International Journal of Minerals, Metallurgy and Materials https://doi.org/10.1007/s12613-020-2214-x

International Journal of Minerals, Metallurgy and Materials Accepted manuscript, https://doi.org/10.1007/s12613-020-2214-x © University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2020

Two Refractory High-entropy Alloys CrHfNbTaTi and CrHfMoTaTi: Phase,

Microstructure and Compressive Properties

Jiaojiao Yia, Fuyang Cao b, Mingqin Xua, Lin Yangc*, Lu Wangc, Long Zengd aSchool of Mechanical Engineering, University of Technology, 1801 Zhongwu RD, Zhonglou ,

Changzhou, 213001, PR

Email: Jiaojiao Yi: [email protected]

Mingqin Xu: [email protected] bSchool of Materials Science and Engineering, Jiangsu University, No.301, Xuefu RD, ,

Zhenjiang, 212013, PR China

Email: [email protected] cSchool of Materials Science and Engineering, Jiangsu University of Technology, 1801 Zhongwu RD, Zhonglou

District, , 213001, PR China

Email: Lin Yang*: [email protected]

Lu Wang: [email protected] dSchool of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan RD. Minhang

District, Shanghai, 200240, PR China

Email: [email protected]

Abstract

Two new refractory high-entropy alloys, CrHfNbTaTi and CrHfMoTaTi, derived from the well-known

HfNbTaTiZr alloy by principal element substitutions, were prepared by vacuum arc-melting. Their phase components, microstructures, and compressive properties in the as-cast state were investigated intensively, and the results showed that bothAccepted alloys were mainly Manuscript composed of BCC and Not cubic laves Copyedited phases. In terms of the mechanical properties, the yield strength increased remarkably from 926 MPa for HfNbTaTiZr to 1258 MPa for

CrHfNbTaTi, while retaining a promising ductility of around 24.3 % fracture strain for CrHfNbTaTi. The morphology and composition of the network-shaped interdendritic regions were closely related to the improved mechanical properties due to elemental substitution. Dendrites were surrounded by an incompact interdendritic shell after Mo incorporation, which deteriorated the yield strength and increased the brittleness. Keywords: Refractory; High-entropy alloys; Phase structure; Microstructure; Yield strength; Elongation

1. Introduction

Due to the development of the aerospace, energy generation, and petrochemical industries, there is a high demand for structural materials that can be applied in high-temperature environments. To cater to the corresponding rigorous operation conditions of these industries, Ni-based superalloys have been proposed, but they have an incipient melting point near 1300 ℃, which greatly limits their applications to those which have temperatures between 1160 ℃ and 1277 ℃. [1-3] In addition to most of conventional refractory metals or alloys, the design of Ni-based superalloys is based on one simple idea: one element with attractive properties is used as the principal element, and minor amounts of other elements are added to improve these properties.[4] This has sometimes been referred to as a mature technology because it offers finitude modification possibilities.[5]

To overcome inherent drawbacks, high-entropy alloys (HEAs) have been designed, which consist of at least

5 principal elements with the atomic concentrations of each element between 5 and 35 at.%. [6, 7] Soon after the introduction of HEAs, refractory HEAs, in which the most commonly used elements were Cr, Ti, Zr, Hf, Nb, Ta,

Mo, W, and V, were first proposed by Senkov et al.[8, 9] The first two reported refractory HEAs, NbMoTaW and

VNbMoTaW, both possessed a simple body-centered cubic (BCC) microstructure and pronounced solution hardening.[8, 9] The two alloys also have high yield strengths of 405 MPa and 477 MPa at 1600 ℃, respectively; however, their drawbacks are noticeable, including a density much higher than 9.0 g/cm3, room temperature (RT) brittleness (fracture strain of about 1.7%), and a low RT yield strength of about 1000 MPa. Subsequently, many new refractory alloys were tried and reported.[10-13] Thereinto, the most well-known refractory alloy,

