Fabrication of the Casting Products in Cu–Zn–Mn–Ni Medium-Entropy Brasses

Abstract

The ingots of Cu_xZnMnNi (x = 1, 2) medium-entropy (ME) brasses were fabricated using metallic mold-casting process without a vacuum chamber. The molten metal was obtained by high-frequency melting of the mixture of pure Cu, pure Ni, and pre-alloy ingots of Mn–Cu and Zn–Ni using silica-based crucible in Ar flow. The metallic mold-casting ingots were obtained using centrifugal casting in air atmosphere. The composite of body-centered-cubic (BCC) and face-centered-cubic (FCC) phases were obtained in the ingots of equiatomic CuZnMnNi ME brass, while a near-single FCC phase was obtained in the ingots of non-equiatomic Cu₂ZnMnNi ME brass, where the identification of the constituent phases was mainly performed by XRD analysis. The ingots showed superior deformability and high 0.2% proof stress during compression test conducted at room temperature.

 

This Paper was Originally Published in Japanese in J. Japan Institute of Copper 59 (2020) 24–31. Minor corrections in abstract, main text, Figure and Table captions were performed with the translation from Japanese to English and proofreading by native speakers. Reference 43) was updated. An appendix containing figures based on the original and new experimental results has been added to the paper.

1. Introduction

High-entropy alloys (HEAs) are widely studied as a new category of structural and functional materials.^1–17) In general, HEAs are multi-component alloys that exhibit the following characteristics: the presence of more than five constituent elements, composition similar to equiatomic ratio, and the presence of a solid solution phase. Some of the important effects in HEAs are suggested as the followings:^8,9)

1. (1)    High entropy effect,
2. (2)    Severe lattice distortion effect,
3. (3)    Sluggish diffusion effect, and
4. (4)    Cocktail effect.

Considering that HEAs are cast materials, they exhibit the following advantages:^14,17–20)

1. (i)    the ingots can be fabricated using conventional casting method without special techniques and equipment,
2. (ii)    ductile materials can be obtained due to the formation of a solid solution,
3. (iii)    high-strength materials are formed as a result of severe lattice distortion, and
4. (iv)    solidification shrinkage was reduced by the formation of graphite in graphite-formation type HEAs.

Although HEAs are defined based on their various properties, the entropy-based definition of HEAs, using the mixing entropy (configurational entropy) of the ideal solution and regular solution,^8,9) is widely accepted. It is written by the following equation:   

\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 mole fraction of the i-th element. HEAs, medium-entropy alloys (MEAs), and low-entropy alloys (LEAs) are defined as follows:   

\begin{equation} \Delta S_{\textit{mix}} \geq 1.5R\ (\mathit{HEA}) \end{equation} (2)

  

\begin{equation} 1.0R \leq \Delta S_{\textit{mix}} \leq 1.5R\ (\mathit{MEA}) \end{equation} (3)

  

\begin{equation} \Delta S_{\textit{mix}} \leq 1.0R\ (\mathit{LEA}) \end{equation} (4)

In this study, the alloys were classified according to the aforementioned entropy-based equations. Cu–Zn-based alloys with ΔS_mix ≥ 1.5R are defined as High-Entropy Brasses (HE Brasses) and those with 1.0R ≤ ΔS_mix ≤ 1.5R are defined as Medium-Entropy Brasses (ME Brasses). Further, Brasses with ΔS_mix ≤ 1.0R are Low-Entropy Brasses (LE Brasses).

The concept of HE Brasses and HE Bronzes was suggested by Law et al. in 2015.²¹⁾ They reported Cu–Mn–Ni–Al alloys, Cu–Mn–Ni–Sn alloys, Cu–Mn–Ni–Zn alloys, Cu–Mn–Ni–Al–Sn alloys, and Cu–Mn–Ni–Zn–Sn alloys, which were designed as the combination of the equiatomic CuMnNi alloy and Cu–Zn alloys, and Cu–Sn alloys and Cu–Al alloys, where Cu–Zn, Cu–Sn, and Cu–Al was considered as the basic alloy systems of brass, bronze and aluminum-bronze, respectively. We demonstrate that the ingots of HE and ME Brasses can be fabricated using conventional casting process with silica-based crucible melting in an inert gas and metallic mold casting in Cu_xZnMnNi (x = 1, 2, 3 and 4) ME Brasses and Cu_xZnMnNiSn_y (x = 1 and y = 0.2, x = 2 and y = 0.45) HE Brasses.²²⁾ The metallic mold-casting ingots in Cu_xZnMnNi (x = 1, 2, 3 and 4) ME Brasses and Cu_xZnMnNiSn_y (x = 1 and y = 0.2, x = 2 and y = 0.45) HE Brasses showed high mechanical strength and ductility.²²⁾

