Formation of amorphous phase with crystalline globules in Fe–Cu–Nb–B immiscible alloys

Journal of Alloys and Compounds 619 (2015) 267–274

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Formation of amorphous phase with crystalline globules in Fe–Cu–Nb–B immiscible alloys Takeshi Nagase a,b,⇑, Masanori Suzuki b, Toshihiro Tanaka b a b

Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, 7-1, Mihogaoka, Ibaraki, Osaka 567-0047, Japan Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1, Yamada-Oka, Suita, Osaka 565-0871, Japan

a r t i c l e

i n f o

Article history: Received 25 June 2014 Received in revised form 6 August 2014 Accepted 7 August 2014 Available online 9 September 2014

a b s t r a c t The microstructure of arc-melted ingots and rapidly solidified melt-spun ribbons of quaternary Fe–Cu–Nb–B immiscible alloys was investigated, with a focus on amorphous-phase formation and solidification structure. A continuous melt-spun ribbon with an Fe–Nb–B-based amorphous matrix and 10–100 nm diameter fcc-Cu crystalline globules was obtained for the (Fe0.75Nb0.10B0.15)80Cu20 alloy. Ó 2014 Elsevier B.V. All rights reserved.

Keywords: Rapid solidification Iron alloys Amorphous material TEM HREM

1. Introduction Binary Fe–Cu alloys are well known to show a metastable liquid miscibility below the liquidus line [1]. The liquid phase separation behavior in binary Fe–Cu and multicomponent Fe–Cu-based alloys has been of great interest to the materials science fields [2–16], industrial engineering, such as recycling Cu from Fe-based alloy scraps [17,18], for precursor preparation of porous materials [19], and as a dual-layer melt-extraction wire composed of a soft, magnetic Fe-based amorphous alloy core with a crystalline Cu cover layer [20]. Recently, new types of immiscible metallic materials, called two-phase metallic amorphous alloys, were developed from a variety of immiscible alloys [21–24]. The formation of two amorphous phases originates from a positive heat of mixing (DHmix) between the two major elements and a high glass-forming ability (GFA) of the separated liquids. In the case of Fe–Cu-based alloys, the formation of a two-phase metallic amorphous material has not yet been reported; whereas the simultaneous occurrence of liquid phase separation and Fe-based amorphous phase formation has been well studied [25–36]. Spherical crystalline precipitates tend to form in the Fe-rich, Fe–Cu-based multicomponent immiscible alloys with amorphous phase formation, indicating that their

⇑ Corresponding author at: Research Center for Ultra-High Voltage Electron Microscopy, Osaka University, 7-1, Mihogaoka, Ibaraki, Osaka 567-0047, Japan. Tel.: +81 6 6879 7941; fax: +81 6 6879 7942. E-mail address: [email protected] (T. Nagase). http://dx.doi.org/10.1016/j.jallcom.2014.08.229 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

liquid phase separation results in the formation of composites with a spherical crystalline phase and an Fe-based amorphous matrix. The importance of this concept for practical applications of metallic glasses is described in detail in Appendix. In this study, the formation of an amorphous phase with a spherical crystalline phase (crystalline globules) was investigated in Fe–Cu–Nb–B alloys. 2. Alloy design Fe–Cu-based alloys, which show simultaneous amorphous phase formation and liquid phase separation, are difficult to design due to competition between the atomic pairs of the constituent elements. Large negative values of DHmix for the atomic pairs favor amorphous phase formation by liquid state stabilization [37]; whereas large positive DHmix values for atomic pairs lead to liquid destabilization, resulting in liquid phase separation. In this study, a combination map of DHmix for binary atomic pairs in quaternary Fe–Cu-based alloys, in addition to a predicted quaternary phase diagram based on ab initio calculations, shown in Fig. 1, was suggested for alloy design. Fig. 1a shows the combination map of DHmix for the constituent element binary atomic pairs in the quaternary Fe–Cu–Nb–B alloys, where the value of DHmix was estimated by modified Miedema method [38,39]. The large negative DHmix values for Fe–Nb, Fe–B, and Nb–B, as well as the large DHmix positive values for Cu–Fe, Cu–Nb, and Cu–B, imply a liquid phase separation of Fe–Nb–B

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Fig. 3. XRD patterns of (Fe0.75Nb0.10B0.15)100xCux (x = 0, 20, 30, and 50) arc-melted alloy ingots.