HfNbTaTiZr, with a BCC structure revealed excellent ductility exceeding 50% but much lower strength than the first two refractory HEAs.[1, 10, 13, 14]

To improve the strength of refractory HEAs based on HfNbTaTiZr, J.W. Yeh et al. subsequently investigated the mechanical behavior ofAccepted HfNbTaTiZr alloys withManuscript different grain sizes.[15] Not The resultCopyedited showed that the alloy with the smallest grain size of 38 μm possessed the highest yield strength of 958 MPa, which was just slightly higher than the initial value of 929 MPa. Modifying the composition of alloying elements, such as minor element addition and element substitution, is a conventional and effective way to alter material properties. Senkov et al. documented that the partial substitution of Al for Hf in HfNbTaTiZr reduced the density of the alloy by 9% and increased the RT hardness by 29% and the enhanced yield strength by 98%.[16] Chun-Ming Lin et al. investigated the effect of Al addition on the mechanical properties of HfNbTaTiZr and found that the yield strength of

2 equiatomic HfNbTaTiZrAl alloy was enhanced to about 1500 MPa under compression.[1] Inspired by these, L.

Wang et al. reported a new refractory HEA, HfMoNbTiZr, in which all Mo replaced by Ta in comparison to

HfNbTaTiZr.[4] This new alloy showed a significantly enhanced yield strength up to 1575 MPa as well as excellent ductility.[4] Besides, in a recent report, E ́. Fazakas et al. studied the effect of Cr on the mechanical properties of Ti20Zr20Hf20Nb20Cr20 refractory HEA, and attributed the enhanced strength and hardness to the segregation of Cr-containing laves phases (Cr2Hf, Cr2Nb) during casting.[17]

Accordingly, it suggested that Cr tends to form laves phase with Hf, Nb or Ta, which promotes the strength, and Cr is also more cost-efficient than other component elements (Hf, Nb, Ta, Ti, and Zr). Thus, in the present work, a new CrHfNbTaTi refractory HEA was designed by first replacing Zr with Cr in HfNbTaTiZr. Besides, the modulus of Mo was more than three times over Nb, even though they have a similar densities,[18] due to which the substitution between them is expected to greatly increase the mechanical properties. Hence, stemming from the alloy CrHfNbTaTi, another new CrHfMoTaTi was derived as well. The phases, microstructures, and mechanical properties of the new CrHfNbTaTi and CrHfMoTaTi refractory HEAs, at RT were investigated and discussed.

2. Materials and methods

CrHfNbTaTi and CrHfMoTaTi alloys were prepared using vacuum arc melting of equimolar mixtures in a

Ti-getter high-purity argon atmosphere. The elements of Cr, Hf, Nb, Mo, Ta, and Ti all were in the bulk form slugs with a purity of 99.99%, 99.9%, 99.95%, 99.95%, 99.95%, and 99.9% (at.%), respectively. During melting, ingots were melted for 3 min in a water-cooled copper crucible and turned over once after each melting process with electromagnetic stirring. This was repeated at least four times to ensure that the alloys were in a well-mixed state. The dimension of the solidified button sample was about 8 mm in thickness and 15 mm in diameter. The weight loss of each alloy during arc melting was lower than 0.2 percent, as there were no volatile elements in the alloys. The 10 g ingot piecesAccepted were melted by arc meltingManuscript and then sucked into Not an RT cylinder-shaped Copyedited copper mold with a diameter of 4 mm and a length of 60 mm. Table 1 shows the as-prepared alloy compositions of CrHfNbTaTi and CrHfMoTaTi cylindrical specimens, whose actual compositions were very close to the designed compositions.

The crystal structure of the as-cast cylindrical samples was identified via X-ray diffractometry (XRD) using a PANalytical X'Pert Powder diffractometer with CuKα radiation operating at 40 kV/40 mA and a scanning rate of 3 °/min with a 2휃 range of 20 to 120°. The microstructure was analyzed using a field emission scanning electron microscope (FESEM, Zeiss sigma 500). The compressive tests were carried out on a computer-controlled Instron

3 (Instron, Norwood, MA) mechanical testing machine using Φ3.7× 5.6 cylindrical samples at a constant ramp speed of 5.6× 10-3 mm/s and an initial strain rate of 10-3 s−1.

