Solidication Microstructure of AlCoCrFeNi2.1 Eutectic High Entropy Alloy Ingots*

Takeshi Nagase1,2, Mamoru Takemura3, Mitsuaki Matsumuro3 and Toru Maruyama4

1Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, Ibaraki 567–0047, Japan2Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, Suita 565–0871, Japan3Osaka Research Institute of Industrial Science and Technology, Izumi 594–1157, Japan4Faculty of Chemistry, Materials and Bioengineering, Kansai University, Suita 564–8680, Japan

AlCoCrFeNi2.1 eutectic high-entropy alloy (EHEA) ingots were successfully obtained by high-frequency melting and centrifugal met-al-mold casting under an Ar ow. The microstructure of the ingots was investigated by trans-scale observations using optical microscopy (OM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), and scanning transmission electron microscopy (STEM). The constituent phases of the ingots were identied as fcc and bcc phases by X-ray diffraction (XRD) analysis, and were not depen-dent on the position of the ingot. The microstructure was observed to have a primary fcc dendrite and fcc+bcc eutectic structure at the in-ter-dendrite region, regardless of the position of the ingots. The size of the solidication structure was affected by the cooling rate. Faster cooling rates resulted in ner solidication structures. TEM observations claried the development of L12 ordering structures in the primary fcc dendrite phase, while the ordering peak could not be detected by XRD analysis.　[doi:10.2320/matertrans.F-M2017851]

(Received September 6, 2017; Accepted November 2, 2017; Published December 22, 2017)

Keywords:　 metal, alloy, high-entropy alloy, microstructure, solidication

1.　 Introduction

A new class of metallic materials called high-entropy al-loys (HEAs) were developed and are presently attracting considerable interest in the area of materials science and en-gineering1–7). The abbreviation “HEA” is roughly used to re-fer to the metallic materials that are constructed with equal or nearly equal quantities of ve or more metals and whose constituent phase is solid solution. Several HEA denitions were suggested, which gives rise to some confusion. Recently, an entropy-based denition was suggested and has been widely adopted, because the phrase “high entropy” sug-gests a denition based on the magnitude of entropy, as de-noted in the literature5–7). The entropy-based denition is as follows:

∆Smix = −Rn

i=1

xi ln xi (1)

∆S mix ≥ 1.5R (HEA) (2)

where ΔSmix is the mixing entropy of an ideal solid solution (and regular solid solution), R is the gas constant (8.314 J/K mol), and xi is the mole fraction of the element i. Based on the ΔSmix, a medium-entropy alloy (MEA) and low-entropy alloy (LEA) can be dened as follows:

1.0R ≤ ∆S mix ≤ 1.5R (MEA) (3)

∆S mix ≤ 1.0R (LEA) (4)

The mixing entropy represented by eq. (1) shows the maxi-mum value in the case of an equi-atomic ratio, namely, xi = 1/n. The maximum mixing entropy in alloys is expressed by eq. (5), and this value increases with an increase in the num-ber of elements.

(∆Smix)max = −R ln n (5)

The value of (ΔSmix)max in a quaternary (4-component) alloy system is 1.39R, and that in a 5-component alloy system is 1.61R. Based on the entropy-based denition of eqs. (1) and (2), only multicomponent alloys with equal and nearly equal quantities of ve or more constituent elements can be classi-ed as HEAs. Table 1 lists the mixing entropies calculated for typical traditional alloys and equi-atomic alloys, focus-ing on cast iron and steel, with some of the data being taken from the literature6). A number of previously developed tra-ditional and common alloys cannot attain a value of 1.5R in ΔSmix. This indicates that a value of “5” is not a critical num-ber for an HEA, but a metallic material that is constituted of equal or near-equal quantities of “5” or more metals is im-portant for obtaining alloys for which ΔSmix ≥ 1.5R. Recent progress in the taxonomy of HEAs7) indicates that not only multicomponent alloys with a single solid-solution phase, but also composites with a main solid-solution phase and minor intermetallic compounds, as well as composites with two solid solutions, and so on, can also be called HEAs. In the present study, the entropy-based denition was adopted for HEAs.

HEAs exhibit four effects that differentiate them from other, conventional alloys. In other words, four basic core ef-fects differentiate conventional alloys (namely, LEAs and MEAs) from HEAs2,5–7). These four core effects are (1) high entropy, (2) severe lattice distortion, (3) sluggish diffusion, and (4) “cocktail.” These four core effects differentiate HEAs from previously developed conventional alloys, inter-metallic compounds, and amorphous alloys/metallic glass. Considering that HEAs are cast materials, they exhibit the following three interesting effects: (1) a solid-solution phase

* This Paper was Originally Published in Japanese in J. JFS 89 (2017) 119–129. In order to more precisely explain the background, the purpose of the study, the experimental procedures, and the results, some parts of the contents were revised. The Refs. 21, 22), and 23) were added to clar-ify the source of pure elements. The Refs. 12) and 35) were added to dis-cuss the solidication microstructure in more detail. Figure 1 was slightly modied to make the casting process and the position of the in-gots clearly understandable. Figure 7 was slightly modied to discuss the solidication microstructure in more detail.

