Solidification Microstructures in 3d-Transition Metal High Entropy Alloys with Cu Element

Abstract

In this study, high-entropy alloys (HEAs) with Cu as the main constituent element were investigated, focusing on the distribution of Cu in the ingots. Based on the taxonomy of HEAs, those with Cu as the casting material were classified as (1) HEAs whose main constituent elements were 3d transition metals, such as Co, Cr, Fe, Mn, Ni, and Cu (3d-HEAs), and (2) high-entropy (HE) brasses based on the Cu–Zn alloy system and HE bronzes based on Cu–Sn and/or Cu–Al alloy systems. In the case of 3d-HEAs with Cu, the distribution of Cu in the ingots exhibited the following tendency: (1-1) segregation from the dendrite to the residual liquid, resulting in the formation of Cu-rich interdendritic regions in the ingots; (1-2) liquid-phase separation resulting in the formation of a Cu-rich liquid, which formed a macroscopically phase-separated structure; and (1-3) the dispersion of fine Cu precipitates embedded in the solid solution matrix.

 

This Paper was Originally Published in Japanese in J. Japan Inst. Copper 60 (2021) 167–175. Some figures were corrected for the revision from gray-contrast images to color-contrast images. Reference 50 and Fig. 4 were slightly modified.

Fig. 2 Relationship between the main constituent elements and periodic table for various HEAs including 3d-transition-metal HEAs (3d-HEAs), refractory HEAs (RHEAs), HEAs for metallic biomaterials (BioHEAs), and lightweight HEAs (LW-HEAs), focusing on Cu elements. (a) 3d-HEAs, RHEAs, BioHEAs and LW-HEAs, (b) HE brasses, and HE bronzes.

1. Introduction

Recently, new metallic materials called high-entropy alloys (HEAs) have been rapidly developing.^1–6) HEAs are generally multicomponent alloys that consist of five or more constituent elements. Their compositions are close to the equiatomic composition or approximately equal to the equiatomic composition to increase the configuration entropy term of the alloy. From the viewpoint of casting alloys, HEAs are regarded as novel cast alloys with the following features.^7–10) (1) Ingots can be obtained using conventional melting and casting apparatus. In other words, it is not necessary to use a specific apparatus for fabricating the ingots. (2) Solid solution ingots can be obtained. A solid solution is favorable for obtaining ductile ingots. (3) High strength can be expected even for the as-cast specimens because of the severe lattice distortion effect. (4) In graphite-formation-type HEAs such as HE cast irons, casting defects can be suppressed via volume expansion owing to graphite formation during solidification. Miracle et al. investigated the trends of constituent elements used in multicomponent alloys (multi-principal element alloys, MPEAs) containing HEAs.¹¹⁾ Figure 1 shows the frequency of elements used as constituents in 408 MPEAs focusing on Cu, as reported in the literature¹¹⁾ in 2017. Cu is frequently used as a constituent element of MPEAs and HEAs. Figure 2 shows the relationship between the periodic table and main constituent elements for various HEAs. Figure 2(a) shows a diagram focusing on 3d-transition-metal HEAs (3d-HEAs), whose constituent elements are mainly 3d-transition-metal elements (a typical 3d-HEA is the Cantor alloy of CoCrFeMnNi¹⁾), refractory HEAs (RHEAs),^12–18) HEAs for metallic biomaterials (BioHEAs),^19–24) and lightweight HEAs (LW-HEAs).^25–29) Cu does not tend to be a main constituent element in RHEAs, BioHEAs, and LW-HEAs. In a previous review,¹⁶⁾ a list of RHEAs was shown, which indicated that there are very few RHEA alloy systems containing Cu. Lightweight elements, such as Al, Ti, Mg, and Li, are typically used as the main constituent elements of LW-HEAs, whereas Cu with a large atomic weight is rarely used as a main constituent element. In contrast, Cu is one of the main constituent elements (Co, Cr, Fe, Mn, Ni, Cu, and Al) of 3d-HEAs. The characteristics of the Cu element distribution in the ingots of 3d-HEAs include ① segregation during solidification and enrichment at the interdendritic regions, including the enrichment of Cu and Mn element through the decomposition of the remaining liquid phase at the interdendritic region via peritectic reactions,^30–33) ② liquid-phase separation and formation of a Cu-rich liquid,^33–44) and ③ dispersion of Cu-rich fine precipitates embedded in the solid solution matrix.^45–47) Figure 2(b) shows a diagram focusing on high-entropy (HE) brasses and HE bronzes.^48–50) HE brasses are multicomponent alloys based on Cu–Zn alloy systems, whereas HE bronzes are multicomponent alloys based on Cu–Sn tin-bronzes and/or Cu–Al aluminum-bronzes. The 3d HEAs with Cu are a combination of Cu and elements on the left side of Cu in the periodic table (Fig. 2(a)). However, HE brasses and HE bronzes are multicomponent alloys consisting of a combination of Cu and elements on the right side of Cu in the periodic table (Fig. 2(b)). In this study, the following alloy systems were evaluated based on the distribution of Cu in their ingots (Fig. 2(a)). (A) CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs,^33,44) which are typical examples of alloys exhibiting the segregation and enrichment of Cu at the interdendritic regions, including the enrichment of Cu and Mn element through the decomposition of the remaining liquid phase at the interdendritic region via peritectic reactions, were investigated. (B) CoCrFeMnNiCu_yB_0.2 (y = 1, 2, 3) HEAs,^33,44) which are typical examples of alloys exhibiting liquid-phase separation, and (C) AlCoFeNiCu_0.3 HEAs,⁵¹⁾ which are typical examples of alloys in which Cu-rich nanoscale precipitates are embedded in a solid solution matrix, were evaluated.