Table 1. Nominal and actual compositions (at.%) of the studied alloys in this work

Composition Cr Hf Nb Mo Ta Ti

Nominal 20 20 20 - 20 20 CrHfNbTaTi Actual 19.4 19.9 21.7 - 18.6 20.3

Nominal 20 20 - 20 20 20 CrHfMoTaTi Actual 20.5 21.6 - 18.2 17.7 22

3. Results and discussion

3.1 Phase analysis

Figure 1. XRD patterns of CrHfNbTaTi (a) and CrHfMoTaTi (b) alloys in the as-cast condition together with that of the as-cast HfNbTaTiZr alloy (c) reported in Ref. [7].

Figure 1 shows the XRD patterns of the as-cast cylindrical CrHfNbTaTi and CrHfMoTaTi samples in this work (a, b) and as-solidifiedAccepted HfNbTaTiZr alloy Manuscriptextracted from Ref. [19] Not(c). The referenceCopyedited HfNbTaTiZr alloy consists of a single BCC phase (Figure 1(c)). In comparison, the alloy CrHfNbTaTi possesses more complicated phases even though only one principal element, Zr, was replaced by Cr. This implies that Cr plays a vital role in determining the phases of this type of alloys. More specifically, the as-cast CrHfNbTaTi alloy contained three phases, two BCC (BCC1 and BCC2) and one cubic laves phase (Figure 1(a)). The much weaker peaks of BCC2 phase than the BCC1 phase implies that BCC1 is the dominant constituent of the BCC phases. Further, if all peaks are considered, it can be concluded that the BCC phases are the major phases in the CrHfNbTaTi refractory HEA.

4 In addition, according to the JCPDS cards, the BCC2 phase was enriched with Ti and Cr, and depleted by Hf, Nb, and Ta, which matches well with TiCr. The laves phase is consistent with a Cr-rich laves phase (Cr2(Hf, Ti, Ta,

Nb)), as well as that in CrHfNbTiZr alloy reported in Ref. [17]. The formation of this type of laves phase has been documented by A. Inoue et. al.[20].

Hereby, the atomic size ratio of the largest and smallest elements in a composition larger than 1.225 was considered beneficial to laves phase formation. Thus, the presence of the Cr-rich laves phase in CrHfNbTaTi alloy is spontaneously anticipated due to the large atomic size difference between Hf (157.8 pm), Nb (142.9 pm) and

Cr (124.9 pm). Compared with the XRD patterns of the HfNbTaTiZr alloys, the diffraction peaks of the BCC1 phase in the as-cast CrHfNbTaTi alloy shifted towards higher angles, indicating the lattice parameter of the primary BCC phase was reduced compared with that of BCC phases in HfNbTaTiZr. According to the XRD patterns, the lattice parameters of the BCC phase in the CrHfNbTaTi and HfNbTaTiZr were determined to be a =

334.1 pm and a = 340.4 pm, respectively. This change in the lattice parameter might stem from the small atomic size of Cr (124.9 pm) relative to the substituted element Zr (160.3 pm).[17] Moreover, the lattice parameters of

BCC2 and laves phases of the studied CrHfNbTaTi alloy were determined to be 311.6 pm and 720.1 pm, respectively.

Table 2. The crystal structures, atomic radius (r), and Young’s modulus (E), of the elements in the studied alloys.