Figure 1 shows 0.2% proof stress of the various copper-based casting alloys, including Cu_xZnMnNi ME Brasses and Cu_xZnMnNiSn_y HE Brasses, as a function of ΔS_mix.¹⁷⁾ The black cross-mark (×), open triangle (△), and filled upside-down triangle (▼) indicate conventional copper castings,²³⁾ brass castings,²³⁾ and bronze casting,²³⁾ respectively. Filled-square (blue ■) denotes ME Brass and bronze castings reported by Laws et al.²¹⁾ Open squares and circles (blue □, and red ○) represent ME and HE Brasses reported by Nagase et al.²²⁾ The A and B groups indicate Cu–Zn–Mn–Ni ME brass castings^21,22) and Cu–Zn–Mn–Ni–Sn HE brass castings,²²⁾ respectively. Almost all copper, brass, and bronze castings used as the industrial materials belong to LEAs. A few brasses, including high strength brass such as CAC304C (Japanese Industrial Standard, JIS), correspond to MEAs. The proof stress tends to increase with an increase in the ΔS_mix of copper, brass, and bronze castings. This corresponds to the increase in the number and composition of non-Cu constituent elements in Cu alloys. For example, the proof stress of Cu-based alloy castings increased from copper castings to brass castings to high strength brass castings. The Cu-based alloy castings whose ΔS_mix was more than 1.5R were not found in Cu-based casting alloys defined by JIS.²³⁾ ME bronzes of CuMnNiSn_0.33 and CuMnNiAl_0.3 (blue, ■),²¹⁾ and HE brasses of Cu_xZnMnNiSn_y (x = 1 and y = 0.2, x = 2 and y = 0.45) (red ○)²²⁾ show higher strength than the conventional copper castings (×),²³⁾ brass castings (△),²³⁾ and bronze castings (▼).²³⁾ An increase in ΔS_mix in Cu–Zn-based brasses and Cu–Sn-based and/or Cu–Al-based bronzes, i.e., the development of ME and HE brasses and/or bronzes is a new strategy for the development of high strength Cu-based alloy castings.

Fig. 1

0.2% proof stress of the various copper-based casting alloys, including Cu_xZnMnNi ME Brasses and Cu_xZnMnNiSn_y HE Brasses, as a function of the mixing entropy ΔS_mix. The black cross-mark (×), open triangle (△), and filled upside-down triangle (▼) indicate conventional copper castings,²³⁾ brass castings,²³⁾ and bronze casting,²³⁾ respectively. Filled-square (■) denotes ME Brass and bronze castings reported by Laws et al.²¹⁾ Open squares and circles (□ and ○) represent ME and HE Brasses reported by Nagase et al.²²⁾ The groups indicated by indexes A and B indicate Cu–Zn–Mn–Ni ME Brass castings^21,22) and Cu–Zn–Mn–Ni–Sn HE Brass castings,²²⁾ respectively. A part of the figure was referred from the literature.¹⁷⁾

Cu–Zn–Mn–Ni ME Brasses, which belong to the alloys indicated by index A in Fig. 1, exhibit superior ductility.^21,22) HE Brasses of Cu_xZnMnNiSn_y (x = 1 and y = 0.2, x = 2 and y = 0.45), indicated by index B in Fig. 1, were obtained by adding Sn to Cu_xZnMnNi (x = 1, and 2) ME Brasses.²²⁾ The characteristics of the constituent elements of ME and HE Brasses were different from 3d-transition metal type HEAs (3d-HEAs) ^1,2,4,24) including HE cast irons,^14,17,20) refractory HEAs (RHEAs) composed with 4 group · 5 group · 6 group elements,^25–31) HEAs for metallic biomaterials (bio-HEAs),^32–37) and Light-Weight HEAs (LW-HEAs) · LW-MEAs composed with lightweight elements.^38–43) The primary constituent elements of ME and HE Brasses were Cu and Zn, where Zn has high vapor pressure. The development of the casting process in multi-component alloys with Zn is important for the development of ME and HE Brasses. In this study, the development of the metallic mold-casting process without using any special equipment and vacuum chamber of Cu_xZnMnNi (x = 1, and 2),^21,22) and the solidification microstructure of the ingots obtained by melting with silica-based crucibles in an inert gas and metallic mold casting in air atmosphere were studied.