Fig. 1

Frequency of using elements as constituents in 408 multi-principal element alloys, reported in Ref. 11) in 2017.

Fig. 2

Relationship between the main constituent elements and periodic table for various HEAs including 3d-transition-metal HEAs (3d-HEAs), refractory HEAs (RHEAs), HEAs for metallic biomaterials (BioHEAs), and lightweight HEAs (LW-HEAs), focusing on Cu elements. (a) 3d-HEAs, RHEAs, BioHEAs and LW-HEAs, (b) HE brasses, and HE bronzes.

2. Materials and Methods

To clarify the enrichment of Cu in the interdendritic region of the ingots of 3d-HEAs with Cu, the CoCrFeMnNiCu_x HEAs were investigated. CoCrFeMnNiCu_x HEAs were designed as a combination of equiatomic CoCrFeMnNi (Cantor alloy), with a considerably high face-centered cubic (FCC) solid-solution formation ability,¹⁾ and Cu. Table 1 shows the atomic composition ratios of the constituent elements and the x-dependence of the mixing entropy (ΔS_mix) of the CoCrFeMnNiCu_x HEAs. The parameter ΔS_mix reflects the mixing entropy of the ideal and regular solutions and is expressed as follows:   

\begin{equation} \Delta S_{\textit{mix}} = -R \sum\nolimits_{i=1}^{n} x_{i}\ln (x_{i}), \end{equation} (1)

where R is the gas constant, and x_i is the molar fraction of the i^th component. In the classification of alloys based on the mixing entropy ΔS_mix,¹⁾ HEAs are defined as solid solution alloys with ΔS_mix ≥ 1.5R, and is satisfied at x < 5 in CoCrFeMnNiCu_x alloys. In this study, ingots of the alloys were prepared with x = 1 as the equiatomic-, and x = 2 and 3 as the nonequiatomic-composition-ratio alloys. Ingots of CoCrFeMnNiCu_x alloys were fabricated from a mixture of pure Co, Cr, Fe, Mn, Ni, and Cu fragments using an arc-melting method. The cooling rate in the arc-melting process was estimated to be approximately 2000 K/s by measuring the secondary dendritic arm spacing in the Al–Cu alloys.⁹⁾

Table 1 Alloy abbreviations, nominal atomic ratio, and mixing enthalpy (ΔS_mix) of CoCrFeMnNiCu_x Alloys. R is the gas constant.

CoCrFeMnNiCu_yB_0.2 (y = 1, 2, 3) HEAs as a combination of CoCrFeMnNiCu_x (x = 1, 2, 3) and B³³⁾ were designed as Cu-containing 3d-HEAs with liquid-phase separation. B is typically added to enhance the liquid-phase separation in binary Fe–Cu⁵¹⁾ and Co–Cu⁵²⁾ alloy systems. Ingots of CoCrFeMnNiCu_xB_0.2 (x = 1, 2, 3) HEAs were also fabricated by arc melting a mixture of pure Co, Cr, Fe, Mn, Ni, Cu, and B fragments. Table 2 shows the atomic composition ratios and ΔS_mix of the CoCrFeMnNiCu_xB_0.2 (x = 1, 2, 3) HEAs.