Fig. 1. The prediction method for Fe–Cu-based alloys, presenting simultaneous amorphous phase formation and liquid phase separation, shows (a) the DHmix combination map of the constituent elemental binary atomic pairs in quaternary Fe–Cu–Nb–B alloys; and (b) the predicted Fe–Cu–Nb–B quaternary phase diagram from ab initio calculations.

and Cu liquids. The ternary Fe–Nb–B alloy system is well known to exhibit a high glass-forming ability (GFA) [40,41]. This indicates that the quaternary Fe–Cu–Nb–B alloys have a strong tendency for simultaneous liquid phase separation and amorphous phase formation. The predicted Fe–Cu–Nb–B quaternary phase diagram is useful for designing new metallic materials. Fig. 1b shows the quaternary phase diagram at 0 K for the Fe–Nb–B–Cu alloy constructed using first principles calculations from the Materials Project [42–45]. In the predicted phase diagram, there is no information about a liquid phase; however, the possibility of

intermetallic compounds can be discussed, as they are important for liquid phase separation and amorphous phase formation. There are no binary Fe–Cu, Nb–Cu, and B–Cu; ternary Fe–Nb–Cu, Fe–B–Cu, and Nb–B–Cu; or quaternary Fe–Cu–Nb–B intermetallic compounds. This indicates that the liquid phase separation of Fe–Nb–B-based and Cu-based liquids is unlikely to be interrupted by the formation of intermetallic compounds during Fe–Cu–Nb–B alloy cooling. It should be noted here that intermetallic compounds such as FeB, FeB2, and FeNbB are not seen in the predicted phase diagram. This is the result of the missing compounds’ formation energy not being calculated by the ab initio calculations. In spite of the limited data obtained by ab initio calculations, the predicted phase diagram remains useful for discussing the existence of intermetallic compounds in the quaternary Fe–Nb–B–Cu alloy. There is no discrepancy between the forecast from the DHmix combination map of binary atomic pairs based on the modified Miedema model [38,39], and that from the predicted phase diagram based on ab initio calculations rather than empirical rules. The DHmix

Fig. 2. Cross section images of (Fe0.75Nb0.10B0.15)100xCux arc-melted alloy ingots: (a) x = 20, (b) x = 30, and (c) x = 50; (b0 ) magnified image of the macroscopically phaseseparated x = 30 ingot interface.

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Fig. 6. XRD patterns of the (Fe0.75Nb0.10B0.15)100xCux (x = 0, 20, 30) melt-spun alloy ribbons.

3. Experimental procedure

Fig. 4. EPMA analysis of the (Fe0.75Nb0.10B0.15)70Cu30 arc-melted alloy ingot.

Master ingots of (Fe0.75Nb0.10B0.15)100xCux alloys (x = 0, 20, 30, 50 at.%) were prepared from Fe, Cu, Nb, B, and an Fe–B pre-alloy in a water-cooled Cu hearth by arc melting under a purified Ar atmosphere. Rapidly quenched ribbons were produced from the master ingots by means of a single-roller melt-spinning method. A fused quartz nozzle was used, 14 mm in diameter with a 1.0 mm orifice, where the heating of the master ingot was performed by radio-frequency. The roller surface velocity was approximately 42 m s1. The structure of the melt-spun ribbon was examined by X-ray diffraction (XRD) using Cu Ka radiation. Scanning electron microscopy (SEM)-backscattered electron image (BEI) observation and electron probe microanalysis (EPMA)-wavelength dispersive X-ray spectrometry (WDX) analysis were performed using the JEOL JEM-5600 and JXA-8800R systems, respectively. Transmission electron microscopy (TEM) and high-resolution electron microscopy (HREM) observations were carried out using the Hitachi H-800 and Hitachi HF-2000 systems, respectively. Thermal analysis was performed using differential scanning calorimetry (DSC) with a Mac Science DSC-3100S. The thin films for the TEM and HREM analyses were prepared by an ion-thinning method using Gatan’s precision ion-polishing system (PIPS, model 691).

4. Results

Fig. 5. The outer appearance and magnified OM images of the free and wheel-side surfaces of (Fe0.75Nb0.10B0.15)100xCux melt-spun alloy ribbons: (a) x = 0, (b) x = 20, and (c) x = 30.

combination map for the constituent element binary atomic pairs and the predicted quaternary phase diagram, both of which are readily available, indicate that the quaternary Fe–Nb–B–Cu alloy could produce simultaneous amorphous phase formation and liquid phase separation, resulting in the creation of an Fe-based amorphous alloy with Cu-based crystalline globules.