2. Experimental Procedure

Table 1 shows the nominal alloy composition of Cu_xZnMnNi (x = 1, 2) ME Brasses expressed in mass%. Figure 2 illustrates the raw materials used for the fabrication of the ingots of Cu_xZnMnNi ME Brasses by centrifugal metallic mold casting via high frequency melting in silica-based crucibles. Commercial elemental wire-cuts of Cu (Mitsuwa Pure Chemicals Co. Ltd., Japan, purity: 99.99%; as shown in Fig. 2(a)), wire-cuts of Ni (Mitsuwa Pure Chemicals Co. Ltd., Japan, 99.99%; as shown in Fig. 2(b)), lumps of Mn–Cu pre-alloy ingots (Osaka Asahi Co. Ltd., Cu_45.62Mn_54.38 (at%); as shown in Figs. 2(c1), 2(c2)), and lumps of Zn–Ni pre-alloy ingots (Osaka Asahi Co. Ltd., composition: Ni_27.81Zn_72.19 (at%); as shown in Figs. 2(d1), 2(d2)) were used as the starting materials. The size of the lumps of pre-alloy ingots were below 10 mm. The Mn-rich Mn–Cu pre-alloy ingots and Zn-rich Zn–Ni ingots were used instead of pure Zn and Mn lumps. The fabrication of ingots from the mixture of pure element lumps, including Zn lumps, was difficult because of the white fumes that occur during the oxidation of Zn element. The use of Zn–Ni pre-alloy was effective to suppress the white fumes during the melting and casting processes, resulting in the successful fabrication of the metallic mold-casting ingots using conventional casting process in Cu_xZnMnNi (x = 1, 2) ME Brasses.

Table 1 Nominal alloy composition of Cu_xZnMnNi (x = 1, 2) ME Brasses expressed in mass%.

Fig. 2

Raw materials for the fabrication of the Cu_xZnMnNi ME Brasses ingots via high frequency melting in silica-based crucible and centrifugal metallic mold casting. (a) pure-Cu wire cuts, (b) pure-Ni wire cuts, (c1) Mn–Cu pre-alloy ingots, (c2) lumps of Mn–Cu pre-alloy ingots for melting in silica-based crucible, (d1) Zn–Ni pre-alloy ingots, (d2) lumps of Zn–Ni pre-alloy ingots for melting in silica-based crucible.

Figure 3 shows the schematic illustration (Fig. 3(a)) and the pictures of the centrifugal metallic mold-casting equipment (Fig. 3(b)) (Denki Kogyo Co. Ltd., High Frequency Induction Casting Machines, MD-201).⁴⁴⁾ The magnified image of the cross-section of the silica-based crucible obtained using X-ray computed tomography is shown in Fig. 3(c). The ingots were fabricated in the following sequence:

Fig. 3

Schematic illustration (a) and outer appearance (b) of the centrifugal metallic mold-casting equipment, and magnified image of the cross-section of the silica-based crucible (c).

1. (1)    The mixture of raw materials was set in the silica-based crucible (Index A in Fig. 3); The raw materials of pure Cu, pure Ni, Zn–Ni pre-alloy, and Mn–Cu pre-alloys were set randomly in the crucible.
2. (2)    The crucible and mixture of raw materials were set in the coil of high frequency melting as indicated by index B in Fig. 3.
3. (3)    The mixture of raw materials was heated in Ar flow; it is indicated by index C in Fig. 3.
4. (4)    The molten metal was directly stirred several times using a quartz rod.
5. (5)    After the complete melting of raw materials, the molten metal was cast into the metallic mold by rotating the crucible via centrifugal casting. The metallic mold is indicated by index D in Fig. 3. BN-based release was coated at the surface of the metallic mold.