Table 2 Alloy abbreviations and nominal atomic ratio of CoCrFeMnNiCu_yB_0.2 Alloys. R is the gas constant.

The solidification microstructure of an ingot obtained by crucible melting of an AlCoFeNiCu_0.3 HEA⁵³⁾ was investigated; the ingot was considered as a typical alloy bulk specimen composed of Cu-rich nanoscale precipitates and a solid solution matrix. Table 3 shows the atomic composition and empirical alloy parameters of ΔS_mix, ΔH_mix,^5,6) and δ(ΔH_mix)²⁹⁾ related to the mixing enthalpy, δ parameter reflecting the distribution of atomic radii among the constituent elements, and dimensionless parameter Ω^5,6) as the ratio of mixing entropy and enthalpy in the AlCoFeNiCu_0.3 alloy.⁵³⁾ The parameters ΔH_mix, δ(ΔH_mix), δ, and Ω are expressed as follows:   

\begin{equation} \Delta H_{\textit{mix}} = \sum\nolimits_{i=1}^{n} \sum\nolimits_{j,j\neq i}^{n} 4 \cdot x_{i} \cdot x_{j} \cdot \Delta H_{i\text{–}j}, \end{equation} (2)

  

\begin{equation} \delta (\Delta H_{\textit{mix}}) = \sum\nolimits_{i=1}^{n} \sum\nolimits_{j,j\neq i}^{n} 4 \cdot x_{i} \cdot x_{j} \cdot |\Delta H_{i\text{–}j} - \Delta H_{\textit{mix}}|, \end{equation} (3)

  

\begin{equation} \delta = 100\times \sqrt{\sum\nolimits_{i=1}^{n} x_{i} \cdot \left(1 - \frac{r_{i}}{\bar{r}} \right)^{2}}, \end{equation} (4)

  

\begin{equation} \varOmega = \frac{\overline{T_{m}} \cdot \Delta S_{\textit{mix}}}{|\Delta H_{\textit{mix}}|}, \end{equation} (5)

where ΔH_i–j is the mixing enthalpy of the binary i–j atomic pair, and r_i is the atomic radius of the i^th element. The values of ΔH_i–j and r_i were obtained from the literature.⁵⁴⁾ $\bar{r}$ and $\overline{T_{m}}$ are the compositional average values of the atomic radius and the melting point of the pure elements, respectively. ΔH_mix is an empirical alloy parameter used to evaluate the interactions of the constituent elements. The closer ΔH_mix is to zero, the smaller is the interaction between the constituent elements in an alloy and the higher is the ability to form a solid solution. δ(ΔH_mix) is an empirical alloy parameter used to evaluate the dispersion of ΔH_i–j. The closer δ(ΔH_mix) is to zero, the smaller is the dispersion of ΔH_i–j among the constituent elements and the higher is the ability to form a solid solution through the suppression of the liquid-phase separation. δ is an empirical parameter for examining the atomic radius difference between the constituent elements. The closer δ is to zero, the smaller is the atomic radius difference between the constituent elements, the lower is the tendency to form intermetallic compounds and ability to form glass, and the higher is the ability to form a solid solution. Ω is a dimensionless parameter obtained from the entropy-to-enthalpy ratio. A higher Ω implies a higher ability to form a solid solution. The empirical parameters indicate a high solid-solution formation tendency in the AlCoFeNiCu_0.3 HEA. Crucible melting (1873 K for 5 min) of the AlCoFeNiCu_0.3 HEA was performed using a mixture of pure element fragments by high-frequency melting with alumina crucibles. The cooling rate during the melting-cooling process in the crucible was estimated from thermocouple measurements to be approximately 4 K/s immediately before solidification started, and the average cooling rate during solidification was approximately 0.7 K/s. X-ray diffraction (XRD) measurements were performed using a Rigaku RINT2500 diffractometer with Cu-Kα radiation. Specimens for scanning electron microscopy (SEM; JEOL JEM-5600 (W-filament thermionic gun)) and electron probe microanalysis (EPMA)-wave dispersive spectroscopy (WDS; JEOL JXA-8800R (W-filament thermionic gun)) were prepared by polishing using a SiC sandpaper, followed by buffing using an Al₂O₃ polishing agent. Transmission electron microscopy (TEM; Hitachi H-800 (W-filament thermionic gun)) and scanning transmission electron microscopy (STEM; JEOL JEM-2100F (field-emission gun)) specimens were prepared by Ar ion milling.