Fig. 2 shows the cross section images of the arc-melted (Fe0.75Nb0.10B0.15)100xCux, (x = 20, 30, 50) alloy ingots. Fig. 2b0 is the magnified image of Fig. 2b. A macroscopic phase-separated interface between the metallic silver region and the copper-colored region, indicated by the white arrows in Fig. 2a–c and the black arrows in Fig. 2b, was observed regardless of the Cu quantity. A metallic silver colored spherical phase embedded in the copper-colored matrix as well as a copper colored spherical phase embedded in the metallic silver-colored matrix, indicated by the red arrows, can also be seen. These features are typical for a solidification structure with a liquid phase separation. Fig. 3 shows the XRD patterns of the arc-melted (Fe0.75Nb0.10B0.15)100xCux, (x = 0, 20, 30, 50) ingots. The main constituent phases in the ternary Fe–Nb–B alloy ingot were identified as bcc-Fe and an FeNbB intermetallic compound (space group P62m, Pearson symbol hP9). In the quaternary Fe–Cu–Nb–B alloys, sharp diffraction peaks corresponding to the fcc-Cu can be seen in the XRD patterns together with those corresponding to the bcc-Fe and the FeNbB, indicating that Fe and Cu were separated in the arc-melted ingots. Fig. 4 shows the EPMA analysis of the arc-melted (Fe0.75Nb0.10B0.15)70Cu30 alloy ingot, with its macroscopically phaseseparated interface. In the bright-field image (BFI, Fig. 4a); the gray and white contrast phases are identified as Fe-rich and Cu-rich phases, respectively. The Nb and B tend to participate in the Fe-rich phases. These results indicate that the Fe–Nb–B- and Cu-based

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←End do. ΔH (a.u.)

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liquid phases solidified separately in the Fe–Cu–Nb–B immiscible alloys, in good agreement with the prediction based on the DHmix combination map and the predicted phase diagram shown in Fig. 1. Fig. 5 displays the outer appearance and magnified optical microscope (OM) images of the free and wheel-side surfaces of the melt-spun (Fe0.75Nb0.10B0.15)100xCux, (x = 0, 20, 30) alloy ribbons. A continuous melt-spun ribbon with a metallic silver color was obtained for the Fe75Nb10B15 (Fig. 5a) and (Fe0.75Nb0.10B0.15)80Cu20 (Fig. 5b) alloys; whereas flake-like samples formed in the (Fe0.75Nb0.10B0.15)70Cu30 (Fig. 5c) alloy. A mirror-surface image of the free-surface side, a typical feature of amorphous phase formation, can be seen for the Fe75Nb10B15 (Fig. 5a) and (Fe0.75Nb0.10B0.15)80Cu20 (Fig. 5b) alloys. In the (Fe0.75Nb0.10B0.15)70Cu30

(Fig. 5c) alloy, a macroscopically phase-separated interface can be seen between the metallic silver region and the copper-colored region, indicated by the black arrows, in the magnified OM images of the free and wheel-side surfaces. Fig. 6 shows the XRD patterns of the (Fe0.75Nb0.10B0.15)100xCux, (x = 0, 20, 30) alloy melt-spun ribbons. The XRD patterns do not have a significant difference between the free-surface side (free) and the wheel-side surface (wheel) for all specimens. A broad peak typical of an amorphous single phase can be seen in the ternary Fe75Nb10B15 alloy. Both a sharp diffraction peak corresponding to the fcc-Cu crystalline phase and a broad halo peak can be seen in the (Fe0.75Nb0.10B0.15)80Cu20 alloy. The intensity of the fcc-Cu crystalline phase peaks increased with increasing Cu content in the alloy. In the (Fe0.75Nb0.10B0.15)70Cu30 alloy, sharp diffraction peaks corresponding to the fcc-Cu and bcc-Fe were observed; a broad peak was not apparent due to the low intensity and overlap of the peak positions. To clarify the formation of an amorphous phase in the Fe–Cu–Nb–B alloys, a DSC analysis of the melt-spun ribbons were performed. Fig. 7 shows the DSC curves of the melt-spun (Fe0.75Nb0.10B0.15)100xCux, (x = 0, 20, 30) alloy ribbons. The ternary Fe75Nb10B15 alloy shows an endothermic reaction with its onset indicated by Tg and a sharp exothermic peak with its onset indicated by Tx. These peaks can be explained as an amorphous phase glass transition and a thermal crystallization, respectively. An amorphous phase showing a glass transition during heating can be obtained in melt-spun ribbons of the ternary Fe–Nb–B alloy. The Cu addition affected these DSC curves, with the endothermic glass transition not seen in the quaternary Fe–Cu–Nb–B alloys. The amount of heat released for the exothermic peaks decreased with increasing Cu content. In the (Fe0.75Nb0.10B0.15)80Cu20 alloy, a broad exothermic peak appeared, having an onset temperature similar to Tx in the ternary Fe–Nb–B alloy, implying the crystallization of an amorphous phase. Exothermic peaks caused by crystallization of an amorphous phase cannot be seen in the (Fe0.75Nb0.10B0.15)70Cu30 alloy.

