Journal of Materials Science & Technology, 2020, 51(0): 173-179 DOI: 10.1016/j.jmst.2020.01.066

Research Article

Spheroidization behaviour of a Fe-enriched eutectic high-entropy alloy

Yu Yina, Damon Kenta,b, Qiyang Tana, Michael Berminghama, Ming-Xing Zhang,a,*

a School of Mechanical and Mining Engineering, the University of Queensland, St Lucia, QLD 4072, Australia

b School of Science and Engineering, University of the Sunshine Coast, Sippy Downs, QLD 4556, Australia

Corresponding authors: Mingxing.Zhang@uq.edu.au(M. - X. Zhang)

Received: 2019-12-1   Accepted: 2020-01-28   Online: 2020-08-15

Abstract

A cost-effective Fe-enriched eutectic high-entropy alloy (EHEA), Fe35Ni25Cr25Mo15, was designed and prepared to avoid the use of expensive Co that is commonly used in HEAs. However, the as-cast Fe-enriched EHEA was associated with brittleness. The present work aims to evaluate the possibility and feasibility of spheroidization of the lamellar structure of the EHEA in order to improve the ductility. Due to the high cooling rate of arc-melting, the as-melted Fe35Ni25Cr25Mo15 EHEA was found to be a pseudo eutectic alloy comprised of alternant σ phase (Cr0.22Mo0.18Fe0.6-type intermetallic) and face centred cubic (FCC) phase. The lamellar structure in the Fe-enriched EHEA remained stable up to 800 °C. The instability of the lamellar structure occurred at temperatures over 800 °C, which was resulted from migration of high-density faults (i.e. lamellar termination and ledges in the lamellae). However, the Fe35Ni25Cr25Mo15 EHEA still exhibited brittleness even after spheroidization at 1100 °C for 168 h due to the formation of the hard and brittle σ matrix in the pseudo Fe35Ni25Cr25Mo15 EHEA as a result of decomposition of the lamellar structure. Therefore, in contrast to the softening of traditional eutectic alloys, spheroidization treatment was considered as invalid to improve the ductility of pseudo-eutectic HEA with high fraction of intermetallic phase. The present work provides a valuable reference for those who aim to improve the ductility of brittle EHEA through spheroidization.

Keywords: Eutectic high-entropy alloys ; Instability ; Lamellar structure ; Spheroidization ; Mechanical properties

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Cite this article

Yu Yin, Damon Kent, Qiyang Tan, Michael Bermingham, Ming-Xing Zhang. Spheroidization behaviour of a Fe-enriched eutectic high-entropy alloy. Journal of Materials Science & Technology[J], 2020, 51(0): 173-179 DOI:10.1016/j.jmst.2020.01.066

1. Introduction

Unlike traditional alloys containing one or two principal elements, high entropy alloys (HEAs) normally consist of 3 or more principal elements [[1], [2], [3]]. For as-cast HEAs, research over the last decade has shown that the most non-equiatomic HEAs with multi-phases exhibit superior mechanical properties compared with the most equiatomic single-phase HEAs [[4], [5], [6], [7]] and dual-phase (DP) eutectic high-entropy alloys (EHEAs) in particular [[8], [9], [10]]. For example, Lu et al. [9] firstly proposed to use eutectic alloy concept to design HEAs, aiming at good castability and composite structure to resolve the strength-ductility trade-off. The directly casting AlCoCrFeNi2.1 DP EHEA exhibit both high strength and ductility in a wide temperature range, which outperform most of conventional cast alloys and also HEAs reported previously [9]. In addition, based on a pseudo binary strategy proposed by Jin et al. [11], three groups of AlCoCrFeNi-based EHEAs with tensile fracture strengths of ∼1.0 GPa and elongations of over 10% were developed. Similarly, a quaternary Fe20Co20Ni41Al19 EHEA with nano-lamellar structure presented an ultimate tensile strength of 1.3 GPa and an elongation of 17.1% at room temperature, which is superior to most reported as-cast HEAs [12].