Materials Transactions, Vol. 59, No. 2 (2018) pp. 255 to 264 ©2017 Japan Foundry Engineering Society

can be obtained using a conventional casting process with-out any need for special techniques, given the high entropy from a thermodynamic aspect and sluggish diffusion from a kinetic aspect, (2) ductile materials can be obtained because of the formation of a solid solution, and (3) high-strength materials are formed as a result of the severe lattice distor-tion effect. The main route adopted for the fabrication of an HEA specimen is melting and casting because of the need to mix the constituent elements5–7). While some refractory HEAs8,9) and bio-HEAs10), which consist of 4–6 group ele-ments including Ti and others with high melting points, are tried to be fabricated by powder-metallurgy processes with solid-state mixing, melting and casting is the most com-monly used technique for fabricating HEAs. When adapting melting and casting processes for the fabrication of HEAs, the following problems should be considered and overcome prior to further development and practical application; (1) the method for melting a mixture of constituent elements, considering the widely varying melting points, the degree of oxidation, the reactivity between the molten metal and the crucible, and the vapor pressure, (2) the general tendency of the heterogeneous properties of the constituent phases and the microstructure, as well as the solidication segregation, has not been claried for HEAs, and (3) the grain size in HEA specimens obtained by melting and solidication ex-hibits a tendency to be coarser rather than that of HEA spec-imens obtained by powder metallurgy. Regarding problem (2), to the best of our knowledge, there have been no sys-tematic studies addressing the solidication segregation in HEAs. Furthermore, the constituent phases and microstruc-ture evaluation were performed mainly by XRD analysis without any consideration of position-dependence. The com-

plexity of the combination of constituent elements in HEAs points to cast products having various type of heterogeneous properties, although little attention has been afforded this topic.

In the present study, the solidication microstructure of AlCoCrFeNi2.1 eutectic high-entropy alloys (EHEAs) was investigated by trans-scale observation12) to clarify the posi-tion dependence of the constituent phases and the micro-structure of cast products. The present study set out to estab-lish a means of evaluating the solidication microstructure of HEA casting products.

2.　 Materials and Methods

2.1　 Alloy composition and alloy parametersIn is well known that HEAs exhibit the following general

tendencies in terms of their mechanical properties: sin-gle-phase, bcc-structured HEAs exhibit a relatively high strength but tend to be brittle, while fcc-structured HEAs are ductile but not strong enough to be superior to common al-loys developed to date. EHEAs were developed based on the concept of combining bcc- and fcc-structured HEAs to achieve a balance of strength and ductility. The composite structure could be obtained from the eutectic reaction occur-ring in the thermal melt during solidication11,12). The fabri-cation of industrial-scale ingots of AlCoCrFeNi2.1 EHEAs was attempted to research future applications for engineer-ing cast materials11,12). However, the possibility of solidica-tion segregation and the position dependence of the constitu-ent phases, as well as the microstructure of the ingots of AlCoCrFeNi2.1 EHEAs, were not claried. An AlCoCrFeNi2.1 EHEA was selected for the investigation be-

Table 1　Mixing entropies calculated for typical traditional alloys and equiatomic alloys. HEA, MEA, and LEA denote high-entropy alloy, medium-entropy alloy, and low-entropy alloy, respectively.

Alloys ΔSmix LEA MEA HEA

Cast IronFe-C-Si　(Fe-3.6C-2.4Si [mass.%])

0.58R

　(Fe-3.3C-2.7Si [mass.%]) 0.58R Austenitic cast iron, FCDA-NiMn 13 7　(Fe-3C-3Si-7Mn-13Cr)

0.88R

Cast stainless steel, SCS16　(Fe-0.03C-1.5Si-2Mn-12Ni-20Cr-2.5Mo, [mass%])

1.13R

Heat-resistant cast steel, SCH22　(Fe-0.45C-1.75Si-1.5Mn-22Ni-27Cr, [mass%])

1.29R

Stainless spheroidal carbide cast iron, KSH-C　(Fe-2.84C-0.95Si-0.61Mn-9.2Ni-17.3Cr-9.34V [mass.%])

1.40R

Stainless steelSUS-304　(Fe-18Cr-8Ni [mass%])

0.74 R

SUS-316　(Fe-18Cr-14Ni-3Mo [mass%])

0.94 R

Al alloy A7075 0.43R Mg alloy AZ91D 0.35R Cu alloy 7-3 brass 0.61R Ni alloy Inconel 718 1.31R

Alnico alloyAlnico2　(Fe-10Al-19Ni-13Co-3Cu [mass%])

1.33R

Alnico9　(Fe-7Al-15Ni-35Co-4Cu-5Ti [mass%])

1.55R

Equiatomic alloy TiNbTaZr 1.39R CoCrMnFeNi 1.61R

256 T. Nagase, M. Takemura, M. Matsumuro and T. Maruyama

cause the solidication microstructure can be evaluated eas-ily compared with single-phase HEAs and the applicability of this alloy to engineering applications.