Element Cr Hf Nb Mo Ta Ti

Crystal BCC HCP BCC BCC BCC HCP structure

r (pm) 124.91 157.75 142.9 136.2 143 146.15

E (GPa) 279 78 105 329 186 116

To unveil the effect of the introduction of an element with significantly different intrinsic features on the phase component, in this work,Accepted Mo was used to replace Manuscript Nb to change CrHfNbTaTi Not to CrHfMoTaTi.Copyedited Both elements (Nb and Mo) possessed a BCC structure at RT, a relatively low density, and similar melting temperature, but the young’s modulus of Mo is almost three times over than that of Nb. The related element properties of the constituent elements in the studied alloys are listed in Table. 2. After mutual substitution, despite the presence of a few unknown peaks in the XRD patterns, the number of phases dramatically increased from three to six ( Figure

1(b)). These six phases were determined to be four BCC phases (BCC1, BCC2, BCC3, and BCC4) and two laves phases. According to the JCPDS cards, the BCC1, BCC2, BCC3, and BCC4 phases matched well with TaMo,

5 Hf9Mo, MoTi/TiCr, and MoCr, respectively. This suggests that the BCC1, BCC3 and BCC4 phases were mainly enriched with Mo, while the BCC2 phase was enriched with Hf. Their lattice parameters in the former order were determined as a =325.3 pm, 339.9 pm, 316.6 pm, and 300.3 pm, respectively. Their lattice parameters were determined by the atomic size of the predominant and secondary elements (Hf: 157.8 pm, Mo: 136.2 pm, Ta:

143.0 pm, Ti: 146.2 pm, and Cr: 124.9 pm). Besides, as a Cr-containing alloy, a Cr-rich C15 laves phase, Cr2(Hf,

Ta), was also identified with a lattice parameter of a = 704.3 pm, which is as well as the structure of the C15 laves phase in the CrHfNbTaTi but different in the compositional constituents (apparently lacking Ti and Nb). This is in good agreement with the former BCC phase analysis, in which the incorporation of many Ti and Nb atoms formed BCC phases; however, another laves phase, was also interestingly matched well with the C15 laves phase,

Mo2Hf, which was unrelated to Cr. Its lattice parameter was up to 746.0 pm. In general, the substitution of Mo for Nb facilitated the formation of Mo-containing BCC phases by affecting the formation of the Cr-containing laves phase and by directly participating in the formation of the BCC phases.

3.2 Microstructure

Figure 2. SEM micrographsAccepted of the microstructure Manuscript of CrHfNbTaTi (a) Not and CrHfMoTaTi Copyedited (c) in the as-cast condition. (b) and (d) are the magnification of the local region of (a) and (c), respectively.

The microstructure of the as-cast CrHfNbTaTi and CrHfMoTaTi alloys are shown in Figure 2. A typical dendrite structure with an average primary arm size of around 2 μm was observed in the as-cast CrHfNbTaTi alloy, as shown in Figure 2(a, b). The dendritic regions (bright regions marked as D) were irregularly surrounded by a continuous and corridor interdendritic regions (grey part marked as ID). The volume fraction of the dendritic regions was larger than the remaining interdendritic regions, which implies that the dendritic regions should

6 correspond to the BCC phases. In contrast, the interdendritic regions were likely Cr-rich laves phases according to the phase analysis, which showed that the BCC phases were the majority phase in CrHfNbTaTi. The chemical compositions of both regions determined from EDS analysis are summarized in Table 3. It shows that the interdendritic regions were mainly enriched in Cr and Hf, further confirming that the interdendritic regions were composed of Cr-rich laves phases. Additionally, Cr had the lowest melting temperature, which is the reason for the slim and continuous morphology of the interdendritic regions. For the dendritic regions, the main enriched elements were Nb, Ta and Ti. According to the JCPDS cards, the primary XRD peaks for both BCC phases approximately correspond to compounds mainly composed of Nb and Ta, consistent with the elemental distribution of the dendritic regions. This also implies that the dendritic regions corresponded to the BCC phases.

Table 3. Quantitative chemical analysis of as-cast CrHfNbTaTi and CrHfMoTaTi samples (at.%).

