Effects of Co and Ti on microstructure and mechanical behavior of Al0.75FeNiCrCo high entropy alloy prepared by mechanical alloying and spark plasma sintering

Author’s Accepted Manuscript Effects of CO and Ti on microstructure and mechanical behavior of Al 0.75FeNiCrCo High entropy alloy prepared by mechanical alloying and spark plasma sintering Zhen Chen, Weiping Chen, Bingyong Wu, Xueyang Cao, Lusheng Liu, Zhiqiang Fu www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(15)30307-5 http://dx.doi.org/10.1016/j.msea.2015.08.056 MSA32690

To appear in: Materials Science & Engineering A Received date: 1 July 2015 Revised date: 15 August 2015 Accepted date: 18 August 2015 Cite this article as: Zhen Chen, Weiping Chen, Bingyong Wu, Xueyang Cao, Lusheng Liu and Zhiqiang Fu, Effects of CO and Ti on microstructure and mechanical behavior of Al 0.75FeNiCrCo High entropy alloy prepared by mechanical alloying and spark plasma sintering, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.08.056 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effects of Co and Ti on microstructure and mechanical behavior of Al0.75FeNiCrCo high entropy alloy prepared by mechanical alloying and spark plasma sintering Zhen Chen, Weiping Chen*, Bingyong Wu, Xueyang Cao, Lusheng Liu, Zhiqiang Fu† School of Mechanical and Automotive Engineering, South China University of Technology, Guangzhou, Guangdong 510640, China

Abstract The effects of Co removal and Ti addition on the microstructure and mechanical behavior of a high-entropy alloy (HEA), Al0.75FeNiCrCo, were studied systematically. Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 were successfully prepared by the combination of mechanical alloying (MA) and spark plasma sintering (SPS). During the MA process, a primary body-centered cubic (BCC) supersaturated solid-solution phase and a face-centered cubic (FCC) supersaturated solid-solution phase were formed in each of the three investigated alloys. Removing Co from the Al0.75FeNiCrCo HEA or adding Ti into the Al0.75FeNiCrCo HEA could facilitate increasing the content of BCC phase both during the MA process and after the SPS process. Following SPS, bulk Al0.75FeNiCrCo was composed of a major FCC phase (~79 vol.%) and a minor BCC phase (~21 vol.%). However, with Co removal, bulk Al0.75FeNiCr alloy exhibited a main BCC phase (~60 vol.%) and an FCC phase (~40 vol.%). With Ti addition, bulk Al0.75FeNiCrCoTi0.25 alloy consisted of a major FCC phase (~77 vol.%) and a minor BCC phase (~23 vol.%). The EDS/TEM

*

Corresponding author. Tel.:+86-20-87113832; fax:+86-20-87112111; E-mail address: [email protected] (W. Chen) † Corresponding author. E-mail address:[email protected] (Z. Fu) 1

results of the bulk Al0.75FeNiCrCo alloy revealed that the BCC phase was enriched in Al-Ni, while the FCC phase was enriched in Fe-Cr-Co. Meanwhile, a small fraction of nanoscale twins were present in the bulk Al0.75FeNiCrCo alloy. The bulk Al0.75FeNiCrCo alloy exhibited high strength and high hardness, mainly attributed to solid-solution strengthening, twin-boundary strengthening and grain-boundary strengthening. With Co removal, the bulk Al0.75FeNiCr alloy showed lower strength, hardness and ductility in comparison with the Al0.75FeNiCrCo alloy. With Ti addition, compared to the Al0.75FeNiCrCo alloy, the bulk Al0.75FeNiCrCoTi0.25 alloy exhibited higher strength and hardness with a slightly lower ductility. Keywords: high entropy alloys; mechanical alloying; spark plasma sintering; Cobalt; Titanium; microstructure. 1. Introduction Traditional alloys are mostly composed of one principal element with some minor elements added into the alloy systems, in order to tailor the microstructures and properties, for example, Fe-, Al- and Mg-based alloys [1-3]. As the number of alloying elements increases, intermetallic compounds (IMC) and/or other complicated compounds exhibiting complex microstructures that would further degrade mechanical properties can readily be formed in the traditional alloy systems. Recently, a novel class of alloys that defined as high entropy alloys (HEAs) by Yeh et al. have broken the traditional alloy design concept since 2004 [4]. Normally, these new series of alloys consist of at least four principal elements, and atomic concentration for each constituent element is between 5 and 35 at.% [4,5]. Due to the multi-principal-element composition, HEAs have more pronounced entropic effects than conventional alloys, i.e. “high-entropy effect” [4,5]. Therefore, most of HEAs reported in the previous literature display solid-solution phases with FCC and/or 2