Table 2 lists the nominal alloy composition and the alloy parameters for AlCoCrFeNi2.1 EHEA. As shown in the nom-inal alloy composition (Table 2(a)), the atomic composition ratio of Ni is larger than that of the other elements (Al, Co, Cr, and Fe). Based on the weight percentage, the weight ra-tio of (Ni + Co) is over 50 mass%, and that of (Ni + Co + Cr) is nearly 75%. Various parameters were suggested to predict the formation of the solid-solution phase in multi-compo-nent alloys. Table 2(b) lists the parameters for HEAs in AlCoCrFeNi2.1 EHEA, and the following ve parameters (A)–(E) were used in the present study.

(A) Mixing entropy, ΔSmix

The mixing entropy was easily calculated using eq. (1). The value of ΔSmix of AlCoCrFeNi2.1 EHEA is 1.55R (Table 2(b)).

(B) Mixing enthalpy, ΔHmix

Mixing enthalpy was used to predict the tendency to form a solid-solution phase, together with the glass forming abil-ity (GFA) in multicomponent alloys2,3,5,6,14–16). The mixing enthalpy was expressed as:

∆Hmix =

n

i=1

n

j i

4∆Hi jxix j (6)

where ΔHij is the mixing enthalpy of an X-Y binary system at an equi-atomic composition (X50Y50) for the liquid phase. The value of ΔHij is taken from the literature16).

(C) Delta parameter, δThe atomic size difference in a multi-component alloy is

also an important parameter for predicting the formation of the solid-solution and amorphous phases. The δ parameter for discussing the atomic size differences is as follows3,5,6,17):

δ =

n

i=1

xi 1 − ri

r× 100 (7)

r =n

i=1

xiri (8)

where ri is the atomic radius of i-th element and r is the av-erage atomic radius. The value of ri was again taken from the literature16).

(D) Omega parameter, ΩThe Ω parameter17) contains both ΔSmix and ΔHmix for pre-

dicting the formation of a solid solution, and is expressed as follows:

Ω =Tm · ∆S mix

|∆Hmix| (9)

Tm =

n

i=1

xi · (Tm)i (10)

where (Tm)i is the melting temperature of the element i, Tm is the average melting temperature which is derived from the rule governing a mixture of pure elements and their melting temperatures.

(E) Valence electron concentration parameter, VECVEC is a frequently used parameter for predicting the

phase stability of fcc and bcc structures in HEAs, and is ex-pressed as follows7,18,19),

VEC =n

i=1

ci · VECi (11)

where VECi is the value of VEC for a pure element. The lit-erature discusses the use of VEC or the electrons per atom ratio (e/a) for predicting the phase stability of HEAs in de-tail7). We opted to use VEC in the present study. The fcc phases occur at VEC > 8.0, and the bcc phases at VEC ≦ 6.87, with a mixture of bcc and fcc phases being obtained at 6.87 ≦ VEC < 8.0. There are some exceptions to the VEC rule, however, which provide a convenient means of design-ing fcc- and bcc-based HEAS, as well as those which are a composite of bcc and fcc HEAs.

Zhang et al. reported that a solid-solution phase tended to form when −20 ≦ ΔHmix ≦ 5 [kJ/mol], δ ≦ 6.4, 0.012 ≦ ΔSmix ≦ 0.0175 [kJ/mol]3). Guo et al. suggested that the fol-lowing conditions promote the formation of solid solutions: −11.6 ≦ ΔHmix ≦ 3.2, δ ≦ 6.620). Yang et al. suggested that δ ≦ 6.4 and Ω ≧ 1.1 should be used to predict the formation of the solid-solution phases17). Although some inconsisten-cies are observed between the above-mentioned criteria, they are helpful for predicting the formation of solid-solution phases in multicomponent alloys. Table 2(b) lists the values of the ΔSmix, ΔHmix, δ, and Ω parameters for an AlCoCrFeNi2.1 alloy. The 0.0129 kJ/mol of ΔSmix corre-sponds to 1.55R. All the values exhibit a strong solid-solu-tion formation tendency in an AlCoCrFeNi2.1 alloy. The VEC value of an AlCoCrFeNi2.1 alloy is 7.7, indicating the formation of a mixture of fcc and bcc phases.

2.2　 Experimental procedureFigure 1 is a schematic illustration of the metal mold cast-

ing process and the position of the ingot. Figure 1(a), (b) shows the crucible and a mixture of pure element blocks, and the centrifugal-casting equipment, respectively. The schematic illustration of the centrifugal-casting equipment and the position of the ingot for microstructure analysis are shown in Fig. 1(c), (d), respectively. A mixture of pure ele-ment blocks, with a total weight of about 80 g was set in the silica-based crucible, as indicated by A in Fig. 1(a). The al-loy composition is listed in Table 1(a). Pure Al shots (Purity 99.99%, Mitsuwa Chemical Co., Ltd., Japan21)), Co akes

Table 2　(a) Alloy composition and (b) alloy parameters of AlCoCrFeNi2.1 eutectic high-entropy alloys.