Table 3 Alloy parameters for AlCoFeNiCu_0.3 HEA alloy: ΔS_mix, ΔH_mix [kJ/mol], δ(ΔH_mix) [kJ/mol], δ, and Ω. R is the gas constant.

3. Results and Discussion

3.1 3d-HEAs with Cu-rich interdendritic regions

Figure 3 shows XRD patterns obtained from the central part of the arc-melted ingots of the CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. The sharp diffraction peaks originated from two FCC phases (● and ○) with different lattice parameters, while no diffraction peaks are observed corresponding to intermetallic compounds or chemically ordered structures based on the FCC structure.

Fig. 3

XRD patterns of the arc-melted ingots in CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. Figure adapted from Ref. 33).

Figure 4 shows SEM-back scattering electron (BSE) images of the arc-melted ingots of the CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. The dendrite structure composed of the gray-contrast dendritic regions and white-contrast interdendritic regions is observed regardless of the alloy composition. The area ratio of interdendritic/dendritic regions increased with the increase in x in the CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs.

Fig. 4

SEM-BSE images of the arc-melted ingots with CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. (a) CoCrFeMnNiCu (x = 1), (b) CoCrFeMnNiCu₂ (x = 2), (c) CoCrFeMnNiCu₃ (x = 3).

Figure 5 shows EPMA-WDS element maps of the arc-melted ingots as the elemental distributions of the CoCrFeMnNiCu (x = 1) HEAs as the equiatomic HEAs and CoCrFeMnNiCu₃ (x = 3) HEAs as the typical examples of the nonequiatomic HEAs. Table 4 shows the chemical alloy composition analysis results of the dendritic and interdendritic regions obtained by WDS. Co, Cr, and Fe were enriched at the dendritic regions, while Mn and Cu concentrated at the interdendritic regions. In the CoCrFeMnNiCu_x HEAs, the solubility of Cu in the dendritic regions did not increase monotonously when x increased. The increase in the concentration of Cu in the CoCrFeMnNiCu_x alloys led to an increase in the area ratio of interdendritic/dendritic regions. Based on the composition analysis results (Table 4) and lattice constants in Materials Project^55,56) regarding pure Co,⁵⁷⁾ Cr,⁵⁸⁾ Fe,⁵⁹⁾ Mn,⁶⁰⁾ Ni,⁶¹⁾ and Cu⁶²⁾ with FCC structures, the lattice constants of the dendritic and interdendritic regions were calculated using the Vegard’s law. The lattice constant of the interdendrite was estimated to be larger than that of the dendrite region. In Fig. 3, the FCC phases indicated by the black-filled circles (●) and black-open circles (○) were considered to correspond to Cu-rich interdendritic regions and Cu-poor dendritic regions, respectively. Table 5 shows the distribution coefficient (k) between the FCC and liquid phases at the liquidus temperature (T_L) in the CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. The values of k and T_L were estimated by thermodynamic calculations using FactSage,⁶³⁾ with SGTE2017⁶⁴⁾ as the thermodynamic database. k_i was defined as the ratio of the concentration of the i^th element in the solid phase (C_S,i) to that in the liquid phase (C_L,i), that is, (C_S,i)/(C_L,i). Liquid separation was not considered in the thermodynamic calculations. The values of k for Cu and Mn were lower than unity, while those for Co, Cr, and Fe were higher than unity. The enrichment of Cu and Mn at the interdendritic regions rather than at the dendritic regions (Figs. 3, 4, Table 4) can be explained qualitatively by the thermodynamic calculations (Table 5).

Fig. 5

EPMA-WDS element mapping images of the arc-melted ingots with CoCrFeMnNiCu_x (x = 1, 3) HEAs. (a) CoCrFeMnNiCu (x = 1) alloy as the equiatomic HEAs. (b) CoCrFeMnNiCu₃ (x = 3) alloy as the typical example of the nonequiatomic HEAs.

Table 4 Results of the chemical composition analysis of the dendritic (D) and interdendritic (ID) regions of the alloys by WDS. (a) CoCrFeMnNiCu (x = 1) alloy as the equiatomic HEAs. (b) CoCrFeMnNiCu₃ (x = 3) alloy as the typical example of the nonequiatomic HEAs.