BCC structures attributed to the “high-entropy effect” [4,6-12]. Inspection of the published literature reveals that HEAs possess high hardness and high strength [13,14], excellent wear resistance [15], high thermal stability [16], distinctive electrical and magnetic properties [17] and outstanding oxidation resistance [18], etc., and accordingly they are suitable for a wide range of engineering applications. To date, a variety of experimental techniques, such as arc-melting [19,20], fluxed water quenching [21], melt spinning [22], injection casting [23], laser cladding [24,25], mechanical alloying (MA) [26,27], etc., have been explored to synthesize HEAs. Inspection of the published literature reveals that the most widely used method is arc-melting (casting), and composition segregation and inhomogeneous microstructures are usually observed in as-cast HEAs [8,28]. Hence, a combination of MA that can produce nanocrystalline materials with good homogeneity from elemental powders and SPS which can consolidate powders to high density rapidly by applying pressure and passing electric pulse current, has been developed to prepare bulk HEAs [29-31]. In AlFeNiCrCox (expressed in molar ratio) HEA system, increasing concentration of Co element can lead to an increment in ductility and a decrement in strength [32]. On the other hand, target phases and mechanical properties of Al-containing and/or Ti-containing HEAs are obtainable depending on the concentrations of Al and Ti elements heavily, for instance, AlxFeNiCrCo, AlxFeNiCrCoTi and AlFeNiCrCoTix HEA systems [31,33-36]. In the well-studied AlxFeNiCrCo HEA system, Al0.75FeNiCrCo HEA possessing a mixture of FCC+BCC solid-solution phases has a potential balance of ductility and strength, which may result in excellent mechanical properties, and this alloy produced by MA and SPS has heretofore never been studied [19,36]. Taking the above discussion into account, in the present work, Al0.75FeNiCrCo HEA was 3

fabricated by MA and SPS, and consequently alloying behavior, microstructure and mechanical behavior of this alloy were studied in detail. Moreover, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 alloys were designed and synthesized by MA and SPS, with the aim of investigation of the influence of Co removal and Ti addition on the microstructure and mechanical properties of the Al0.75FeNiCrCo HEA. 2. Experimental details Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 mechanically alloyed powders were prepared by 45 h dry milling without process-control-agent (PCA) followed by 5 h wet milling using ethanol as PCA. The starting materials were Al, Fe, Ni, Cr, Co and Ti elemental powders with high purity (>99.7 wt%) and particle size of ~45mm (-325 mesh). The MA process was carried out using stainless-steel vials and under an argon (Ar) atmosphere, and then it was performed on a high energy planetary ball milling machine (QM-3SP4 Planetary Ball Mill) operated at 300 rpm. Tungsten carbide balls were used as milling medium and the ball-to-powder weight ratio was 10:1. To study alloying behavior, powder samples were taken out following milling for 0, 5, 15, 30, 45, 50 h, respectively. The MA process was finished after 50 h of milling, and the powders were allowed to dry in a vacuum oven. Then a 75 μm sieve was used for filtering the ultimate powders. Following drying and filtering, the as-milled powders were consolidated by Dr. Sinter Model SPS-825 Spark Plasma Sintering System (Sumitomo Coal Mining Co. Ltd., Japan) in a graphite die with 20.4 mm inner diameter under a constant pressure of 30 MPa. In order to remove the sample after sintering easily, the punches and the powders were separated by graphite foils of 0.2 mm in thickness. The samples were sintered at 1273 K for 8min, and the heating rates were 80 K min–1 from room temperature to 373 K and 100K min–1 from 373 K to 1273K, respectively. 4

Powders and bulk alloys after SPS were analyzed by a Bruker D8 ADVANCE X-ray diffractometer (XRD) with a Cu Kα radiation. Scanning electron microscopy (SEM, NOVA NANOSEM 430, USA) with an energy dispersive spectrometer (EDS) was used for observing microstructures of bulk alloys which were polished and etched with aqua regia solution ( HCl and HNO3 in a volume ratio of 3:1) approximately 10 s. Thin foil specimens were prepared by mechanical thinning followed by ion milling at room temperature, and subsequently were analyzed by a transmission electron microscopy (TECNAI G2 S-TWIN F20, FEI, USA) with selected area electron diffraction (SAED) analysis and an energy dispersive spectrometer. According to ASTM E9-09 [37], the room-temperature compressive properties of the cylindrical samples (ϕ3mm×4.5 mm in size) were measured using an Instron 5500 testing system at a strain rate of 1×10-3 s-1. Hardness measurement was conducted on a Digital Micro Hardness Tester HVS-1000 Vickers hardness instrument under a load of 300 gf (~2942 mN). The reported hardness value for each bulk alloy is an average of 10 measurements. 3. Results and discussion 3.1 Influence on Alloying behavior Fig.1 shows the XRD patterns of Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 HEA powders under different milling durations. Fig.1(a) shows the XRD patterns of the mechanically alloyed powders of Al0.75FeNiCrCo for different times. All of the pure elements can be predominantly observed in the initial blending powders (0 h). As the milling time increases to 5 h, drastic decrement of diffraction intensity is observed simultaneously with disappearance of peaks corresponding to Al. As the milling time is increased to 15 h, the peaks of Co disappear and the rest of peaks are similar to the peaks of the 5 h milled powders except for a minor decrement in 5

diffraction intensity. After milling for 30 h, the XRD pattern shows peaks corresponding to a primary BCC solid-solution phase and an FCC solid-solution phase without any peaks associated with pure elements, suggesting the complete formation of solid-solution phases. When the milling time reaches up to 45 h and 50 h, the XRD patterns are quite similar to that of 30 h except peak broadenings. Note that the 45 h milled powders were subjected to wet ball milling (in ethanol) for 5 h with the aim to further reduce the grain and particle sizes for being conducive to attaining dense bulk samples. Fig.1(b) shows the XRD patterns of the mechanically alloyed powders of Al0.75FeNiCr for different times. As the milling time is increased to 5 h, Al still exists, however, has significant decrement in diffraction intensity. As the milling time extends to 15 h, the diffraction peaks of Al disappear completely. Unlike the Al0.75FeNiCrCo HEA, following 30 h of milling, the diffraction peaks of Fe cannot be detected while the Cr is still faintly existent, suggesting that the alloying process has not been finished. After 45 h of milling, a major BCC solid-solution phase with a minor FCC solid-solution phase are observed, in the mean time, there are no diffraction peaks corresponding to pure elements. Similarly, the XRD patterns of 50 h milled powders were similar to that of 45 h milled powders except peak broadening. The XRD diffractograms of Al0.75FeNiCrCoTi0.25 HEA powders under different milling durations are shown in Fig. 1(c). After 15 h of milling, the additive element of Ti still can be observed in the XRD diffraction patterns. Except diffraction peaks of Ti, the alloying behavior of Al0.75FeNiCrCoTi0.25 HEA is quite similar to that of Al0.75FeNiCrCo HEA. It is worth pointing out that when the solid-solution phases are completely formed