Except for the above EHEAs, most other reported EHEA systems exhibited ultra-high strength but substantial brittleness. This includes the as-cast CoFeNi2V0.5Nb0.75 and Co2Mo0.8Ni2VW0.8 EHEAs. [[14], [15], [16]]. These alloys showed high ultimate compression strength (UCS) of around 2 GPa, but low compression fracture strain (εc) below 15%. In addition, almost all the current EHEAs (e.g. AlCoCrFeNi2.1, Fe20Co20Ni41Al19, CoFeNi2V0.5Nb0.75 and Co2Mo0.8Ni2VW0.8) contain expensive metals, such as Co or Mo, which increased the overall cost of the alloys. Hence, the critical challenges to applicate this type of EHEA are to enhance their ductility and decrease the cost. To address this problem, we previously designed a low-cost Fe-enriched dual-phase EHEA - Fe35Ni25Cr25Mo15 [13]. But, as-cast Fe35Ni25Cr25Mo15 alloy also showed high brittleness at room temperature.

Adopting the concept of spheroidization from traditional eutectic alloys, the present work aims to evaluate the feasibility of spheroidization of the lamellar structure of EHEAs in order to improve ductility. The brittle Fe-enriched EHEA (Fe35Ni25Cr25Mo15) [13] was used as an example. The instability of the lamellar structure in this EHEA under high temperature annealing was investigated and the influence of spheroidization on the room-temperature mechanical properties was evaluated.

2. Alloy design and experimental procedure

To lower the cost of HEAs, the overall design strategy was to avoid the use of expensive Co that is commonly used in HEAs. Based on Yu and co-workers’ previous work [13], a new cost-effective Fe35Ni25Cr25Mo15 EHEA was designed through increasing the ratio of Fe to Mo in an equiatomic Fe25Ni25Cr25Mo25 HEA to produce more ductile face centred cubic (FCC) phase because the latter was predominated by brittle intermetallic phase. It has been commonly considered that Fe and Ni are FCC stabilizers, while Cr and Mo are intermetallic forming elements [13]. Hence, the EHEA can be designed through adjusting the ratio of Fe/Ni to Cr/Mo. In addition, to ensure the high oxidation and corrosion resistance, Cr content was retained at 25 at.%. As Fe is much cheaper than other elements while Mo is the most expensive one in the FeNiCrMo alloy system, the ratio of Fe to Mo was accordingly increased to obtain the dual phase (DP) eutectic structure consisting of FCC and intermetallic phase. The proposed Fe-enrich EHEA exhibited high strength at room temperature. Due to its composite structure and high Cr content, this new alloy should possess potential applications as high-temperature materials. But, reducing its room-temperature brittleness remains a great challenge.

The designed Fe35Ni25Cr25Mo15 EHEA was prepared in an arc-melting furnace using a water-cooled copper mould in an argon atmosphere. Before arc-melting, commercial-grade metal powders (Fe, Ni, Cr and Mo) with 99.9 wt.% purity were mixed and then compacted into small cylinders of 14 mm in diameter and 20 mm in length. Then, the compacted cylinders were melted in the arc furnace. To ensure the chemical homogeneity, each ingot was remelted four times with processing parameters of 250 A current and 28.5 V voltage. Mild oxidation was notified on the top surface of the ingots. This was attributed to the high specific surface area of the powders used. However, since samples for mechanical property tests and microstructure characterization were cut from the middle of the bar ingots, the effect of surface oxidation can be reasonably ignored. Rectangular bar ingots with dimension of about 120 mm × 12 mm × 8 mm were cut to various sizes for heat treatment, microstructural characterization and mechanical property tests. Spheroidization treatments were conducted in an air furnace within a temperature ranging from 600 °C to 1100 °C for 24, 72, and 164 h followed by air-cooling.

The phase constituents were identified through X-ray diffraction (XRD; Bruker D8) using Cu radiation at a scanning rate of 1° min-1, and a 2θ angle ranging from 20° to 100°. Metallographic samples with a thickness of around 2 mm were mounted with electroconductive resin. The mounted samples were then mechanically ground and polished followed by etching using an aqua regia etchant (HNO3 + 3 HCl). The microstructure was examined in a scanning electron microscopy (SEM; JEOL-6610) and the chemical compositions were analysed using energy-dispersive spectrometry (EDS). Transmission electron microscopy (TEM) samples were prepared using a Focused Ion Beam (FIB) method in a FEI Scios Dual Beam system and examined in a TECNAI F20 TEM with operating voltage of 200 kV.

