**A Novel Low-Activation VCrFeTa***x***W***x* **(***<sup>x</sup>* **= 0.1, 0.2, 0.3, 0.4, and 1) High-Entropy Alloys with Excellent Heat-Softening Resistance**

#### **Weiran Zhang 1, Peter K. Liaw <sup>2</sup> and Yong Zhang 1,3,\***


Received: 15 November 2018; Accepted: 6 December 2018; Published: 11 December 2018

**Abstract:** The microstructure, Vickers hardness, and compressive properties of novel low-activation VCrFeTa*x*W*<sup>x</sup>* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) high-entropy alloys (HEAs) were studied. The alloys were fabricated by vacuum-arc melting and the characteristics of these alloys were explored. The microstructures of all the alloys exhibited a typical morphology of dendritic and eutectic structures. The VCrFeTa0.1W0.1 and VCrFeTa0.2W0.2 alloys are essentially single phase, consisting of a disordered body-centered-cubic (BCC) phase, whereas the VCrFeTa0.2W0.2 alloy contains fine, nanoscale precipitates distributed in the BCC matrix. The lattice parameters and compositions of the identified phases were investigated. The alloys have Vickers hardness values ranging from 546 HV0.2 to 1135 HV0.2 with the x ranging from 0.1 to 1, respectively. The VCrFeTa0.1W0.1 and VCrFeTa0.2W0.2 alloys exhibit compressive yield strengths of 1341 MPa and 1742 MPa, with compressive plastic strains of 42.2% and 35.7%, respectively. VCrFeTa0.1W0.1 and VCrFeTa0.2W0.2 alloys have excellent hardness after annealing for 25 h at 600–1000 ◦C, and presented compressive yield strength exceeding 1000 MPa with excellent heat-softening resistance at 600–800 ◦C. By applying the HEA criteria, Ta and W additions into the VCrFeTaW are proposed as a family of candidate materials for fusion reactors and high-temperature structural applications.

**Keywords:** low-activation high-entropy alloys (HEAs); high-temperature structural alloys; microstructures; compressive properties; heat-softening resistance

#### **1. Introduction**

With the rapid development of human civilization, the demand for energy is increasing and the fossil fuel sources are running out. As nuclear energy produces more energy with less pollution, in the long run, nuclear energy will be the next major energy source after fossil fuels, such as coal and oil, to meets human needs [1–3]. However, with fast-growing nuclear power technology, people have higher requirements for the reliability and safety of the nuclear-power [4,5]. The structural materials for new commercial fusion nuclear reactors operate in a harsh environment that is high temperature and chemically reactive, and experiences time-varying stress and intense neutron radiation [4,6], while required to be environmentally-friendly (reduced activation properties). This has motivated worldwide research and development (R&D) on advanced nuclear power systems. Therefore, the exploitation of novel and advanced materials that meet the requirements of these severe conditions will be a key issue for the development of new commercial reactors in the future [4].

Reduced activation or low-activation materials means that the main source of radioactivity after neutron irradiation is short- or medium-lived radioactive elements [7]. The challenge is managing the radioactive waste after shutting down the reactor, and fusion will lose its advantage of being a cleaner energy. Hence, it is important to choose specific materials for the reliable operation of these reactors [8]. The reduced activation ferritic/martensitic (RAFM) steels [9–11], such as Eurofer 97, China low activation martensitic (CLAM), and F82H, are considered to be the original candidate blanket structural materials and/or first wall for future fusion-power devices due to their excellent thermophysical properties, high thermomechanical capabilities, low-activation property, and resistance to neutron irradiation. RAFM steels have been developed using modified (8–12)CrMoVNb type ferritic martensitic steels by replacing Nb, Mo, and Ni with W, Mn, and Ta to achieve the low activation properties [1,7]. However, the operating temperature limit of RAFM steels is currently about 550 ◦C, which limits the overall thermodynamic efficiency of the power plant [12]. In order to widen the operating temperature window for fusion reactors, several alternative advanced materials options are being pursued. These alternatives include oxide-dispersion-strengthened (ODS) ferritic steels [5], vanadium alloys, and silicon carbide fiber-reinforced silicon carbide matrix composites [13,14]. Existing materials struggle to meet the requirements of fusion reactors operating in extreme environments, such as higher temperatures and stronger neutron irradiation [13,15–17]. Therefore, the first task for the development of fusion energy is to develop high-performance materials. According to the requirements of fusion reactors for short- or medium-lived radioactive materials and existing alloys, such as RAFM steels, ODS steels, and vanadium alloys, we summarize the low activation elements and high activation elements in Table 1.

**Table 1.** Low activation elements and high activation elements.


High-entropy alloys (HEAs) are new materials developed in the field of metals in the past decade [18–20]. The term HEAs signifies unconventional alloy systems composed of at least four principal elements, and the atomic percent of each composed element is between 5 at. % (atomic percent) and 35 at. %, which benefits the formation of single-phase solid-solution on the simple underlying face-centered-cubic (FCC), body-centered-cubic (BCC), and hexagonal-close-packing (HCP) structures compared with intermetallics [19,21,22]. HEAs are strongly contrasted with conventional alloys, which are usually based on one or two major elements, and the addition of trace amounts of alloying elements mostly leads to the formation of new phases [4,23]. HEAs can have high hardness [24–27], great creep resistance [28,29], good irradiation resistance [4,6,16,30], good structural stability [28,31–34], and excellent high-temperature strength [28,31,35]. These advantages make HEAs specifically suitable for high-temperature [13] and irradiation applications [14,30]. Zhang et al. [15] reported that Al*x*CoCrFeNi (*x* = 0.1, 0.75, and 1.5) shows great phase stability and swelling resistance under heavy ion irradiation at room temperature to high displacement per atom (dpa). They thought that this trend is due to the severe lattice distortion and sluggish diffusion, which are unique to HEAs. MoNbHfZrTi [36] shows a high compressive yield strength of 1719 MPa and 1575 MPa in as-cast and as-homogenized states at room temperature, respectively. This alloy has high compressive yield strength at elevated temperatures (825 MPa at 800 ◦C and 187 MPa at 1200 ◦C) and shows a drop in flow stress after yielding.

The study of HEAs has expanded from the central region to the surroundings of the phase diagram, which means that the research is changing from examining equiatomic single-phase solid-solution alloys to non-equimolar multi-phase solid-solution alloys [9,22,37,38]. According to the concept of HEAs, the reduced activation elements of Fe, Cr, V, W, and Ta were chosen to form several non-equimolar and an equiatomic TaWFeCrV HEAs. However, W is a candidate element for plasma-oriented components in commercial energy-fusion reactors, as it can increase strength and reduce the brittle transition temperature [3,39–41]. Cr reinforces the corrosion resistance of the alloys [5,42,43]. V and Ta can improve creep properties, reduce grain size, and enhance the toughness and strength of the alloys [44,45]. Fe has excellent ductility [46] and is inexpensive. In this study, the V–Cr–Fe–Ta–W alloys were prepared by vacuum-arc melting. Their alloying behaviors, microstructures, and mechanical properties were investigated in detail.

#### **2. Experimental Procedures**

Alloy ingots with the nominal composition of VCrFeTa*x*W*<sup>x</sup>* (*x*: molar ratio; *x* = 0.1, 0.2, 0.3, 0.4, and 1 denoted by T0.1, T0.2, T0.3, T0.4, and T1, respectively) were prepared by vacuum arc melting with high purity elements (the purity of each elements was better than 99.9 wt. %) in a Ti-gettered high-purity argon atmosphere. The nominal chemical compositions of the obtained five alloys are listed in Table 2. These ingots were re-melted at least 5 times in order to achieve compositional homogeneity, and each sample weight was around 50 g. The produced alloys were annealed at 800 ◦C for 25 h in a high-purity argon furnace and cooled down in water.

**Table 2.** Nominal compositions of experimental alloys (Nominal compositions, at. %).


The analyzed samples were cut from the middle part of the as-cast alloys to attain flat surfaces for the microstructure study. The exposed surfaces were subsequently ground and polished with the standard polishing process. The crystal structures of as-cast and annealed samples were identified by X-ray diffraction (XRD) (Bruker D8, Karlsruhe, Germany.) with Cu Kα radiation generated at 40 kV and 40 mA and with the scanning angles (2θ) ranging from 30◦ to 100◦ at a step of 0.0102◦ and 10 s dwell time per step. A scanning electron microscope (SEM) (ZEISS SUPRA 55, Jena, Germany.) was used for microstructural analyses, and experimental compositions were analyzed by the energy-dispersive spectroscopy (EDS) (ZEISS SUPRA 55, Jena, Germany.).

Mechanical properties were studied in terms of compression and micro-hardness in air at room temperature. The cylindrical test specimens with a diameter of 3 mm and height of 6 mm were cut from the middle of the ingots for compressive tests, which were performed on a computer-controlled electronic universal testing machine. The high-temperature compressive performances were studied on a Gleeble machine. The heating rate was set to 20 ◦C/s, and the holding time was five minutes, then air cooled. The initial strain rate for all compressive tests was 10−<sup>3</sup> s−1. The Vickers micro-hardness tests were carried out with a load of 200 g and a 15 s dwelling time with at least 12 tested points for each test specimen. SEM was used to observe the fracture surfaces of the samples after compressive tests.

#### **3. Results**

#### *3.1. Structural Characterization*

The XRD patterns of as-cast alloys are shown in Figure 1. Only one BCC1 crystal structure (the Fe-, Cr-, and V-rich phase) had lattice parameters of 0.2935 nm and 0.2937 nm, which were identified in T0.1 and T0.2 alloys in the as-cast state according to Bragg's law, respectively. With the increase in Ta and W content, the structures became rather complex. Both BCC1 and BCC2 (W-rich phase) solid-solution structures, together with Laves (Fe2Ta-type) phases, were found in these alloys. When the molar ratio of Ta and W increased to 0.3, BCC2 appeared with a lattice parameter of 0.3174 nm. The lattice parameter of the BCC1 phase was 0.2947 nm with a peak width much wider than those of

the T0.1 and T0.2 alloys. However, there was a weaker peak of the Laves phase around 47◦. The phase composition of T0.4 is similar to T0.3, with the only difference in the strength. The peak (110) for BCC1 is broader than that of the T0.3, and the peak intensity is weaker. The peak of BCC2 phase with the lattice parameter of 0.3178 nm is stronger, both in the peak width and intensity compared with T0.3. Despite this trend, T0.4 is mainly composed of BCC1 with a lattice parameter of 0.2963 nm. When the alloy reaches an equimolar ratio, VCrFeTaW, which is denoted T1, is composed of two BCC phases with lattice parameters of 0.2962 nm for BCC1 and 0.3166 nm for BCC2.