Figure 4 shows the cooling rate of the metallic mold-casting ingots evaluated by secondary dendrite arm spacing in Al–Cu alloys.^18,19) The cooling rate at the metallic mold contacted regions ①②③⑥ was estimated at approximately 200–550 K/s, and that at the inner regions ④⑤ was estimated at approximately 100 K/s. The cooling rate during the metallic mold casting in this study was approximately an order of magnitude slower than the cooling rate of the general arc melting method (∼2000 K/s),^35,45) and approximately two orders of magnitude faster than the cooling rate when the molten metal was cooled in the silica-based crucible (∼1 K/s).²⁰⁾

Fig. 4

The cooling rate of the metallic mold-casting ingots evaluated by secondary dendrite arm spacing in Al–Cu alloys. A part of the figure was referred from the literature.^18,19)

The phase identification was mainly performed using XRD analysis. The solidification microstructure was investigated using optical microscopy (OM), scanning electron microscopy (SEM), electron probe micro analyzer (EPMA) – wavelength dispersive X-ray spectrometry (WDS), and transmission electron microscopy (TEM). OM images were obtained from the specimens etched by the corrosion fluid composed with Iron chloride (III)-hydrochloric acid-ethyl alcohol (FeCl₃: 1 g, HCl: 5 ml, C₂H₅OH: 200 ml) at room temperature. TEM specimens were synthesized by ion-slicer.

3. Results and Discussions

Figure 5 shows the XRD patterns of the ingots of Cu_xZnMnNi ME Brasses; they were obtained from the regions ② and ⑤ in Fig. 4. The peaks of the ingot in the equiatomic CuZnMnNi ME Brass was identified as FCC (○) and BCC (●) phases. The main constituent phase in the ingot of non-equiatomic Cu₂ZnMnNi ME Brass was identified as the FCC phase. Peaks corresponding to intermetallic compounds were not observed in the XRD patterns.

Fig. 5

XRD patterns of the Cu_xZnMnNi ME Brasses ingots. The XRD patterns of the ingots were obtained from the region ② and ⑤ in Fig. 4.

Figure 6 shows the macroscopic and microscopic solidification structure analysis results of the ingots of the equiatomic CuZnMnNi ME brass (Fig. 6(a)) and non-equiatomic Cu₂ZnMnNi ME Brass (Fig. 6(b)). Figures 6(a1) and 6(b1) show that the shape of ingots obtained using the centrifugal metallic mold casting process was similar to the metallic mold. Figures 6(a2) and 6(b2) show the schematic illustration of the position of the ingots. The macroscopic solidification structure in OM images demonstrates the formation of coarse columnar grains from the mold-contacted part to the center of the ingot (①→④, ②→⑤, ⑥→⑤ in Fig. 4), and that of equiaxed grain at the central region (④ and ⑤ in Fig. 4) in the equiatomic CuZnMnNi ME brass (Fig. 6(a3)) and non-equiatomic Cu₂ZnMnNi ME Brass (Fig. 6(b3)) regardless of the alloy composition. The formation of the macroscopic casting defects of the cavity reached the central region (⑤ in Fig. 4), and that of the blow holes were not observed in OM images (Figs. 6(a3) and 6(b3)). In accordance with the conventional micro solidification structure in the ingots, the inner regions of X in Fig. 6(a3) and P in Fig. 6(b3) were investigated; the regions X and P corresponded with the region ② in Fig. 4. The equiaxis dendrite structure was observed in the OM image (Fig. 6(a4)) and SEM-Back Scattering Electron (BSE) image (Fig. 6(a5)) of the region X in the equiatomic CuZnMnNi ME Brass. In Cu–Zn–Mn–Ni ME Brasses, the observation of the solidification microstructure by SEM-BSE image was difficult.²²⁾ The SEM-BSE image (Fig. 6(a5)) was a high-contrast image for an enhanced view of the characteristics of the solidification microstructure. The equiaxis dendrite structure was also observed in the region P in the non-equiatomic Cu₂ZnMnNi ME Brass (Fig. 6(b4)). The area fraction of the interdendrite region in the non-equiatomic Cu₂ZnMnNi ME Brass was much lower than in the equiatomic CuZnMnNi ME Brass. The particular solidification microstructure formation with the formation of inclusions was not observed at the metallic mold contacted regions indicated by indexes Y (Fig. 6(a6)) and Z (Fig. 6(a7)) in the equiatomic CuZnMnNi ME Brass and by index Q (Fig. 6(b5)) in the non-equiatomic Cu₂ZnMnNi ME Brass.

Fig. 6

Solidification microstructure analysis of the ingots of the (a) equiatomic CuZnMnNi ME Brass and (b) non-equiatomic Cu₂ZnMnNi ME Brass. (a1) and (b1) outer appearances, and (a2) and (b2) schematic illustrations of the position of the ingots, (a3) and (b3) macro solidification structure of the cross section of the ingots. (a4), (a6), (a7), (b4), and (b5) micro solidification structure obtained from OM images, and (a5) micro solidification structure obtained from SEM-BSE image.

