeps the original amorphous matrix. As a result, the plasticity of the alloy is rarely reduced after the addition of Ag. After the addition of Nd, the compressive strength decreases to 465.5 MPa, which is related to the formation of a large number of fragile lamellar eutectic. When Yb is added, the alloy’s compressive strength increases to 606.2 MPa. Such enhancement of compressive strength may be related to two factors. For one thing, when a Yb element is added, a large number of small quasicrystal particles are formed in the Mg matrix. Since there is no specific slip surface in the quasicrystal system, the dislocations cannot move easily at room temperature. Therefore, these hard and brittle quasicrystal particles act as a strengthening phase and improve the alloy’s strength. For another thing, a multi-layer swirling solidified structure composed of α-Mg matrix and densely-distributed small quasicrystal particles forms during the solidification process. This tough, interlaced, dense multi-layer structure leads to increasing compressive strength. In addition, the compressive strength of samples with Dc = 3 mm is worse than those of samples with Dc = 2 mm, which may be related to the larger size of precipitated phases in the 3 mm samples.
The size and volume fraction of quasicrystal phases formed in Nd- and Yb-containing alloys can be seen from Table 4. When the rod diameter decreases from 3 mm to 2 mm, the quasicrystal size decreases and the volume fraction of quasicrystals increases obviously in both the Nd- and Yb-containing alloys. Moreover, due to a higher content of rare earth elements, the critical stable radius [37] for forming spherical quasicrystals decreases. As a result, the quasicrystal size in the Yb-containing alloy is smaller than that in the Nd-containing alloy. The amount and volume fraction of quasicrystal phases in the Yb-containing alloy are higher than in the Nd-containing alloy.
In this study, when Nd and Yb were introduced into the original alloy, the quasicrystal phase induces brittleness in the case of the Nd-containing alloy while it has a beneficial effect on the strength of the Yb-containing alloy. The microstructure of the Nd-containing alloy is mainly composed of petal-like quasicrystals with large size and lamellar eutectics. Every petal-like quasicrystal is surrounded by a lamellar eutectic network. The existence of these fragile eutectics induces obvious brittleness in the Nd-containing alloy. At the same time, the microstructure of the Yb-containing alloy is composed of quasicrystals and α-Mg. Fragile eutectics are not formed in this alloy. In particular, the multi-layer swirling structure is formed by close packing of a large number of small quasicrystals and α-Mg. Such special structure provides the possibility of improvement in the compressive strength. For another aspect, the content of rare-earth element Yb is higher than that of Nd, which induces more significant constitutional supercooling and further decreases the critical stable radius for forming spherical quasicrystals. As a result, a large amount of petal-like and spherical quasicrystals with small sizes instead of large petal-like quasicrystals are created. It was found that the smaller the quasicrystal size, the greater its microhardness [36]. So, the strengthening effect of the Yb-containing alloy is improved by these small quasicrystals.
Figure 6 shows the fracture morphologies of the samples with a diameter of 2 mm after compression testing. Figure 6a,b show the compression fracture morphologies of the Mg62.9Zn32.3Ca4.8 amorphous alloy. It can be seen that the fracture surface shows a large number of herringbone fracture patterns, indicating that the overall fracture mode of the alloy is brittle. In addition, obvious vein patterns can be observed within some local areas of the fracture, showing the characteristics of local plastic fracture. These fracture pattern characteristics correspond to the good mechanical properties exhibited by the amorphous alloy. Figure 6c shows the fracture morphology of the Mg63Zn30.2Ca4.5Ag2.3 alloy. Both samples with diameters of 3 mm and 2 mm exhibit obvious cleavage steps, indicating that the fracture mode is a typical brittle fracture. Figure 6d and the inset show the fracture morphologies of the Mg59.8Zn33.1Ca4.7Nd2.4 alloy. The macroscopic fracture shows a river pattern, indicating that the fracture mode is brittle. It can be further observed from the magnified fracture structure that the quasicrystal phase and matrix in the alloy are separated during the compression process, resulting in obvious gaps and pores, which can explain the decline in compressive strength compared to the original alloy. Figure 6e,f shows the fracture morphologies of the Mg59.3Zn32.4Ca4.8Yb3.5 alloy. The macroscopic fracture exhibits a river pattern, indicating that the fracture mode is also brittle. Several quasicrystal particles can be observed in the magnified fracture structure, which indicates that the crack mainly propagates along the grain boundaries at the quasicrystal edge. The dense grain distribution makes an important contribution to the higher compressive strength.