**Figure 1.** X-ray diffraction (XRD) patterns of the as-cast VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys.

Comparing the X-ray diffraction parameters of each sample, the lattice parameter of BCC1 first rises and then falls, as shown in Figure 1. When the contents of Ta and W reached 0.4 molar, the lattice constants of BCC1 and BCC2 were the largest. The lattice-constant changing trend of BCC2 is similar to that of BCC1. The difference is that the decreasing degree of BCC2 is greater than BCC1. It can be seen from the XRD pattern that the Laves phase has obvious diffraction peaks when *x* reaches 0.4. The intensity of the diffraction peak of BCC1 decreased significantly but still dominated. The phase composition and lattice constant of each alloy are listed in Table 3. We explain this phenomenon in two ways. Considering the atomic radius, the atomic radii of Ta and W are 1.48 Å and 1.41 Å, respectively, which are larger than the atomic radii of other elements (Fe: 1.24 Å, Cr: 1.25 Å, and V: 1.32 Å). When the contents of Ta and W were less than 0.4, they could be dissolved in the solid solution of BCC1 to a certain degree. Since the atomic radii of Ta and W are larger than those of other elements, the lattice distortion of the solid-solution matrix, BCC1, and the lattice constant of BCC1 increased, which inevitably appeared. Thereby, the solid-solution lattice-strain energy increased. When the content of Ta exceeds 0.4, the solid solubility of the Ta element in the solid-solution matrix of BCC1 has reached saturation. The addition of excess Ta induces the precipitation of the Laves phase. The precipitate of the second phase mitigates the lattice distortion of the solid-solution matrix to a certain extent. Hence, the lattice strain energy of the solid solution is released. An interesting phenomenon can be seen from the XRD pattern: when *x* is 1, the peak of the Laves phase is not enhanced by the previous analysis, but is slightly weakened. This feature may be related to the phase of BCC2, which is rich in W.


**Table 3.** Phase compositions of the VCrFeTa*x*W*x* alloy system and lattice constants of the solid-solution phase.

When *x* increased from 0.3 to 0.4, due to the precipitated Laves phase, the lattice constant of BCC2 significantly increased. When x increased to 1, the lattice constant of W drastically reduced, which may be related to the fact that W is a dominant phase, and elements, such as Fe, Cr, and V, having a small atomic radius, are dissolved in W.

#### *3.2. Microstructures and Chemical Compositions*

The microstructures of the as-cast and as-polished VCrFeTa*x*W*<sup>x</sup>* samples captured using SEM exhibited multiple phases. Figure 2 presents the microstructures of the samples. All the alloys exhibited a typical cast dendrite (DR) structure. The EDS component analyses of the alloys are provided in Table 4. W and V were mainly present in the DR regions, Fe and Ta were concentrated in the inter-dendritic (IR) regions, whereas Cr was more evenly distributed. The distribution of elements could mainly be explained by the mixing enthalpy between elements and the difference in melting points. During the solidification of the alloys, W and V, with higher melting points, first solidified to form DR. The mixing enthalpy (Table 5) between Ta and Fe, −15 KJ/mol, which is the minimum for the selected alloy system, indicates that Ta and Fe combined together more easily than other elemental pairs in the process of solidification due to their great intermetallic compounding ability. Due to the low melting point of Fe, Fe and Ta formed in IR after the formation of DR. Ta and Fe formed a Laves phase in the solid-solution matrix composed of W and V. The microstructures of the alloy system will be described in detail later.

**Figure 2.** Scanning electron microscope (SEM) backscatter electron images of the as-cast alloys. (**a**,**b**) VCrFeTa0.1W0.1; (**c**,**d**) VCrFeTa0.2W0.2; (**e**,**f**) VCrFeTa0.3W0.3; (**g**,**h**) VCrFeTa0.4W0.4; and (**i**,**j**) VCrFeTaW.


**Table 4.** Chemical compositions in different regions of various alloys by energy-dispersive spectroscopy (EDS) (at. %).

**Table 5.** The formation enthalpies between elements (KJ/mol).


Figure 2a–d correspond to the microstructures of the T0.1 and T0.2 alloys. From these figures, we found that there is a small amount of the second phase in the white regions in the figures, which is the Laves phase according to the EDS results. The volume fraction of the Laves phase is lower than the XRD detection limit, as no obvious Laves phase peak was observed in the XRD pattern, as presented in Figure 1. It can be seen from the figures that they exhibit a similar DR microstructure. The IR structure of the Laves phase combined with the primary BCC1 phase in the DR region formed eutectic structures and is composed of a network-like appearance. We found that the DR size of T0.2 is less than T0.1, indicating that the addition of Ta and W can effectively reduce the size of the DR crystal grains, which is consistent with the width of the diffraction peak of T0.2, which is wider than T0.1 in the XRD pattern. The eutectic structure regions of the T0.2 alloy are significantly larger than those of the T0.1 alloy. The plate-like Laves phase thickens with increasing Ta content, meaning that the volume fraction of the Laves phase is increasing with the Ta content. This trend is consistent with the appearance of a less pronounced Laves phase around 36◦ on the XRD curve of T0.2 (Figure 1). From the magnified view of T0.2 in Figure 2b, there is a black area in the middle of the gray and white regions. Based on the EDS results, we found that the area is enriched by V and Fe, which was the BCC1 solid solution. On the gray DR, there are many fine precipitate particles, which is consistent with the increase in the BCC1 lattice constant.

When the content of Ta and W reached 0.3, the microstructure of the alloy still exhibited a typical DR structure (Figure 2e,f). Combined with the EDS results, the DR structure is formed by gray regions of solid-solution phases and precipitated particles of the Laves phase. The Laves phase in the IR region, with the BCC1 primary phase of the black regions in the V-rich area, formed a eutectic structure. The Laves-phase volume fraction increased significantly, which is consistent with the XRD results (Figure 1).

For the microstructure of the T0.4 alloy, shown in Figure 2g,h, the Laves phase volume fraction increased relative to T0.1, T0.2, and T0.3 alloys shown in Figure 2a–c, respectively. The Laves phases grew into a bar-like shape distributed in the matrix (Figure 2h). Combined with XRD and EDS results, the DR regions consist of a gray W-rich area of BCC2 and a black Fe–Cr–V-rich area of BCC1. The eutectic structure formed by the precipitated Laves peak of the IR structure and the primary BCC1. The distortion of the matrix was released, and the lattice constant of BCC1 reached the maximum owing to the complete precipitation of the Laves phase, which was also observed for BCC2.

The T1 alloy in Figure 2i,j had a similar appearance to T0.4, and the microstructure was a typical DR crystal structure. The white region of DR is distinctly fishbone-like shaped and is enriched in W, whereas a portion of V is a solid dissolved in the W matrix. Due to the smaller atomic radius of the V solid solution in the W matrix, the lattice constant of W decreased, which is consistent with the XRD pattern in Figure 1. The IR structure presents a hypereutectic structure, and the primary phase is not the BCC1 in the black regions, but the Laves phase in the gray regions, whereas, in the highlighted region, Laves and BCC1 phases are present in eutectic structure. These trends indicate that the additions of Ta and W not only influence the phase composition of the alloy system, but also change the microstructure of the alloy system. The addition of Ta promoted regular changes in the microstructures of the BCC1 alloy system, and the transition from the hypoeutectic structure (*x* = 0.4) to hypereutectic structures (*x* = 1) occurred.

#### *3.3. Mechanical Properties*

#### 3.3.1. Mechanical Properties at Room Temperature

Figure 3a shows a histogram of the microhardness of the alloys at room temperature and the average microhardness values are listed in Table 6. It can be seen from Figure 3a that the microhardness of the alloy system increases linearly with the increase in the Ta and W contents, from 564 HV0.2 of T0.1 to 1135 HV0.2 for the equimolar alloy, T1. The hardness enhanced with increasing *x* due to: (1) the Laves phase belongs to the intermetallic compound, and the hardness is high and (2) the addition of Ta and W caused the lattice distortion of the alloy matrix to increase continuously, and the hardness of the alloy also improved. The relationship between Ta and W content, *x*, and microhardness (yHV) is fitted in Figure 3b, which can be expressed as yHV = 553.1 + 609.2*x*, where y is the microhardness of the alloy. The linear correlation coefficient, R, is 0.942, meaning that relationship is linear.


**Table 6.** Mechanical properties of the as-cast VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys.

σ0.2: yield strength; σbc: ultimate compressive strength; and εp: plastic-strain limit.

**Figure 3.** (**a**) Hardness histogram and standard deviation of the as-cast VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys at room temperature; (**b**) The hardness curve of the as-cast VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys as a function of the Ta and W content, and hardness test process.

The compressive engineering stress-strain curves are shown in Figure 4, and the compressive properties, such as yield strength σ0.2, fracture strength σbc, and plastic strain limit εp, are summarized in Table 6. Evidentially, the Ta and W content had a very pronounced effect on the compressive behavior of the alloys. It can be seen that the T0.1 and T0.2 alloys showed excellent compressive properties, with yield strength, fracture strength, and plastic strain values in T0.1 and T0.2 of 1341 MPa, 2917 MPa, and 42.2%; and 1742 MPa, 3265 MPa, and 35.7%, respectively. Compared with the compressive properties of T0.1 and T0.2, both the yield strength and fracture strength increased, which is a trade-off with plastic strain. This trend is mainly due to the presence of a large amount of dispersed fine precipitates in the T0.2 alloy and an increase in the lattice distortion, resulting in an increase in the strength of the alloy. The plastic strain decreased due to the increase in the volume fraction of the second phase. It can be seen that the compressive properties substantially decrease with the Ta and W content changing from 0.3 to 1 mainly because of the increase in the volume fraction and size of the Laves phase as the Ta content increases. The compressive performance of T0.3 is worse than T0.4 and T1, probably due to the precipitated particles growing in the matrix.

**Figure 4.** The compressive stress-strain curves of the as-cast VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys at room temperature with a diameter of 3 mm.

The fracture-surface morphologies of the as-cast alloys after compressive tests at room temperature are depicted in Figure 5. Figure 5a,b show the side sections of T0.1 and T0.2 alloys, respectively. It can be seen from the figures that the alloys are not crushed, and the macroscopic appearance are waist-drum shaped. The picture of the T0.1 alloy in Figure 5a shows that the fracture surface presents severe plastic deformation, indicating its superior room-temperature plasticity and exhibiting a certain plasticity fracture. In Figure 5b, the T0.2 alloy exhibits a fracture angle of nearly 45◦, and approximately parallel sliding steps of unevenness can be observed in the fracture profile. The fracture surface of the T0.3 alloy is shown in Figure 5c, which has a distinct cleavage step due to the cleavage fracture. T0.4 and T1 alloys have similar fracture patterns, and the fracture surfaces are relatively flat, demonstrating obvious tear and river patterns, which are typical quasi-cleavage fractures, in Figure 5d,e, respectively. This trend is consistent with the compressive strength of T0.4 and T1 being better than T0.3.

**Figure 5.** SEM micrographs of the fracture surfaces of VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys at room temperature: (**a**) VCrFeTa0.1W0.1; (**b**) VCrFeTa0.2W0.2; (**c**) VCrFeTa0.3W0.3; (**d**) VCrFeTa0.4W0.4; and (**e**) VCrFeTaW.

