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Article

Microstructural Evolution from Dendrites to Core-Shell Equiaxed Grain Morphology for CoCrFeNiVx High-Entropy Alloys in Metallic Casting Mold

by
Leigang Cao
*,
Lin Zhu
,
Hongde Shi
,
Zerui Wang
,
Yue Yang
,
Yi Meng
,
Leilei Zhang
and
Yan Cui
*
Department of Materials Science and Engineering, North China University of Technology, Beijing 100144, China
*
Authors to whom correspondence should be addressed.
Metals 2019, 9(11), 1172; https://doi.org/10.3390/met9111172
Submission received: 30 September 2019 / Revised: 24 October 2019 / Accepted: 29 October 2019 / Published: 30 October 2019
(This article belongs to the Special Issue Rapid Solidification Processing)
Browse Figures
Versions Notes

Abstract

:
The CoCrFeNiVx (x = 0, 0.25, 0.5, 0.7, 0.8, 0.9, and 1.0) high-entropy alloys (HEAs) were fabricated by the copper mold casting process. The microstructure, phase constitution, and mechanical properties were investigated by using X-ray diffraction, scanning electron microscopy, transmission electron microscopy analyses and compressive testing. It revealed that, when x ≤ 0.25, the alloys solidified into a single fcc phase. When 0.5 ≤ x ≤ 0.8, the alloys solidified into a dendritic structure of the fcc phase with the formation of the σ phase in the interdendrite region. Interestingly, when x exceeded 0.9, the alloys presented a typical core-shell equiaxed grain morphology. The core region consisting of a mixture of fcc + σ phases was surrounded by the shell of the single σ phase and the interdendrite region solidified into the single fcc phase. The dual-phase “eutectiod” structure in the core region of the equiaxed grain might be formed from the decomposition of the unidentified metastable phase. As the V fraction increased, the compressive yield strength of the CoCrFeNiVx alloys gradually increased from 164 MPa (x = 0) to 458 MPa (x = 0.8), and then sharply increased to 722 MPa (x = 0.9) and 1493 MPa (x = 1.0).