(a) Alloy composition

Al Co Cr Fe Ni

atomic ratio 1 1 1 1 2.1

at% 16.39 16.39 16.39 16.39 34.43

mass％ 8.51 18.59 16.40 17.62 38.88

(b) Alloy parameters

ΔSmix ΔHmixδ Ω VEC

[kJ/mol K] [kJ/mol]

0.0129 −11.9 5.3 1.8 7.70

257Solidication Microstructure of AlCoCrFeNi2.1 Eutectic High Entropy Alloy Ingots

(Purity 99.9%, Osaka Asahi Co., Ltd., Japan22)), Cr granules (Purity 99.9%, Mitsuwa Chemical Co., Ltd., Japan21)), Fe akes (Purity 99.9%, Toho Zinc Co., Ltd., Japan23)), and Ni grains (Purity 99.9%, Osaka Asahi Co., Ltd., Japan22)), were used as the pure metal blocks in the present study. The cruci-ble (A) and the Cu mold (B) were set in the centrifugal-cast-ing equipment as shown in Fig. 1(b). Thermal melt was ob-tained by high-frequency melting under an Ar ow. The Ar supply is indicated by C in Fig. 1(b), (c). The ingots were obtained by two different processes to give an “air-cooled in-got” and a “metal-molded ingot.” In the case of the air-cooled specimens, the thermal melt was cooled in the cruci-ble with a slow cooling rate. In contrast, the metal-molded specimens were obtained by centrifugal casting. The metal mold was rst heated to remove any moisture. No mold-re-lease agent was used.

To clarify the position dependence of the constituent phases and the microstructure, the following method was ap-plied in the present study. (1) The ingot was cut into nine part by electro-discharge machining as shown in Fig. 1(d). (2) The specimens were wet-polished by SiC abrasive grind-ing paper and/or diamond polishing paper. (3) The speci-mens of No. ① (metal-mold-contacted, top part), No. ② (metal-mold-contacted, middle part), No. ③ (met-al-mold-contacted, bottom part), No. ④ (non-contact with metal mold, top part), No. ⑤ (non-contact with metal mold, middle part), No. ⑥ (metal-mold-contacted, bottom part) were selected as typical samples. (4) The constituent phases and the microstructures of specimens ①, ②, ③, ④, ⑤, and ⑥ were investigated by XRD analysis and trans-scale obser-vation13) using OM, SEM, EPMA, TEM, and STEM. The thin lms required for the TEM and STEM measurements were prepared by mechanical polishing and ion-milling at room temperature. The Micro-Vickers hardness test was per-formed with a load setting of 1 kgf (9.807 kN).

The cooling rate during the solidication of the ingot pre-pared by air cooling and metal-mold casting was estimated from the secondary dendrite arm spacing in the Al95.5Cu4.5 alloy24,25). For the air-cooled ingot, the cooling rate was in the order of 1 K/s, regardless of the position of the ingot.

The cooling rates in the ingot prepared by metal-mold cast-ing are shown in Fig. 2. The cooling rate in the metal mold ingot points to the position dependence and was estimated to be 100–500 K/s. The cooling rates at positions ③ and ⑥ were higher than those in other positions. Positions ④ and ⑤, which were not in contact with the metal mold, exhibit a relatively low cooling rate. We could consider that the cool-ing rate shown in Fig. 2 is higher than the typical value cal-culated by the ratio of the volume and surface area in the in-gots fabricated by conventional metal-mold casting26). The difference may be due to the application of the centrifugal casting process, non-use of a mold-release agent, and the use of a Cu mold in the present study.

3.　 Results

Figure 3 shows the outer appearance of the AlCoCrFeNi2.1 EHEA ingots prepared by (a) metal-mold casting and (b) the air-cooling process. The metal-mold in-got exhibits a characteristic metallic luster without any oxide or casting defects such as open cavities or blowholes. In contrast, the air-cooled ingot is covered by an oxide layer. Figure 4 shows a macroscopic view of the solidication structures of the AlCoCrFeNi2.1 EHEA ingots prepared by (a) metal-mold casting and (b) the air-cooling process. The development of the columnar grains from the surface to the central area, and that of the equi-axial structure in the central region can be seen in the metal mold ingot. In the air-cooled

Fig. 2　Position dependence of cooling rate of ingot prepared by met-al-mold casting. The cooling rate was estimated from the spacing of the secondary dendrite arms in the Al95.5Cu4.5 (at%) alloy.

Fig. 3　Appearance of the AlCoCrFeNi2.1 eutectic high-entropy alloy in-gots prepared by (a) metal-mold casting and (b) air-cooling process.

Fig. 1　Schematic illustration of metal-mold casting process and ingot po-sition. (a) Schematic illustration of crucible and mixture of pure element blocks, (b) appearance of centrifugal-casting equipment, (c) schematic illustration of centrifugal-casting equipment, and (d) schematic illustra-tion of ingot position for microstructure analysis.