Figure 7 shows the EPMA-WDS element mapping of the equiaxis dendrite structure corresponding to the Figs. 6(a4) and 6(a5) in the metallic mold casting ingots in equiatomic CuZnMnNi ME Brasses. Figure 8 shows the EPMA-WDS element mapping of the equiaxis dendrite structure corresponding to Fig. 6(b4) in the metallic mold-casting ingots in non-equiatomic Cu₂ZnMnNi ME Brasses. Ni was enriched in the dendrite region, while Zn was enriched in the interdendrite region in the equiatomic CuZnMnNi and non-equiatomic Cu₂ZnMnNi ME Brasses, regardless of the alloy composition in Cu–Zn–Mn–Ni ME Brasses. Table 2 shows the chemical composition analysis of the dendrite (Dend.) and interdendrite (Interdend.) regions in the equiatomic CuZnMnNi ME Brass ingots.²²⁾ Dendrite and interdendrite regions contained Cu, Zn, Mn, and Ni at a ratio close to the equiatomic composition, indicating the formation of multicomponent phases in both the regions. The maximum difference in atomic composition between the dendrite and interdendrite regions was shown in Ni. The difference in Ni atomic percent between dendrite and interdendrite regions was significantly small and at most 8.5%. This indicates that the partition coefficient during the solidification of the equiatomic CuZnMnNi ME Brass is close to 1, and the element distribution and formation of the segregation region occurred with the solidification. However, the degree of element distribution was small resulting in the formation of the composite of dendrite and interdendrite regions with a small difference in the chemical composition.

Fig. 7

EPMA-WDS element mapping of the equiaxis dendrite structure in the metallic mold-casting ingots in equiatomic CuZnMnNi ME Brasses.

Fig. 8

EPMA-WDS element mapping of the equiaxis dendrite structure in the metallic mold-casting ingots in non-equiatomic Cu₂ZnMnNi ME Brasses.

Table 2 Chemical composition analysis of dendrite (Dend.) and interdendrite (Interdend.) regions in the equiatomic CuZnMnNi ME Brass ingots. A part of the data was referred from literature.²²⁾

Figure 9 shows TEM observation results of dendrite phase in the equiaxis dendrite structure of metallic mold-casting ingots of equiatomic CuZnMnNi ME Brass. In the bright field image (Fig. 9(a)), the existence of fine precipitates embedded in the matrix was not observed. The selected area diffraction patterns (Figs. 9(b1), 9(b2), and 9(b3)) observed from various direction can be indexed as fundamental diffraction spots of the FCC phase. The ordering spots based on the FCC-based ordering phase was not observed. Electron Back-Scatter Diffraction (EBSD) was reported to be effective for the identification of the constituent phases in ME Brass.²¹⁾ Based on the XRD analysis (Fig. 5) and the solidification microstructure analysis (Fig. 6, 7, 8, 9, S1, S2), the dendrite and interdendrite regions corresponded to the FCC and BCC regions, respectively.

Fig. 9

TEM observation results of the dendrite phase of the equiaxis dendrite structure in the metallic mold-casting ingots in equiatomic CuZnMnNi ME Brasses. (a) Bright field image, (b1)–(b3) selected area electron diffraction patterns.

Figure 10 shows nominal stress–nominal strain curves of the metallic mold-casting ingots in ME Brasses; the detailed information was reported in the literature.²²⁾ The metallic mold-casting ingots exhibit superior compression ductility. The equiatomic CuZnMnNi ME Brass with FCC + BCC dual phase structure (0.2% proof stress, 300 MPa) showed higher yield stress than non-equiatomic Cu₂ZnMnNi ME Brass with FCC near-single phase (0.2% proof stress, 168 MPa). The BCC phase in ME Brass and HE Brass showed the particular structure with nano-scale precipitates and BCC-based ordering structure.²²⁾ The existence of BCC phase may lead to the increase in the yield stress of the ingots in Cu–Zn–Mn–Ni ME Brasses.

Fig. 10

Nominal stress–nominal strain curves of the metallic mold-casting ingots in ME brasses.