1. Introduction

Alloying is used as an effective route to improve the mechanical properties of structural metallic materials over the long term via well-established strengthening mechanisms such as solid-solution strengthening by inducing lattice distortion and precipitation strengthening with the formation of second phases with different properties [1,2,3]. Traditionally, researchers have mainly focused on the corners of a phase diagram for alloy design. It is because, according to the Gibbs phase rule [4], the complex element constitution usually gives rise to the increasing kinds and fraction of the addition phase (e.g., intermetallics). This will, in many cases, weaken the comprehensive mechanical properties. Conversely, it is interesting that multi-component alloys such as the fcc-type CoCrFeNiMn alloy [5], bcc-type AlCoCrFeNi alloy [6], hcp-type YGdTbDyLu alloy [7], or a mixture of them, are prone to form a simple phase constitution, instead of multi-phase constitution. The phenomenon is against the theoretical prediction. The early ground-breaking experiments of multi-component alloys with equal or near-equal elemental constituents were undertaken intermittently in 1981–2004 by Cantor et al. [4,8], Ranganathan [9], and Yeh et al. [9,10,11]. This kind of alloy system was termed by Yeh as the widely accepted “high-entropy alloys (HEAs)”, which contains at least five metallic elements, with the composition of each element being 5–35 at.%.
Aside from their interesting solidification behaviors, high-entropy alloys also present great potential to be used as structural materials due to their excellent properties such as high hardness [12,13], high strength [14,15,16], good oxidation resistance [17], soft-magnetic properties [18], softening resistance at high temperature [3], and recently revealed, perfect irradiation resistance [19]. Considering the existence of dozens of metallic elements, this new type of alloy implies that a large number of new multi-component metallic materials with excellent properties, which was neglected previously, might be fabricated to serve humanity. Therefore, high-entropy alloys have become a hot topic in materials science. In the past 15 years, many investigations have been performed to study the fundamental feature of multi-component HEAs and many valuable achievements have been obtained. First, the variation of certain constitutional elements can strongly change the crystalline structure, microstructural morphology, and the subsequent mechanical properties. For example, the microstructure of the AlxCoCrFeNi system [20] presents a complex transition from a columnar cellular structure to dendritic grain structures as the Al composition is increased. Correspondingly, the fraction of the bcc phase is increased gradually and reaches 100% when x exceeds 0.9, giving rise to an increasing trend in the hardness. A similar variation trend of the crystalline structure and tensile strength was also found in AlxCoCrFeNiMn HEAs [1]. Recently, laser additive manufacturing provides a high-throughput characterization method to assess the effect of the element on the microstructure and properties of the HEAs by fabricating the compositionally graded HEAs [21,22,23]. The designed composition gradient can give rise to the continuous change of the mechanical properties of the alloy, which can also be used as an approach for surface treatment or gradient materials. Furthermore, some effective approaches that are usually used in traditional metallic materials can also be adopted to strengthen the HEAs. For example, He et al. [15] successfully developed the (CoCrFeNi)94Ti2Al4 alloys with enhanced tensile properties by incorporating the nanosized coherent reinforcing L12–Ni3(Ni, Al) precipitates. As a consequence, the yield strength, tensile strength, and plasticity of the water-quenched alloys after (1) cold rolling (30%), annealing (1273 K + 2 h), and aging (1073 K + 18 h), and (2) cold rolling (70%) and aging (973 K + 2 h), could be improved to 645 MPa, 1094 MPa, and 39% and 1005 MPa, 1272 MPa, and 17%, respectively. In comparison, the tensile strength and plasticity of the as-homogenized control sample were only about 503 MPa and 68%, respectively. Gludovatz et al. [24] successfully prepared a grain refined CoCrMnFeNi alloy with an ~6 µm grain by cold forging, cross rolling, and full recrystallization. The prepared alloy presents excellent mechanical properties even in cryogenic conditions. In addition, the transformation-induced plasticity (TRIP) effect, which has been developed to improve the plasticity of the steels, titanium alloys [25], and amorphous alloys [26], has also been successfully introduced into the Fe50Mn30Cr10Co10 [27] and TaHfZrTi [28] HEAs.
It seems that many obvious improvements can be achieved in the increase of the mechanical properties of the HEAs by means of the processing techniques. However, we must realize that the fabricating process is actually based on the microstructure and phase constitution of the as-solidified alloys, namely that the feature of the as-solidified HEAs is rather important for either direct use as structural materials or for further modification of the mechanical property by subsequent processing techniques like rolling and heat treatment. Therefore, it is still necessary to experimentally investigate the alloying effects on the microstructure evolution and phase selection, and their effects on the subsequent mechanical properties, especially under the condition of lack of the thermodynamic phase diagram. In addition to bcc and hcp precipitates, the tetragonal phase can also be an effective reinforcing phase in fcc-matrix HEAs with the addition of the additional component. One such is the equal-atomic five-component CoCrFeNiV alloy [29], the strength of which could be strongly enhanced by the formation of the σ phase with dramatic loss of the plasticity. Conversely, equal-atomic four-component CoCrFeNi alloy usually solidified into a typical fcc structure with excellent plasticity, but low strength [20,30]. This indicates that the comprehensive mechanical properties of the CoCrFeNi alloy could be improved by controlling the fraction and distribution of the σ phase reinforcement as the V content is increased. However, the solidification behavior of the CoCrFeNiVx HEAs with the V addition is still unclear. Therefore, in this study, as-cast CoCrFeNiVx HEAs were fabricated by the copper mold casting method to investigate microstructure evolution, solidification behavior, and mechanical properties with the addition of the V element. The results can give access to understanding and controlling the fundamental solidification process of the CoCrFeNiVx HEAs for the development of the high-performance materials.