#### 3.3.2. Mechanical Properties at High-Temperature

Since T0.1 and T0.2 alloys exhibit excellent room-temperature strength, we also examined their high-temperature properties. Annealing at 600 ◦C, 800 ◦C, and 1000 ◦C for 25 h did not cause alloy softening. The hardness of the annealed T0.1 alloys are 605 HV0.2/ 600 ◦C and 621 HV0.2/800 ◦C, and T0.2 alloys are 721 HV0.2/600 ◦C and 762 HV0.2/800 ◦C, respectively, obviously higher than those of as-cast alloys presented in Table 7, indicating that the T0.1 and T0.2 alloys have the great softening resistance. Table 7 and Figure 6 show the compressive engineering stress-strain curves of T0.1 and T0.2 alloys from room temperature to 1000 ◦C. The strength increased and the ductility decreased with the yield strength at 600 ◦C. The strength decreased and the ductility increased at 800 ◦C. At 1000 ◦C, the strength of T0.1 and T0.2 alloys dropped significantly, indicating a certain degree of softening, which is also reflected in the hardness at this temperature listed in Table 7. Surprisingly, the yield strength, fracture strength, and plastic strain at 800 ◦C for T0.1 and T0.2 are 1019 MPa, 1289 MPa, and 50% and 1033 MPa, 1260 MPa, and 40.6%, respectively.

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**Table 7.** Mechanical properties of T0.1 and T0.2 alloys at 600–1000 ◦C.

Note: σ0.2: yield strength, σbc: ultimate compressive strength, and εp: plastic-strain limit.

**Figure 6.** The compressive stress-strain curves of (**a**) VCrFeTa0.1W0.1 and (**b**) VCrFeTa0.2W0.2 alloys at different temperatures with a diameter of 3 mm.

#### **4. Discussion**

#### *4.1. Phase Selection*

From the results, the phase structures of the VCrFeTa*x*W*x* HEAs are not as simple as presented in the XRD patterns in Figure 1. The various criteria were analyzed in order to further understand the phase formation of the alloy in this system.

Studies have shown that HEAs are prone to form solid solutions with an FCC structure or a BCC [19,21,47]. Hume-Rothery rules [21,48,49] govern the criteria for the formation of solid solutions in the binary alloy system, which include the crystal structure factor, atomic-size factor, valence-electron-concentration factor, and chemical electronegativity factor. As a special kind of the solid-solution alloy, HEAs have many components, complicating distinguishing solutes or solvents. Hence, it is difficult to study them using traditional methods [22,50]. Recent investigations extended the Hume-Rothery rules for explaining the criteria for the formation of the solid-solution structure in the HEA area with the aid of empirical relationships.

Zhang et al. [48,51] proposed three parameters affecting the formation of the HEA solid-solution phase: atomic-size difference (delta, δ), mixing enthalpy (ΔHmix), and mixing entropy (ΔSmix), to predict the phase formation in HEAs, amorphous metallic glasses, and intermetallic compounds. These calculation methods are detailed as follows:

$$
\Delta \mathbf{H}\_{\text{mix}} = \sum\_{i=1, i \neq j}^{n} \Omega\_{ij} \mathbf{c}\_i \mathbf{c}\_j
$$

$$
\Omega\_{ij} = 4 \Delta \mathbf{H}\_{\text{AB}}^{\text{mix}}
$$

$$
\Delta \mathbf{S}\_{\text{mix}} = \text{klnw} = -\mathbb{R} \sum\_{i=1}^{n} (\mathbf{c}\_i \ln \mathbf{c}\_j)
$$

$$
\delta = \sqrt{\sum\_{i=1}^{n} \mathbf{c}\_i (\mathbf{r}\_i - \overline{\mathbf{r}})^2}
$$

$$
\overline{\mathbf{r}} = \sum\_{i=1}^{n} \mathbf{c}\_i \mathbf{r}\_i
$$

where *n* is the number of the involved elements in an alloy, ΔHmix AB is the mixing of the enthalpy of binary equiatomic AB alloys, Ω*ij* is the regular melt-inter-action parameter between the *i*-th and *j*-th elements, R is the gas constant, c*<sup>i</sup>* and r*<sup>i</sup>* are the atomic percentage and atomic radius of the *i*-th element, respectively, and r is the average atomic radius. They concluded that solid solutions tend to form in the region delineated by δ > 6.6%, −15 KJ/mol ≤ ΔHmix ≤ 5 KJ/mol, and 11 J/(K·mol) ≤ ΔSmix ≤16.5 J/(K·mol).

To be better understand the criteria of HEAs, a new parameter Ω [21,51] was proposed to correlate the relative contribution of the change in ΔHmix and ΔSmix, expressed as:

$$
\Omega = \frac{\text{T}\_{\text{m}} \Delta \text{S}\_{\text{mix}}}{|\Delta \text{H}\_{\text{mix}}|}
$$

$$
\text{T}\_{\text{m}} = \sum\_{i=1}^{n} \mathbf{c}\_{i} (\mathbf{T}\_{\text{m}})\_{i}
$$

where (Tm)*<sup>i</sup>* is the melting temperature of the *i*-th component. By analyzing the phase formation using the parameters, Ω and δ, of various reported multicomponent alloys, new parameters for forming solid-solution phases in HEAs were suggested [48]: Ω ≥ 1.1 and δ ≤ 6.6%.

HEAs often form solid solutions of an FCC or BCC phase, and the above criteria can effectively predict whether the alloy can form a solid-solution structure, but it is impossible to predict whether the solid-solution structure of the alloy is an FCC or BCC phase. Guo et al. [52] proposed the

relationship between the valence-electron concentration (VEC) and solid-solution stability. The VEC in a multi-component system is:

$$\text{VEC} = \sum\_{i=1}^{n} \mathbf{c}\_{i} (\text{VEC})\_{i}$$

where (VEC)*<sup>i</sup>* is the VEC of the *i*-th element. From published experimental results [37,52], the limitation that they suggested were: FCC phases occur at VEC ≥ 8.0, BCC phases at VEC < 6.87, and a mixture of FCC and BCC phases at 6.87 ≤ VEC < 8.

For the VCrFeTa*x*W*x* alloys in the present work, we studied the phase-formation rules according to the above parameters. The specific results are shown in Table 8. It can be seen from the table that as the contents of Ta and W increase, the δ of the alloy system grows from 3.59% to 5.41%, as plotted in Figure 7a. This indicates that the degree of the lattice distortion caused by the atomic arrangement is increasing, which is consistent with the lattice constants of the two phases when x is between 0.1 and 0.4 (the lattice constant decreases in the case of the equimolar HEA, which is related to the Laves phase being the primary phase, and V is distributed into a solid-solution with W). Despite the change in δ and ΔHmix with the increase in the Ta and W content, their values meet the requirement for forming solid solutions, as plotted in Figure 7a,b, indicating that the mixing entropy effect of the alloy is stronger than that of the mixing enthalpy. The tendency of the alloy to form solid-solutions is improved, indicating that the effect of the mixing of entropy on the solid solution formation strengthens. The VEC in the studied alloys is shown in Figure 7c. With increasing Ta and W addition, the value of VEC decreased from 6.28 to 6, which meets the BCC-forming requirement (6.87 ≤ VEC < 8), demonstrating that the BCC phase is stable in the VCrFeTa*x*W*x* alloy system.

**Figure 7.** (**a**) The curves of δ, Ω, and ΔSmix as a function of the Ta and W content for VCrFeTa*x*W*<sup>x</sup>* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys; (**b**) The relationship between parameters δ and ΔHmix for the as-cast VCrFeTa*x*W*x*; (**c**) The relationship between the valence-electron concentration (VEC), and Ta and W content of VCrFeTa*x*W*<sup>x</sup>* alloys; (**d**) The relationship between the parameters δ and Ω, for the as-cast VCrFeTa*x*W*x*. (SS: Solid-Solutions; I: Intermetallics compound; SS + I: Solid-Solutions + Intermetallics compound).


**Table 8.** ΔHmix, ΔSmix, Ω, δ, valence-electron concentration (VEC), the theoretical density, and the melting points for VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) alloys.

With the increase in the Ta and W content, although both ΔSmix and Ω (≥ 1.1) increased (Figure 7a), from Figure 7d, T0.3, T0.4, and T1 HEAs are in the SS + I (SS: Solid-solution, I: Intermetallics compound) region. The mixing of the enthalpy promotes the formation of intermetallic compounds, whereas the Ta–Fe binary system has the most negative mixing of enthalpy (−15 kJ/mol) listed in Table 5, indicating that the bonding force between these two elements is the strongest. The Fe2Ta phase formation occurs, which is consistent with the XRD patterns in Figure 1.

#### *4.2. Ta and W Effects at Room Temperature*

The microstructure analysis of the VCrFeTa*x*W*<sup>x</sup>* alloys with different Ta and W contents, performed in the current research, revealed several features. First, an increase in the Ta and W content substantially decreased the BCC1 matrix phase of the alloys (Figure 1). Conversely, the volume fraction of BCC2 and Laves phases increase. Second, an addition of Ta and W resulted in the formation of intermetallic phases, namely the Laves phase of Fe2Ta-type. The Laves phase can be associated with a highly negative-enthalpy of the intermetallic-phases formation [19,37] (Table 6). Third, with increasing Ta and W contents, the change of precipitation was the most obvious and most significant phenomenon. When *x* = 0.1, the Laves precipitation phase has a lamellate shape and forms an eutectic structure with the matrix. When *x* = 0.2, the Laves-precipitation phase is thicker, and still forms an eutectic structure with the matrix, and fine particles, most of which are the Laves phase precipitated on the matrix. When *x* = 0.3–1, the precipitation phase gradually aggregates and grows. These results indicate that the addition of Ta and W not only changes the phase composition of the alloy system, but also varies the microstructure of the alloy system. However, the addition of Ta promoted the regular Laves phase microstructure change of the BCC1 alloy system, and a transition from the hypoeutectic structure (*x* ≤ 0.4) to hypereutectic structure (*x* = 1) occurred. This indicates that Ta can promote the eutectic transformation of the Fe–Cr–V alloy system.

The Ta and W contents have a prominent effect on the mechanical properties of the VCrFeTa*x*W*x* alloys, as expected given the pronounced changes in the microstructure. With the addition of Ta and W, the hardness of the alloy increases linearly, while the plasticity and strength of the studied alloys exhibited a complex dependence. In the HEA solid-solution phase, various atoms randomly occupy the lattice position of the crystal. Each atom is surrounded by other kinds of atoms, and all atoms can be regarded as solute atoms or solvent atoms [37,53]. In addition, the types of atoms vary in size, causing severe lattice distortion in the solid solution, which in turn leads to high solid-solution strengthening [46,54,55], thereby increasing the strength and hardness of the alloy, especially the HEAs of the BCC structure [24,56–58]. Conversely, W obviously promotes the formation of BCC2, and Ta prefers forming Laves phases. The mixing enthalpy between Ta and Fe is −15 kJ/mol (Table 5), which is the lowest for the VCrFeTa*x*W*<sup>x</sup>* system, meaning that Ta and Fe combine together quite easily compared with the other element pairs during solidification owing to their great compatibility. The hardness values of the VCrFeTa*x*W*x* alloys increase with increasing *x*. As the content of Ta and W increased, the yield strength increased from T0.1 to T0.2. However, in the VCrFeTa*x*W*<sup>x</sup>* (*x* = 0.3, 0.4, and 1) alloys, the yield strength deteriorated quickly with an increase in *x* (Figure 4). This phenomenon can be associated with a large volume fraction of Laves phases, and the size of Laves particles increases

with the change in *x*. Notably, the strengthening contribution of the precipitation particles to the BCC phase alloys was prominent, which has been previously studied [59]. In the present study, the fine second phase particles of the Laves phase even had a positive effect on the strength, as the T0.2 alloy has an excellent yield strength and plasticity.