Differential Scanning Calorimetry (DSC; NETZSCH thermische analyse) analysis was carried out over a temperature ranging from 20 °C to 1450 °C at a heating and cooling rate of 10 °C min-1 in an argon atmosphere. The DSC samples with dimension of 4.3 mm in diameter and 3 mm in length were cut from the bar ingots, and then ground and polished before testing. After the DSC, microstructural characterization of these samples was also performed using the JEOL-6610 SEM. Hardness was measured using a Struers Vickers hardness tester with a load of 5 kg and a dwell time of 12 s. Due to the brittleness of the alloy, mechanical properties were measured through compression tests. Test samples were machined to the size of 5 mm in diameter and 10 mm in length. The testing was conducted on a 100 kN Shimadzu AGS-X machine fitted with video extensometer at a strain rate of 5 × 10-4 s-1 at room temperature, and representative compressive true stress-strain curves were presented.

3. Results and discussion

3.1. Microstructure and phase constitution of the as-cast EHEA

Fig. 1(a, b) show the backscattered electron (BSE) morphology of the arc-melted Fe35Ni25Cr25Mo15 EHEA. The as-cast alloy exhibited a typical lamellar eutectic (LE) structure (non-faceted/non-faceted) with small amounts of anomalous eutectic (AE) structure. From the high magnification image in Fig. 1(b), a high-density of faults can be identified in the lamellar structure, including lamellar terminations (marked as blue arrows) and some ledges in the lamellae (red arrows). The EDS mapping in Fig. 1(c) shows chemical contrast of two phases, including a Fe and Ni enriched phase and a Mo enriched phase in the lamellae. The XRD spectrum in Fig. 1(d) confirms the dual-phase structure of the as-cast Fe35Ni25Cr25Mo15 EHEA. The phases are an FCC solid solution and a Cr0.22Mo0.18Fe0.6-type body-centred tetragonal phase (σ phase). TEM was performed to further verify the phase constituents in the alloy.

Fig. 1.   BSE images (a, b), EDS mapping (c) and XRD spectra (d) of the as-cast Fe35Ni25Cr25Mo15 EHEA; blue arrows in (b) indicate the lamellar termination and red arrows indicate ledges in the lamellae.


To understand the formation sequence of the LE and AE revealed in Fig. 1(a, b), DSC analysis was conducted. Fig. 2 shows the DSC heating (a) and cooling curves (b) of the as-cast EHEA. Although the heating curve shows one endothermic peak, indicating simultaneous melting of the two phases in both the LE and AE regions at ∼1318 °C, three exothermic peaks can be identified in the cooling curve between 1350 °C and 1288 °C, corresponding to different phase transformations during solidification. To understand the phase transformations occurred during DSC cooling, the microstructure of the sample after DSC test was examined and is shown in Fig. 3. Three distinguishable microstructures, pro-eutectic σ phase (marked as P), LE structure and AE structure can be observed, which differ from the microstructure of the as-cast samples. This difference in microstructure results from the different cooling rates during the DSC versus the arc-melting used to prepare the samples. In contrast to the slow cooling rate during the DSC test (10 °C min-1), the typical cooling rate for arc-melting is much higher (103-104 °C s-1) [29,30], which inhibited the formation of the pro-eutectic σ phase in the as-cast sample. Slow cooling during DSC enabled primary σ phase to form. This result indicates that the Fe35Ni25Cr25Mo15 alloy is a hyper-eutectic alloy. According to previously published results on traditional eutectic alloys, the AE structure tends to form at high undercooling as a result of irregular or partial remelting [[17], [18], [19], [20]].

Fig. 2.   DSC heating (a) and cooling curves (b) of as-cast Fe35Ni25Cr25Mo15 EHEAs.


Fig. 3.   BSE image of DSC-tested EHEA.