during MA,

the contents

of FCC

phases

in

Al0.75FeNiCr and

Al0.75FeNiCrCoTi0.25 are slightly lower in comparison with that of Al0.75FeNiCrCo. 6

As described above, Al dissolves most rapidly while Ti dissolves most slowly. In conclusion, the order of decreasing alloying rate for the pure elements in the investigate alloys are: Al→Co→Ni → Fe → Ti → Cr. Thus, compared to the Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 alloys, the formation of completed solid-solution phases in the Al0.75FeNiCr alloys with a higher concentration of Cr takes a longer time. Alloying rate is dominated by melting point and by ductility when pure elements have similar melting points [38]. It can be explained as the following two factors: (a) An element with a lower melting point could show a higher intrinsic diffusion coefficient compared to that with higher melting point, (b) Co and Ni have very similar coefficients and their alloying rates are dominated by ductility, therefore the brittle Co with HCP structure dissolves a bit more rapidly than the ductile Ni with FCC structure [38]. Apparently, removing Co from the Al0.75FeNiCrCo HEA or adding Ti into the Al0.75FeNiCrCo HEA are both in favor of increasing the relative volume fraction ratio between BCC phase and FCC phase after MA, i.e., increasing the content of BCC phase and decreasing the content of FCC phase. During MA, Al, Fe, Cr and Ti tend to form BCC structured phases, whereas Co and Ni favors forming FCC structured phases [14,38-42]. It is obvious that with Co removal and with Ti addition would lead to concentrations of elements which are in favor of forming BCC structured phases. Hence, both removing Co and adding Ti result in increasing the relative volume fraction ratio between BCC phase and FCC phase. As the milling time increases, the occurrence of peak broadenings in the XRD patterns of these three investigated alloys are caused by the reduction in crystallite size and the increment in lattice strain [43-46]. After deducting the instrumental contribution, the crystal size and lattice strain of the BCC and FCC solid-solution phases for the investigated alloys with different milling times were calculated based on the XRD patterns using Viogt peak profile analyses. The calculated results are listed in Table 1. 7

It can be seen that crystallite size decreases while lattice strain increases as extending the milling time, confirming peak broadenings are attributed to the crystalline refinement and the increment in lattice strain [43-46]. 3.2 Influence on microstructure, phase composition and evolution After densification by SPS at 1273K, the relative density of bulk Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 are 98.85%, 97.28% and 98.90%, respectively. Fig.2 shows the XRD patterns of the bulk Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 alloys after densification by SPS at 1273 K. It indicates that the bulk Al0.75FeNiCrCo HEA consists of a primary FCC phase (~79 vol.%) with a lattice parameter of ~0.3592 nm and a BCC phase (21 vol.%) with a lattice parameter of ~0.2880 nm. After removing Co from the Al0.75FeNiCrCo HEA, however, the bulk Al0.75FeNiCr alloy displays a main BCC phase (~60 vol.%) with a lattice parameter of ~0.2867 nm and an FCC phase (~40 vol.%) with a lattice parameter of ~0.3598 nm. After adding Ti, the bulk Al0.75FeNiCrCoTi0.25 alloy contains a primary FCC phase (~77 vol.%) with a lattice parameter of ~0.3603 nm and a BCC phase (~23 vol.%) a lattice parameter of ~0.2870 nm. As described in section 3.1, the main phases of these alloys after MA are BCC phases with small volume fractions of FCC phases, suggesting that phase transformations took place during SPS. As we known, the MA process is a non-equilibrium route, which can facilitate the formation of metastable supersaturated solid solutions [14,39,47]. In the mean time, the enthalpy stored in nano-sized grain boundaries with significant volume fraction could be regarded as excess energy, which would lower down the driving force for phase transformation, thus phase transformations occurred readily during the SPS [47,48]. At a high temperature of 1273 K, reordering can be resulted from the annihilation of defects that introduced into metastable phases that produced during 8

the MA process. As a result, metastable phases produced after MA evolved into stable equilibrium phases after SPS. In conclusion, phase transformations occurred in these three as-milled powders after densification by SPS at 1273 K. Compared to the bulk Al0.75FeNiCrCo HEA, removing Co from the Al0.75FeNiCrCo HEA shows a significantly higher volume fraction of BCC phase, whereas adding Ti into the Al0.75FeNiCrCo HEA leads to a slightly higher volume fraction of BCC phase. Normally, Ti favors formation of BCC phase in HEAs, in contrast, Co is in favor of forming FCC phase [49]. Therefore removing ~21.05 at.% of Co element causes the significant decrement of FCC phase, that is, the significant increment of BCC phase. In addition, adding ~5 at.% of Ti accounts for the slightly increment of BCC phase. Fig. 3 shows low-magnified SEM micrographs with high-magnified images in the insets of the bulk Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 alloys consolidated by SPS at 1273 K. Fig. 3(a) shows the SEM image of the bulk Al0.75FeNiCrCo alloy, indicating two kinds of regions, i.e., featureless regions (region 1) and etched regions (region 2) that have been etched by aqua regia solution. EDS/SEM results shown in Table 2 demonstrate that region 1 and region 2 are Fe-Cr-Co-rich phase and Al-Ni-rich phase, respectively. This suggests that featureless regions and etched regions correspond to the main FCC phase and the BCC phase, respectively, since Al-Ni-rich phase have BCC structure whereas Fe-Cr-Co-rich phase display FCC structure in HEAs [49-51]. After removing Co, the different phases in the bulk Al0.75FeNiCr are not readily distinguished, displaying larger featureless regions (region 3) and smaller featureless regions (region 4) with some pitting corrosion regions. It is postulated that the phases in the bulk Al0.75FeNiCr have similar resistance to etching by the aqua regia solution, which is consistent with that of Al0.6FeNiCr [52]. 9