Fig. 8. TEM microstructures of the (Fe0.75Nb0.10B0.15)100xCux (x = 20, 30) melt-spun alloy ribbons. (a) x = 20: (a1) BF image, (a2) magnified BF image, and (a20 ) corresponding SAD pattern; (b) x = 30: (b1) BF image, (b2) magnified BF image of an amorphous matrix region (indicated by A in Fig. 8b1), (b20 ) corresponding SAD pattern of (b2); (b3) magnified BF image of a crystalline region (indicated by B in Fig. 8b1), and (b30 ) magnified BF image of (b3).

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Fig. 9. HREM images and nanobeam EDS results of nanocrystalline globules embedded in an amorphous matrix of (Fe0.75Nb0.10B0.15)80Cu20 melt-spun alloy ribbon. (a) Typical nanoglobule example: (a1) HREM image, (a2) FFT pattern obtained from a nanocrystalline globule (A), and (a3) FFT pattern obtained from the amorphous matrix (B). (b) Typical example of a nanoglobule with many stacking faults: (b1) HREM image, (b2) nanobeam EDS results obtained from a nanocrystalline globule (C), and (b3) nanobeam EDS results obtained from a nanocrystalline globule (D).

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confirmed by TEM observation. In the (Fe0.75Nb0.10B0.15)70Cu30 alloy (Fig. 8b1), the dispersion of crystalline globules in the amorphous matrix, which is indicated by the letter A, and the crystalline region, indicated by the letter B, was observed. The regions indicated by A and B were confirmed as the Fe-rich and Cu-rich regions, respectively, by TEM-EDS analysis. The smooth interface between the A and B regions indicates that the interface was formed through liquid phase separation. Fig. 8b2 and b3 are the magnified images of regions A and B, respectively. Fig. 8b2, showing the BFI, reveals that the crystalline globules are clearly embedded in the amorphous matrix; the crystalline globule density in the (Fe0.75Nb0.10B0.15)70Cu30 alloy (Fig. 8b2) was higher than that in the (Fe0.75Nb0.10B0.15)80Cu20 alloy (Fig. 8a2). The amorphous phase formation in region A was confirmed by its corresponding SAD pattern (Fig. 8b20 ). As shown in Fig. 8b3, spherical phases can also be seen in the Cu-rich crystalline part (region B). An amorphous phase formation in the crystalline B region could not be confirmed by its corresponding SAD pattern (Fig. 8b30 ). Fig. 9 shows the HREM images and nanobeam EDS results of the nanocrystalline globules embedded in the amorphous matrix of the (Fe0.75Nb0.10B0.15)80Cu20 alloy. The spherical crystalline phase formation was also confirmed by the HREM images (Fig. 9a1 and b1). The nanocrystalline globule indicated by the letter A in Fig. 9a1 was identified as a fcc-Cu phase from the fast Fourier transform (FFT) pattern (Fig. 9a2). The matrix phase shows a salt-and-pepper contrast in the HREM image (Fig. 9a1), and the FFT pattern obtained from this, indicated by the letter B, shows a halo ring (Fig. 9a3) indicative of amorphous phase formation. In the (Fe0.75Nb0.10B0.15)80Cu20 alloy, TEM-BF images (Fig. 8a1 and a2), crystalline globules with a number of stacking faults can be seen, and a typical HREM image of the globules is shown in Fig. 9b1. The existence of stacking faults in the fccCu crystalline globules was confirmed by the HREM and TEM. Fig. 9b2 and b3 shows the nanobeam EDS analyses of a crystalline globule, indicated by the letter C, and the matrix region, indicated by the letter D, respectively. Only the Cu intensity can be seen in the crystalline globule (Fig. 9b2); whereas the Fe and Nb enrichment is seen in the amorphous matrix region (Fig. 9b3). The TEM, HREM, and nanobeam-EDS analyses clarified that the fcc-Cu crystalline globules dispersed in the Fe–Nb-enriched amorphous phase formed in the (Fe0.75Nb0.10B0.15)80Cu20 alloy. In the (Fe0.75Nb0.10B0.15)70Cu30 alloy, the amorphous phase was confirmed by TEM as a minor phase, with an emulsion structure similar to that of the (Fe0.75Nb0.10B0.15)80Cu20 alloy observed in this region. This report is the first finding of an amorphous phase formation with a unique solidification structure in which fcc-Cu crystalline globules were dispersed in an Fe-based amorphous matrix of Fe–Cu– Nb–B immiscible alloys.