As engineering materials, the yield strength of HEAs is an important parameter for the design of a component. The dependence of the room-temperature compressive yield strength of the reported HEAs [13,24,29,56–72] and the studied alloys in the present work are plotted in Figure 8. The T0.1 and T0.2 HEAs possess relatively-high yield strength, whereas they all exhibit excellent compressive strains. T0.1 and T0.2 have a yield strength of 1341 MPa and 1742 MPa with plastic strains of 42.2% and 35.7%, respectively. They exhibit excellent plasticity, as they could not be broken in the compressive tests. The refractory HEAs reported in papers exhibit high yield strength, while their plastic strains are lower than 15%; or the reported refractory HEAs show remarkable plastic strains more than 50%, while their yield strengths are not remarkable. If we set the coordinates of the alloys in Figure 8 at (10%, 1600 MPa), only the T0.2 alloy is in the first quadrant. Therefore, we concluded that the BCC HEAs developed so far still have insufficient compressive strain at room temperature, or exhibit low strength. It is noteworthy that T0.2 studied in the current work not only possesses a compressive strain up to 35.7%, but also has a yield strength as high as 1742 MPa. In other words, T0.2 may have excellent ductility, which makes this kind of HEA better than other HEAs for engineering applications. More significantly, the T0.2 alloy, as a low-activation HEA, has obvious performance advantages over the previously-reported low-activation alloys, indicating that it has potential as a candidate material for fusion reactors.

**Figure 8.** The map of compressive yield strength and ductility combinations of various refractory high-entropy alloys [13,24,29,56–72] at room temperature. Initial strain rates range from 10−<sup>4</sup> to 10−<sup>3</sup> s<sup>−</sup>1.

#### *4.3. Heat-Softening Resistance*

Typically, annealing leads to alloy softening due to the rapid diffusion of atoms, resulting in internal stress relaxation caused by alloy defects at high temperatures [28,29]. However, due to the high mixing entropy and the obvious sluggish diffusion effect, HEAs have good high-temperature stability and great resistance to high temperature softening [28,35,37]. According to the Gibbs free energy formula, with the increase in temperature, the effect of entropy is more obvious than that of enthalpy [21]. The high mixing-entropy increases the solid solubility of the alloy, which is conducive to the formation of the solid-solution structure. Because of the sluggish diffusion effect, the alloy easily forms a supersaturated solid solution and a fine precipitated phase. Annealing at elevated temperatures

can partially release the internal stress and dissolve atoms. For the studied alloys in the present work, the increase in hardness after annealing from 600 ◦C to 800 ◦C is due to solid-solution strengthening. Although the microhardness of the T0.1 and T0.2 alloys decreased at 1000 ◦C, the reduction was not large compared to the hardness values at room temperature. For traditional alloys, tempering after quenching tends to result in significant softening. However, HEAs have significant advantages in this regard. HEAs, such as NbMoTaW [56], NbMoTaWV [56], and AlNbTiV [68], with a typical BCC structure have been reported to show that, although the alloys have poor compressive strains at room temperature, the plasticity of the alloys increases as the temperature rises. As shown in Figure 9, from room temperature to 800 ◦C, the yield strength of the T0.1 and T0.2 alloys are significantly higher than those of superalloys, such as Inconel 718 and Haynes 230. Although the strength of MoNbHfZrTi is higher than those of the T0.1 and T0.2 alloys at room temperature, its compression plasticity is only 10.1%, far lower than the T0.1 and T0.2 alloys. However, the strengths of T0.1 and T0.2 alloys exceed the reported high-entropy alloys include MoNbHfZrTi at 600–800 ◦C [36,68]. Thereby, the HEAs exhibit excellent resistance to high-temperature softening. As an alternative alloy for future fusion reactors with the low activation, the target alloys in the present work have certain advantages at high temperatures compared to CLAM, ODS, and V–4Cr–4Ti, as presented in Figure 9.

**Figure 9.** Temperature dependence of the compressive yield strength of superalloys [37], reported HEAs [36,68], and low activation alloys [13].

#### **5. Summary**

Novel low-activation HEAs VCrFeTa*x*W*x* (*x* = 0.1, 0.2, 0.3, 0.4, and 1) were fabricated by vacuum arc melting, and their microstructure and mechanical properties were studied in the present work. The investigated alloys exhibited a relatively simple microstructure and promising properties. Based on the obtained results and discussions, our conclusions are as follows:

(1) The microstructures of all investigated alloys exhibited a typical dendritic and eutectic structure, with VCrFeTa0.1W0.1 and VCrFeTa0.2W0.2 presenting a mainly BCC1 (a VCrFe-rich region) solid solution and Laves phases (an Fe2Ta-type). VCrFeTa0.3W0.3, VCrFeTa0.4W0.4, and VCrFeTaW contain BCC1 and BCC2 (W-rich region) solid solutions and Laves phases.

(2) The Vickers hardness of the alloys increased with increasing Ta and W contents. The hardness values of VCrFeTa0.1W0.1, VCrFeTa0.2W0.2, VCrFeTa0.3W0.3, VCrFeTa0.4W0.4, and VCrFeTaW are 546 HV0.2, 673 HV0.2, 726 HV0.2, 886 HV0.2, and 1135 HV0.2, respectively. This feature is attributed to the solid-solution strengthening and the increased amount of Laves and BCC2 phases.

(3) The augmented Ta and W contents increased the compressive strength but decreased the plastic strain of T0.1 and T0.2 alloys. The T0.1 and T0.2 alloys exhibited compressive yield strengths of 1341 MPa and 1742 MPa, with plastic strains of 42.2% and 35.7%, respectively. The solid-solution strengthening of the BCC matrix and the formation of hard Laves phases precipitated in particles are two main factors contributing to alloy strengthening. With the addition of Ta and W, the compressive performance deteriorated sharply due to the increase in the volume fraction and the growth of the Laves phase.

(4) By applying the atomic ratios strategy and the criteria for disordered solid solutions, the alloy system of VCrFeTa0.1W0.1, VCrFeTa0.2W0.2, VCrFeTa0.3W0.3, VCrFeTa0.4W0.4, and VCrFeTaW fully met the VEC, δ-Ω, and δ-Hmix criteria. Studies have shown that HEAs maintain stable phase structures and properties after high-temperature annealing. This means that the design of high-temperature structural materials using the HEA concept could promote the development of these alloys for use in extreme high-temperature extreme such as blades, engines, aerospace, fusion reactors and other applications.

(5) The high-temperature mechanical properties of T0.1 and T0.2 alloys were examined. After annealing 25 h at 600–1000 ◦C, the T0.1 and T0.2 alloys maintained high hardness. The compressive yield strengths of the T0.1 and T0.2 alloys are promising with heat-softening resistance at 600–800 ◦C. The yield strengths of T0.1 and T0.2 alloys were much higher than 1000 MPa at 800 ◦C. This study provides guidance for further development of the current concepts to produce refractory high-temperature structural materials, which are candidate alloys for fusion reactors.

**Author Contributions:** W.Z. prepared the high-entropy alloys and finished the manuscript. Y.Z. and P.K.L. offered the theoretical guidance. All authors contributed to the general discussion.

**Funding:** Y.Z. would like to thank the financial supports from the National Science Foundation of China (NSFC, Granted Nos. 51471025 and 51671020). P.K.L. would like to acknowledge the National Science Foundation (DMR-1611180 and 1809640), and Department of Energy (DOE) Office of Fossil Energy, National Energy Technology Laboratory (NETL) (DE-FE-0011194), and the U.S. Army Office Project (W911NF-13-1-0438).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Magnetic Properties and Microstructure of FeCoNi(CuAl)0.8Sn***x* **(0** *<sup>≤</sup> <sup>x</sup> <sup>≤</sup>* **0.10) High-Entropy Alloys**

#### **Zhong Li 1, Chenxu Wang 1, Linye Yu 1,2, Yong Gu 1,3, Minxiang Pan 1, Xiaohua Tan 1,\* and Hui Xu 1,\***


Received: 9 October 2018; Accepted: 8 November 2018; Published: 13 November 2018

**Abstract:** The present work exhibits the effects of Sn addition on the magnetic properties and microstructure of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) high-entropy alloys (HEAs). The results show all the samples consist of a mixed structure of face-centered-cubic (FCC) phase and body-centered-cubic (BCC) phase. The addition of Sn promotes the formation of BCC phase, and it also affects the shape of Cu-rich nano-precipitates in BCC matrix. It also shows that the Curie temperatures (*Tc*) of the FCC phase and the saturation magnetization (*Ms*) of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs increase greatly while the remanence (*Br*) decreases after the addition of Sn into FeCoNi(CuAl)0.8 HEA. The thermomagnetic curves indicate that the phases of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs will transform from FCC with low *Tc* to BCC phase with high *Tc* at temperature of 600–700 K. This work provides a new idea for FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs for their potential application as soft magnets to be used at high temperatures.

**Keywords:** high-entropy alloys (HEAs); phase constitution; magnetic properties; Curie temperature; phase transition

#### **1. Introduction**

Since the first report of high-entropy alloys (HEAs) in 2004 [1], researchers have shown an increased interest in the study of HEAs. HEAs are the definition of alloys that are typically composed of more than 5 principal elements, which have broken the traditional alloy design concept based on 1 or 2 principal elements [2]. In contrast with the conventional alloys, HEAs predominantly trend to form an amorphous structure [3,4] or a simple solid solution with body-centered-cubic (BCC) phase [5,6], face-centered-cubic (FCC) phase [7–9] or a mixture of them [10,11], which is attributed to the high mixing entropy of HEAs [12,13]. The unique design concept and the significant mixing entropy effect of HEAs give them potential application in many high-entropy structural and functional materials. For example, HEAs have huge potential for use in jet-engine turbines, thin-film resistors, heat- or wear-resistant parts, functional coatings, and electronic products [13,14]. Recently, the good magnetic properties of HEAs capture the increasing interest in this new field of materials [15,16]. It is worth noting that many HEAs [17–20] consist of several ferromagnetic elements, such as Fe, Co, and Ni, and they also have a good comprehensive mechanical properties, which make them have great application potential in soft magnetic materials. HEAs with good magnetic and mechanical properties are expected

to be used in electric motors, electromagnets, and magnetic recording. Liu et al. [21,22] reported the FeCoNi0.25Al0.25 HEA exhibits a high saturation magnetization (*Ms* = 101.0 emu/g) and a low coercivity (*Hc* = 268 A/m). Zuo et al. [23] found CoNiMnGa HEA shows a low saturation magnetostriction coefficient and a high Curie temperature (*Tc*). In our previous work [24], the FeCoNi(CuAl)0.8 HEA consisting of BCC and FCC phases shows good magnetic and mechanical properties, and it was found that BCC phases show a higher *Ms* than that of FCC phases for the FeCoNi(CuAl)0.8 HEA. The other work of our group [25] found a minor amount of Ga addition into FeCoNi(CuAl)0.8 HEA can promote the formation of BCC phase and improve the *Ms* of the alloy, and the value of remanence (*Br*) and coercivity (*Hc*) also increases. It was reported that the addition of Sn can hinder the formation of FCC phase [26] and promote the formation of BCC phase [27]. Therefore, in this work, a minor amount of Sn was added into the FeCoNi(CuAl)0.8 HEA hoping to get a higher volume fraction of BCC phase.