4. Conclusion

This paper reported the fabrication of the ingots using metallic mold casting process and the solidification microstructure of the ingots in Cu_xZnMnNi (x = 1, 2) ME Brasses. The results are summarized as following:

1. (1)    The 0.2% proof stress-mixing entropy ΔS_mix map in various copper-based casting alloys, including Cu_xZnMnNi ME Brasses and Cu_xZnMnNiSn_y HE Brasses, was constructed. The increase in the ΔS_mix is effective to increase in the mechanical strength in the Cu-based casting alloy ingots.
2. (2)    The ingots of Cu_xZnMnNi ME Brasses can be obtained from a mixture of pure-Cu, pure-Ni, Mn–Cu pre-alloy, and Zn–Ni pre-alloy lumps using conventional casting process via crucible melting in an inert gas and metallic mold casting in air atmosphere without any special equipment and vacuum chamber.
3. (3)    The dual-phase structure composed of the major FCC and minor BCC structures was obtained in the equiatomic CuZnMnNi ME Brass ingots. A near-single FCC structure was obtained in the non-equiatomic Cu₂ZnMnNi ME brass ingots.
4. (4)    The macroscopic solidification microstructure in the ingots exhibited a columnar dendrite structure, and the microscopic solidification microstructure in the ingots showed an equiaxes dendrite structure. The element distribution was not homogeneous, and a small distribution of the chemical composition corresponding to the dendrite structure was observed.

Acknowledgments

This work was partially supported by JSPS KAKENHI (grant number 18K04750, 19H05172) and the scientific grants received from the Japan Copper and Brass Association. A part of this work was conducted under the interuniversity cooperative research program of (Proposal number 18G0036, 19G0035, 20G0020) the Cooperative Research and Development Center for Advanced Materials, Institute for materials Research, Tohoku University. The authors are grateful to Mr. I. Narita at Tohoku University for his help with the EPMA measurements.