5. Discussion The outer appearance (Fig. 5), XRD (Fig. 6), and DSC (Fig. 7) patterns show the formation of an amorphous phase and a fcc-Cu crystalline phase in the (Fe0.75Nb0.10B0.15)80Cu20 alloy. To investigate the microstructure of the rapidly solidified Fe– Cu–Nb–B alloys and the amorphous-phase formation as a minor phase in the (Fe0.75Nb0.10B0.15)70Cu30 alloy, TEM and HREM observations of the melt-spun ribbons were performed. Fig. 8 shows the TEM microstructures of the (Fe0.75Nb0.10B0.15)80Cu20 (Fig. 8a) and (Fe0.75Nb0.10B0.15)70Cu30 (Fig. 8b) alloy melt-spun ribbons. The BFI (Fig. 8a1) and the magnified BFI (Fig. 8a2) of the (Fe0.75Nb0.10B0.15)80Cu20 alloy show 10–100 nm crystalline precipitates dispersed in the featureless matrix. The corresponding selected area diffraction (SAD) pattern (Fig. 8a20 ) displays a broad halo ring corresponding to an amorphous matrix, along with discontinuous Debye rings generated by the fcc-Cu crystalline phase. The dispersion of the fcc-Cu crystalline globules in the amorphous phase was

Discussing the liquid phase separation behavior in Fe–Cu–Nb–B quaternary alloys is difficult, primarily as there is no systematic experimental data available for such a complicated alloy system, although the occurrence of a liquid phase separation in Fe–Cu– Nb–B quaternary alloys can be predicted by a DHmix map (Fig. 1a) and phase diagram (Fig. 1b) constructed using ab initio calculations. In this section, the occurrence of a liquid phase separation and partition of the constituent elements into separated liquids are discussed, based on the thermodynamic calculations by FactSage (ver. 6.4) using the thermodynamic databases for alloy systems from the Scientific Group Thermodata Europe (SGTE) [46]. The database was not constructed by a simple regular solid solution model with DHmix estimated on the Miedema model; however, the excess Gibbs energy of mixing in the liquid phase was described by Redlich–Kister–Muggianu polynomial expression

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Fig. 10. Thermodynamic calculations for the quaternary Fe–Cu–Nb–B alloy: (a) Gibbs free energy of a single liquid phase and a mixture of separated liquids in the (Fe0.75Nb0.10B0.15)100xCux alloys as a function of x. The solid line and broken line denote the single liquid phase and mixture of separated liquids, respectively. (b) Calculated miscibility gap in the (Fe0.75Nb0.10B0.15)100xCux alloys as a function of x, (c) constituent element mole fractions in the Fe-rich liquid, and (d) constituent element mole fractions in the Cu-rich liquid in the (Fe0.75Nb0.10B0.15)100xCux alloys as a function of x.

under regular solution model, and the critical assessment of the polynomial parameters was done to reproduce most of available experimental data on thermodynamic properties and phase equilibria. Fig. 10 shows the thermodynamic calculation results for the quaternary (Fe0.75Nb0.10B0.15)100xCux alloys as a function of x. The Gibbs free energy of the single liquid phase (indicated by the solid line) of the (Fe0.75Nb0.10B0.15)100xCux alloys at 2500 K shows a downward curve. The Gibbs free energy curve of a single liquid changes from a downward shape to an upward one with a peak composition at approximately x = 60 by decreasing the temperature. At 1500 K and 1000 K, the Gibbs free energy of the separated liquid mixture, indicated by the broken line, is considerably lower than that of the single liquid phase, indicating the liquid phase separation tendency from a single liquid phase to Fe–Nb–B-rich and Cu-rich liquid phases during cooling. Fig. 10b shows the calculated miscibility gap in the (Fe0.75Nb0.10B0.15)100xCux alloy as a function of x. A clear liquid miscibility gap with an approximate 2000 K peak temperature can be seen. Fig. 10c and d shows the constituent element mole fractions in the Fe-rich and Cu-rich liquids, respectively, of the (Fe0.75Nb0.10B0.15)100xCux alloys as a function of x. The Nb and B show a clear tendency of enriching the Fe-rich liquid rather than the Cu-rich liquid. In the arc-melted ingots, the Nb and

B enrichment occurred in the Fe-based crystalline phase rather than in the Cu-based crystalline phase, confirmed by EPMA analysis (Fig. 4), in good agreement with the thermodynamic calculations (Fig. 10). The Nb enrichment in the Fe-based amorphous phase rather than in the fcc-Cu crystalline globule of the rapidly solidified melt-spun ribbon was confirmed by nanobeam EDS (Fig. 9). The Nb enrichment in the Fe-based amorphous phase of the melt-spun ribbon is similar to that of the Fe-based crystalline phase region in arc-melted ingots evaluated by EPMA–WDX. In contrast, the detection of B was significantly more difficult using EDS, as it is a light element. Based on the GFA in Fe-based alloys [33,34,40], it is reasonable that B enriched the Fe-rich liquid rather than the Cu-rich liquid during rapid solidification in the melt-spinning process due to the difficulty of an amorphous phase forming in the binary Fe–Nb alloy system [47]. When the constituent phases of the melt-spun ribbon are hypothesized for the pure Cu phase and the ternary Fe75Nb10B15 amorphous phase in the (Fe0.75Nb10B15)80Cu20 alloy, the area fraction of the fcc-Cu crystalline region with respect to the Fe–Nb-based amorphous region (SCu:SAm.) was evaluated as 75.7:24.2 from the TEM microstructure analysis. The mole fraction of the fcc-Cu crystalline region and Fe–Nb-based amorphous region (MCu:MAm.) was obtained by