Table 5 Distribution coefficient (k) between the FCC phase and the liquid phase at the liquidus temperature (T_L) in CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs. The value of k and T_L were estimated by thermodynamic calculations completed using FactSage and SGTE2017.

3.2 3d-HEAs with macroscopically phase-separated Cu-rich regions via liquid-phase separation

The tendency of Cu segregation was stronger when a small amount of B was added to the 3d-HEAs with Cu.³³⁾ The addition of B enhanced the liquid-phase separation to form Cu-rich liquids in 3d-HEAs with Cu, resulting in the formation of macroscopically phase-separated structures. Figure 6 shows the XRD patterns of the arc-melted ingots of CoCrFeMnNiCu_xB_0.2 (x = 1, 2, 3) HEAs.³³⁾ The XRD patterns were obtained from the central regions of the arc-melted ingots. The sharp diffraction peaks in Fig. 6 can be assigned to the two FCC phases with different lattice constants (● and ○). Figure 7 shows the SEM-BSE images of the arc-melted ingots of the CoCrFeMnNiCu_xB_0.2 (x = 1, 2, 3) HEAs. The macroscopically phase-separated structure is composed of a gray-contrast matrix (upper side) and white-contrast matrix (lower side) in the CoCrFeMnNiCu_xB_0.2 (x = 2, 3) HEAs (Figs. 7(b), 7(c)), which indicates liquid-phase separation. Figure 8 shows the EPMA-WDS elemental maps of the CoCrFeMnNiCu_xB_0.2 (x = 3) alloys as typical examples of a macroscopically phase-separated structure formed via liquid separation in CoCrFeMnNiCu_xB_0.2 HEAs. The Cu-rich region with white contrast in the BSE image were rich in Mn, whereas the Cu-poor regions with gray contrast in the BSE image were rich in Co, Cr, and Fe. B tends to concentrate in plate-like Cr-rich phases dispersed in the Cu-poor phase and forms Cr–B-rich compounds.³³⁾

Fig. 6

XRD patterns of the arc-melted ingots in CoCrFeMnNiCu_xB_0.2 (x = 1, 2, 3) HEAs. Figure adapted from Ref. 33).

Fig. 7

SEM-BSE images of the arc-melted ingots in CoCrFeMnNiCu_yB_0.2 (y = 1, 2, 3) HEAs. (a) CoCrFeMnNiCuB_0.2 (y = 1), (b) CoCrFeMnNiCu₂B_0.2 (y = 2), (c) CoCrFeMnNiCu₃B_0.2 (y = 3).

Fig. 8

EPMA-WDS element mapping images of the arc-melted ingots with CoCrFeMnNiCu₃B_0.2 (y = 3) HEA.

Figure 9 shows the values of the mixing enthalpy of the i–j atomic pairs (ΔH_i–j) [kJ/mol] for Co, Cr, Fe, Mn, Ni, and Cu, which are the constituent elements of the CoCrFeMnNiCu_x HEAs. The values of ΔH_i–j were obtained from the literature.⁵⁴⁾ The values of ΔH_i–j (i = Cu, j = Cr, Mn, Fe, Co, Ni) are positive, which indicates that Cu exhibits a repulsive nature against the other constituent elements of Co, Cr, Fe, Mn, and Ni in the CoCrFeMnNiCu_x HEAs. In the ΔH_i–j relationship diagram (Fig. 9(b)), the value for Co–Cr–Fe–Ni is zero or negative, while those for Cu and Mn are positive with respect to Co–Cr–Fe–Ni. The two groups, Co–Cr–Fe–Ni and Cu–Mn, can be considered to correspond to the dendrite structures with Co–Cr–Fe–Ni-rich dendritic and Cu–Mn-rich inter-dendritic regions (Table 4, Figs. 3, 4, 5) and the macroscopically phase-separated structure consisting of the Co–Cr–Fe–Ni-rich and Cu–Mn-rich phases with liquid-phase separation (Figs. 6, 7, 8), respectively.