In addition, the formation of the various microstructures during solidification are marked near the corresponding exothermic peaks on the DSC cooling curve. During solidification of the EHEA, the primary σ phase forms first, corresponding to the smallest peak in Fig. 2(b). From Fig. 3, it is observed that the volume fraction of AE is higher as compared with the primary σ phase and LE region, hence it corresponds to the highest (3rd) exothermic peak in the DSC cooling curve and thereby the second peak is related to the formation of the LE region.

3.2. Instability of lamellar structure in EHEA at elevated temperature

Fig. 4(a-g) show optical microscopy (OM) images revealing the morphological evolution of the EHEA when annealing at various temperatures from 600 °C to 1100 °C for 24 h. The lamellar eutectic structure remained stable up to 800 °C. With annealing at temperatures over 800 °C, the lamellar eutectic structure started to pinch off, coarsen and spheroidize. The lamellar structure completely decomposed at 1000 °C and 1100 °C, though the microstructure was not fully spheroidized. This result indicates that, in the current as-cast EHEA, although the lamellar structure was destabilized over 800 °C, full spheroidization required much longer annealing time. Fig. 5 shows the hardness variation of the EHEA with annealing temperature. The hardness of the samples annealed below 800 °C remained similar to the as-cast alloy, which is consistent with the similar microstructures observed in these samples (Fig. 4(a-d)). Whereas, due to the decomposition of the lamellar structure, the hardness of the samples annealed at temperatures over 900 °C is gradually reduced. As spheroidization is regarded as one of the most effective approaches to increase the ductility of eutectic or eutectoid alloys, longer time annealing at 1100 °C was conducted to maximize spheroidization of the LE in the EHEA.

Fig. 4.   (a-g) OM images of Fe35Ni25Cr25Mo15 EHEA after annealing at different temperatures for 24 h.


Fig. 5.   Hardness variation of as-cast and as-treated Fe35Ni25Cr25Mo15 EHEAs.


Fig. 6 shows XRD spectra of the Fe35Ni25Cr25Mo15 alloy after annealing at 1100 °C for different time, indicating the same phase constituents (dual FCC and σ phase) as the original as-cast alloy. Thus, the annealing did not cause phase transformation. Fig. 7(a-d) are BSE images of the Fe35Ni25Cr25Mo15 alloys annealed at 1100 °C, which show two distinguishable phase features (a white region and a grey region) with the contrast due to their differing chemical compositions. EDS analysis indicated that the grey areas were Fe and Ni enriched FCC phase and the white areas were Cr and Mo enriched σ phase. As aforementioned, although the lamellar structure decomposed after 24 h annealing, complete spheroidization did not occur. An increased annealing time up to 72 h significantly improved the spheroidization. However, further annealing up to 168 h did not promote further spheroidization of the alloy. For longer annealing times above 72 h, the microstructure consisted of the σ phase as the matrix with spheroidized FCC phase.

Fig. 6.   XRD spectra of Fe35Ni25Cr25Mo15 EHEA after annealing at 1100 ℃ for various time.


Fig. 7.   BSE images of Fe35Ni25Cr25Mo15 EHEA after annealing at 1100 ℃ for various time: (a) as cast, (b) annealed for 24 h, (c) annealed for 72 h, (d) annealed for 168 h.


To verify the phase constituents shown in Fig. 6, TEM characterization is presented in Fig. 8, Fig. 9. The TEM bright field image and corresponding selected diffraction patterns in Fig. 8(a-c) confirm the two phases to be FCC and σ phase in the as-cast alloy. The TEM-EDS mapping (Fig. 8(d)) indicates that the FCC phase is Fe and Ni enriched (grey area in Fig. 7(a)), while the σ phase is rich in Cr and Mo (white area in Fig. 7(a)). Fig. 9(a-d) indicates that the spheroidized alloy is comprised of spheroidized FCC phase (grey area in Fig. 7(d)) and the σ phase matrix (white area in Fig. 7(d)) after annealing at 1100 °C for 168 h.

Fig. 8.   TEM image (a); SAD patterns (b, c); and EDS mapping (d) of as-cast Fe35Ni25Cr25Mo15 EHEAs (FCC phase is marked by a red dot and s phase marked by a blue dot).