In this work, the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs were studied from phase constitutions to microstructure and magnetic properties. It was found that these HEAs show high *Ms*, low *Br*, and high *Tc*, which indicate their potential application as soft magnetic materials. This paper offers a good method for designing future high-performance soft magnetic materials.

#### **2. Materials and Methods**

The FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs were prepared via arc-melting the constituent elements of 99.99% purity using a water-cooled Cu crucible. The alloys were sucked into a 100 × 10× 2 mm water-cooled Cu mold after remelting four times. X-ray diffraction (XRD, D/max-2500 V, Rigaky Corporation, Tokyo, Japan) was used to characterized the crystal structures of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs at a scan speed of 1◦/min. Scanning electronic microscopy (SEM, Hitachis-3400N) was used to observe the morphology of the samples. Transmission electron microscope (TEM, JEM-2100F, JEOL, Ltd., Tokyo, Japan) was employed for the microstructure of HEAs. Dual-beam focused ion beam (FIB, FEI Helios 600i, Hillsboro, OR, USA) was used to prepare the TEM samples. The high angle annular dark field (HAADF) images were performed by a scanning transmission electron microscope with energy dispersive spectrometer (STEM/EDS, JEM-2100F, JEOL, Ltd., Tokyo, Japan). The *Ms* and thermomagnetic curves was obtained from vibrating sample magnetometer (VSM, Lakeshore 7407, Westerville, OH, USA). The coercivity (*Hc*), hysteresis losses (*Pu*), remanence (*Br*), initial permeability (*μi*), and maximum permeability (*μmax*) were obtained from hysteresis curves (DC) test system (HCTS, FE-2100SD, Forever elegance, Hunan, China) using a 27 × 9 × 1.7 mm rectangular sample at a magnetic field of 25 kA/m. The thermal stability was analyzed by differential scanning calorimeter (DSC, DIAMOND) at a heating rate of 10 K/min.

#### **3. Results and Discussion**

#### *3.1. X-ray Diffraction*

Figure 1 shows the XRD patterns of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs. It is found that all of these HEAs consist of a mixed structure of FCC phase and BCC phase. A few diffraction peaks of unknown phases appear in the XRD patterns for x ≥ 0.02 and it is especially obvious for x ≥ 0.06. Here, I(111)FCC and I(110)BCC were used to denote the diffraction intensity of the strongest peak of (111) for FCC phase and (110) for BCC phase, respectively. Therefore, the relative content of FCC and BCC phases can be expressed as I(110)BCC/I(111)FCC. Table 1 shows the ratio of I(110)BCC/I(111)FCC. For x=0, the ratio of I(110)BCC/I(111)FCC is 0.38. This means the FCC phase in the FeCoNi(CuAl)0.8 HEA is a dominant phase. When the content of Sn increases from 0.02 to 0.10, the ratio of I(110)BCC/I(111)FCC increases from 0.69 to 10.53, suggesting that the content of BCC phase increases rapidly. Based on XRD patterns, the lattice parameters of FCC and BCC phases can be calculated and shown in Table 1. It is seen that the lattice parameters of FCC and BCC phases all increase as *x* increases from 0 to 0.04, then they decrease and remain almost stable as *x* further increases. The decrease of lattice parameters of FCC and BCC phases may be due to the precipitation of these unknown phases.

**Figure 1.** XRD patterns of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) high-entropy alloys (HEAs).

**Table 1.** The ratio of I(110)BCC/I(111)FCC and lattice parameters of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs.


#### *3.2. Magnetic Properties*

Figure 2 shows hysteresis loops of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs measured by VSM. Figure 3 shows the magnetization curves and hysteresis loops of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs measured by HCTS. The corresponding magnetic parameters obtained from Figures 2 and 3 are shown in Table 2. As can be seen from Table 2, the value of *Ms* increases from 78.6 Am2/kg to 88.8 Am2/kg as *x* increases from 0 to 0.10, which increases almost 13 percent. To make it easier to compare, the corresponding magnetic parameters as well as the ratio of I(110)BCC/I(111)FCC as a function of *x* are shown in Figure 4. It is obvious that the *Ms*, coercivity (*Hc*), and hysteresis losses (*Pu*) increase while the initial permeability (*μi*) and maximum permeability (*μmax*) have a generally decreasing trend with increasing *x*. These results may be due to the increase of the volume fraction of BCC phase and the decrease of the volume fraction of FCC phase, which is in agreement with that of reported FeCoNi(CuAl)0.8Ga*<sup>x</sup>* (0 ≤ *x* ≤ 0.08) HEAs [25]. However, the value of remanence (*Br*) decreases after the addition of Sn into FeCoNi(CuAl)0.8 HEA, and this is completely different from the effect of Ga, which contributes to the increase of *Br* [25]. A careful analysis of these magnetic parameters shows the decrease of *Br* may be due to the rapid decrease of permeability.

**Figure 2.** Hysteresis loops of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs measured by VSM.

**Figure 3.** Magnetization curves and hysteresis loops of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs measured by HCTS.

**Figure 4.** The ratio of I(110)BCC/I(111)FCC and magnetic properties as a function of *x* for FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs.

**Table 2.** Magnetic parameters of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs.


The temperature dependence of magnetization for FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs measured at an applied magnetic field of 1 T is shown in Figure 5. It can be seen that the magnetization of the FeCoNi(CuAl)0.8 HEA decreases first with increasing temperature, and it has hit bottom of 28.9 Am2/kg when the temperature is 630.4 K, then it increases quickly and reaches a constant value of about 52 Am2/kg at 706 K. It is worth noting that the magnetization does not reduce to 0 at its lowest point. That means a magnetic phase with higher Curie temperature exists in the alloy. As can be seen from Figure S1, there is one phase transformation peak [25] for FeCoNi(CuAl)0.8 HEA, at which the phase transforms from FCC to BCC phase. Therefore, it can be concluded that the Curie temperatures of BCC phase are obvious higher than that of FCC for FeCoNi(CuAl)0.8 HEA. That means the FCC phase exhibits paramagnetic behavior and a disordered magnetic structure while the BCC phase still shows ferromagnetic behavior with the increase of temperature. It also agrees well with our previous study [25]. Therefore, the magnetization of the alloy will increase when the temperature is above 630.4 K and below 706 K. The Curie temperature (*Tc*), corresponding to the ferromagnetic to paramagnetic state transition of FCC phase for FeCoNi(CuAl)0.8 HEA, is indicated by arrow in Figure 5. For ease of comparison, the Curie temperatures of FCC phase for other FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0.02 ≤ *x* ≤ 0.10) HEAs are also shown in Figure 5. For *x* = 0.04, the outline of the curve is similar to that of the Sn-free alloy, but the Curie temperature of FCC phase increases significantly and reaches 634.9 K. For *x* ≥ 0.08, the Curie temperature of FCC phase continues to increase. However, there is only little change for the value of magnetization as the temperature continues to increase. This is because only a small number of FCC phase transform to BCC phase according to the results of XRD in Figure 1 and DSC curves in Figure S1. Therefore, it can be concluded that the addition of a minor amount of Sn can obviously increase the Curie temperature of FCC phase. The phases of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs will transform from FCC with low *Tc* to BCC phase with high *Tc* at temperature of 600–700 K, which leads to the increase of magnetization. This provides a new idea for FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs for their potential application as soft magnets to be used at high temperature.

**Figure 5.** The thermomagnetic curves of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs.

#### *3.3. Microstructure*

The SEM backscattered-electron (SEM-BSE) microstructures of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs are displayed in Figure 6. In Figure 6a, two obviously identifiable contrasts are found in the Sn-free alloy, which can be identified as dendritic regions and interdendritic regions (marked as DR and IR, respectively). According to our previous studies [24,25], the DR and IR region can be confirmed to be FCC and BCC phase, respectively. For *x* ≥ 0.02, regions with strong contrast appear between the DR and IR regions, namely, the phase boundary regions, which can be marked as PB region. Moreover, the volume fraction of DR gradually decreases while the volume fraction of IR and PB increase with the increase of *x*. In addition, it is worth noting that the DR phases are almost invisible in the FeCoNi(CuAl)0.8Sn0.10 HEA (Figure 6f). The above results are in good agreement with the XRD results. Therefore, we can conclude that the addition of Sn in FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs can promote the formation of BCC phases and PB phases.

**Figure 6.** Typical SEM-BSE images of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs.

In order to get more details of the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs, the structures of the FeCoNi(CuAl)0.8 and FeCoNi(CuAl)0.8Sn0.10 HEAs were further analyzed by TEM and the results are shown in Figure 7. Figure 7a1 confirms that FeCoNi(CuAl)0.8 HEA consists of two kinds of phases which are named as DR (dendritic region) and IR (interdendritic region) according to the results revealed by the typical SEM-BSM images (Figure 6). The selected-area-electron-diffraction (SAED) patterns of Figure 7a2,a3 suggest the FCC crystal structure of DR and the BCC crystal structure of IR in FeCoNi(CuAl)0.8 HEA. Meanwhile, it can be seen from Figure 7a1 that the surface of FCC phase is very smooth while the surface of BCC phase is much harsh. The high-magnification bright-field image of BCC phase for FeCoNi(CuAl)0.8 HEA is shown in Figure 7a4 and displays that the BCC phase contains a large number of nanoscale precipitates which distribute homogeneously in the BCC matrix. Moreover, the average size of nanoscale precipitates is 20 ± 5 nm. After the addition of Sn into the FeCoNi(CuAl)0.8 HEA, DR and IR regions can also be found in the FeCoNi(CuAl)0.8Sn0.10 HEA in Figure 7b1. Similarly, the former one can be indexed as FCC phase while the latter one is BCC phase according to the diffraction calibration in Figures 7 and 7. In addition, it sees that two new regions (marked as A and B, respectively) are observed in Figure 7b1. We can infer that they are the source of the unknown phase peak in the XRD pattern. As shown in Figure 7b4, the high-magnification bright-field image of BCC phase for FeCoNi(CuAl)0.8Sn0.10 HEA is apparently different from that of the Sn-free alloy. The shape of the nano-precipitates is rod-like, and their density is lower than that in FeCoNi(CuAl)0.8 HEA. At the same time, the nanoprecipitates with an average length of about 100 nm and a width of about 30 nm are larger than that in FeCoNi(CuAl)0.8 HEA.

**Figure 7.** TEM images of FeCoNi(CuAl)0.8 HEA: (**a1**) bright-field image; (**a2**) SAED pattern of DR region; (**a3**) SAED pattern of IR region; (**a4**) high-magnification bright-field image of IR region. TEM images of FeCoNi(CuAl)0.8Sn0.10 HEA: (**b1**) bright-field image; (**b2**) SAED pattern of DR region; (**b3**) SAED pattern of IR region; (**b4**) high-magnification bright-field image of IR region.