258 T. Nagase, M. Takemura, M. Matsumuro and T. Maruyama

ingot, a large shrinkage cavity can be seen. The size of the solidication structure in the air-cooled ingot was larger than that in the metal-mold ingot. These results indicate that ingots without casting defects can be obtained by high-fre-quency melting under an Ar ow and by centrifugal casting using a metal mold. It should be emphasized here that the present study focused on the position dependence of the constituent phases and the microstructure in the ingots. The castability and the formation of macroscopic cast defects are an important topic affecting the practical application of HEAs and EHEAs. The formation of a shrinkage cavity in the air-cooled specimens may be prevented by the control of the casting process. However, a detailed study of the forma-tion of shrinkage cavities was not a concern in the present study, but will be addressed in our future work.

To clarify the position dependence of the constituent phases and the microstructure in the ingot, XRD analysis was performed for specimens ①–⑥. The results for the metal-mold ingot and air-cooled ingot are shown in Figs. 5 and 6, respectively. In the XRD patterns, the intensity is ex-pressed as a linear scale in the left-hand gures and by a log-arithmic scale by the right-hand gures to enhance minor peaks. In the metal-mold ingots, the sharp peaks in the XRD patterns of the specimens ①–⑥ can be identied as fcc or bcc phases. The lattice constant of the fcc and bcc was esti-mated to be 0.358 ± 0.001 [nm] and 0.286 ± 0.001 [nm], re-spectively, and no position dependence of the lattice con-stants was observed. The peak intensity of the fcc phase is higher than that of the bcc phase, indicating that the volume fraction of the fcc phase is larger than that of the bcc phase. In the logarithmic scale plot (right side), no minor peaks corresponding to intermetallic compounds were observed. Lu et al. reported the existence of minor peaks correspond-ing to the B2 ordering phase in the XRD pattern obtained from an weighing 2.5 Kg, while the B2 ordering was not ob-served in the present study. The peak intensity ratio in the fcc phase is position dependent between (111) and (200). For specimens ③ and ⑥, from the bottom of the met-al-mold contacted region, exhibit a peak intensity for (200)

Fig. 4　Macroscopic view of solidication structures of AlCoCrFeNi2.1 eu-tectic high-entropy alloy ingots prepared by (a) metal-mold casting and (b) air-cooling process.

Fig. 5　Position dependence of XRD patterns of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by metal-mold casting. (a) Positions ①, ②, and ③; and (b) positions ④, ⑤, and ⑥. The intensity is ex-pressed using a linear scale in the left-hand gures and with a logarith-mic scale in the right-hand gures to enhance minor peaks.

Fig. 6　Position dependence of XRD patterns of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by air-cooled casting. (a) Positions ①, ②, and ③; and (b) positions ④, ⑤, and ⑥. The intensity is ex-pressed using a linear scale in the left-hand gures and with a logarith-mic scale in the right-hand gures.

259Solidication Microstructure of AlCoCrFeNi2.1 Eutectic High Entropy Alloy Ingots

that is much higher than that for (111). This can be ex-plained by the crystallographic orientation due to the prefer-ential growth of the primary fcc-dendrite from the met-al-mold contacted surface, as described above. The XRD patterns for the air-cooled ingot exhibit the following ten-dencies which are similar to those of the metal-mold ingot: (1) A composite of the fcc and bcc phases was formed, and the position-dependence of the constituent phases was not observed. (2) The lattice constants of the fcc and bcc phases exhibit no position dependence. (3) The volume fraction of the fcc phases tends to be larger than that of the bcc phase. (4) An ordering peak cannot be seen in the XRD patterns with linear and logarithmic intensity plots of the XRD patterns.

Figure 7 shows SEM back scattering electron (BSE) im-ages of the solidication structures of the AlCoCrFeNi2.1 EHEA ingots prepared by (a) metal-mold casting and (b) the air-cooling process. Specimens ④, ⑤, and ⑥ were selected as typical examples. Specimen ⑥ exhibits a fast cooling rate due to its being in contact with the metal mold, while speci-mens ④ and ⑤ exhibit a lower cooling rate because neither were in contact with the metal mold. The SEM microstruc-tures were observed in the central area of the partitioned

specimens, with the lower parts of the SEM images corre-sponding to the lower parts of the ingots. Figure 7(c) shows the SEM microstructure, focusing on the metal-mold/ingot interface. The black contrast region, a typical example of which is indicated by D, was claried as being not an inclu-sion formed through the solidication process but as an arti-fact introduced during the polishing. The randomly dis-persed black particles can be assumed to be SiC and/or SiC-related particles from the SiC abrasive grinding paper. In the metal-mold ingots, the development of the primary dendrite and the eutectic structure at the inter-dendrite re-gion can be seen (Figs. 7(a1)–(a3)). The size of the primary dendrite indicates a position dependence, as follows, ④ > ⑤ > ⑥ (Fig. 7(a1) > Fig. 7(a2) > Fig. 7(a3)). In the air-cooled ingots, the primary dendrite and the eutectic structure in the inter-dendrite region can be seen in the inset of the SEM images for specimens ⑤ and ⑥ (Figs. 7(b2), (b3)), while the development of the dendrite structure was not ob-served in specimen ④ (Fig. 7(b1)). The area ratio of (pri-mary dendrite)/(eutectic structure in the inter-dendrite re-gion) in the specimens also shows the position dependence as follows: ⑤ > ⑥. The full lamellar structure can be seen for specimen ④. No signicant difference in the size of the