2. Materials and Methods

CoCrFeNiVx (x = 0, 0.25, 0.5, 0.7, 0.8, 0.9, and 1.0) alloys were fabricated by an arc-melting furnace. The purity of raw materials was higher than 99.99 wt%. The furnace was pumped to a pressure of 2 × 10−3 Pa and then back-filled with Ar gas to a pressure of 1.5 × 103 Pa. Pure titanium ingot was melted first to remove the residual oxygen and oxygen-containing substance in the evacuated chamber, and then the alloys were arc-melted five times to ensure chemical homogeneity.
The arc-melted alloy was loaded into a quartz crucible with a hole (about 1 mm diameter) in the center of the base, which was then placed into a single roll melt-spinner. The top end of the crucible was connected with the high-pressure gas line. The copper roller for melt-spinning was removed and replaced by the copper mold. During the casting process, the crucible had two fixed working positions for melting and ejecting the liquid alloys. Therefore, before evacuating the chamber, the crucible was adjusted to the appropriate position to make sure that during the heating step, the section of the crucible with the alloy was surrounded by the radio-frequency (RF) coil and, during the ejecting step, the bottom of the crucible was right above the copper mold with the distance being about 0.3 mm. The system was then evacuated to a pressure of 3 × 10−3 Pa for induction heating. A reflective mirror, in line with the axis of the crucible, was fixed on the top of the instrument to determine the melting process of the alloy. Once the alloy was melted, the crucible was instantly dropped down to the “ejecting position” and the melt was rapidly ejected into the mold immediately with high-pressure argon. A series of cylindrical alloy rods with a diameter of 5 mm and a length of 10 mm were fabricated, as is shown in Figure 1.
The cross-sections of all alloys were ground using SiC papers (320–2000 grit) for phase identification by X-ray diffractometer (XRD, Rigaku Ultimate-IV, Rigaku, Tokyo, Japan) with Cu Kα radiation. The scans focused on the entire cross-section area of the casting rods using a continuous mode with a scan speed of 10 °/min. Then, the samples were mounted in conductive resin and polished using 2.5 µm and 1 µm diamond paste and 0.05 µm alumina suspension (Buehler, Lake Bluff, IL, USA). The samples were washed manually using dilute detergent and by ultrasonic cleaning using an ethanol medium during each polishing step. The microstructure of the alloys (central region of the cross-section) and the average element composition of the alloys and the expected areas/phases were investigated using a backscattered electron detector (BSD) and an energy dispersive spectrometry detector (EDX, Bruker Quantax XFlash 6|60, Bruker, Berlin, Germany) mounted on a Zeiss sigma 300 scanning electron microscope (SEM, Carl Zeiss Ltd, Cambridge, UK) with an accelerating voltage of 15 kV. TEM analyses also focused on the central region of the cross-section for microstructure observation and phase identification using a transmission electron microscope (TEM, FEI Tecnai G2 F30, FEI Co. Ltd, Hillsboro, OR, USA) with an accelerating voltage of 300 kV. The TEM specimen was fabricated by the ion milling of a disk-shaped foil with the thickness and diameter being about 30 µm and 5 mm, respectively, by using a Gatan 691 precision ion polishing system (Gatan, Pleasanton, CA, USA). XRD and TEM diffraction patterns were annotated for phase identification using the powder diffraction pattern file database (2014). Room-temperature compressive properties were performed on the cylindrical samples (ø 5 × 10 mm2) with a strain rate of 1 × 10−3 s−1. It should be ensured that there are no casting holes on the surface and cross section of the cylindrical samples. A compression test was carried out one time for the CoCrFeNiVx (x ≤ 0.8) alloys. For the V0.9 and V1.0 alloys, the fracture surface of the fragments after compression test was also checked and the test needs to be repeated if a casting hole is found.

3. Results and Discussion

3.1. Alloy Composition

The EDX measurements were carried out to confirm the compositions of the CoCrFeNiVx alloys (denoted as Vx hereafter). The average composition of each alloy was determined by scanning at least 10 random areas of the cross-section with the rectangular scanning range of 500 × 350 μm2 (Table 1). It is clear that the measured element compositions of the casting rods were almost the same as the expected values. It followed that the samples met the requirement of the experimental investigation.