EDS/SEM results shown in Table 2 reveal that region 3 and region 4 are Fe-Cr-rich phase which is associated with the main FCC phase and Al-Ni-rich phase corresponding to the BCC phase, respectively. Adding Ti into the Al0.75FeNiCr alloy, the bulk Al0.75FeNiCrCoTi0.25 alloy also exhibit featureless regions (region 5) and etched regions (region 6). EDS/SEM results in Table 2 show that region 5 is Fe-Cr-Co-rich phase and region 6 is Al-Ti-rich phase, while Ni distributes uniformly in these two phases. In other words, featureless regions and etched regions of the bulk Al0.75FeNiCrCoTi0.25 alloy are associated with the main FCC phase and the BCC phase, respectively [49,50]. Fig.4 shows TEM micrograph and corresponding SAED patterns of the bulk Al0.75FeNiCrCo HEA after SPS. In order to confirm the phase compositions of the bulk Al0.75FeNiCrCo HEA, EDS/TEM analyses of all the phases have been conducted in detail. It is worth pointing out that based on the EDS/TEM results and the corresponding SAED, bright-field image in Fig. 4(a) are marked by the symbols of “BCC and FCC”. The corresponding SAED pattern of BCC phase along [111] zone axis and the corresponding SAED pattern of FCC phase along [011] zone axis are showed in Figs. 4(b) and (c), respectively. According to the SAED patterns, the lattice parameters of the BCC and FCC phases are calculated as 0.290 and 0.360 nm, respectively, which are pretty similar to the values calculated from the XRD pattern in Fig. 2(a). The chemical composition results listed in Table 2 indicate that the BCC phase is Al-Ni-rich phase and the FCC phase is Fe-Cr-Co-rich phase, consistent with the EDS/SEM results and previous studies [49-51]. Note that the accuracy of EDS/TEM is better than that of EDS/SEM. The phases of the bulk Al0.75FeNiCrCo alloy are mainly ultra-fine grained, ranging from ~100 nm to several nanometers. As discussed above, the Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 alloys consist 10

of only two solid-solution phases, i.e., one FCC phase and one BCC phase, which are fewer than expected. Originally, the formation of solid-solution phases in HEAs is attributed to high entropy of mixing (ΔSmix) of ≥13.38 J K-1mol-1 [4]. High entropy of mixing can be estimated by the equation n

S mix   R (ci ln ci ) , where ci is the atomic fraction of the each element, R is the gas constant i 1

equal to 8.314 JK-1mol-1, and n is the number of component elements. Therefore, ΔSmix for the Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 alloys are 13.32, 11.47 and 14.32 JK-1mol-1, respectively. Thus sole consideration of entropy of mixing is insufficient to dominate the solid-solution phase formation in these three alloys. Yang et al. [53] proposed two important factors for the stable solid-solution formation in multi-component HEAs based on the calculation of mostly reported HEAs, i.e., the conditions of the thermodynamic parameter Ω≥1.1 and atomic size difference δ < 6.6% should be met simultaneously; and larger Ω and smaller δ are beneficial to the formation of solid solution. These two parameters Ω and δ can be calculated by Eqs.(2) and (3) as follows:



Tm Smix H mix



 C (1  r )

n

i 1

(2)

ri

2

(3)

i

Where ci is the atomic percentage of the each component element; r   i 1 ci ri is the average n

atomic radius and ri is the atomic radius of the each component. Tm and ΔHmix are calculated as follows: n

Tm   ci (Tm )i i 1

11

(4)

ΔH mix 

n



i 1,i  j

Ωij ci c j

(5)

Where (Tm)i is the melting point for each component element in HEAs, and Ωij  4ΔH ijmix is the regular solution interaction parameter between the ith and jth component, ΔH ijmix is the enthalpy of mixing between the ith and jth components. Table 3 presents the values of mixing enthalpy of atom-pairs of these three investigated alloys. Moreover, average valence electron concentration n