Fig. 7　SEM-back scattering electron (BSE) images of solidication structures of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by (a) met-al-mold casting and (b) air-cooling process. (a1) Metal-mold casting–position ④, (a2) ⑤, and (a3) ⑥; and (b1) Air cooling–position ④, (b2) ⑤, and (b3) ⑥. (c) Metal-mold casting–position ⑥ focusing on the mold wall.

260 T. Nagase, M. Takemura, M. Matsumuro and T. Maruyama

primary dendrite of specimens ④ and ⑤ was observed. The size of primary dendrite in the air-cooled ingots is much larger than that in the metal-mold ingots, regardless the posi-tion in the ingots. A TEM analysis claried that the primary dendrite phase is an fcc phase, while the eutectic structure is composed of fcc and bcc phases. These are described in de-tail below. The development of the dendrite was not seen at the metal-mold/ingot interface, and the formation of the eu-tectic structure in the inter-dendrite region was observed in the region from about 100 μm from the metal-mold/ingot in-terface. The columnar grain was observed at the met-al-mold/ingot interface, which can be explained by the for-mation of a chill layer and the development of columnar grains at the metal-mold/ingot interface27–29). The inclusions formed only at the metal-mold/ingot interface region were not seen in the AlCoCrFeNi2.1 EHEA ingots.

To quantitatively evaluate the size of the solidication mi-crostructure, we investigated the relationship between the secondary dendrite arm spacing (DAS) and the cooling rate (K) of the ingots. The results are shown in Fig. 8. The DAS was generally expressed as a function of K using the follow-ing equation24,25,27–31):

DAS = C · KA (12)

where A and C are constants. The linear curve in Fig. 8 is the approximate curve based on eq. (12), and the value of A was estimated to be -0.23. The DAS for the AlCoCrFeNi2.1 EHEA ingots can be evaluated using eq. (12). These results indicate that the solidication microstructure in the AlCoCrFeNi2.1 EHEA consists of primary fcc dendrite and fcc+bcc eutectic in the inter-dendrite region, while the size of the solidication microstructure increased as the cooling rate fell.

We performed SEM-EPMA and STEM-EDS mapping to investigate the distribution of the constituent elements during the solidication. The results are shown in Fig. 9. The EPMA-WDS mapping results (Fig. 9(a)) exhibit the fol-lowing tendencies: (1) the Al and Ni elements are enriched in the inter-dendrite region, (2) the Co, Cr, and Fe elements are enriched in the primary fcc dendrite. The STEM-EDS mapping results, focusing on the inter-dendrite region, are shown in Fig. 9(b). In the STEM-HAADF image, the white and black contrast regions correspond to the bcc and fcc phases, respectively. The bcc phase exhibits a tendency to be

enriched by Al and Ni, while the fcc phases exhibit the op-posite tendency in the eutectic region. Table 3 lists the ex-perimental analysis results for the chemical composition of the primary fcc dendrite as evaluated by EPMA-WDS, while the fcc and bcc phases in the eutectic region were evaluated by STEM-EDS. The fcc dendrite appears to be lacking in Al, relative to the nominal composition. For the eutectic region, the following two tendencies can be observed: (1) the Al ele-ment is lacking in the fcc phase, relative to the nominal composition, which is similar to the primary fcc dendrite. (2) The Al and Ni elements are enriched in the bcc phase, relative to the primary fcc dendrite and the fcc phase in eu-tectic region. In other words, the bcc phase in the eutectic re-gion is an Ni- and Al-rich phase. This indicates that the so-lidication of the AlCoCrFeNi2.1 EHEA progresses according to the following sequence: (1) fcc dendrite forma-tion and the ejection of Al to the residual liquid, (2) fcc + bcc eutectic structure formation in the inter-dendrite region, resulting in the formation of an Ni/Al-rich bcc phase.

XRD observations (Figs. 5 and 6), SEM observations (Fig. 7), and EPMA/STEM analyses (Fig. 9) claried that the constituent phases in an AlCoCrFeNi2.1 EHEA ingot are the primary fcc dendrite and the fcc+bcc eutectic structure in the inter-dendrite region. To investigate the microstructure of the primary dendrite fcc phase, TEM analysis of specimen ⑤ (a typical metal-mold ingot sample) was performed. The results are shown in Fig. 10. In the bright eld (BF) image, no inclusions or precipitates were observed. The inset is a BF image focusing on the triple points of the grain bound-ary. No precipitates can be seen. Figures 10(b1)–(b3) show the [001], [011], and [111] selected area diffraction (SAD) patterns, where Figs. 10(c1)–(c3) and Fig 10(d) are the [001], [011], and [111] key diagrams for the XY3 compound with an L12 structure, and a schematic illustration of the bi-nary XY3 compound with a L12 structure, respectively. In the key diagrams (Figs. 10(b1)–(b3)), and represent the fundamental diffraction of the fcc structure and the su-per-lattice diffraction based on the L12 structure, respec-tively. The SAD patterns (Fig. 10(c1)–(c3)) can be explained as being those of the L12 structure. This indicates that the primary fcc phase is not a chemically random fcc solid-solu-tion phase but an fcc phase with weak chemical ordering based on the L12 ordering structure, where the chemical or-dering was too weak to be detected by XRD analysis.