3.2. Microstructure Evolution

It was revealed by the SEM and TEM analyses that the microstructure of the alloys transformed from dendrite to an equiaxed structure morphology with the V addition, as shown in Figure 2, Figure 3 and Figure 4. First, the V0 and V0.25 alloys present a single fcc phase, according to the SEM micrographs (Figure 2a,b) and the TEM images and diffraction patterns (Figure 2d,e,g,h). The formation of the additional σ phase was found in the V0.5 alloy (Figure 2c). Then, the CoCrFeNiVx with a high V content (0.7 ≤ x ≤ 1.0) presents a distinct transition of microstructure morphologies, although they possessed the same phase constitution of the fcc and σ phases. The V0.7 and V0.8 alloys displayed a typical dendrite structure of the fcc phase, with remaining liquid in the interdendritic region solidifying into the σ phase (Figure 3a,c). This indicates a simple solidification behavior of the dendrite growth of the fcc phase with the residual liquid in the interdendrite region forming the σ phase. A similar microstructural morphology was also observed in the copper-mold solidified (FeNiCoCu)95V5 HEAs [31].
However, the V0.9 and V1.0 alloys presented a typical equiaxed grain structure, with the grains being surrounded by a light-gray shell, as can be seen in Figure 4a,b. The detailed microstructure information was further studied by SEM at high magnification as well as TEM analysis (Figure 4c–f). The results showed that the matrix of the equiaxed grain was a continuous σ phase with the dispersion of fcc lamellae in the central region of the grain, presenting a typical core-shell structure. The interdendritic phase was proven to be a single fcc phase.
Figure 5 is the XRD diffraction patterns of the CoCrFeNiVx alloys. It indicated that, for the alloys with a low concentration of V element (x ≤ 0.5), it is mainly composed of the single fcc phase. For the alloys with x ≥ 0.7, the σ phase with a tetragonal crystal structure was observed. This is consistent with the studies by Salishchev et al. [29], who also observed the existence of the σ phase in the as-solidified equal-atomic CoCrFeNiV HEA. Figure 5 shows that the peak intensity for the σ phase was enhanced along with the increase in the V concentration. Moreover, there was a significant shift of the diffraction peak of {111} crystal plane from 43.64° (CoCrFeNiV0.25 alloy) to 43.50° (CoCrFeNi alloy), implying a lattice expansion of the fcc phase caused by the V element. In addition, the peak intensity of the {111} crystal plane of the fcc phase in the V0.9 alloy was not the strongest one in the diffraction patterns, which might be due to the existence of the preferential grain orientation in the as-solidified bulk metallic materials [32]. In fact, the σ phase was formed first in the V0.5 alloy, according to the microstructure presented in Figure 2. The deviation of phase identification by XRD analysis is due to the observed low fraction of the σ phase in the V0.5 alloy.
It is difficult to measure the cooling curves of the melt in the copper mold directly, which could be calibrated by means of indirect methods. Kozieł [33] estimated the cooling rate of the melt experienced in the copper mold by using Fe-25 wt% Ni as the “standard thermometers”. The results indicated that for the ø 5 mm casting rod, the minimum values of the superficial and axial cooling rates were about 1400 K s−1 and 11 K s−1, respectively. Therefore, it could be accepted that, in the casting process, the melt in the center region of the mold experiences a certain undercooling, which can give access to a free growth of grains into the surrounding undercooled melt, usually resulting in the equiaxed microstructural morphology. Therefore, it seems acceptable that the V0.9 and V1.0 melts followed the described solidification path with the equiaxed crystal growth. A similar result was also observed in the copper-mold cast AlxCoCrFeNiTi HEAs, wherein large equiaxed grains solidified first from the melt and grew up [34]. However, it is rather difficult to understand the nature of the distinct core-shell structure. Previously, Ahamd et al. [35] investigated the undercooling behaviors on a Ni-25.3 at.