VEC=  ci (VEC )i ( (VEC)i is the valence electron concentration of the ith element) is proposed to i 1

predict the crystal structures of solid-solution phases in HEAs [54,55]. The calculated values for the parameters δ, ΔHmix, ΔSmix, Tm, Ω and VEC for these alloys are shown in Table 4. The ΔHmix, ΔSmix, Ω and δ for Al0.75FeNiCrCo are -10.88 kJ mol-1, 13.32 JK-1mol-1, 1.76 and 4.73, respectively. Removing Co from the Al0.75FeNiCrCo alloy reduces ΔHmix, ΔSmix and Ω, while increases δ; However, adding Ti into the Al0.75FeNiCrCo alloy increases ΔSmix and δ, while reduces ΔHmix and Ω. It is obvious that the parameters Ω and δ agree well with the solid-solution formation rules proposed by Yang et al., indicating the formation of stable solid-solution phases is possible for these three investigated alloys, consistent with our experimental results. In addition, the values of VEC for these three alloys meet the proposition that FCC+BCC structured solid-solution phases will be formed when 6.87 ≤ VEC < 8. Thus these three investigated alloys shows a mixture of FCC+BCC solid-solution phases [54,55]. Tables 2 and 3 are combined to investigate phase formation in the Al0.75FeNiCrCo alloy. For the Al0.75FeNiCrCo alloy, the mixing enthalpy of Al and Ni (-22 kJ mol-1) stays the most negative level, which could lead to Ni atoms preferring Al sites [51,56] (as shown in Table 3), and Co has a high concentration (17.9±1.3 at%) in Ni-Al rich phase might be attributed to the high negative 12

mixing enthalpy between Co and Al (-19 kJ mol-1). In addition, the similar atomic size of Co and Ni and 0 kJ mol-1 mixing enthalpy between them, could lead to Co and Ni form solid solution readily. According to Table 3, Fe, Cr and Co present small atomic size difference and mixing enthalpies for each atom-pair among these three elements is very close to 0 kJ mol-1, which is in favor of formation of solid-solution phase, therefore the Fe-Cr-Co-rich solid-solution phase could be readily formed. For the Al0.75FeNiCr alloy, the mixing enthalpy between Al and Ni is the most negative after removing Co from the Al0.75FeNiCrCo alloy, and therefore the presence of Al-Ni-rich phase is also preferred. Similarly, the mixing enthalpy of Fe and Cr is -1 kJ mol-1 and their atomic radii are very similar, thus the Fe-Cr-rich solid-solution phase can be formed easily. After adding Ti into the Al0.75FeNiCrCo alloy, since Ti has similar atomic radius and electronegativity with Al, thus the situation in the Al0.75FeNiCrCoTi0.25 alloy is pretty similar to that of the Al0.75FeNiCrCo alloy. Nano-twinning occurred in some FCC structured phase of the bulk Al0.75FeNiCrCo HEA, the similar nanoscale twins in HEAs prepared by MA and SPS have also been observed in the previous studies [26,27,47]. The lamella thickness of the nanoscale twins of grain A in Fig.5 (a) is ~50 nm. _ _

Fig. 5(b) shows the SAED pattern of grain A along [011]M (matrix) and [011]T (twin) zone axes, confirming the nanoscale twins belong to the FCC phase. Actually, nanoscale twins are only observed in the FCC phase, which have also been confirmed by EDS/TEM. The formation of twins in FCC phase might be attributed to the low stacking fault energy anticipated in the FCC that can be in favor of twining [57]. 3.3 Influence on mechanical properties Fig.6 shows the room-temperature compressive stress-strain curve of bulk Al0.75FeNiCr, Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 alloys after SPS. The yield strength (σy), compressive 13

strength (σmax) and strain-to-failure (εf) and Vickers hardness values of the studied alloys are given in Table 5. The compressive strength of Al0.75FeNiCrCo, Al0.75FeNiCr, and Al0.75FeNiCrCoTi0.25 are 2221, 2184 and 2376 MPa, respectively. The average Vickers hardness values of Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 have been measured as 577, 556 and 609 Hv, respectively. The high strength and hardness of these investigated alloys is possibly attributed to solid-solution strengthening caused by Al or/and Ti atoms with larger atomic radius than the other constituent elements, twin-boundary strengthening, as well as grain-boundary strengthening owing to ultra-fine grains [58,59]. With Co removal, the compressive strength of Al0.75FeNiCr decrease slightly compared to that of Al0.75FeNiCrCo, and the strain-to-failure of the former is slightly lower than the later. With Ti addition, Al0.75FeNiCrCoTi0.25 displays a higher compressive strength and a slightly lower strain-to-failure that that of Al0.75FeNiCrCo. Furthermore, Al0.75FeNiCrCoTi0.25 shows the highest Vickers hardness among the three investigated alloys. As analyzed in Section 3.2, Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 exhibits a major FCC phase and a minor BCC phase, whereas Al0.75FeNiCr is composed of a main BCC phase and an FCC phase. In other words, the volume fraction of the BCC phase in the Al0.75FeNiCr alloy is evidently higher than that of the BCC phase in the Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 alloys. Generally, FCC phase tends to possess better ductility than BCC phases in HEAs, whereas BCC phases tends to show higher strength and hardness [43,48]. Nevertheless, Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 with a higher volume fraction of the FCC phases exhibit a bit higher yield strength in the present work. After removing Co, compared to the Al0.75FeNiCrCo alloy, the Al0.75FeNiCr alloy shows more evidently strain hardening according to Fig. 6 and Table 5 [58]. It means that the Al0.75FeNiCrCo alloy has finer grains than the Al0.75FeNiCr alloy. According to the 14