REFERENCES

• 1)   B.  Cantor,  I.T.H.  Chang,  P.  Knight and  A.J.B.  Vincent: Mater. Sci. Eng. A 375–377 (2004) 213–218.
• 2)   J.W.  Yeh,  S.K.  Chen,  S.J.  Lin,  J.Y.  Gan,  T.S.  Chin,  T.T.  Shun,  C.H.  Tsau and  S.Y.  Chang: Adv. Eng. Mater. 6 (2004) 299–303.
• 3)   S.  Ranganathan: Curr. Sci. 85 (2003) 1404–1406. https://www.currentscience.ac.in/Downloads/article_id_085_10_1404_1406_0.pdf
• 4)   Y.  Zhang,  Y.J.  Zhou,  J.P.  Lin,  G.L.  Chen and  P.K.  Liew: Adv. Eng. Mater. 10 (2008) 534–538.
• 5)   J.W.  Yeh: JOM 65 (2013) 1759–1771.
• 6)   M.H.  Tsai and  J.W.  Yeh: Mater. Res. Lett. 2 (2014) 107–123.
• 7)   Y.  Zhang,  T.T.  Zuo,  Z.  Tang,  M.C.  Gao,  K.A.  Dahmen,  P.K.  Liaw and  Z.P.  Lu: Prog. Mater. Sci. 61 (2014) 1–93.
• 8)  B.S. Murty, J.-W. Yeh and S. Ranganathan: High-Entropy Alloys, first ed., (Elsevier, Amsterdam, 2014).
• 9)  M.C. Gao, J.-W. Yeh, P.K. Liaw and Y. Zhang: High-Entropy Alloys, Fundamentals and Applications, first ed., (Springer, Berlin, 2016).
• 10)   A.  Takeuchi: J. Jpn. Inst. Metals and Materials 79 (2015) 157–168.
• 11)   A.  Takeuchi: J. Jpn. Soc. Powder Metall. 63 (2016) 209–216.
• 12)   Y.F.  Ye,  Q.  Wang,  J.  Lu,  C.T.  Liu and  Y.  Yang: Mater. Today 19 (2016) 349–362.
• 13)   D.B.  Miracle and  O.N.  Senkov: Acta Mater. 122 (2017) 448–511.
• 14)   T.  Nagase: J. Soc. Mech. Eng. 121 (2018) 8–11.
• 15)   W.  Zhang,  P.K.  Liew and  Y.  Zhang: Sci. China Mater. 61 (2018) 2–22.
• 16)   D.B.  Miracle: Nat. Commun. 10 (2019) 1805.
• 17)   T.  Nagase: KINZOKU 89 (2019) 899–910.
• 18)   T.  Nagase,  M.  Takemura,  M.  Matsumuro and  T.  Maruyama: J. JFS 89 (2017) 119–129.
• 19)   T.  Nagase,  M.  Takemura,  M.  Matsumuro and  T.  Maruyama: Mater. Trans. 59 (2018) 255–264.
• 20)   T.  Nagase,  T.  Kakeshita,  K.  Matsumura,  K.  Nakazawa,  S.  Furuya,  N.  Ozoe and  K.  Yoshino: Mater. Des. 184 (2019) 108172.
• 21)   K.J.  Laws,  C.  Crosby,  A.  Sridhar,  P.  Conway,  L.S.  Koloadin,  M.  Zhao,  S.  Aron-Dine and  L.C.  Bassman: J. Alloy. Compd. 650 (2015) 949–961.
• 22)   T.  Nagase,  A.  Shibata,  M.  Matsumuro,  M.  Takemura and  S.  Semboshi: Mater. Des. 181 (2019) 107900.
• 23)  Dou Goukin Imono No Hyojyun, (Japan Foundry Society, 2011).
• 24)   B.  Gludovatz,  A.  Hohenwarter,  D.  Catoor,  E.H.  Chang,  E.P.  George and  R.O.  Ritchie: Science 345 (2014) 1153–1158.
• 25)   O.N.  Senkov,  G.B.  Wilks,  D.B.  Miracle,  C.P.  Chuang and  P.K.  Liaw: Intermetallics 18 (2010) 1758–1765.
• 26)   O.N.  Senkov,  G.B.  Wilks,  J.M.  Scott and  D.B.  Miracle: Intermetallics 19 (2011) 698–706.
• 27)   O.N.  Senkov,  J.M.  Scotta,  S.V.  Senkova,  D.B.  Miracle and  C.F.  Woodward: J. Alloy. Compd. 509 (2011) 6043–6048.
• 28)   O.N.  Senkov,  J.M.  Scott,  S.V.  Senkova,  F.  Meisenkothen,  D.B.  Miracle and  C.F.  Woodward: J. Mater. Sci. 47 (2012) 4062–4074.
• 29)   O.N.  Senkov,  D.B.  Miracle,  K.J.  Chaput and  J.-P.  Couzinie: J. Mater. Res. 33 (2018) 3092–3128.
• 30)   T.  Nagase,  S.  Anada,  P.D.  Rack,  J.H.  Noh,  H.  Yasuda,  H.  Mori and  T.  Egami: Intermetallics 38 (2013) 70–79.
• 31)   T.  Nagase,  S.  Anada,  P.D.  Rack,  J.H.  Noh,  H.  Yasuda,  H.  Mori and  T.  Egami: Intermetallics 26 (2012) 122–130.
• 32)   M.  Todai,  T.  Nagase,  T.  Hori,  A.  Matsugaki,  A.  Sekita and  T.  Nakano: Scr. Mater. 129 (2017) 65–68.
• 33)   S.P.  Wang and  J.  Xu: Mater. Sci. Eng. C 73 (2017) 80–89.
• 34)   T.  Nagase,  M.  Todai,  T.  Hori and  T.  Nakano: J. Alloy. Compd. 753 (2018) 412–421.
• 35)   T.  Nagase,  K.  Mizuuchi and  T.  Nakano: Entropy 21 (2019) 483.
• 36)   T.  Hori,  T.  Nagase,  M.  Todai,  A.  Matsugaki and  T.  Nakano: Scr. Mater. 172 (2019) 83–87.
• 37)   T.  Nagase,  Y.  Iijima,  A.  Matsugaki,  K.  Ameyama and  T.  Nakano: Mater. Sci. Eng. C 107 (2020) 110322.
• 38)   K.M.  Youssef,  A.J.  Zaddach,  C.  Niu,  D.L.  Irving and  C.C.  Koch: Mater. Res. Lett. 3 (2015) 95–99.
• 39)   R.  Feng,  M.C.  Gao,  C.  Lee,  M.  Mathes,  T.  Zuo,  S.  Chen,  J.A.  Hawk,  Y.  Zhang and  P.K.  Liaw: Entropy 18 (2016) 333.
• 40)   Y.  Qiu,  Y.J.  Hu,  A.  Taylor,  M.J.  Styles,  R.K.W.  Marceau,  A.V.  Ceguerra,  M.A.  Gibson,  Z.K.  Liu,  H.L.  Fraser and  N.  Birbilis: Acta Mater. 123 (2017) 115–124.
• 41)   R.  Feng,  M.C.  Gao,  C.  Zhang,  W.  Guo,  J.D.  Poplawsky,  F.  Zhang,  J.A.  Hawk,  J.C.  Neuefeind,  Y.  Ren and  P.K.  Liaw: Acta Mater. 146 (2018) 280–293.
• 42)   T.  Nagase,  A.  Terayama,  T.  Nagaoka,  N.  Fuyama and  T.  Sakamoto: J. JFS 91 (2019) 717–729.
• 43)   T.  Nagase,  A.  Terayama,  T.  Nagaoka,  N.  Fuyama and  T.  Sakamoto: Mater. Trans. 61 (2020) 1369–1380.
• 44)  https://www4.hp-ez.com/hp/haradadenki/page2, (assessed, 05/30, 2019).
• 45)   T.  Nagase,  M.  Matsumoto and  Y.  Fujii: J. Alloy. Compd. 738 (2018) 440–447.
• 46)   A.  Yasuhara: JEOL News 40 (2005) 46–49. https://www.jeol.co.jp/download_jeolnews.html
• 47)  JEOL Ltd.: EM-09100IS Ion Slicer, https://www.jeol.co.jp/en/products/detail/EM-09100IS.html, (assessed, 01/16, 2020).