Fig. 9.   TEM image (a); SAD patterns (b, c); and EDS mapping (d) of Fe35Ni25Cr25Mo15 EHEA after annealing at 1100 ℃ for 168 h (FCC phase is marked by a red dot and s phase marked by a blue dot).


The above microstructure characterization indicates that the lamellar structure of the as-cast Fe35Ni25Cr25Mo15 EHEA started to destabilize after annealing at 800 °C for 24 h. Generally, a perfectly infinite lamellae with flat faces can be regarded to be intrinsically stable since there is no primary radius to drive perturbation growth [21]. However, most lamellar structures in eutectic alloys contain defects, such as lamellar terminations, ledges in the lamellae and colony/grain boundaries, which induce coarsening of the lamellar structure at higher temperatures. Depending on the defects, coarsening can occur through a continuous mode (fault migration) or a discontinuous mode (grain boundary migration) [[21], [22], [23]]. Currently, the fault migration theory is widely accepted for understanding the instability of lamellar structures at high temperatures. As shown in Fig. 1(b), high-density faults (i.e. lamellar terminations and ledges in the lamellae) can be clearly identified in the alloy. The difference in curvature between the fault tip and the adjacent flat interface is related to a chemical potential gradient. Capillary forces derived from the curvature difference promote atomic diffusion from the tips to the adjacent flat interfaces in the lamellae [21]. Accordingly, the fault tip is dissolved and the adjacent flat surface is coarsened, resulting in partial spheroidization of the lamellar structure.

In addition, increasing the temperature from 800 °C to 1100 °C promoted atomic diffusion and fault migration, therefore fostering the spheroidization as shown in Fig. 4(a-g). According to the model developed by Cline et al. [23,24], the fault migration rate in binary eutectic can be expressed as follows:(1)VF=4DV¯αγCβλα2RT(Cα-Cβ)ln(2λ/λα)where VF is the fault migration rate, D is the diffusion coefficient, γ is the interfacial energy between the α and β phases, Cα and Cβ are the concentrations of solutes in both phases (Cα > Cβ), λ is the inter-lamellar spacing, λα is lamellar thickness of solute rich phase, T is temperature and V¯α is the partial molar volume of solute rich phase. For a given eutectic alloy, the parameters in Eq. (1) can be assumed to be invariant except the temperature T and diffusion coefficient D. Thus, the fault migration rate VF is mainly defined by the value of D/T [25]. According to Gupta et al. [26], the diffusion coefficient D in Pb-Sn eutectic alloys increased up to several orders of magnitudes when the temperature T surpassed a critical value. Therefore, the high temperature annealing treatments over 800 °C may significantly increase rates of diffusion in the Fe35Ni25Cr25Mo15 EHEAs and thereby the velocity of fault migration. Combined with the high-density faults in the eutectic structure of the Fe35Ni25Cr25Mo15 alloy, instability is substantially increased at elevated temperature.

Fig. 7 shows that spheroidization in the current alloy was not fully completed at 1100 °C even after 168 h (Fig. 7(a-d)). As the microstructure after annealing for 72 h was similar to that for 168 h, the potential to achieve full spheroidization at 1100 °C is limited. Higher temperatures are necessary to complete the spheroidization. A similar result was reported by He et al. [25] in a CoCrFeNiNb0.65 EHEA. Partial spheroidization occurred in this alloy at 900 °C after 24 h, but the majority of the spheroidization process was not complete. In addition, using deep etching and tilted sample observation with SEM, Lei et al. reported that the observed spheroids in the NiAl-Cr(Mo) eutectic alloys that were annealed at 1150 °C for 400 h were actually cylindrical in shape [27]. Higher annealing temperatures were required to break the cylindrical phase into spheres. For instance, spheroidization of a NiAl-Mo eutectic alloy was only completed at around 1575 °C for 100 h [28]. Thus, it is reasonable to believe that full spheroidization of lamellar eutectic in EHEAs is challenging, and generally requires very long time and/or high temperatures.