Hall-Petch effect, alloys with finer grains will exhibit higher strength and hardness. It can be inferred that the effect of grain size on mechanical properties may be more significant than that of the BCC content. [52]. Decrement of ductility in the Al0.75FeNiCr alloy results from the increased content of BCC phase. The bulk Al0.75FeNiCr alloy shows a lower relative density than the bulk Al0.75FeNiCrCo alloy, which also can decrease the ductility and strength. In addition, removing such a high atomic concentration of ~21.05 at.% of Co element may change the intrinsic nature of the BCC and FCC phases in the Al0.75FeNiCr HEA, which may ultimately influence its mechanical properties. Hence, with Co removal, strength, hardness and strain-to-failure are decreased. With Ti addition, solid-solution strengthening can be enhanced in the Al0.75FeNiCrCoTi0.25 because of higher concentration of Ti contributing to solid-solution strengthening, and accordingly results in that strength and hardness are enhanced while strain-to-failure is decreased. 4. Conclusions Bulk Al0.75FeNiCrCo, Al0.75FeNiCr and Al0.75FeNiCrCoTi0.25 HEAs were successfully synthesized by MA followed by consolidation via SPS. Following MA, a main BCC phase with an FCC phase present was observed in each of these three investigated alloys. Removing Co from the Al0.75FeNiCrCo HEA or adding Ti into the Al0.75FeNiCrCo HEA are in favor of increasing the relative volume fraction ratio between BCC phase and FCC phase both in the MA and SPS stages. After SPS, both Al0.75FeNiCrCo and Al0.75FeNiCrCoTi0.25 consist of a primary FCC phase and a minor BCC phase. The volume fraction of FCC phase in the former is slightly higher than that in the latter. While a main BCC phase and an FCC phase are present in Al0.75FeNiCr. TEM results suggest that the bulk Al0.75FeNiCrCo HEA consist of an FCC phase and a BCC phase. The BCC phase is enriched in Al-Ni, while the FCC phase is enriched in Fe-Cr-Co. Meanwhile, a small 15

fraction of nanoscale twins were observed in the bulk Al0.75FeNiCrCo alloy. The three investigated alloys exhibit high strength and hardness. With removal of Co from Al0.75FeNiCrCo decreases strength, hardness and ductility. Adding Ti into the Al0.75FeNiCrCo alloy can increase strength and hardness, but decrease ductility. Acknowledgements The authors acknowledge the financial support from Fundamental Research Funds for the Central Universities, SCUT (2013ZZ014), from Specialized Research Fund for the Doctoral Program of Higher Education (20130172120027), and from National Natural Science Foundation of China (51271080).

16

References [1] W. Wang, C. Zhang, P. Xu, M. Yasir, L. Liu, Mater. Des. 73 (2015) 35–41. [2] Z.D. Zhao, Q. Chen, S.H. Huang, H.Y. Chao, Mater. Des. 31 (2010) 1906–1916. [3] L. Wang, K.Q. Qiu, Y.L. Ren, J. Alloys Compd. 612 (2014) 5–9. [4] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, et al., Adv. Eng. Mater. 6 (5) (2004) 299–303. [5] J.W. Yeh. Ann. Chim. Sci. Mat. 31 (2006) 633–648. [6] Y.P. Wang, B.S. Li, M.X. Ren, C. Yang, H.Z. Fu., Mater. Sci. Eng. A 491 (2008) 154–158. [7] T.T. Shun, Y.C. Du, J. Alloys Compd. 479 (2009) 157–160. [8] S. Guo, C. Ng, C.T. Liu, J. Alloys Compd. 557 (2013) 77–81. [9] W.H. Liu, Y. Wu, J.Y. He, T.G. Nieh, Z.P. Lu, Scr. Mater. 68 (2013) 526–529. [10] H.F. Sheng, M. Gong, L.M. Peng, Mater. Sci. Eng. A 567 (2013) 14–20. [11] M.H. Tsai, H. Yuan, G.M. Cheng, W.Z. Xu, K.Y. Tsai, C.W. Tsai, et al., Intermetallics 32 (2013) 329–336. [12] M.H. Tsai, J.W. Yeh, Mater. Res. Lett. 2 (3) (2014) 107–123. [13] S. Varalakshmi, G. A. Rao, M. Kamaraj, B.S. Murty, J. Mater. Sci. 45 (19) (2010) 5158–5163. [14] S. Varalakshmi, M. Kamaraj, B.S. Murty, Mater. Sci. Eng. A 527 (4–5) (2010) 1027–1030. [15] J.M. Wu, S.J. Lin, J.W, Yeh, S.K. Chen, Y.S. Huang, Wear 261(5-6) (2006) 513–519. [16] C.J. Tong, M.R. Chen, S.K. Chen, J.W. Yeh, T.T. Shun, S.J. Lin, et al., Metall. Mater. Trans. A 36 (5) (2005) 1263–1271. [17] C.Z. Yao, P. Zhang, M. Liu, G.R. Li, J.Q. Ye, P. Liu, et al., Electrochi. Acta 53 (28) (2008) 8359–8365. 17