Appendix

We attempted to identify the constituent phases of dendrite and interdendrite regions in the equiatomic CuZnMnNi ME Brass by EBSD measurement and TEM observation. Figure S1 shows an image quality (IQ) map (a) and phase map (PM) (b) by EBSD analysis for the equiatomic CuZnMnNi ME Brass fabricated by metallic mold-casting. The specimen for EBSD analysis was prepared by mechanical-polishing and then ion polishing at 4.5 kV for 300 s at room temperature. We could obtain clear electron back-scattering patterns from the dendritic region as indicated by bright contrasts in IQ map of Fig. S1(a), and find that the dendritic region had an FCC structure as shown in Fig. S1(b). The phase of the interdendrite region was alternatively analyzed to be BCC structure of CuZn (Fig. S1(b)), although the EBSD patterns from the interdendritic region was weak and unclear as represented by dark region in Fig. S1(a). This is probably because the interdendrite regions were deeply etched by ion-polishing. A detailed experiments and analysis should be required to clarify that the phase structure has BCC structure.

Fig. S1

(a) image quality (IQ) map, (b) phase map (PM) by EBSD analysis for the metallic-mold casting ingots of the equiatomic CuZnMnNi ME Brass with an equiaxis dendrite structure. In PM of (b), the Kikuchi pattern was analyzed based on the data for pure FCC-Cu (indicated in red) and that for the BCC structure of CuZn (green).

Figure S2 shows TEM images of the interdendrite region for equiatomic CuZnMnNi ME Brass. The TEM specimens were prepared using the ion-slicer equipment^46,47) as opposed to the ion milling equipment for the wide-area observation. In the bright-field image shown in Fig. S2(a), we can find that the interdendrite grain as indicated by the arrow X was composed of single phase (including no precipitates and inclusions). The selected area electron diffraction (SAD) pattern taken from the interdendrite grain (X) (Fig. S2(b)) shows that the fundamental diffraction spots were indexed by the BCC structure. Some of weak diffraction spots which cannot be indexed as BCC phase with A2 structure were seen in the SAD pattern, however, these cannot be identified because of the low intensity. Further discussion focusing on BCC structure formation will be reported in the other works. XRD patterns (Fig. 5 in main text), OM and SEM microstructures (Fig. 6 in main text), EPMA element mapping (Fig. 7 in main text), TEM observation focusing on dendrite regions (Fig. 9 in main text), EBSD analysis (Fig. S1 in the appendix) and TEM observation focusing on interdendrite regions (Fig. S2 in the appendix) can be explained by the formation of the equiaxis dendrite structure composed with dendrite regions with FCC structure and inderdendrite regions with BCC structure without any discrepancy.

Fig. S2

(a) bright-field TEM image and (b) selected area electron diffraction (SAD) pattern for the metallic-mold casting ingots of the equiatomic CuZnMnNi ME Brass with an equiaxis dendrite structure. The SAD pattern of (b) was taken from the interdendrite grain indicated by X in (a). The fundamental diffraction spots in (b) was corresponded to that viewed from [111] direction to a BCC structure.

In the metallic mold casting ingots of the non-equiatomic Cu₂ZnMnNi ME Brass, the interdendrie region was considered to have the BCC structure, and the lack of sharp peaks corresponding to the BCC structure in XRD patterns (Fig. 5) was due to the significantly small quantity of the interdendrite region. The quantity of the interdendrite region with BCC structure decreased as the x value in Cu_xZnMnNi ME Brasses increased, resulting in a near-single FCC structure in the non-equiatomic Cu₂ZnMnNi ME brass ingots.