Alloy Regions Phases Cr Hf Nb Mo Ta Ti

D BCC1, 2 5.2±0.6 20.8±0.3 22.2±0.2 - 25.3±1.1 26.6±0.3 CrHfNb -TaTi ID Laves1 44.9±5.4 21.4±0.5 6.7±1.6 - 15.8±1.7 11.2±2.6

D BCC1, 3, 4 11.2±0.4 14.4±0.3 - 35.8±0.2 17±0.4 21.6±0.1 CrHfMo ID1 BCC2,4+Laves1,2 30.2±0.3 27.8±0.1 - 19±0.6 5.9±0.2 17.1±0.9 -TaTi ID2 BCC3 17.6±1.9 24.8±0.5 - 13.3±0.4 5.3±0 38.9±2.9

D is short for dendrite, and ID for interdendrite.

After replacing Nb with Mo in CrHfNbTaTi, expect the typical two distinct regions observed as well as that in the CrHfNbTaTi alloy, the difference between them is recognizable (see Figure 2(c, d)). In detail, for the as- cast CrHfNbTaTi alloy, theAccepted microstructure showed Manuscript one typical fishbone-like Not dendritic Copyedited grain (brighter part) with blunt edges tightly embedded into a continuous matrix (grey part) in Figure 2(a, b). In contrast, the dendritic regions of CrHfMoTaTi marked as D appeared more similar to equiaxed grains, and the interdendritic regions contained two parts: a bright part marked as ID1 and a grey part marked as ID2. Moreover, the dendritic regions that surrounded the interdendritic shell regions are often incompact, and the shell thickness was around 1 μm.

According to the EDS analysis, all regions' chemical compositions were determined and are summarized in Table

3. The interdendritic regions (ID1 and ID2) were rich in Cr, Hf, and Ti, while the dendritic regions (D) were rich

7 in Mo and Ta. Apparently, as with CrHfNbTaTi, the interdendritic regions were also ascribed to a laves phase, while the dendritic regions was BCC phases.

Figure 3. SEM images of as-cast CrHfNbTaTi (a) and CrHfMoTaTi (b) high entropy alloys, EDS mapping images of the corresponding elements of Cr, Hf, Nb, Ta and Ti in (a) and Cr, Hf, Mo, Ta and Ti in (b).

To figure out the overall distribution of each element, EDS mapping was also carried out (Figure 3).

According to mapping, 1~2 μm thick transition layers (marked with white arrows in the Ti map in Figure 3(a)) enriched with Ti were observed at the interface of the bright dendritic regions in CrHfNbTaTi. This was also observed with the BCC crystal structures in the hot isostatically pressed NbCrMo0.5Ta0.5TiZr alloy.[21] Similarly, part of the ID2 regions with 2~3 μm in size (marked with red arrows in the bottom of the Ti map in Figure 3b) rich in Ti were observed in CrHfMoTaTi. Combined with the XRD analysis, they correspond to the BCC2 TiCr phase of CrHfNbTaTi and the BCC3 TiCr phase of CrHfMoTaTi, respectively. The overall experimental compositions were very similar to the nominal compositions, as shown in Table. 1, which further confirmed the reliable melting of this composition.

3.3 Mechanical properties

The engineering stress, 휎 , vs. engineering strain, ε curves for the as-cast CrHfNbTaTi and CrHfMoTaTi alloys obtained during compression testing at RT, together with those of solidified HfNbTaTiZr alloy reported in

Ref. [19] are shown in Figure 4(a). Note that the compressive curve of the alloy CrHfMoTaTi was offset along the axis to display clearly. The compressive properties such as yield strength 휎0.2 , ultimate compressive strength

휎b, and fracture strain 휀f ofAccepted these alloys are given Manuscriptin Table. 4. For CrHfNbTaTi, Not the yield Copyedited strength, 휎0.2 , was 1258

± 15 MPa. After yielding, the as-solidified HfNbTaTiZr and CrHfNbTaTi alloys showed obvious work- hardening behavior, which is commonly to be evaluated by the instantaneous work hardening index, n*. This

d(ln휎T) parameter is defined as n*= , based on the common power-law relationship 휎 = k 휀 n, where 휎 and 휀 are d(ln휀T) T T T T the true stress and true strain [22].