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converting SCu:SAm based on the densities of the fcc-Cu (8.96 kg m3) and the Fe0.75Nb10B15 (7.54 kg m3) resulting in a MCu:MAm value of 81.8:18.2. An experimental result for the Fe0.75Nb10B15 amorphous alloy density is not available, so the density of Fe66Nb4B30 metallic glass (7.54 kg m3) [48] was used for this study. Despite the estimation, the MCu:MAm = 81.8:18.2 mol fraction indicates that B enriched the Fe-rich liquid rather than the Cu-rich liquid during the melt-spinning rapid solidification process. The liquid phase separation into Fe- and Cu-based liquids, which was detected in this experimental study, was supported by the thermodynamic calculations. It should be emphasized that the alloy design based on the DHmix combination map of

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constituent element binary pairs and the predicted quaternary phase diagram based on ab initio calculations were useful for the development of new Fe-based amorphous alloy systems. Fcc-Cu crystalline globules dispersed in an Fe-based amorphous phase formed in the melt-spun ribbons of the Fe–Cu–Nb–B immiscible alloys. The measurement of the actual ejection temperature during the melt-spun ribbon preparation is vital for clarifying the mechanism of the unique solidification structure formation; this will be done in our future work. Without the actual ejection temperature, the formation of a unique solidification structure can be explained using mechanisms based on the liquid phase separation and amorphous phase formation. For the (Fe0.75Nb0.10B0.15)80Cu20 alloy, the upper limit temperature for liquid immiscibility is approximately T = 1820 K, shown in Fig. 10b. It is reasonable to consider that a single liquid in the fused quartz tube was generated by the radio-frequency heating, which then separated to the Fe– Nb–B-rich and Cu-rich liquids during the rapid cooling in the melt-spinning process, followed by the Fe–Nb–B-rich and Cu-rich liquids cooling separately. The high GFA of the Fe–Nb–B-based alloy liquid showed a liquid-to-glass transition, resulting in an Fe-based amorphous phase. The crystallization of a Cu-rich liquid phase with a spherical shape occurred during cooling, resulting in the crystalline globule formation. The emulsion-type structure formed through the liquid phase separation was frozen in place by the liquid-to-glass transition of the Fe-rich liquid matrix. One may consider another mechanism for the fcc-Cu crystalline globule dispersion in the Fe-based amorphous phase: an amorphous single phase forms initially, followed by fcc-Cu crystalline phase precipitation from the amorphous single phase through thermal crystallization. However, the model without the liquid phase separation is unlikely as the Cu partition from the amorphous single phase to the fcc-Cu crystalline globules does not have sufficient time to form during the considerably brief melt-spinning process. The simultaneous occurrence of a liquid phase separation and amorphous phase formation in the Fe–Nb–B-based alloy liquids is the key factor for the emulsion-type structure formation. It is well known that Cu plays an important role in Fe-based amorphous alloys at even a small amount (approximately 1 at.%) for alloys such as Fe–Zr–B–Cu (i.e., NANOPERM) [49] and Fe–Cu– Si–Nb–B (i.e., FINEMET) [50]. Cu clustering during annealing in the Fe-based amorphous phase leads to the nano-crystallization of bcc-Fe in FINEMET alloys [51]. In this study, the possible phase separation in the aforementioned Fe-based amorphous alloys was not considered, as this was outside the scope of this study. The formation of core–shell type fcc-Cu crystalline globules embedded in a Fe-based amorphous alloy was reported in Fe–Zr– B–Cu alloys [26]; whereas the core–shell type nano-crystalline globules were not observed in the Fe–Nb–B–Cu alloys. The difference in the glass forming ability and microstructure among Fe–Si–B–Cu, Fe–Nb–B–Cu, and Fe–Zr–B–Cu [26–28] alloys will be discussed in a future paper.

6. Conclusions The microstructure of arc-melted ingots and melt-spun ribbons of Fe–Cu–Nb–B immiscible alloys was investigated, showing the formation of an amorphous phase and a rapidly solidified structure. The obtained results are summarized as follows:

Fig. A1. Crack propagation along Cu crystalline globules embedded in the Fe–Nb–B amorphous matrix, represented here by the Cu globules embedded in a (Fe0.75Nb0.10B0.15)80Cu20 melt-spun alloy ribbon. (a) BF image, (b1) BF image, and (b2) the results of the crack propagation analysis of b1.