[18] Y. Yu, W.M. Liu, T.B. Zhang, J.S. Li, J. Wang, H.C. Kou, et al., Metall. Mater. Trans. A 45(1) (2013) 201–207. [19] W.R. Wang, W.L. Wang, S.C. Wang, Y.C. Tsai, C.H. Lai, J.W. Yeh, Intermetallics 26 (2012) 44–51. [20] A. Munitz, M.J. Kaufman, J.P. Chandler, H. Kalaantari, R. Abbaschian, Mater. Sci. Eng. A 560 (2013) 633–642. [21] A. Takeuchi, N. Chen, T. Wada, Y. Yokoyama, H. Kato, A. Inoue, et al., Intermetallics 19 (2011) 1546–1554. [22] A. Cunliffe, J. Plummer, I. Figueroa, I. Todd, Intermetallics 23 (2012) 204–207. [23] H.Y. Ding, K.F. Yao, J. Non-Cryst. Solids 364 (2013) 9–12. [24] X.W. Qiu, C.G. Liu, J. Alloys Compd. 553 (2013) 216–220. [25] T.M. Yue, H. Xie, H.O. Yang, G.H. Meng, J. Alloys Compd. 587 (2014) 588–593. [26] S.C. Fang, W.P. Chen, Z.Q. Fu, Mater. Des. 54 (2014) 973–979. [27] Z.Q. Fu, W.P. Chen, Z. Zhen, H. Wen, E.J. Lavernia, Mater. Sci. Eng. A 619 (2014) 137–145. [28] Y. Dong, K.Y. Zhou, Y.P. Lu, X.X. Gao, T.M. Wang, T.J. Li, Mater. Des. 57 (2014) 67–72. [29] Z.A. Munir, U. Anselmi-Tamburini, M. Ohyanagi, J. Mater. Sci. 41(3) (2006) 763–777. [30] R. Sriharitha, B.S Murty, R.S. Kottada, J. Alloys Compd. 583 (2014) 419–426. [31] W. Ji, Z.Y. Fu, W.M. Wang, H. Wang, J.Y. Zhang, Y.C. Wang, et al., J. Alloys Compd.589 (2014) 61–66. [32] Z.Q. Fu. Research on microstructure and properties of AlCrFeNi-M high entropy alloys and its composites, (M.S. thesis), Harbin Institute of Technology, 2011. [33] Y.J. Zhou, Y. Zhang, Y.L. Wang, G.L. Chen, Appl. Phys. Lett. 90 (2007) 181904. 18

[34] K.B. Zhang, Z.Y. Fu, J.Y. Zhang, W.M. Wang, H. Wang, Y.C. Wang, et al., Mater. Sci. Eng. A 508 (2009) 214–219. [35] Y. Yu, J. Wang, J. Li, H. Kou, W. Liu, Mater. Lett. 138 (2015) 78–80. [36] W.R. Wang, W.L. Wang, J.W. Yeh, J.Alloys Compd. 589 (2014) 143–152. [37] ASTM E9-09. Standard Test Methods of Compression Testing of Metallic Materials at Room Temperature, 2009. [38] Y.L. Chen, Y.H. Hu, C.A. Hsieh, J.W. Yeh, S.K. Chen, J. Alloys Compd. 481 (1–2) (2009) 768–775. [39] S. Varalaksmi, M. Kamaraj, B.S. Murty, Metall. Mater. Trans. A 41(10) (2010) 2703–2709. [40] R. Sriharitha, B.S Murty, R.S. Kottada, Intermetallics 32 (2013) 119–126. [41] S. Mridha, S. Samal, P.Y. Khan, K. Biswas, Govind, Metall. Mater. Trans. A 44(10) (2013) 4532–4541. [42] S. Praveen, B.S. Murty, R.S. Kottada, JOM 65(12) (2013) 1797–1804. [43] Z.Q. Fu, W.P. Chen, S.C. Fang, D.Y. Zhang, H.Q. Xiao, D.Z. Zhu, J. Alloys Compd. 553 (2013) 316–323. [44] C. Suryanarayana, Prog. Mater Sci. 46 (2001) 1–84. [45] F.A. Mohamed, Acta Mater. 51 (2003) 4107–4119. [46] K.B. Zhang, Z.Y. Fu, J.Y. Zhang, J. Shi, W.M. Wang, H Wang, et al., J. Alloys Compd. 485 (2009) L31–L34 . [47] Z.Q. Fu, W.P. Chen. S.C. Fang, X.M. Li, Mater. Sci. Eng. A 597 (2014) 204–211. [48] S. Praveen, B.S. Murty, R.S. Kottada, Mater. Sci. Eng. A 534 (2012) 83–89. [49] B.S. Murty, J.W. Yeh, S. Ranganathan, High-Entropy Alloys, Butterworth-Heinemann, 1st 19

edition, June, 2014. [50] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, et al., Prog. Mater Sci. 61 (2014) 1–93. [51] W.P. Chen, Z.Q. Fu, S.C. Fang, H.Q. Xiao, D.Z. Zhu, Mater. Des. 51 (2013) 854–860. [52] Z.Q. Fu, W.P. Chen, H.M. Wen, Z. Chen, E.J. Lavernia, J. Alloys Compd. 646 (2015) 175–182. [53] X. Yang, Y. Zhang, Mater. Chem. Phys. 132 (2-3) (2012) 233–238. [54] S. Guo, C.T. Liu, Prog. Nat. Sci. Mater. Int. 21(2011) 433–446. [55] S. Guo, C. Ng, J. Lu, C.T. Liu, J. Appl. Phys. 109 (2011) 103505. [56] A. Takeuchi, A. Inoue, Mater Trans. 46 (12) (2005) 2817–2829. [57] Y.T. Zhu, X.Z. Liao, X.L. Wu, Prog. Mater Sci. 57 (1) (2012) 1–62. [58] H.M. Wen, T.D. Topping, D. Isheim, D.N. Seidman, E.J. Lavernia, Acta Mater. 61 (2013) 2769–2782. [59] K.K. Ma, H.M. Wen, T. Hu, T.D. Topping, D. Isheim, D.N. Seidman, et al., Acta Mater. 62 (2014) 141–155.