8 Figure 4. Compressive engineering stress-strain curves of the as-cast CrHfNbTaTi, CrHfMoTaTi and the reported

d(ln휎T) HfNbTaTiZr[19] alloys (a) and instantaneous work hardening index, n*,= vs. true strain curves of d(ln휀T)

CrHfNbTaTi and HfNbTaTiZr alloys (b). To display clearly, the compressive curve of the alloy CrHfMoTaTi was offset along the axis.

Table 4. Compressive yield strength 휎0.2 , ultimate compressive strength 휎b , fracture strain 휀f of CrHfNbTaTi,

CrHfMoTaTi and HfNbTaTiZr.[19]

Alloy 휎0.2 (MPa) 휎b (MPa) 휀f (%)

CrHfNbTaTi 1258 2061 24.3

CrHfMoTaTi 1051 1051 7.8

HfNbTaTiZr 929 - >50

The instantaneous work hardening as a function of true strain was calculated for the as-cast samples of

HfNbTaTiZr and CrHfNbTaTi using compression test data. The results presented in Figure 4(b) show that the instantaneous work hardening of HfNbTaTiZr and CrHfNbTaTi both exhibited two stages. First, the instantaneous work hardening sharply decreasedAccepted from true strain Manuscript 휀T =0.7% - 3.8% and 휀NotT =2.6% -Copyedited 13.8% (stage I). This was followed by a gradual increase of the HfNbTaTiZr alloy and a stable stage from true strain 휀T =14.7%-21.1% of

CrHfNbTaTi (Stage II). The classical decrease in the instantaneous work hardening index in Stage I was primarily attributed to dislocation gliding during the initial deformation stage. In Stage II, the instantaneous work hardening index tended to remain stable at about n*= 0.47 for CrHfNbTaTi. The contribution of dynamic strain aging on work hardening involved the suppression of microvoids/microcracks at this stage.[23] Thus, its strength eventually reached a maximum value of 2061 MPa before fracture. However, by replacing Nb with Mo,

9 CrHfMoTaTi showed typical brittle fracture at 1051 MPa with no work hardening behavior during compression tests.

Compared with the previously reported HfNbTaTiZr alloy, the yield strength of CrHfNbTaTi alloys dramatically increases from 926 MPa to 1258 MPa, while their promising plasticity was retained. Moreover, if the alloys such as CrHfNbTiZr,[17] CrNbTiVZr,[24] HfMoNbZrTi,[4] HfMoTaTiZr[25] and MoNbTaVW,[9] were taken into account, this alloy possessed a better strength-ductility trade-off. There is a close relationship between the microstructure and mechanical properties of an alloy. According to the above phase analysis, by substituting Zr with Cr, the alloy changed from a single BCC phase to multiple phases, including both BCC and

Cr-containing laves phases. Since the BCC phase is the major/only phase in CrHfNbTaTi/HfNbTaTiZr alloys, dislocation movements are sensitive to the lattice plane distance d of the BCC phase because a wider lattice distance d is beneficial to the dislocation movement and subsequent formation of stress fields from dislocations.[26] When there are more stress fields, there is more resistance to dislocation movement and better solid-solution strengthening behavior. The largest lattice plane distance in a BCC structure is {110}, which can

푎 be calculated from the actual lattice parameters a by d= . However, the d of CrHfNbTaTi (which is ℎ2 + 푘2 + 푙2 110

236.2 pm) is smaller than that of HfNbTaTiZr (which is 240.7 pm) due to the substitution of Zr with a larger atomic radii and Cr with a smaller one. This indicates that the other Cr-containing laves phases contributed to the improved yield strength of CrHfNbTaTi. Moreover, Cr-containing laves phases are inherently brittle and strong,