% Si alloy (hypereutectic composition) using a melt-fluxing technique. They observed in the experiments that at all undercoolings, the melts were always solidified into a lamellar eutectic structure of single-phase γ-Ni31Si12 lamellae and a supersaturated Ni-rich phase, with the latter undergoing a eutectoid decomposition. The formation of the metastable phase was attributed to the possible low atomic mobility (sluggish diffusion), which was supported by the favorable experimental observations. Sluggish diffusion effect has long been treated as an intrinsic characteristic of the multi-component HEAs [11,36]. A number of recent studies [37,38,39,40,41,42,43] on both the solid and liquid state of the HEAs indicate the existence of the sluggish diffusion effect in HEAs. However, it can be influenced by the variation of the element component. As a result, the metastable phase might be retained easier in the as-solidified HEAs in comparison to that of the binary or ternary alloys (low-entropy alloy), according to the experimental evidences of the disordered bcc phase in the AlCoCrFeNi alloy [3] by traditional casting process (45 × 45 × 10 mm3) and the metastable phase in the Ni-25.3% Si and Ni-Fe-Si melts by severe non-equilibrium solidification processes [35,44,45].
By analogy, the observed duplex structure in the core region of the grain might be from the eutectoid decomposition of the unidentified metastable phase like a supersaturated solid solution. It is clear that this “eutectoid” structure does not fill the entire grain, with a thin layer of the single σ phase in the surface region of the grain. Considering that the σ phase is the continuous matrix phase, we may speculate that the elemental composition of the grain varies along the radial direction of the equiaxed grain during the solidification process due to the effect of the recalescence phenomenon. For further discussion, the elemental compositions of the local regions for the V0.7, V0.8, V0.9, and V1.0 alloys were also analyzed by EDX, as can be seen in Table 2. To obtain the average composition, the EDX measurement was repeated randomly at least 10 times for the fcc and σ phases in the V0.7 and V0.8 alloys, and at least five times for the core, shell, and interdendritic regions in the V0.9 and V1.0 alloys. It is clear that the composition of the fcc phase in the V0.7 and V0.8 alloys was quite close to the average composition of the corresponding alloys, which could be possible since the fcc phase is the primary phase solidified from the melt. The core region of the equiaxed grains contains a high composition of the Cr and Fe elements and a low composition of Co and Ni elements, in comparison to the average elemental composition of the V0.9 alloy. The EDX results of the V1.0 alloy also present a similar result, as shown in Table 2. Therefore, the origin of the core-shell equiaxed grain structure might be explained as follows. Initially, the undercooled melt facilitates the nucleation and subsequent growth of the unidentified metastable phase with a depletion of Cr and Ni elements in the surrounding areas. Grain growth is usually accompanied by the process of releasing latent heat, giving rise to continuous warming of the surrounding undercooled liquid. It might change the atomic diffusion behavior in the local region close to the solidified grain. In the early stage, a low fraction of solid phase, indicating low quantity of the released latent heat, will not give rise to a significant temperature increase of the surrounding melt. In this case, the atomic diffusion will not change too much and, therefore, the unidentified metastable phase could grow to a large size (Figure 5a,c). Then, as the fraction of the solid phase is increased, a high quantity of the latent heat would be expected and, in turn, the surrounding liquid could be heated to a high temperature, which can accelerate the atomic mobility and favor the formation of the thermodynamic equilibrium phase. Therefore, due to the available atomic diffusion rate in the “warm” melt, the surface region of the equiaxed grains and the interdendrite region solidified into the single σ and fcc phases, respectively. The elemental composition of the single σ phase was still close to that of the core region of the equiaxed grain. However, the obvious variation of the composition was observed in the fcc phase (V0.9 and V1.0 alloy) with a higher Ni composition, which is probably because of the enhanced atomic diffusion ability in the “warm” interdendritic melt. A similar result can be seen in the V0.7 and V0.8 alloys, namely that the σ phase in the interdendrite region, solidified in the final stage, possesses high compositions of Cr and V elements due to the sufficient diffusion time. The primary solidified unidentified metastable phase decomposes into the σ and fcc phases in the solidification process, with the presence of the eutectoid structure in the core region of the equiaxed grains.