Fig. 9

Value of the mixing enthalpy of i–j atomic pairs (ΔH_i–j) [kJ/mol]. (a) ΔH_i–j matrix, (b) ΔH_i–j relationship map. The values of ΔH_i–j were found in the literature.⁶⁴⁾

Figure 10 shows the liquid miscibility gap in the CoCrFeMnNiCu_yB_z (0 ≤ y ≤ 4) (z = 0, 0.2, 0.4) alloys, as estimated by thermodynamic calculations. Thermodynamic calculations were performed using FactSage⁶³⁾ and SGTE2017⁶⁴⁾ thermodynamic databases. Non-B, 02 B, and 04 B denote z = 0, 0.2, and 0.4 alloys, respectively. Thermodynamic calculations showed that the liquid-phase separation temperature increased with increasing B concentration regardless of the Cu concentration in the CoCrFeMnNiCu_yB_z HEAs. The occurrence of liquid-phase separation in the CoCrFeMnNiCu_yB_0.2 (y = 2, 3) HEAs (Figs. 7, 8) roughly corresponds to the thermodynamic calculation results showing that the addition of B enhances liquid-phase separation in 3d-HEAs with Cu (Fig. 10).

Fig. 10

Liquid miscibility gap in CoCrFeMnNiCu_yB_z (0 ≦ y ≦ 4) (z = 0, 0.2, 0.4) alloys calculated by FactSage and SGTE2017. 0B, 0.2B, and 0.4B denote z = 0, z = 0.2, and z = 0.4 alloys, respectively.

3.3 3d-HEAs with fine Cu-rich precipitates

In 3d-HEAs with Cu, the enrichment of Cu at the interdendritic region during solidification^30–33) (typical examples are CoCrFeMnNiCu_x HEAs shown in Figs. 3, 4, and 5, and Table 4) and/or the formation of Cu-rich separated liquids via liquid-phase separation^34–44) (typical examples are CoCrFeMnNiCu_xB_0.2 HEAs shown in Figs. 6, 7, and 8) can lead to macroscopic and/or microscopic-scale segregation in the ingots. This severe segregation tendency is not suitable for the development of 3d-HEAs with Cu as the casting material. The values of ΔH_i–j (Fig. 9) indicate repulsive interactions between Cu (i = Cu) and Co, Ce, Fe, Mn, and Ni (j = Co, Ce, Fe, Mn, Ni), which are the typical constituent elements of 3d-HEAs. The difficulty in the formation of a solid solution phase with a high Cu concentration in the ingots without macroscopic and microscopic segregation and/or liquid-phase separation was predicted by ΔH_i–j. The formation of nanoscale Cu-rich phase precipitates embedded in a solid solution matrix by controlling the constituent elements and Cu composition can enhance the mechanical strength of ingots of 3d-HEAs with Cu, even if a solid solution with a high Cu concentration is not obtained. In this study, the solidification microstructures of AlCoFeNiCu_0.3 HEAs⁵³⁾ in ingots prepared by conventional crucible melting and casting are reported as typical examples of composites in which Cu-rich nanoscale precipitates are embedded in a solid solution matrix.

Figure 11 shows the XRD pattern of the AlCoFeNiCu_0.3 HEA ingot obtained from the central region, indicated by the arrow in the inset cross-sectional optical microscopy image. The sharp diffraction peaks can be indexed to the main body-centered cubic (BCC) phase (■) and the minor FCC phase (○). The XRD pattern of the ingot of the AlCoFeNiCu_0.3 HEA exhibited a strong cooling rate dependence, and the peak intensity ratio of the FCC and BCC phases was dependent on the cooling rate, even in the same metallic mold casting ingots.⁵³⁾

Fig. 11

XRD patterns of AlCoFeNiCu_0.3 HEA ingots obtained from the central region indicated by the arrow in the inset cross-section optical microscopy image of the ingot. The ingots were prepared by high-frequency induction melting using silica-based crucibles.

Electron microscopy observations were carried out using a sample of the central part of the ingot, as indicated by the black arrow in Fig. 11. In the SEM-BSE image acquired using the conventional SEM with a W-filament thermionic gun, a typical equiaxial dendrite structure similar to that shown in Fig. 4 and the typical liquid-phase separation structure similar to that shown in Fig. 6 are not observed. Figure 12 shows the EPMA-WDS element maps of the ingot of the AlCoFeNiCu_0.3 HEA. Open spaces corresponding to the interdendritic regions were observed in the secondary electron (SE) images. A dendrite structure developed in the central area of the ingots. In the BSE image, the microstructure corresponding to segregation based on the dendrite structure was not observed. The elemental distribution corresponding to the development of the dendrite structure was not observed in the WDS element maps. No macroscopic- and/or microscopic-scale Cu-rich phases were observed by SEM and EPMA at the central regions of the ingots.