[18] and they filled in the interdendritic regions and formed a network, due to which a great number of dislocations that play a vital role in deformation might pile-up at the grain boundaries, which significantly enhanced the strength. Additionally, the interdendritic regions containing laves phases enriched with Cr also contained other elements. This implies that the phases in the interdendritic regions were relatively more complex, especially the thin transition layers that coated the dendritic phase, which may play an important role in the combination of strength and plasticity. This kind of complexity might decrease the stiffness of the framework, resulting in outstanding plasticity. For CrHfMoTaTi,Accepted the deterioration Manuscript of the mechanical Not strength Copyeditedafter replacing Nb with Mo in CrHfNbTaTi might be closely related to the incompact coupling between the dendritic and interdendritic regions, which facilitated the formation of poles. The poles are generally considered to be one kind of crack initiators, playing an important role in determining the brittle failure behavior at the initial deformation stage, especially in the brittle and strong laves phase.

4. Conclusions

10 In summary, CrHfNbTaTi was mainly composed of two BCC and one Cr-containing laves phases, while

CrHfMoTaTi consisted of six phases: four BCC and two laves phases. The laves phases in both alloys possessed a C15 structures. In addition, compared with the reference alloy HfNbTaTiZr, replacing Zr with Cr increased the yield strength from 926 MPa to 1258 MPa, while retaining an elongation of around 24.3%. The enhanced strength might stem from the constraining effect of the strong laves phase framework. Furthermore, the component complexity of the framework decreased the stiffness, which was responsible for the promising ductility. Besides, the alloy CrHfMoTaTi showed a typical brittle fracture with a strength of 1051 MPa, which might stem from the presence of multiple phases and because the dendritic regions surrounded by an incompact interdendritic shell after the incorporation of Mo.

Declaration of Interest Statement

We declared that the present work has no conflicts with other researchers or organizations.

References

[1] C.M. Lin, C.C. Juan, C.H. Chang, C.W. Tsai, and J.W. Yeh, Effect of Al addition on mechanical properties and microstructure of refractory AlxHfNbTaTiZr alloys, J. Alloys Compd., 624(2015), p. 100.

[2] T.M. Pollock and S. Tin, Nickel-based superalloys for advanced turbine engines: chemistry, microstructure and properties, J. Propul. Power., 22 (2006), No. 2, p. 361.

[3] E. Karaköse and M. Keskin, Microstructure evolution and mechanical properties of intermetallic Ni–xSi (x=5,

10, 15, 20) alloys, J. Alloys Compd., 528(2012), p. 63.

[4] N.N. Guo, L. Wang, L.S. Luo, X.Z. Li, Y.Q. Su, J.J. Guo, and H.Z. Fu, Microstructure and mechanical properties of refractory MoNbHfZrTi high-entropy alloy, Mater. Des., 81(2015), p. 87.

[5] D.B. Miracle, High entropy alloys as a bold step forward in alloy development, Nat. Commun., 10(2019), No.

1, p. 1805.

[6] 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, Nanostructured high-entropy alloys with multipleAccepted principal elements: Manuscript novel alloy design concepts Not and outcomes,Copyedited Adv. Eng. Mater., 6 (2004), p. 299.

[7] B. Cantor, I.T.H. Chang, P. Knight, and A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A., 375-377(2004), p. 213.

[8] O.N. Senkov, G.B. Wilks, D.B. Miracle, C.P. Chuang, and P.K. Liaw, Refractory high-entropy alloys.

Intermetallics, 18 (2010), No. 9, p. 1758.

11 [9] O.N. Senkov, G.B. Wilks, J.M. Scott, and D.B. Miracle, Mechanical properties of Nb25Mo25Ta25W25 and

V20Nb20Mo20Ta20W20 refractory high entropy alloys, Intermetallics, 19(2011), No. 5, p. 698.

[10] R.R. Eleti, T. Bhattacharjee, A. Shibata, and N. Tsuji, Unique deformation behavior and microstructure evolution in high temperature processing of HfNbTaTiZr refractory high entropy alloy, Acta. Mater., 171(2019), p. 132.

[11] K.C. Lo, Y.J. Chang, H. Murakami, J.W. Yeh, and A.C. Yeh, An oxidation resistant refractory high entropy alloy protected by CrTaO4-based oxide, Sci. Rep., 9 (2019), p. 7266.