3.3. Mechanical Property

Figure 6 shows the compressive mechanical properties of CoCrFeNiVx rods with different V content at a strain rate of 1.0 × 10−3 s−1. The results of the yield stress and fracture strength are given in Table 3. First, when x ≤ 0.5, the V0.0, V0.25, and V0.5 alloys were finally compressed into a drum shape without fracture, indicating an excellent ductility. The yield stress of the alloys, being at a low level, increased slightly from 164 MPa (V0.0 alloy) to 291 MPa (V0.5 alloy). This can be attributed to the solid solution strengthening, although the V0.5 alloy contained a low fraction of the σ phase. A similar result was also observed by Gwalani et al. [22] in the compositionally graded AlCrFeMoVx (0 ≤ x ≤ 1) high-entropy alloy fabricated by laser additive manufacturing. The addition of the V element does not change the phase constitution, but it results in the increase of the hardness from 485 HV to 581 HV. Meanwhile, as the V element was increased, the fraction of the second phase (σ phase in the interdendritic region) increased. It is known that the σ phase, with a tetragonal crystal structure, was stronger than the fcc matrix phase due to the inherent characteristics of the crystal structure [46]. The enhanced yield stress of the V0.7 and V0.8 alloys, being about 392 MPa and 458 MPa, was due to the increasing fraction of the σ phase.
The further increase of the mechanical properties was observed in the V0.9 and V1.0 alloys, with the yield strength being about 722 MPa and 1493 MPa, respectively. Generally, the precipitation of the additional phase could improve the strength first and then weaken the strength when its fraction exceeded a certain level, which can be worse if the additional phase is of a large size. Although the microstructure of the equiaxed grains in the V0.9 and V1.0 alloys was similar to a eutectic structure, the continuous matrix phase was the brittle σ phase. Therefore, to a certain degree, the microstructure of the V0.9 and V1.0 alloys could be treated as the hard particle reinforced HEA composite, resulting in a significant increase in the strength with a decrease in plasticity. A significant increase of the yield strength in the V1.0 alloy, in comparison with that of the V0.9 alloy, can be attributed to the higher fraction of the hard grains and the lower fraction of the fcc phase in the interdendrite region.

4. Conclusions

(1)
The addition of the V element facilitates the formation of the σ phase from the fcc matrix for the CoCrFeNiVx alloys, with a microstructure transition from a single fcc phase structure (x ≤ 0.25) to a dendrite structure (0.5 ≤ x ≤ 0.8), and finally to the distinct equiaxed grain structure (0.9 ≤ x ≤ 1.0).
(2)
The formation of the σ phase can effectively enhance the mechanical properties of the alloys. As the V content increased, the yield strength of the alloys first increased gradually from 164 MPa (x = 0) to 458 MPa (x = 0.8), and then increased dramatically from 722 MPa (x = 0.9) to 1493 MPa (x = 1.0).
(3)
The core region of the equiaxed grain in the V0.9 and V1.0 alloys presents the fcc + σ eutectoid microstructure, which is derived from the eutectoid decomposition of the unidentified metastable phase.

Author Contributions

Conceptualization, L.C. and Y.C.; Funding acquisition, L.C. and Y.C.; Methodology, L.Z. (Lin Zhu), H.S. and Z.W.; Writing—original draft, L.C.; Writing—review & editing, L.C., Y.Y., Y.M., and L.Z. (Leilei Zhang).

Funding

This research was funded by the National Key R&D Program of China (2017YFB0703102), the Beijing Natural Science Foundation (2194074), the Science and Technology Project of Beijing Municipal Education Commission (KM201910009005), the Fundamental Scientific Research Funding of Beijing Municipal Education Commission (110052971921/040), the Yuqing Talent Support Program (18XN012-081), and the Undergraduate Scientific Research and Entrepreneurship Action Plan (18XN130) of North China University of Technology.