(1) Amorphous phase formation was confirmed in the meltspun ribbon despite the liquid-phase separation in quaternary Fe–Cu–Nb–B immiscible alloys. (2) An Fe–Nb–B-based amorphous phase with 10–100 nm fccCu crystalline globules formed in the melt-spun ribbon of the (Fe0.75Nb0.10B0.15)80Cu20 alloy. A liquid phase separation

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to Fe–Nb–B- and Cu-based liquids, in addition to the high glass-forming ability of the Fe–Nb–B alloy system, produced a unique solidification structure in the Fe–Cu–Nb–B immiscible alloys. (3) A prediction method was suggested for Fe–Cu-based immiscible alloys with an amorphous phase formation.

[5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

Acknowledgment This work was supported by JSPS KAKENHI Grant Number 50362661. Appendix A The brittleness of metallic glasses poses a significant problem for their practical applications. A dispersion of crystalline globules is effective in stimulating crack deflection, which may lead to the improvement of the glass deformability; a typical example of such a dispersion is shown in Fig. A1. Fig. A1 shows the TEM-BF image of a crack in a (Fe0.75Nb0.10B0.15)80Cu20 alloy melt-spun ribbon, with the crack having been introduced during the TEM specimen preparation. In Fig. A1a, the letters X1 and X2 indicate typical examples of interfacial delamination between the Fe-based metallic glass matrix and the crystalline Cu globules. The crack clearly cuts through the crystalline Cu globules, indicated by the letters Y1 and Y2. These can direct the crack deflection of the metallic glass matrix along the Cu crystalline globules, indicated by the letter Z. In Fig. A1b1, the letter A displays a zigzag crack rather than a straight crack. To evaluate the effect of the Cu globules on crack propagation in detail, an analysis was performed. The number of Cu globules per unit length (C) was defined as follows:

C ¼ N=L

ðA1Þ

where N is the number of Cu globules either on a line along the crack or on a randomly selected straight line, and L denotes the length of the line in nm. The analyses results are shown in Fig. A1b2, with the line A running along the crack, and the lines B–G are arbitrary lines drawn to evaluate the average Cu globule density embedded in the Fe–Nb–B-based metallic glass matrix. The values of CA–CG were as follows: CA(crack) = 0.00565 nm1, while CB = 0.00238, CC = 0.00156, CD = 0.00162, CE = 0.00298, CF = 0.00204, and CG = 0.00301 nm1. The value of CA is significantly larger than CB–CG, indicating that the crack propagated along the Cu globules. The crack deflection achieved by the Cu crystalline globule dispersion can be considered effective for improving the deformability of Fe-based metallic glasses.

[15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40]

[41] [42] [43]

[44] [45] [46]

[47] [48]

References [1] Y. Nakagawa, Acta Metall. 6 (1958) 1086. [2] P.A. Kindqvist, B. Uhrenius, On the Fe–Cu phase diagram, Calphad 4 (1980) 193–200. [3] T.Y. Chuang, R. Schmid, Y.A. Chang, Metall. Mater. Trans. A (1921) 15. [4] A. Munitz, Metall. Trans. 18B (1987) 565.

[49] [50] [51]