[12] V. Soni, O.N. Senkov, B. Gwalani, D.B. Miracle, and R. Banerjee, Microstructural design for improving ductility of an initially brittle refractory high entropy alloy, Sci. Rep., 8 (2018), No. 1, p. 8816.

[13] O.N. Senkov, A.L. Pilchak, and S.L. Semiatin, Effect of cold deformation and annealing on the microstructure and tensile properties of a HfNbTaTiZr refractory high entropy alloy, Metall. Mater. Trans. A.,

49(2018), No. 7, p. 2876.

[14] M. Wang, Z. Ma, Z. Xu, and X. Cheng, Microstructures and mechanical properties of HfNbTaTiZrW and

HfNbTaTiZrMoW refractory high-entropy alloys, J. Alloys Compd., 803(2019), p. 778.

[15] S. Chen, K.K. Tseng, Y. Tong, W. Li, C.W. Tsai, J.W. Yeh, and P.K. Liaw, Grain growth and Hall-Petch relationship in a refractory HfNbTaZrTi high-entropy alloy, J. Alloys Compd., 795(2019), p. 19.

[16] O.N. Senkov, S.V. Senkova, and C. Woodward, Effect of aluminum on the microstructure and properties of two refractory high-entropy alloys, Acta. Mater., 68 (2014), p. 214.

[17] É. Fazakas, V. Zadorozhnyy, L.K. Varga, A. Inoue, D.V. Louzguine-Luzgin, F. Tian, and L. Vitos,

Experimental and theoretical study of Ti20Zr20Hf20Nb20X20 (X=V or Cr) refractory high-entropy alloys, INT. J.

REFRACT. MET. H. 47(2014), p. 131.

[18] D.B. Miracle and O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta. Mater.,

122 (2017), p. 448. [19] O.N. Senkov, J.M. Scott,Accepted S.V. Senkova, D.B.Manuscript Miracle, and C.F. Woodward,Not Copyedited Microstructure and room temperature properties of a high-entropy TaNbHfZrTi alloy, J. Alloys Compd., 509 (2011), No. 20, p. 6043.

[20] A. Inoue, Stabilization of metallic supercooled liquid and bulk amorphous alloys, Acta. Mater., 48(2000), p.

279.

[21] O.N. Senkov and C.F. Woodward, Microstructure and properties of a refractory NbCrMo0.5Ta0.5TiZr alloy,

Mater. Sci. Eng. A., 529(2011), p. 311.

12 [22] J. Chiang, B. Lawrence, J.D. Boyd, and A.K. Pilkey, Effect of microstructure on retained austenite stability and work hardening of TRIP steels, Mater. Sci. Eng. A., 528 (2011), No. 13-14, p. 4516.

[23] S. Liu, Z. Xiong, H. Guo, C. Shang, and R.D.K. Misra, The significance of multi-step partitioning:

Processing-structure-property relationship in governing high strength-high ductility combination in medium- manganese steels, Acta. Mater., 124 (2017), p. 159.

[24] O.N. Senkov, S.V. Senkova, D.B. Miracle, and C. Woodward, Mechanical properties of low-density, refractory multi-principal element alloys of the Cr–Nb–Ti–V–Zr system, Mater. Sci. Eng. A., 565(2013), p. 51.

[25] C.C. Juan, M.H. Tsai, C.W. Tsai, C.M. Lin, W.R. Wang, C.C. Yang, S.K. Chen, S.J. Lin, and J.W. Yeh,

Enhanced mechanical properties of HfMoTaTiZr and HfMoNbTaTiZr refractory high-entropy alloys,

Intermetallics, 62(2015), p. 76.

[26] S. Chen, X. Yang, K. Dahmen, P. Liaw, and Y. Zhang, Microstructures and Crackling Noise of

AlxNbTiMoV High Entropy Alloys, Entropy 16(2) (2014), p. 870.

Accepted Manuscript Not Copyedited

13