The high angle annular dark field (HAADF) image and element mappings of Fe, Co, Ni, Cu, and Al for FeNiCo(CuAl)0.8 HEA measured by STEM-EDS technique are displayed in Figure 8. As can be seen from Figure 8a, the microstructure of the FeNiCo(CuAl)0.8 HEA is similar to that shown in Figure 7a1. It is worth noting that a phase boundary region with a width of about 13 nm can also be found in Figure 8a. We can confirm that these nanoprecipitates in the BCC region and the phase boundary region are rich in Cu (Figure 8e). Figure 8b–f suggests that the distribution of Fe and Co is very uniform in the alloy, while Ni and Al are enriched in the BCC region. The formation of the Cu-rich phase boundary regions in the as-cast FeNiCo(CuAl)0.8 HEA may be caused by the following three reasons. First, the melting point of Cu is lower than that of Fe, Co, and Ni [28], thus it may solidify after Fe, Co, and Ni when the temperature decreases. Second, the mixing enthalpies [28] (Table S1) between Cu and Fe, Co, and Ni are 13, 6, and 4 kJ/mol, respectively, meaning Cu is more likely to be repelled by other elements to form the Cu-rich phase boundary region. Finally, and most importantly, Cu and other elements are completely soluble at high temperature, but a large amount of Cu will precipitate out due to the rapid decrease of Cu solubility during casting. As for the formation of Cu-rich nano-precipitate in the BCC phase, it is due to the great difference in crystal structure between copper and the BCC phase matrix as well as the decrease of Cu solubility in BCC phase and the positive mixing enthalpies of Cu with Fe, Co, Ni, and Sn.

**Figure 8.** High angle annular dark field (HAADF) image (**a**) and elemental mapping images of FeCoNi(CuAl)0.8 HEA for Fe-Kα (**b**), Co-Kα (**c**), Ni-Kα (**d**) Cu-Kα (**e**) and Al-Kα (**f**).

Figure 9 shows the HAADF image and elemental mapping images of FeCoNi(CuAl)0.8Sn0.10 HEA. Compared with that of FeCoNi(CuAl)0.8 HEA, the microstructure shown in Figure 9a is more complex. From Figure 9a, a large number of Cu-rich nanoprecipitates with rod-like shape can be found in BCC phase. The shape of the precipitates changes to spherical as they approach the phase boundary region, and their size also gradually decreases. Figure 9b–g suggests that the distribution of Fe and Co is uniform in FCC and BCC phase regions, however, little Fe and Co can be found in the region between FCC and BCC phase. Moreover, the distribution of Ni and Al is enriched in BCC regions. One of the interesting things is that Cu and Sn segregate in the region between FCC and BCC phase. Besides, in the region where Sn is enriched, the content of Ni is also very high. Therefore, we can deduce that the unknown phase shown in the XRD patterns is composed of two phases, and one of it is rich in Cu, the other one is rich in Ni and Sn.

**Figure 9.** HAADF image (**a**) and elemental mapping images of FeCoNi(CuAl)0.8Sn0.10 HEA for Fe-Kα (**b**), Co-Kα (**c**), Ni-Kα (**d**) Cu-Kα (**e**), Al-Kα (**f**) and Sn-Kα (**g**).

#### **4. Conclusions**

In this work, the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs were prepared by vacuum arc-melt casting. Effects of Sn content on the phase constitution and magnetic properties were studied. All the samples are composed of FCC and BCC phases, whereas some unknown phases appear with the addition of Sn. The addition of Sn promotes the formation of BCC phase, and it also affects the shape of Cu-rich nanoprecipitates in the BCC matrix. Moreover, the *Ms* increases greatly while the remanence (*Br*) decreases with the increasing of *x* for FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs. The addition of Sn can obviously increase the Curie temperature of the FCC phase. The phase of alloys with a mixture of FCC and BCC will transform from FCC to BCC phase at high temperature, leading to an increase of magnetization. They can be used as new soft magnetic materials at high temperatures.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/1099-4300/20/11/872/ s1, Figure S1: DSC curves of FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs. Table S1: The melting points (K) [28] of different elements and the mixing enthalpies (kJ/mol) [29] between two elements.

**Author Contributions:** H.X. and X.T. conceived and designed the experiments; L.Y. and Z.L. prepared the FeCoNi(CuAl)0.8Sn*<sup>x</sup>* (0 ≤ *x* ≤ 0.10) HEAs; L.Y. and Z.L. performed the microstructural characterization of the samples; Z.L. and C.W. prepared the original draft; Z.L., X.T. and H.X. revised the manuscript; All authors discussed, analyzed the results and approved the final manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (grant number 51471101 and U1531120).

**Acknowledgments:** The authors thank Jianchao Peng, Pengfei Hu, Na Min and Xue Liang of the Instrumental Analysis & Research Center, Shanghai University, China for their assistance in TEM samples preparation and measurements.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Effect of Co and Gd Additions on Microstructures and Properties of FeSiBAlNi High Entropy Alloys**

#### **Sicheng Zhai, Wen Wang, Juan Xu, Shuai Xu, Zitang Zhang and Yan Wang \***

School of Materials Science and Engineering, University of Jinan, Jinan 250022, China; zscshr123321@163.com (S.Z.); ww\_nwc@163.com (W.W.); xujuanMISS@163.com (J.X.); 17862915273@163.com (S.X.); 13796829730@163.com (Z.Z.)

**\*** Correspondence: mse\_wangy@ujn.edu.cn; Tel: +86-531-8276-5473; Fax: +86-531-8797-4453

Received: 10 April 2018; Accepted: 19 June 2018; Published: 22 June 2018

**Abstract:** FeSiBAlNi (W5), FeSiBAlNiCo (W6-Co), and FeSiBAlNiGd (W6-Gd) high entropy alloys (HEAs) were prepared using a copper-mold casting method. Effects of Co and Gd additions combined with subsequent annealing on microstructures and magnetism were investigated. The as-cast W5 consists of BCC solid solution and FeSi-rich phase. The Gd addition induces the formation of body-centered cubic (BCC) and face-centered cubic (FCC) solid solutions for W6-Gd HEAs. Whereas, the as-cast W6-Co is composed of the FeSi-rich phase. During annealing, no new phases arise in the W6-Co HEA, indicating a good phase stability. The as-cast W5 has the highest hardness (1210 HV), which is mainly attributed to the strengthening effect of FeSi-rich phase evenly distributed in the solid solution matrix. The tested FeSiBAlNi-based HEAs possess soft magnetism. The saturated magnetization and remanence ratio of W6-Gd are distinctly enhanced from 10.93 emu/g to 62.78 emu/g and from 1.44% to 15.50% after the annealing treatment, respectively. The good magnetism of the as-annealed W6-Gd can be ascribed to the formation of Gd-oxides.

**Keywords:** high entropy alloys; elemental addition; annealing treatment; magnetic property; microhardness

#### **1. Introduction**

Recently, a new concept was proposed for high entropy alloys (HEAs), which has aroused wide attention and interest [1–3]. Generally, HEAs with equiatomic or near-equiatomic alloying elements mainly consist of face-centered cubic (FCC), body-centered cubic (BCC), or hexagonal closed-packed (HCP) solid solutions, and some intermetallic or amorphous phases. Owing to the special phase structure, HEAs usually possess excellent mechanical properties [4,5] and corrosion resistance [6], especially magnetic properties [7–9]. Several studies have reported that additions of certain elements into HEAs could induce the transformation of crystalline structures and further affect the related properties of HEAs [9–11]. The addition of Al, Ga, and Sn to the CoFeMnNi HEA induced the phase transition from FCC to ordered BCC phases, and further led to the significant improvement of the saturation magnetization (Ms) [9]. The microstructural evolution of (FeCoNiCrMn)100-xAlx HEA system transformed from the initial single FCC structure to final single BCC structure as Al concentration increased from 0 to 20 at. % [10]. Both the tensile fracture and yield strength were enhanced with increasing Al concentration. The HEAs phases are metastable in thermodynamics, therefore, would transform to the stable microstructure after subsequent annealing, which obviously affects the properties of HEAs to some degree [3,8,12–15]. Annealing under a given condition can lead to the occurrence of phase transition from a FCC to BCC phase for FeCoNi(CuAl)0.8 HEAs, resulting in a substantial increase in the Ms from 78.9 Am2/kg to 93.1 Am2/kg [15].

Recently, our group has prepared a series of as-milled FeSiBAlNi-based HEA powders using a mechanical alloying (MA) process [11,14,16]; these displayed an interesting microstructural evolution and magnetic properties. In the present study, the equiatomic FeSiBAlNiM (M = Co, Gd) HEAs were fabricated by a copper-mold spray casting technique. The effects of Co and Gd additions and subsequent annealing treatment on the microstructures, microhardness, and magnetism of the FeSiBAlNi HEAs were systematically investigated.

#### **2. Experimental**

The ingots of FeSiBAlNi, FeSiBAlNiCo, and FeSiBAlNiGd HEAs (denoted as W5, W6-Co, and W6-Gd, respectively) were prepared by an arc melting technique. The melting of these ingots was repeated at least five times to ensure the composition homogeneity in a Ti-gettered high-purity argon atmosphere. Then the ingots were remelted and made into 8 mm diameter rods by copper mold spray casting in an argon atmosphere. Then, the rods were annealed at given temperatures for two hours and cooled inside the furnace in the argon atmosphere. The annealing temperatures were set as two segments denoted as TI and TII in a low and high temperature region, respectively. There are 600 and 1000 ◦C for W5 HEA, 600 and 1000 ◦C for W6-Co HEA, and 650 and 1050 ◦C for W6-Gd HEA. The relatively high annealing temperatures selected for W6-Gd are attributed to the higher melting point (Tm) than that of the other two samples, as shown in the differential scanning calorimetry (DSC) curves.

Microstructural characterization of the as-cast and as-annealed HEAs were conducted by X-ray diffraction (XRD, Rigaku D8 Advance, Bruker, Germany) using Cu Kα radiation, field emission scanning electron microscopy (FESEM, QUANTA FEG 250 operated at 15 kV, Japan) coupled with energy dispersive spectrometry (EDS). The working distance used in SEM measurements was less than 10 mm. The thermal properties were analyzed by differential scanning calorimetry (DSC, TGA/DSC1, Mettler-Toledo, Greifensee, Switzerland) used under a continuous flow (30 mL/min) high-purity argon atmosphere at a heating rate of 10 K/min scanned from room temperature to 1400 ◦C. Microhardness of the tested HEAs was determined by a Vickers hardness tester (HV-10B), with a load of 200 g and a duration time of 15 s. The HV measurement for every tested sample was repeated ten times in order to obtain the average values. The coercive force (Hc), Ms, and remanence ratio (Mr/Ms, Mr: remanence) were determined by an alternating gradient magnetometer (AGM) at room temperature with a maximum applied field of 14000 Oe.

#### **3. Results and Discussion**

The XRD patterns of the as-cast W5, W6-Co, and W6-Gd HEAs are shown in Figure 1a. The as-cast W5 HEA consists of BCC1 (a = 4.475 Å) solid solution and FeSi-rich phase. The XRD pattern of the as-cast W6-Co HEA mainly displays the FeSi-rich phase of solution of other principal elements. In addition, other phase peaks may overlap with FeSi-rich phase peaks. Compared with the W5 and W6-Co HEAs, the effect of Gd addition on phase composition presents an obvious difference. The as-cast products of W6-Gd HEA exhibit the formation of new BCC2 (a = 4.484 Å) and FCC solid solutions. However, the FeSi-rich phase doesn't appear. The phase products of the as-cast W5, W6-Co, and W6-Gd HEAs are presented in Table 1.