Acknowledgments

We would like to thank Fengbin Liu for his revision on the language.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. The ø 5 mm alloy rod solidified in the copper mold.
Figure 1. The ø 5 mm alloy rod solidified in the copper mold.
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Figure 2. (a–c) SEM micrographs of the V0.0, V0.25, and V0.5 rods, respectively. (d–f) TEM micrographs of the V0.0, V0.25 and V0.5 rods, respectively. (g–j) TEM selected area diffraction patterns corresponding to (a), (b), and the phases marked as A and B in (f), respectively. The white flake phase that first formed in the V0.5 alloy (c) was identified as the σ phase. The straight lines in (a–c) are the stretches and the black spots in (a) are the surface contaminants introduced during the polishing process.
Figure 2. (a–c) SEM micrographs of the V0.0, V0.25, and V0.5 rods, respectively. (d–f) TEM micrographs of the V0.0, V0.25 and V0.5 rods, respectively. (g–j) TEM selected area diffraction patterns corresponding to (a), (b), and the phases marked as A and B in (f), respectively. The white flake phase that first formed in the V0.5 alloy (c) was identified as the σ phase. The straight lines in (a–c) are the stretches and the black spots in (a) are the surface contaminants introduced during the polishing process.
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Figure 3. (a,c) SEM micrographs of the V0.7 and V0.8 alloys, showing the typical dendrite morphology. (b,d) TEM results of the V0.7 and V0.8 alloys, respectively. The black spots in (a,c) are the surface contaminants introduced during the polishing process.
Figure 3. (a,c) SEM micrographs of the V0.7 and V0.8 alloys, showing the typical dendrite morphology. (b,d) TEM results of the V0.7 and V0.8 alloys, respectively. The black spots in (a,c) are the surface contaminants introduced during the polishing process.
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Figure 4. (a,b) SEM micrographs of the V0.9 and V1.0 alloys, respectively. (c,d) SEM micrographs at high magnification referring to the white rectangles in (a,b), respectively. (e,f) TEM results of the V0.9 and V1.0 alloys. The regions of the σ phase “shell” and interdendritic fcc phase are highlighted by green dashed lines. The black spots in (a,b) are the surface contaminants introduced during the polishing process.
Figure 4. (a,b) SEM micrographs of the V0.9 and V1.0 alloys, respectively. (c,d) SEM micrographs at high magnification referring to the white rectangles in (a,b), respectively. (e,f) TEM results of the V0.9 and V1.0 alloys. The regions of the σ phase “shell” and interdendritic fcc phase are highlighted by green dashed lines. The black spots in (a,b) are the surface contaminants introduced during the polishing process.
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Figure 5. X-ray diffraction patterns of the studied high-entropy alloys with different V concentrations.
Figure 5. X-ray diffraction patterns of the studied high-entropy alloys with different V concentrations.
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Figure 6. Compressive mechanical properties of the CoCrFeNiVx rods at room temperature.
Figure 6. Compressive mechanical properties of the CoCrFeNiVx rods at room temperature.