Fig. 12

EPMA-WDS element mapping images of AlCoFeNiCu_0.3 HEA obtained from the central region indicated by the arrow in the inset cross-section image of the ingot shown in Fig. 11.

Figure 13 shows TEM results of the ingot of the AlCoFeNiCu_0.3 HEA. The bright-field (BF) image shows the microstructure composed of a matrix and nanoscale precipitates (Fig. 13(a)). Figures 13(b1) and 13(b2) show selected-area electron diffraction (SAED) patterns along the [001] BCC and [011] BCC zone axes, respectively. The SAED patterns (Figs. 13(b1) and 13(b2)) can be explained by key diagrams of the B2 structure shown in Figs. 13(c1) and 13(c2). The matrix in the AlCoFeNiCu_0.3 HEA ingot contained the chemically ordered structure related with B2, where the formation of B2 ordering cannot be identified by the XRD pattern (Fig. 11) because of the significantly low intensity of the ordering diffraction peaks.

Fig. 13

TEM observation of the AlCoFeNiCu_0.3 HEA ingots. The specimens were obtained from the central region indicated by the arrow in the cross-section image of the ingot shown in Fig. 11. (a) BF image, (b1) (b2) SAED patterns, (c1) (c2) key diagrams. (b1) (c1) SAED pattern and key diagram along the [001]BCC zone axis. (b2) (c2) SAED patterns and key diagram along the [011]BCC zone axis.

The development of the B2 ordering structure in the BCC matrix can be detected by the B2 ordering spots in the SAED patterns (● in Figs. 13(c1) and 13(c2)). In the electron diffraction pattern along the [011] BCC direction (Fig. 13(b2)), diffraction spots that could not be explained by the BCC and B2 structures were observed, which were considered to correspond to the nanoscale precipitates. Figure 14 shows STEM images and STEM–energy-dispersive spectroscopy (EDS) element maps of the ingot of AlCoFeNiCu_0.3 HEA. The nanoscale precipitates exhibited white contrast in the high-angle annular dark-field (HAADF) STEM images. The STEM-EDS elemental maps showed that Cu was enriched in the precipitates. Based on the SEM, EPMA (Fig. 12), TEM (Fig. 13), and STEM (Fig. 14) results, the composite structure composed of the main BCC phase and minor FCC phase identified by the XRD patterns (Fig. 11) was considered to correspond to the BCC/B2 matrix with nanoscale Cu-rich precipitates.

Fig. 14

STEM observation and STEM-EDS element mapping images of the AlCoFeNiCu_0.3 HEA ingots. The specimens were obtained from the central region indicated by the arrow in the cross-section image of the ingot shown in Fig. 11. HAADF is the high-angle annular dark-field image.

The main phase of the ingot of the AlCoFeNiCu_0.3 HEA was the BCC phase, whereas that of the CoCrFeMnNiCu_x HEAs was the FCC phase. The addition of Al to 3d-HEAs stabilizes the BCC structure.^2,4–6) The valence electron concentration (VEC) is an effective indicator for predicting the stability of the BCC/FCC phases in HEAs.^5,6)   

\begin{equation} \mathit{VEC} = \sum\nolimits_{i=1}^{n} x_{i}\cdot \mathit{VEC}_{i}, \end{equation} (6)

where VEC_i is the VEC of the pure component i. In 3d-HEAs and RHEAs, the FCC structure tends to form at VEC ≥ 8, the BCC structure tends to form at VEC < 6.87, and the FCC + BCC structure tends to form in the region of 6.87 ≤ VEC < 8. Table 6 shows the VEC values of the CoCrFeMnNiCu_x (x = 1, 2, 3) and AlCoFeNiCu_0.3 HEAs. The VEC values of the CoCrFeMnNiCu_x HEAs (x = 1, 2, 3) were higher than 8. The formation of FCC phases in the ingots of the CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs corresponds to the prediction by the VEC calculation. The formation of the BCC + FCC phases in the ingot of the AlCoFeNiCu_0.3 HEA was also consistent with the structure prediction by the VEC. The prediction of the B2 ordering structure in HEAs has been challenging^65–69) because of the difficult evaluation of the B2 ordering structure by XRD analysis and the lack of a thermodynamic database concerning the ordering structures in multicomponent alloy systems. The prediction of the formation of nanoscale Cu-rich precipitates embedded in the BCC matrix with a B2-ordered structure in the ingot of the AlCoFeNiCu_0.3 HEA and advanced microstructure control are aimed as future tasks.