4. Conclusions

The present paper focused on the composition regions in a Mg‒Zn‒Ca alloy system with a small Dc. The effects of the addition of Ag, Nd, and Yb on the microstructure and mechanical properties of a Mg‒Zn‒Ca alloy were studied. It was found that the compressive strength of the Mg‒Zn‒Ca amorphous alloy can reach 590 MPa, showing the characteristics of overall brittle fracture (herringbone pattern) and local ductile fracture (vein pattern). After adding a small amount of Ag to the alloy, a composite material containing amorphous matrix as well as Mg102.08Zn39.6 and Ag3Mg precipitated phases was obtained. The compressive strength reduced significantly after the addition of Ag. After the addition of Nd and Yb rare earth elements to the alloy, the microstructure with amorphous matrix is replaced by a large amount of quasicrystal phases. The Mg‒Zn‒Ca‒Nd alloy is mainly composed of α-Mg, lamellar eutectic, and petal-shaped quasicrystal phases. The appearance of a large amount of lamellar eutectic significantly reduces the alloy’s mechanical properties. However, the compressive strength of the Mg‒Zn‒Ca‒Yb alloy increases to 606.2 MPa, which is related to its solidified dense multi-layer swirling structure and a large number of small quasicrystals with high microhardness.

Author Contributions

Z.W., X.W., and W.Z. conceived and designed the experiments; Z.L. and L.Y. performed the experiments; Z.L., L.Y., and Y.L. analyzed the data; Z.W. acquired the funding and did the project administration; Y.L., C.Q., and H.Y. contributed reagents/materials/analysis tools; X.W. and W.Z. supervised the work; Z.L. prepared the original draft; Z.W. revised the paper and created the final version.

Funding

This work was financially supported by the Natural Science Foundation of Hebei Province, China (E2015202081) and the Innovation & Entrepreneurship Training Program of Hebei University of Technology (201810080241).