Table 4 lists the position dependencies of the Vickers hardness (Hv) of the AlCoCrFeNi2.1 EHEA ingots prepared by metal-mold casting. The average and standard deviation were evaluated using ve data points at positions ①–⑥. The size of the indentation was approximately 40 μm. The value of Hv was obtained from the region consisting of the pri-mary fcc dendrite and the fcc+bcc eutectic in the inter-den-drite region. No signicant position dependence of the Hv was observed. Figure 10 shows the Hv of the AlCoCrFeNi2.1 EHEA ingots prepared by metal-mold casting as a function of the distance from the metal-mold/ingot interface in the specimen. The Hv exhibits a constant value and no position dependence, while anomalous values in the region near the metal-mold/ingot interface were not observed. The Hv val-ues in the AlCoCrFeNi2.1 EHEA ingots were nearly homo-geneous, while the size of the solidication microstructure

Fig. 8　Secondary dendrite arm spacing of fcc phase as a function of the cooling rate in AlCoCrFeNi2.1 eutectic high-entropy alloy ingots.

261Solidication Microstructure of AlCoCrFeNi2.1 Eutectic High Entropy Alloy Ingots

was found to decrease as the cooling rate increased.

4.　 Discussion

4.1　 Homogeneity in solidication microstructureThe constituent phases in the AlCoCrFeNi2.1 EHEA in-

gots are a composite of fcc and bcc phases, regardless of the casting processes (air cooling or metal-mold casting). The ingot exhibits a microstructure consisting of primary fcc dendrite and an fcc+bcc eutectic structure in the inter-den-drite region, and the development of a L12 ordering structure

at the primary fcc dendrite in the metal-mold ingots. The mi-crostructure of a rapidly solidied melt-spun ribbon of AlCoCrFeNi2.1 EHEA was investigated32), where the cooling rate in the melt-spinning process was estimated to be in the order of 105 K/s33,34). The constituent phases in the melt-spun ribbons are a composite of the fcc and bcc phases, where the microstructure of the ribbon is not a typical lamella structure. This indicates that the constituent phases in the AlCoCrFeNi2.1 EHEA specimens obtained via the so-lidication route are basically fcc+bcc phases with a wide-ranging cooling rate of 1–105 K/s, and exhibit a weak cooling-rate dependence. Lu et al. reported on the lamella structure formation in ingots with a weight of 2.5 Kg11), al-though the position was not mentioned in detail. The full lamella structure was also reported for ingots with a size of 15 mm (width) × 90 mm (length) × 3 mm (thickness)35). In the present study, the development of the primary dendrite structure was not observed and a full lamella structure was seen in position ④ in the air-cooled ingots. This implies that the full lamella structure may be obtained in large specimens

Table 3　Composition analysis results of constituent phases in region ⑤ of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by met-al-mold casting obtained by EPMA-WDS and STEM-EDS.

Al [at%] Co [at%] Cr [at%] Fe [at%] Ni [at%] δ

Dendrite-Fcc 7.7 19.9 20.8 19.3 32.4 3.9

Eutectic-Fcc 5.6 18.0 21.5 19.3 35.7 3.3

Eutectic-Bcc 24.7 11.7 7.2 8.9 47.6 6.0

Fig. 9　Element mapping for position ⑤ of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by metal-mold casting. (a) SEM-EPMA mapping, (b) STEM-EDS mapping, focusing on the eutectic structure in the inter-dendrite region.

262 T. Nagase, M. Takemura, M. Matsumuro and T. Maruyama

with a slow cooling rate. Further discussion of the solidica-tion microstructure in AlCoCrFeNi2.1 EHEAs will be pur-sued in a subsequent paper. The Hv value for the AlCoCrFeNi2.1 EHEA ingots has a very low position depen-dence, as shown in Table 4 and Fig. 11. This can be ex-plained by the lack of position dependence of the constituent phases of the ingots. The properties that change with the size

of the solidication microstructure in the AlCoCrFeNi2.1 EHEA with a cooling rate of 1–500 K/s is not a dominant factor affecting the hardness, and there was no formation of inclusions or precipitates in the region near the metal-mold/ingot interface.