Table 6 VEC values of CoCrFeMnNiCu_x (x = 1, 2, 3) and AlCoFeNiCu_0.3 HEAs.

The constituent phase and solidification structure in the ingots of the 3d-HEAs with Cu were strongly dependent on the constituent elements. The distribution of Cu and morphology of the Cu-rich phases were also strongly dependent on the constituent elements. Cu is frequently used as one of the main constituent elements of 3d-HEAs (Figs. 1 and 2). The values of ΔH_i–j among Cu and other elements of Co, Cr, Fe, Mn, and Ni (i = Cu, j = Co, Cr, Fe, Mn, Ni) were positive (Fig. 9). This corresponds to the formation of macroscopic and/or microscopic phase-separated structures, such as the enrichment of Cu at the interdendritic regions in the dendrite structure (Figs. 1, 2, 3, Tables 4, 5), because of the significant segregation tendency and formation of Cu-rich liquids via liquid-phase separation (Figs. 6, 7, and 8). Various combinations of the constituent elements and alloy compositions are available in HEAs owing to the multiplicity of the constituent elements. Furthermore, the solidification microstructures in HEAs can be controlled by the process parameters, such as the cooling rate. For example, the control of the constituent elements and alloy composition through the addition of Al and a decrease in Cu concentration from the equiatomic composition realizes the solidification microstructure control for the formation of a composite composed of a BCC matrix and nanoscale Cu-rich precipitates without a macroscopic phase-separated structure in the AlCoFeNiCu_0.3 HEA (Figs. 11, 12, 13, 14). This study demonstrates that alloy design with a combination of thermodynamic calculations and alloy parameters of ΔH_i–j and VEC is effective in controlling the solidification microstructure in the ingot, resulting in microstructure customization of the morphology of Cu-rich phases in 3d-HEAs with Cu.

4. Conclusion

The relationship between Cu and HEAs, in particular the distribution of Cu in the ingots of 3d-HEAs with Cu, was investigated. The findings of this study can be summarized as follows:

1. (1)    Cu is frequently used as a constituent element of HEAs.
2. (2)    Cu tends to be the main constituent element of 3d-HEAs, HE brasses, and HE bronzes. However, Cu is not used as the main constituent element in RHEAs, BioHEAs, and LW-HEAs.
3. (3)    In the 3d-HEAs with Cu, Cu exhibit the following tendencies in the ingots. (A) Cu is enriched in the interdendritic region via segregation when the dendrite structure is formed. (B) A macroscopic phase-separated structure with Cu-rich regions was formed when liquid-phase separation occurred in the thermal melt to form a Cu-rich liquid. (C) Nanoscale Cu-rich precipitates embedded in the solid solution matrix were formed when a significant segregation did not occur during solidification.
4. (4)    The ingots of CoCrFeMnNiCu_x (x = 1, 2, 3) HEAs exhibited dual FCC phases with an equiaxial dendrite structure. Cu was enriched in the interdendritic regions. The area ratio of the interdendritic/dendritic regions increased with increasing x. The segregation behavior of Cu could be explained using thermodynamic calculations.
5. (5)    In the ingots of the CoCrFeMnNiCu_yB_z (y = 1, 2, 3) (z = 0, 0.2) HEAs, liquid-phase separation was enhanced by the addition of B. The occurrence of liquid-phase separation corresponds to the mixing enthalpy of ΔH_i–j among the constituent elements and the thermodynamic calculation results.
6. (6)    The constituent phases of the ingot of the AlCoFeNiCu_0.3 HEA were the main BCC phase and minor FCC phase. The Cu-rich phases were considered to be the FCC phase as nanoscale precipitates embedded in the BCC matrix. The SAED pattern obtained from the BCC matrix exhibited B2 ordering spots.

Acknowledgments

This study was partially supported by Japan Society for the Promotion of Science KAKENHI (grant numbers 18K04750, 19H05172, 21H00146, 22H01816) and scientific grants received from the Japan Copper and Brass Association. A part of this study was carried out under the interuniversity cooperative research program (proposal numbers 18G0036, 19G0035, 20G0020, 202012-CRKKE-0009, 202112-CRKKE-0024) of the Cooperative Research and Development Center for Advanced Materials, Institute for Materials Research, Tohoku University.

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