20

Figure Captions: Fig. 1. XRD patterns of HEA powders with different milling time. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25. Fig. 2. XRD pattern of bulk alloys after SPS. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25. Fig. 3. SEM micrographs of bulk HEAs after SPS. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25. Fig. 4. TEM micrograph and corresponding SAED patterns of the bulk Al0.75FeNiCrCo after SPS. (a) Bright-field image, (b) SAED pattern of BCC phase along [111] zone axis, (c) SAED pattern of FCC phase along [011] zone axis. Fig. 5. Twins in the bulk Al0.75FeNiCrCo alloy. (a) Bright-field TEM image showing nanoscale _ _

twins in FCC grain A, (b) SAED pattern of grain A along [011]M (matrix) and [011]T (twin) zone axes. Fig. 6. Engineering stress-strain curve of consolidated alloys under compression at room temperature

21

Table Captions: Table 1. Crystalline size and lattice strain of the investigated alloys during MA. Table 2. Chemical compositions (in at.%) of the bulk alloys analyzed by EDS/SEM or EDS/TEM. Regions 1–6 as marked were analyzed by EDS/SEM, and the FCC and BCC phases of the Al0.75FeNiCrCo were analyzed by EDS/TEM. Table 3. The enthalpy of mixing ( H ijmix , kJ mol-1) of binary equiatomic alloys calculated using Miedema’s approach [56]. Table 4. The values of parameters ΔHmix, ΔSmix, Tm, δ, Ω, and VEC for the investigated alloys. Table 5. Mechanical properties at room temperature of the investigated alloys.

]

22

Fig. 1. XRD patterns of HEA powders with different milling times. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25.

23

Fig. 2. XRD pattern of bulk alloys after SPS. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25.

24

3 1

4 2

5

6

Fig. 3. SEM micrographs of bulk HEAs after SPS. (a) Al0.75FeNiCrCo, (b) Al0.75FeNiCr, (c) Al0.75FeNiCrCoTi0.25.

25

Fig. 4. TEM micrograph and corresponding SAED patterns of the bulk Al0.75FeNiCrCo after SPS. (a) Bright-field image, (b) SAED pattern of BCC phase along [111] zone axis, (c) SAED pattern of FCC phase along [011] zone axis.

26

A

Fig. 5. Twins in the bulk Al0.75FeNiCrCo alloy. (a) Bright-field TEM image showing nanoscale _ _

twins in FCC grain A, (b) SAED pattern of grain A along [011]M (matrix) and [011]T (twin) zone axes.

27

Fig. 6. Engineering stress-strain curve of consolidated alloys under compression at room temperature

Table 1. Crystalline size and lattice strain of the investigated alloys during MA. Alloys

Al0.75FeNiCoCr Al0.75FeNiCr Al0.75FeNiCoCrTi0.25

Milling time (h) 30 50 45 50 30 50

Crystalline size (nm)

Lattice strain (%)

BCC

FCC

BCC

FCC

13 10 14 13 13 11

10 7 11 10 10 9

0.66 0.90 0.57 0.68 0.67 0.82

0.87 1.28 0.88 0.93 0.90 0.99

28

Table 2. Chemical compositions (in at.%) of the bulk alloys analyzed by EDS/SEM or EDS/TEM. Regions 1–6 as marked were analyzed by EDS/SEM, and the FCC and BCC phases of the Al0.75FeNiCrCo were analyzed by EDS/TEM. Alloys

Regions

Al

Fe

Ni

Cr

Co

Ti

Al0.75FeNiCrCo

Nominal composition 1 2 FCC BCC Nominal composition 3 4 Nominal composition 5 6

15.8 14.2 31.9 5.1±0.7 27.1±2.2 20 13.5 34.6 15 7.9 19.3

21.1 25.1 14.6 26.1±1.7 12.4±2.0 26.6 33.9 19.2 20 23.6 19.2

21.1 15.9 26.7 20.2±1.2 36.9±1.6 26.7 23.9 28.1 20 19.7 19.3

21 22.2 9.7 23.2±1.2 5.6±0.5 26.7 28.7 18.1 20 22.8 18.0

21 22.6 17.1 25.4±1.1 17.9±1.3 — — — 20 21.8 18.7

— — —

Al0.75FeNiCr

Al0.75FeNiCrCoTi0.25

29

— — — 5 4.18 5.5

Table 3. The enthalpy of mixing ( H ijmix , kJ mol-1) of binary equiatomic alloys calculated using Miedema’s approach [56]. Element (atomic radii, Å)

Al

Fe

Ni

Cr

Co

Ti

Al (1.43) Fe (1.27) Ni (1.25) Cr (1.28) Co(1.26) Ti (1.46)

— -11 -22 -10 -19 -30

-11 — -2 -1 -1 -17

-22 -2 — -7 0 -35

-10 -1 -7 — -4 -7

-19 -1 0 -4 — -28

-30 -17 -35 -7 -28 —

30

Table 4. The values of parameters ΔHmix, ΔSmix, Tm, δ, Ω and VEC for the investigated alloys Alloys

ΔHmix (kJ mol-1)

ΔSmix (JK-1mol-1)

Tm (K)

δ (%)

Ω

VEC

Al0.75NiFeCrCo Al0.75NiFeCr Al0.75NiFeCrCoTi0.25

-10.88 -12.09 -14.22

13.32 11.47 14.32

1708 1704 1723

4.73 5.11 5.62

1.76 1.36 1.46

7.42 7.00 7.25

31

Table 5. Mechanical properties at room temperature of the investigated alloys. Alloys

σy (MPa)

σmax (MPa)

εf (%)

Hardness (Hv)

Al0.75NiFeCrCo Al0.75NiFeCr Al0.75NiFeCrCoTi0.25

1938 1741 1926

2221 2184 2376

7.6 7.4 6.8

577 556 609

32