Q. Chen, Z. Jin, Metall. Mater. Trans. A 26 (1995) 417. G. Wilde, R. Willnecker, R.N. Sigh, F. Sommer, Z. Metallkd. 88 (1997) 10. G. Wilde, J.H. Perepezko, Acta Mater. 47 (1999) 3009. S.-E. Amara, A. Belhadj, R. Kesri, S. Hamar-Thibault, Z. Metallkd. 90 (1999) 2. X.Y. Lu, C.D. Cao, B. Wei, Mater. Sci. Eng., A 313 (2001) 198. M.A. Turchanin, P.G. Agraval, I.V. Nikolaenko, J. Phase Equil. 24 (2003) 307. X.-M. Zhang, W.-L. Wang, Y. Ruan, B.-B. Wei, Chin. Phys. Lett. 27 (2010) 026401. T. Nagase, A. Yokoyama, Y. Umakoshi, J. Alloys Comp. 509 (2011) 1178. T. Fang, L. Wang, C.X. Peng, Y. Qi, J. Phys.: Condens. Matter 24 (2012) 505103. O.A. Chikova, V.S. Tsepelev, A.N. Konstantinov, V.V. V’yukhin, Russ. Metall. 5 (2013) 13. J. Zhang, X. Cui, Y. Yang, Y. Wang, Metall. Mater. Trans. A 44 (2013) 5544. T. Nagase, Mater. Trans. 55 (2014) 304. K. Taguchi, H. Ono-Nakazato, T. Usui, ISIJ Int. 46 (2006) 633. K. Yamaguchi, T. Ohara, S. Ueda, Y. Takeda, Mater. Trans. 47 (2006) 1864. B.J. Park, Y.M. Chen, T. Ohkubo, Intermetallics 17 (2009) 958. T. Nagase, Y. Umakoshi, J. Alloys Comp. 495 (2010) L1. A.A. Kundig, M. Ohnuma, D.H. Ping, T. Ohkubo, K. Hono, Acta Mater. 52 (2004) 2441. B.J. Park, H.J. Chang, D.H. Kim, W.T. Kim, Appl. Phys. Lett. 85 (2004) 6353. N. Mattern, U. Kuhn, A. Gebert, T. Gemming, M. Zinkevich, H. Wendrock, L. Schultz, Scr. Mater. 53 (2005) 271. D.H. Kim, W.T. Kim, E.S. Park, N. Mattern, J. Eckert, Prog. Mater Sci. 58 (2013) 1103. I. Yamauchi, T. Irie, H. Sakaguchi, J. Alloys Comp. 403 (2005) 211. T. Nagase, A. Yokoyama, Y. Umakoshi, Mater. Trans. 47 (2006) 1105. T. Nagase, A. Yokoyama, Y. Umakoshi, J. Alloys Comp. 494 (2010) 295. T. Nagase, A. Yokoyama, Y. Umakoshi, Scr. Mater. 63 (2010) 1020. T. Koziel, Z. Kedzierski, A. Zielinska-Lipiec, K. Ziewiec, Scr. Mater. 54 (2006) 1991. K. Ziewiec, P. Malczewski, R. Gajerski, A. Ziewiec, J. Non-Cryst. Solids 357 (2011) 73. K. Ziewiec, K. Bryla, A. Blachowski, K. Ruebenbauer, K. Prusik, S. Kac, T. Koziel, J. Microsc. 237 (2010) 232. K. Ziewiec, A. Ziewiec, K. Prusik, J. Achiev. Mater. Manuf. Eng. 37 (2009) 532. T. Koziel, Z. Kedzierski, A. Zielinska-Lipiec, J. Latuche, G. Cieslaki, J. Microsc. 237 (2010) 267. T. Koziel, A. Zielinska-Lipiec, J. Layurh, S. Kac, J. Alloys Comp. 509 (2011) 4981. T. Koziel, J. Latuch, A. Zielinska-Lipiec, Arch. Metall. Mater. 58 (2013) 601. C.D. Cao, S.L. Gong, J.B. Guo, R.B. Song, Z.B. Sun, S. Yang, W.M. Wang, Chin. Phys. B 21 (2012) 086102. A. Inoue, Acta Mater. 48 (2000) 279. F.R. Boer, D.G. Perrifor, Cohesion in Metals, Elsevier Science Publishers B.V., Netherlands, 1988. A. Takeuchi, A. Inoue, Mater. Trans. 28 (2005) 2817. Y. Kawazoe, J.-Z. Yu, A.P. Tsai, T. Masumoto, Nonequilibrium Phase Diagrams of Ternary Amorphous Alloys. Landolt-Bornstein: Numerical Data and Functional Relationships in Science and Technology, Springer, Berlin, 1997. A. Makino, S. Arano, A. Inoue, T. Masumoto, J. Jpn. Inst. Metals 58 (1994) 106. A. Jain, G. Hautier, C. Moore, S.P. Ong, C. Fischer, T. Mueller, K. Persson, G. Ceder, Comput. Mater. Sci. 50 (2011) 2295. S.P. Ong, A. Jain, G. Hautier, M. Kocher, S. Cholia, D. Gunter, D. Bailey, D. Skinner, K. Persson, G. Ceder, The Materials Project, (accessed on 19.03.14). S.P. Ong, L. Wang, B. Kang, G. Ceder, Chem. Mater. 20 (2008) 1798. A. Jain, G. Hautier, S.P. Ong, C. Moore, C. Fischer, K. Persson, G. Ceder, Phys. Rev. B 84 (2011) 045115. C.W. Bale, A.D. Pelton, W.T. Thompson, G. Eriksson, FactSageÒ, Ecole Polytechnique, Montreal, 2001, (accessed on 26.03.14). U. Mizutani, Y. Hoshino, Y. Yamada, Amorphous Gokin Sakusei No Tebiki, Agne, 1986 (in Japanese). M. Stoica, K. Hajlaoui, J. Das, J. Eckert, A.R. Yavari, Rev. Adv. Mater. Sci. 18 (2006) 61. K. Suzuki, A. Makino, N. Kataika, A. Inoue, T. Masumoto, Mater. Trans., JIM 32 (1991) 93. Y. Yoshizawa, S. Oguma, K. Yamauchi, J. Appl. Phys. 64 (1988) 6044. K. Hono, K. Hiraga, Q. Wang, A. Inoue, T. Sakurai, Acta Metell. Mater. 40 (1992) 2137.