Figure 1b shows the DSC curves of the as-cast HEAs. The Tm value of W6-Co HEA is 1129 ◦C, which is lower than that of W5 (1152 ◦C). However, the Gd addition increases the value of Tm, which reaches 1185 ◦C.

XUH **Figure 1.** XRD patterns (**a**) and differential scanning calorimetry (DSC) curves (**b**) of the as-cast W5, W6-Co, and W6-Gd HEAs.

**Table 1.** Phase products of as-cast and as-annealed W5, W6-Co, and W6-Gd high entropy alloys (HEAs) at TI and TII, identified distinctly from XRD patterns.


To further investigate the difference in morphologies and compositions caused by the Co and Gd additions, FESEM coupled with EDS analysis was carried out and is presented in Figure 2 and Table 2. As shown in Figure 2a, a larger number of polygonous light-grey phases distribute dispersedly in the matrix, as well as the irregular black phases with small size. The inset of Figure 2a-1 reveals one rhombic grain with edge sizes less than 8 μm. It needs to be noted that metalloid B as a light element can't be accurately measured. Moreover, the B content of some samples is very small in most regions, therefore, they are omitted in the present study. According to the EDS results, matrix (A) contains more Al and Ni elements, and a certain amount of Fe and Si elements. Furthermore, there is a partial FeSi-rich phase in region (A) and the black region (B) is enriched with Fe and Si elements. This indicates that the FeSi-rich phase mainly exists in region (B), presenting an evenly distribution in the matrix. The rhombic grain (C) is mainly composed of Fe and B elements (instead, region (C) contains more B element above 10 at. %.). The Co addition induces the refinement of the precipitated grains (Figure 2b). The inset in Figure 2b-1 shows that the larger number of dark regions (D) with the smaller size exhibit a uniform distribution state and enrich Fe and Si elements. The gray region (E) is rich in the Ni element, and the bright-grey region (F) is poor in the Al element. Moreover, the component ratio of Fe and Si in regions (E) and (F) is close to 1:1. Although several phases appear in the SEM images, no peaks other than the FeSi-rich phase can be seen in the XRD results of as-cast W6-Co HEA (Figure 1). It suggests that the precipitates probably have a similar lattice constant and crystal structure concerning the matrix [17]. Moreover, it could be a complex compositional fluctuation in the as-cast W6-Co HEA [18]. Figure 2c and inset (Figure 2c-1) present that the as-cast W6-Gd HEA consists of coarse rod-like dendrites as the FCC-matrix phase. They are rich in each principal element with near-equiatomic ratio except for Al element (region (G)). However, the Al element segregates in the interdendritic grains ((H): dark- and deep-grey regions) corresponding to the BCC2 phase with less contents. Usually the precipitation pathways in HEAs can be very complex, and it is a particularly challenging topic, which remains to be studied [19].

XUH **Figure 2.** Field emission scanning electron microscopy (FESEM) micrographs of the as-cast high entropy alloys (HEAs): (**a**) W5; (**b**) W6-Co; and (**c**) W6-Gd. The insets of (a-1), (b-1), and (c-1) are the partial enlargements corresponding to (**a**), (**b**), and (**c**), respectively.


**Table 2.** Chemical compositions of representative regions for W5, W6-Co, and W6-Gd HEAs obtained by energy dispersive spectrometry (EDS).

Figure 3 shows the XRD patterns of the as-annealed W5, W6-Co, and W6-Gd HEAs at different temperatures; their annealing products are also listed in Table 1. After annealing at TI, the annealed products of W5 HEAs consist of a new BCC3 (a = 4.033 Å) solid solution with a FeSi-rich phase. However, the contents of the FeSi-rich phase obviously decrease compared to that of the as-cast state. Moreover, the BCC1 solid solution disappears (Figure 3a). Via annealing at TII the two obtained phases still exist, but the diffraction peak intensity becomes strong, indicating the further growth and

coarsening of the grains. Being distinct from the W5 HEA, no new phase transformation occurs in the W6-Co HEA after annealing at TI and TII, indicating that the W6-Co HEA possesses a good thermal stability (Figure 3b). It suggests that the Co addition leads to the transformation from the metastable characteristic of W5 HEA to a more stable state in thermodynamics. However, the inset presents that the main diffraction peak of the W6-Co HEA shifts to the lower angle with the increased annealing temperature, suggesting a serious lattice distortion caused by the expansion of the lattice. Figure 3c reveals the formation of new phases of AlNi, AlGd, Gd-oxides, besides the primary BCC2 and FCC solid solutions for the as-annealed W6-Gd HEA at TI. The as-annealed products are unchanged at TII, except for the increased phase amounts of Gd-oxides. Moreover, compared with the W5 and W6-Co HEA, the highest Tm value of the W6-Gd HEA may be attributed to the high Tm values of precipitated intermetallic compounds of AlNi (1638 ◦C) and AlGd (1200 ◦C).

**Figure 3.** XRD patterns of the as-annealed HEAs at different temperatures (TI and TII): (**a**) W5, (**b**) W6-Co, and (**c**) W6-Gd.

Figure 4 shows the Vickers hardness (HV) of the as-cast and as-annealed HEAs. The W5 HEA displays the highest HV among the tested HEAs, and the as-cast W5 HEA possesses the highest hardness of 1210 HV. The additions of Co and Gd cause the decline of HV values, and the as-annealed W6-Gd HEA (TII) displays the maximal decline of HV (738). It suggest that the annealing treatment plays a negative effect on the HV of the as-cast samples, which is in agreement with Salishchec's results [20]. With the increased annealing temperature, the internal stress of the as-cast HEAs gradually decreases as well as the microstructural coarsening. The effect of solid solution strengthening became the smaller, and strain softening was revealed in HEAs [21]. Compared with W6-Co and W6-Gd HEAs, the FeSi-rich phase, as the second strengthening phase in the W5 HEAs (especially the as-cast one) evenly distributes in the BCC solid solution matrix, which can contribute to the high HV values.

ĉ Ċ

**Figure 4.** Vickers hardness of the as-cast and as-annealed W5, W6-Co, and W6-Gd HEAs.

The mass magnetization (M) as a function of the magnetic field intensity (H) for the as-cast and as-annealed samples was tested. The Hc, Ms, and Mr/Ms of these HEAs are shown in Figure 5. All Hc values of the tested HEAs are in the range from 10 to 180 Oe (Figure 5a), indicating the soft magnetism nature of these HEAs. It suggests that the annealing treatment induces a weak decrease of Hc values for the W5 and W6-Co HEAs, but the Hc of W6-Gd HEA becomes large after annealing. The Hc is mainly affected by impurity, deformation, crystallite size, and stress, and the subsequent heat-treatment process [22]. Therefore, the Hc values of as-annealed W5 and W6-Co HEAs are slightly lower than those of the as-cast samples, suggesting that the former possess a little larger average crystallite size according to the well-known coercivity-crystal size relationship [8]. Moreover, the origin of the lower Hc can be attributed to the low number density of domain-wall pinning sites [23]. The as-annealed products of W6-Gd HEA contain complex phase compositions, and display the inhomogeneous characteristics, which obviously work towards the pinning effect of domain wall movement. Therefore the Hc values can be enhanced after the annealing treatment.

**Figure 5.** Magnetic properties of the as-cast and as-annealed W5, W6-Co, and W6-Gd HEAs: Hc (**a**), Ms (**b**) and Mr/Ms (**c**).

The variations of Ms are exhibited in Figure 5b. In the as-cast state, there is no distinct difference in Ms for all the tested samples, and W5 HEA emerges with a slightly higher Ms of 12.91 emu/g. After annealing at TI, the Ms of W5 HEA remains nearly unchanged, whereas declines with a reduction of 27.7% at TII. There is no obvious magnetism changes revealed for the as-cast and as-annealed W6-Co HEAs, and the Ms values become stabilized at about 11 emu/g. This stability of Ms is resulted from the stable phase characteristic of W6-Co HEA in the annealing stage. Unlike W5 and W6-Co HEAs, the magnetism of W6-Gd HEA is enhanced during annealing treatment. The Ms value of the as-annealed W6-Gd HEA increased from 10.93 emu/g to 31.91 emu/g at TI, and further up to 62.78 emu/g at TII, suggesting increased soft magnetic properties.

From Figure 5c, the Mr/Ms values of as-cast W5 and W6-Co HEAs are similar to their as-annealed states, which depend on their similar phase compositions. The as-annealed products of W6-Gd HEAs are significantly different from the as-cast one, and the Mr/Ms values are enhanced from 1.44% (as-cast) to 15.5% (at TI). Moreover, the as-annealed W6-Gd HEAs reveal the highest Mr/Ms values among the tested samples, indicating a better soft magnetism.

Residual stress exists in the as-cast HEAs which can deteriorate the soft magnetic properties. Appropriate heat treatment can induce stress relief, which is beneficial to improve the soft magnetic properties [24]. Therefore, except for the as-annealed W5 HEA at TII, the soft magnetic properties of the tested HEAs are properly improved by the structural relaxation through stress-relief annealing [25]. Notably, it suggests that Ms of the as-annealed W6-Gd HEA at TII is about five times higher than that obtained for the as-cast one. Moreover, the magnetic properties are strongly dependent on the microstructure of the materials. The microstructure contribution to magnetism arises from morphology: properties such as magnetic anisotropy, magnetostriction, coercivity, and volume fraction of the precipitates. The decrease in Ms for as-annealed W5 at TII can be related to the enhanced density of grain boundaries and the increase of volumetric fraction of BCC3 solid solutions around FeSi-rich phases, which reduce the magnetic moment. According to the effect of phase compositions on magnetic properties, the increase in Ms for the as-annealed W6-Gd HEA can be ascribed to the formation of Gd-oxides. Moreover, Ms is enhanced by increasing the contents of Gd-oxides after elevating the annealing temperature.

#### **4. Conclusions**

The phase composition, microstructures, microhardness, and magnetic properties of as-cast and as-annealed W5, W6-Co, and W6-Gd HEAs have been investigated. The as-cast and as-annealed W6-Co HEAs maintain the same phase compositions, and are composed of single FeSi-rich phases, indicating the stable phase characteristic. The addition of Gd obviously enhances Tm (1185 ◦C) compared with W5, owing to the exhibition of AlNi and AlGd with high melting points. As-cast W5 possesses the highest hardness of 1210 HV, which is attributed to the uniform distribution of the FeSi-rich phase in the matrix. All the tested HEAs display soft magnetic properties. Moreover, the Ms and Mr/Ms values of W6-Gd were enhanced from 10.93 emu/g to 62.78 emu/g and from 1.44% to 15.50% via the annealing process, respectively. It suggests that Gd-oxides are beneficial to the enhancement of magnetic properties in W6-Gd.