4.2　 Development of L12 structure in primary fcc den-drite phase

The development of the L12 ordering structure in primary fcc dendrite can be detected by TEM-SAD pattern analysis (Fig. 10), while the sharp intensity peaks corresponding to the L12 ordering structure cannot be seen in the XRD pat-terns (Fig. 5). The L12 ordering in the fcc phase is thought to be too weak to be evaluated in detail only by XRD, while the existence of chemical ordering can be detected in certain observation directions in the SAD patterns in TEM observa-tion. In the following, the development of the L12 structure in the fcc phase is discussed from the viewpoint of the alloy parameters for HEAs. The VEC values of the primary fcc dendrite phase, calculated based on an EPMA-WDS analy-sis (Table 3) was 8.0, with this value being the limit value for the formation of the fcc single phase7,18,19). However, the application of VEC theory to the phase stability of the sol-id-solution phase in HEAs focuses on the formation of the fcc phase and bcc phases, not on the formation of the chemi-cal-ordering structure in the fcc and bcc phases. The order-ing phase, such as the bcc-based B2 structure and fcc-based L12, were formed in a number of HEAs, as reported in the literature7), and these are generally called “ordered phases.” In the δ-Ω map suggested by Yang et al.17), the solid-solu-tion phase exhibits a tendency to be formed at δ ≦ 6.6 and Ω ≧ 1.1, while the ordered phase exhibits a tendency to form at 4 ≦ δ ≦ 6.6 in spite of the solid-solution formation range. In the ΔHmix − δ map suggested by Zhang et al.3), the or-dered phase is likely to form in the region of δ ≧ 4.5. These tendencies were derived from the empirical rules, with both rules exhibiting a tendency for the ordered phases to be formed at relatively high δ values. The δ value of the pri-mary fcc dendrite phase, evaluated by the chemical compo-sition analysis results, is 3.9, as shown in Table 2. This was close to the limit value for the formation of an ordered phase in the δ-Ω map7). The δ parameter is used to evaluate the dif-ference in the atomic size effect in HEAs. The L12 ordering formation in the primary fcc dendrite phase may be ex-

Fig. 10　TEM microstructure analysis results for position ⑤ of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by met-al-mold casting. (a) TEM-bright eld (BF) image, ((b1)–(b3)) [001], [011], and [111] selected area diffraction (SAD) patterns, ((c1)–(c3)) [001], [011], and [111] key diagrams of XY3 compound with L12 struc-ture, (d) schematic illustration of binary XY3 compound with L12 structure.

Fig. 11　Vickers hardness (Hv) of AlCoCrFeNi2.1 eutectic high-entropy al-loy ingots prepared by metal-mold casting as a function of the distance from the metal-mold/ingot interface.

Table 4　Position dependence of Vickers hardness (Hv) of AlCoCrFeNi2.1 eutectic high-entropy alloy ingots prepared by metal-mold casting. The average and standard deviation were evaluated using 5 data points.

①　267 ± 6.9 ④　266 ± 8.5

②　272 ± 7.2 ⑤　275 ± 4.4

③　270 ± 10 ⑥　280 ± 8.1

263Solidication Microstructure of AlCoCrFeNi2.1 Eutectic High Entropy Alloy Ingots

plained by the rejection of the Al content from the fcc phase to the residual liquid and the relaxation of the atomic size differences in the constituent elements. It should be empha-sized here that the TEM-SAD pattern analysis is vital to the clarication of the formation of the ordered phase, and the formation of the ordered phase may be predicted by using the alloy parameters for HEAs in AlCoCrFeNi2.1 and Al-Co-Cr-Fe-Ni based EHEAs.

5.　 Conclusions

The constituent phases and microstructure of an AlCoCrFeNi2.1 EHEA were investigated by trans-scale ob-servations covering millimeter-order to nanometer-order scales to evaluate the solidication microstructure in HEAs. The obtained results are summarized as follows: (1) Cast products of AlCoCrFeNi2.1 EHEA can be success-

fully obtained by high-frequency melting under an Ar ow and centrifugal casting using a metal mold.

(2) The solidication microstructure evaluation for HEAs was discussed. The heterogeneity of the constituent phases and microstructure was investigated using nine partitioned parts obtained from an ingot by using XRD (with linear and logarithmic scales) and trans-scale ob-servations using OM, SEM, EPMA-WDS, TEM, and STEM.

(3) The constituent phases in the AlCoCrFeNi2.1 EHEA were a combination of the bcc and fcc phases. This was not dependent on the position in the ingot.

(4) The microstructure of the ingot obtained by centrifugal casting using a metal mold was primary fcc dendrite, with a bcc+fcc eutectic structure in the inter-dendrite re-gions. The size of the microstructure exhibited a cool-ing-rate dependence; as the cooling rate increased, the secondary dendrite space arming and lamellar size of the eutectic structure became ner.

(5) The development of the L12-based ordering structure in the primary fcc dendrite was observed by TEM analysis.

Acknowledgements

This work was partially supported by JSPS KAKENHI (Grant Number 15K06484), by scientic grants from the Japan Foundry Engineering Society (JFS) and by the Organization for Research and Development of Innovative Science and Technology (ORDIST), Kansai University. The authors are grateful to Mr. Kazuki Takahashi of Kansai University for his assistance with the micro Vickers hardness measurements.

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