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Table
Table 1. Average elemental composition of the CoCrFeNiVx alloys measured by energy dispersive spectrometry analysis (EDX).
Table 1. Average elemental composition of the CoCrFeNiVx alloys measured by energy dispersive spectrometry analysis (EDX).
AlloyCo/at.%Cr/at.%Fe/at.%Ni/at.%V/at.%
V0.0Nominal25.0025.0025.0025.00-
Measured25.74 ± 0.2024.97 ± 0.2025.17 ± 0.1924.12 ± 0.25-
V0.25Nominal23.5323.5323.5323.535.88
Measured24.31 ± 0.2623.47 ± 0.2323.78 ± 0.3622.75 ± 0.155.79 ± 0.07
V0.5Nominal22.2222.2222.2222.2211.11
Measured22.92 ± 0.2722.31 ± 0.2222.38 ± 0.2121.42 ± 0.3210.95 ± 0.19
V0.7Nominal21.2821.2821.2821.2814.89
Measured21.93 ± 0.2521.26 ± 0.2921.35 ± 0.1420.67 ± 0.2814.79 ± 0.14
V0.8Nominal20.8320.8320.8320.8316.67
Measured21.46 ± 0.2320.53 ± 0.2621.02 ± 0.1620.39 ± 0.2516.60 ± 0.19
V0.9Nominal20.4120.4120.4120.4118.37
Measured21.01 ± 0.2420.35 ± 0.1920.47 ± 0.2019.93 ± 0.3218.24 ± 0.18
V1.0Nominal20.0020.0020.0020.0020.00
Measured20.71 ± 0.1720.02 ± 0.1720.09 ± 0.2519.46 ± 0.2319.72 ± 0.24
Table
Table 2. Average elemental composition of the CoCrFeNiVx alloys measured by EDX.
Table 2. Average elemental composition of the CoCrFeNiVx alloys measured by EDX.
AlloyCo/at.%Cr/at.%Fe/at.%Ni/at.%V/at.%
V0.7Matrix (fcc phase)22.42 ± 0.2220.75 ± 0.2321.94 ± 0.4921.23 ± 0.2413.65 ± 0.52
ID (σ phase)20.02 ± 0.3725.35 ± 0.5220.51 ± 0.1615.91 ± 0.5718.23 ± 0.33
V0.8Matrix (fcc phase)22.17 ± 0.1519.79 ± 1.2221.63 ± 0.2421.26 ± 0.3214.95 ± 0.24
ID (σ phase)20.07 ± 0.2623.34 ± 0.8020.42 ± 0.3216.57 ± 0.1719.30 ± 0.40
V0.9ID (fcc phase)22.00 ± 0.1417.72 ± 0.1619.99 ± 0.1123.21 ± 0.1317.08 ± 0.11
Shell (σ phase)20.13 ± 0.2222.76 ± 0.5820.75 ± 0.2516.91 ± 0.3819.46 ± 0.37
Core region20.72 ± 0.2621.93 ± 0.2221.09 ± 0.2118.02 ± 0.0918.25 ± 0.26
V1.0ID (fcc phase)21.66 ± 0.1616.69 ± 0.4519.34 ± 0.3623.71 ± 0.6518.60 ± 0.27
Shell (σ phase)19.99 ± 0.2321.80 ± 0.2420.30 ± 0.2017.22 ± 0.2920.69 ± 0.20
Core region20.37 ± 0.2621.63 ± 0.2320.79 ± 0.3117.85 ± 0.1319.36 ± 0.32
Table
Table 3. The yield stress (σy) and fracture strength (σmax) of the CoCrFeNiVx high-entropy alloys.
Table 3. The yield stress (σy) and fracture strength (σmax) of the CoCrFeNiVx high-entropy alloys.
Alloyσy/MPaσmax/MPa
V0.0164-
V0.25234-
V0.5291-
V0.73921600
V0.84581435
V0.97221825
V1.014932032

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Cao, L.; Zhu, L.; Shi, H.; Wang, Z.; Yang, Y.; Meng, Y.; Zhang, L.; Cui, Y. Microstructural Evolution from Dendrites to Core-Shell Equiaxed Grain Morphology for CoCrFeNiVx High-Entropy Alloys in Metallic Casting Mold. Metals 2019, 9, 1172. https://doi.org/10.3390/met9111172

AMA Style

Cao L, Zhu L, Shi H, Wang Z, Yang Y, Meng Y, Zhang L, Cui Y. Microstructural Evolution from Dendrites to Core-Shell Equiaxed Grain Morphology for CoCrFeNiVx High-Entropy Alloys in Metallic Casting Mold. Metals. 2019; 9(11):1172. https://doi.org/10.3390/met9111172

Chicago/Turabian Style

Cao, Leigang, Lin Zhu, Hongde Shi, Zerui Wang, Yue Yang, Yi Meng, Leilei Zhang, and Yan Cui. 2019. "Microstructural Evolution from Dendrites to Core-Shell Equiaxed Grain Morphology for CoCrFeNiVx High-Entropy Alloys in Metallic Casting Mold" Metals 9, no. 11: 1172. https://doi.org/10.3390/met9111172

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