3.3. Influence of spheroidization on mechanical properties of EHEA

Fig. 10(a) shows the hardness of the EHEA after annealing at 1100 °C for various time. Compared with the as-cast alloy, annealing led to a slight decrease in hardness, but the annealing time had marginal impact. The corresponding true compressive stress-strain curves of the Fe35Ni25Cr25Mo15 EHEA are shown in Fig. 10(b). The as-cast alloy exhibited high ultimate strength of 1874 MPa with low strain to fracture (ductility) of 3.7%. Annealing at 1100 °C for 72 h only slightly enhanced the strain to fracture to 6.5% with reduction in ultimate strength to 1570 MPa. Whereas, 168 h annealing decreased both the ultimate strength and strain to fracture.

Fig. 10.   Hardness variation (a), and compressive true stress-strain (b) of Fe35Ni25Cr25Mo15 EHEA with annealing time at 1100 °C.


The decrease in ultimate compression strength of the Fe35Ni25Cr25Mo15 EHEA after annealing at 1100 °C originated from decomposition of the lamellar structure due to its shape instability at elevated temperatures. Fig. 7, Fig. 10 show the close correlation between the eutectic morphology and mechanical properties of the alloy. From Fig. 7(a-d), one can see that annealing at 1100 °C led to coarsening and spheroidization of the lamellae in the Fe35Ni25Cr25Mo15 alloy. The interlamellar spacing was remarkably increased from ∼0.4 μm to ∼4 μm. The bigger interlamellar spacing resulted in a reduction in ultimate compression strength.

Generally, due to the substantial reduction in the volume fraction of interphase boundaries, spheroidization will significantly improve the ductility of eutectic alloys in sacrifice of the strength. However, the Fe35Ni25Cr25Mo15 EHEA still exhibited brittleness after spehroidization. This was attributed to the formation of brittle and hard σ matrix as a result of the spheroidization although the FCC phase was spheroidized (Fig. 7(c, d)). As aforementioned, the Fe35Ni25Cr25Mo15 EHEA is a pseudo eutectic alloy resulting from the fast rate of cooling after arc-melting (Figs. 1(a) and 3). Its hyper-eutectic composition may be associated with higher fractions of the brittle σ phase in the arc-melted EHEA than for equilibrium EHEAs. Thus, the high-proportioned σ phase in the lamellae evolve to the matrix after spheroidization and should be responsible for the brittleness of the spheroidized Fe35Ni25Cr25Mo15 EHEA.

In addition, from Fig. 7, it is reasonable to consider that σ phase acts as the matrix even after longer time or higher temperature annealing, indicating that the brittleness cannot be eliminated by fully spheroidization. Therefore, spheroidization treatment was invalid for the ductility improvement of the brittle pseudo eutectic EHEA with higher fraction of intermetallic phase. This conclusion probably would be true for traditional eutectic alloys provided intermetallic matrix formed after spheroidization of the alloys. However, this is quite unlikely because traditional eutectic alloys do not contain sufficient alloying elements to form intermetallic phases as the matrix. Further studies will provide solid conclusions.

4. Conclusions

(1)The arc-melted Fe-rich Fe35Ni25Cr25Mo15 EHEA is a pseudo eutectic alloy formed due to the high cooling rate during solidification. The microstructure comprises a primarily lamellar eutectic structure with small amount of anomalous eutectic structure. Both consist of Mo-Cr enriched σ phase and Fe-Ni enriched FCC phase.

(2)The lamellar structure of the EHEA was stable up to 800 °C when annealing for 24 h. Above this temperature, morphological instability of the microstructure originating from the migration of high-density faults (i.e. lamellar termination and ledges in the lamellae) occurred. In addition, with increasing annealing temperature, the instability increased significantly due to the higher velocity of fault migration.

(3)The as-cast alloy exhibited high ultimate compression strength of 1874 MPa but low ductility of 3.7%. Unlike traditional eutectic alloys, spheroidization treatment was invalid for the ductility improvement of brittle pseudo-eutectic HEA with higher fraction of intermetallic phase. The brittleness of the spheroidized Fe-enriched EHEA was originated from the the hyper-eutectic composition and thereby the formation of a brittle and hard σ matrix after decomposition of the lamellar structure.

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