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Composition range for the metallic glass formation in Mg‒Zn‒Ca ternary alloy system. Adapted with permission from Springer Nature, Copyright 2009 [9].
Figure 1. Composition range for the metallic glass formation in Mg‒Zn‒Ca ternary alloy system. Adapted with permission from Springer Nature, Copyright 2009 [9].
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Figure 2. X-ray diffraction (XRD) patterns of (a) Mg62.9Zn32.3Ca4.8 and Mg63Zn30.2Ca4.5Ag2.3 alloys; (b) Mg59.8Zn33.1Ca4.7Nd2.4 and Mg59.3Zn32.4Ca4.8Yb3.5 alloys.
Figure 2. X-ray diffraction (XRD) patterns of (a) Mg62.9Zn32.3Ca4.8 and Mg63Zn30.2Ca4.5Ag2.3 alloys; (b) Mg59.8Zn33.1Ca4.7Nd2.4 and Mg59.3Zn32.4Ca4.8Yb3.5 alloys.
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Figure 3. SEM images of experimental alloys: (a) Mg‒Zn‒Ca, Φ 2 mm; (b) Mg‒Zn‒Ca‒Ag, Φ 3 mm; (c) Mg‒Zn‒Ca‒Ag, Φ 2 mm.
Figure 3. SEM images of experimental alloys: (a) Mg‒Zn‒Ca, Φ 2 mm; (b) Mg‒Zn‒Ca‒Ag, Φ 3 mm; (c) Mg‒Zn‒Ca‒Ag, Φ 2 mm.
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Figure 4. SEM images of experimental alloys: (a) Mg‒Zn‒Nd, Φ 3 mm; (b) Mg‒Zn‒Ca‒Nd, Φ 2 mm; (c) Mg‒Zn‒Ca‒Yb, Φ 3 mm; (d–f) Mg‒Zn‒Ca‒Yb, Φ 2 mm.
Figure 4. SEM images of experimental alloys: (a) Mg‒Zn‒Nd, Φ 3 mm; (b) Mg‒Zn‒Ca‒Nd, Φ 2 mm; (c) Mg‒Zn‒Ca‒Yb, Φ 3 mm; (d–f) Mg‒Zn‒Ca‒Yb, Φ 2 mm.
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Figure 5. Compressive curves of Mg‒Zn‒Ca-based alloys.
Figure 5. Compressive curves of Mg‒Zn‒Ca-based alloys.
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Figure 6. SEM images of the fracture morphologies of the experimental alloys: (a,b) Mg‒Zn‒Ca, Φ 2 mm (c) Mg‒Zn‒Ca‒Ag, Φ 2 mm (d) Mg‒Zn‒Ca‒Nd, Φ 2 mm (e,f) Mg‒Zn‒Ca‒Yb, Φ 2 mm.
Figure 6. SEM images of the fracture morphologies of the experimental alloys: (a,b) Mg‒Zn‒Ca, Φ 2 mm (c) Mg‒Zn‒Ca‒Ag, Φ 2 mm (d) Mg‒Zn‒Ca‒Nd, Φ 2 mm (e,f) Mg‒Zn‒Ca‒Yb, Φ 2 mm.
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Table
Table 1. The compositions of the experimental alloys (at %).
Table 1. The compositions of the experimental alloys (at %).
AlloyMgZnCaAgNdYb
Mg‒Zn‒Ca62.932.34.8---
Mg‒Zn‒Ca‒Ag6330.24.52.3--
Mg‒Zn‒Ca‒Nd59.833.14.7-2.4-
Mg‒Zn‒Ca‒Yb59.332.44.8--3.5
Table
Table 2. EDS analysis of areas A, B, and C in Figure 3b (at %).
Table 2. EDS analysis of areas A, B, and C in Figure 3b (at %).
Testing AreaMgZnCaAg
Area A63.230.34.32.2
Area B71.127.21.10.6
Area C49.331.15.514.1
Table
Table 3. EDS analysis of area A and B in Figure 4d (at %).
Table 3. EDS analysis of area A and B in Figure 4d (at %).
Testing AreaMgZnCaYb
Area A56.933.85.43.9
Area B61.730.74.43.2
Table
Table 4. The comparison of the experimental alloys.
Table 4. The comparison of the experimental alloys.
Alloy (at %)Diameter (mm)Compressive Strength (MPa)Quasicrystal Phases
Size (μm)Volume Fraction (%)
Mg62.9Zn32.3Ca4.82590 ± 5.1--
Mg63Zn30.2Ca4.5Ag2.33347.6 ± 8.2--
2506.5 ± 7.5--
Mg59.8Zn33.1Ca4.7Nd2.43298.4 ± 9.3~8.545 ± 3
2465.5 ± 6.4~4.567 ± 4
Mg59.3Zn32.4Ca4.8Yb3.53540.8 ± 5.2~844 ± 5
2606.2 ± 4.9~3.579 ± 3

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MDPI and ACS Style

Liang, Z.; Yang, L.; Li, Y.; Wang, X.; Qin, C.; Zhao, W.; Yu, H.; Wang, Z. Effects of Ag, Nd, and Yb on the Microstructures and Mechanical Properties of Mg‒Zn‒Ca Metallic Glasses. Metals 2018, 8, 856. https://doi.org/10.3390/met8100856

AMA Style

Liang Z, Yang L, Li Y, Wang X, Qin C, Zhao W, Yu H, Wang Z. Effects of Ag, Nd, and Yb on the Microstructures and Mechanical Properties of Mg‒Zn‒Ca Metallic Glasses. Metals. 2018; 8(10):856. https://doi.org/10.3390/met8100856

Chicago/Turabian Style

Liang, Zhuofan, Lianzan Yang, Yongyan Li, Xi Wang, Chunling Qin, Weimin Zhao, Hui Yu, and Zhifeng Wang. 2018. "Effects of Ag, Nd, and Yb on the Microstructures and Mechanical Properties of Mg‒Zn‒Ca Metallic Glasses" Metals 8, no. 10: 856. https://doi.org/10.3390/met8100856

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