**Author Contributions:** Conceived and designed the experiments, S.Z. and S.X.; Reviewed relevant studies and literature, W.W.; Performed the experiments, J.X. and Z.Z.; Analyzed the data, S.X., W.W. and Y.W.; Wrote the paper with the help of the rest of the authors, S.Z. and Z.Z.; Provided feasible advices and critical revision of the manuscript, Y.W.; All authors have read and approved the final manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (No. 51671095).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Evaluation of Radiation Response in CoCrFeCuNi High-Entropy Alloys**

**Yang Wang 1,2, Kun Zhang 1,2,\*, Yihui Feng 2,3, Yansen Li 1,2, Weiqi Tang 1,2 and Bingchen Wei 1,2,\***


Received: 10 October 2018; Accepted: 26 October 2018; Published: 31 October 2018

**Abstract:** CoCrFeCuNi high-entropy alloys (HEAs) prepared by arc melting were irradiated with a 100 keV He+ ion beam. Volume swelling and hardening induced by irradiation were evaluated. When the dose reached 5.0 × 1017 ions/cm2, the Cu-rich phases exhibited more severe volume swelling compared with the matrix phases. This result indicated that the Cu-rich phases were favorable sites for the nucleation and gathering of He bubbles. X-ray diffraction indicated that all diffraction peak intensities decreased regularly. This reduction suggested loosening of the irradiated layer, thereby reducing crystallinity, under He+ ion irradiation. The Nix-Gao model was used to fit the measured hardness in order to obtain a hardness value *H*<sup>0</sup> that excludes the indentation size effect. At ion doses of 2.5 × <sup>10</sup><sup>17</sup> ions/cm<sup>2</sup> and 5.0 × 1017 ions/cm2, the HEAs showed obvious hardening, which could be attributed to the formation of large amounts of irradiation defects. At the ion dose of 1.0 × 1018 ions/cm2, hardening was reduced, owing to the exfoliation of the original irradiation layer, combined with recovery induced by long-term thermal spike. This study is important to explore the potential uses of HEAs under extreme irradiation conditions.

**Keywords:** high-entropy alloy; ion irradiation; hardening behavior; volume swelling

#### **1. Introduction**

High-entropy alloys (HEAs) have recently drawn increased interest because of their distinct compositions, microstructures, and flexible properties. In contrast to conventional alloys, HEAs are composed of more than five principal elements at equal or nearly equal atomic percentages (at.%). HEAs exerts four primary effects: (1) high-entropy effect; (2) sluggish diffusion effect; (3) severe lattice distortion effect; and (4) cocktail effect [1–4]. These effects render HEAs more likely to form a simple solid-solution structure rather than an intermetallic compound, which confers distinct properties on HEAs, including high fatigue strength [5], high hardness [6], good abrasion resistance and corrosion resistance, high breaking strength at low temperatures [7–10], and good softening resistance at elevated temperatures [11]. The use of HEAs as structural materials under extreme environments has been proposed owing to their desirable mechanical properties and thermodynamic stability.

Previous studies on the irradiation effects of HEAs have mostly focused on irradiation resistance and phase stability. Zhang et al. [12] evaluated the effects of irradiation on AlxCoCrFeNi (x = 0.1, 0.75, and 1.5) HEAs under 3 MeV Au<sup>+</sup> ion irradiation at room temperature. Results indicated that compared with conventional nuclear materials, single-phase HEAs based on the face-centered cubic

(FCC) structure in the AlxCoCrFeNi system showed improved radiation resistance; in addition, volume swelling in the AlxCoCrFeNi alloys in the ascending order was FCC < FCC + BCC < BCC (body-centered cubic). Yang et al. [13] recently reported that with temperature increased from 523 K to 923 K, the irradiation-induced defect density of Al0.1CoCrFeNi HEAs decreased, whereas the size of the defect increased. Meanwhile, irradiation led to Ni and Co enrichment as well as Fe, Cr, and Al depletion in dislocation loops and dislocation regions. T. Nagase. et al. [14] showed that as-sputtered CoCrFeMnNi HEAs maintained good irradiation resistance within a wide temperature range from 298 to 773 K without grain coarsening under fast electron irradiation. Jin et al. studied a Ni-based multicomponent alloy under 3 MeV Ni ion irradiation at 773 K [15,16] and found that irradiation-induced volume swelling decreased with an increase in the number of elements in the disordered solid solution under identical irradiation conditions. This finding suggested that irradiation-induced volume swelling was also strongly affected by compositional complexity.

In the last decade, FeCoNi-based HEAs, as one of the successful HEAs, have attracted more and more attention, especially because of its mechanical properties and microstructure evolution [17]. M. Klimova et al. [18] studied the microstructure and mechanical properties evolution of the Al-, C-containing CoCrFeNiMn-type high-entropy alloy during cold rolling. They reported that rolling resulted in an increase in strength and a decrease in ductility of the Al-, C-containing CoCrFeNiMn-type alloy. Feng et al. [19] found that the short range order is positive between Al-Al, Al-Si, Si-Si pairs and negative between Ni-Al, Co-Si, Fe-Co, Ni-Si, and Fe-Si pairs, which leads to an increase in the elastic modulus by sacrificing ductility and isotropy. In addition, the appropriate doping of Y2O3 as a reinforcement phase in CoCrFeMnNi HEAs could increase both the room temperature tensile strength and the wear-resistance [20].

With regard to the irradiation response of HEAs, there are not many researches on the evolution of mechanical properties of FeCoNi-based HEAs after irradiation. Meanwhile, most studies focused on the effect of composition on irradiation-induced swelling and vertical inhomogeneity of radiation damage, and the difference in irradiation response between the phase and phase boundary has rarely been reported. Whether lateral inhomogeneity of irradiation response exists in the HEAs, especially in single-phase FCC structure-based HEAs, has yet to be determined. In the present study, a typical single-phase FCC structure-based CoCrFeCuNi HEA was selected, and the key objective is to determine the irradiation response exerted by He+ ion irradiation. The present study provides an initial examination of the fundamental irradiation behavior of an HEA material, thus offering an insight into the potential of this family of materials for application under extreme environments.

#### **2. Materials and Methods**

CoCrFeCuNi HEAs with a diameter of 5 mm and a length of 100 mm were prepared by arc melting a mixture of pure metal (purity > 99.9% wt %) in a Ti-gettered high-purity argon atmosphere. These ingots were remelted at least four times to prevent chemical heterogeneity and were eventually drop-cast into a copper mold. The cast bar was cut into thin pieces, each with a thickness of 2 mm, and then mechanically polished to a mirror finish. Irradiation experiments were conducted in a BNU-400 kV electrostatic accelerator in a test chamber, applying pressure near or below 10−<sup>5</sup> torr. The polished samples were irradiated with a 100 keV He<sup>+</sup> ion beam at fluences of 2.5 × <sup>10</sup><sup>17</sup> ions/cm2, 5.0 × 1017 ions/cm2, and 1.0 × <sup>10</sup><sup>18</sup> ions/cm<sup>2</sup> at a normal angle at room temperature.

Figure 1 shows the results for ion distribution, irradiation damage (displacement per atom or dpa), and concentration of He atoms with the increases in depth calculated using the Ion Distribution and Quick Calculation of Damage mode in the SRIM (Stopping and Range of Ions in Matter) 2008 code [21]. As shown in Figure 1a, when a large number of incident ions enter the HEAs, they form a spatial influence area rather than a specific path owing to collision cascade. The total number of ions is 99,999, and the effective depth of ion influence is ~560 nm. The dpa is determined by the sum of the predicted Co, Cr, Cu, Fe, and Ni vacancy concentrations and recoil events within 0 nm to ~560 nm from the surface. According to Figure 1b,c, the three curves in each figure linearly change with increased ion doses, implying that the dpa and concentration of He atoms are positively correlated with the total doses of incident ions. Equations (1) and (2) present the method used to extract the irradiation damage (total vacancies produced/atom = dpa) and He atom concentration (in at.%) from SRIM output files [21].

$$
\overbrace{\left(\frac{vacacines}{ions \times \dot{\mathbf{A}}}\right)}^{\text{vacanceics}}\times\left(\frac{10^8\left(\frac{\dot{\mathbf{A}}}{cm}\right) \times \text{Fluence}\left(\frac{ions}{cm^2}\right)}{8.705 \times 10^{22}\left(\frac{atoms}{cm^3}\right)}\right) = \text{dpa},\tag{1}
$$

$$\overbrace{\left(\frac{atoms/cm^3}{atoms/cm^2}\right)}^{\text{range,txt}}\times\left(\frac{\text{Fluence}\left(\frac{ions}{cm^2}\right)}{8.705\times10^{22}\left(\frac{atoms}{cm^3}\right)}\right)\times100=\left(\frac{ions}{atom}\right)\times100=\text{at\% ion,}\tag{2}$$

**Figure 1.** SRIM simulations of CoCrFeCuNi HEAs irradiated with 100 keV He+ ions. (**a**) Ion distribution; (**b**) irradiation damage (dpa); and (**c**) concentration of He atoms with depth.

In addition, the number of atomic displacements reached the peak at a depth of approximately 290 nm, which corresponded to the projected range of the He+ ions in CoCrFeCuNi HEAs. The peak concentrations of the He atoms were approximately 14.9%, 29.8%, and 59.6% under different doses.

The phase structures of pristine and irradiated samples were analyzed by slow-scan X-Ray diffraction (XRD, Philips PW 1050 diffractometer with CuKα radiation, Amsterdam, Netherlands) at 45 kV voltage, 0.01 step size, and 200 mA current for phase analysis with improved accuracy. The microstructure and compositions of pristine and irradiated samples were analyzed by scanning

electron microscopy (SEM, HITACHI S-4800, Tokyo, Japan) taken at a working voltage of 20.0 kV and a working distance of 14.3 and 15.0 mm, with energy-dispersive X-ray spectroscopy (EDX, HORIBA X-Max, Kyoto, Japan).

Nanoindentation experiments on pristine and irradiated samples were conducted using the Agilent Technologies Nano Indenter G200 with a stand Berkovich diamond indenter at room temperature (29.0 ◦C). The hardness measurement parameter included a strain rate of ~0.05 s−1, frequency of ~45 Hz, and harmonic displacement of ~2.0 nm. Each nanoindentation test ran to a maximum of 1000 nm into the indented surface, and all hardness measurement positions were in the matrix phases. Specifically, the sample irradiated at fluences of up to 1.0 × <sup>10</sup><sup>18</sup> ions/cm2, the hardness measurement positions were located apart from the exfoliation of the damaged areas, that is, on the remaining part underneath. The geometry of the tip was standardized from the indentation on fused silica with the same indentation depth, and each sample was subjected to 5 indents with a space wider than 50 μm.

#### **3. Results**

#### *3.1. Microstructural Characterization*

The XRD patterns of the pristine and irradiated CoCrFeCuNi HEA samples at different doses are presented in Figure 2. The figure reveals that the samples remained fully crystalline, and the diffraction peaks of CoCrFeCuNi HEAs in both pristine and irradiated samples at 43◦, 50◦, and 74◦ are from the typical FCC phases corresponding to (111), (200), and (220) lattice planes, respectively. These findings are similar to those in previous studies [22,23]. Figure 2 shows that no obvious phase decomposition occurs under any irradiation condition, which is consistent with high configurational entropy playing a significant role in single-phase stability [24]. A closer examination reveals that each major peak contains two slightly separated peaks. The left peak of the major split peaks is attributed to the Cu-rich phases, whereas the right peak is attributed to the matrix phases. These findings are consistent with the following EDX results. The separated major peaks are hardly differentiated at low angles because no significant difference in average atomic radius between the Cu-rich phases and matrix phases in CoCrFeCuNi HEA is indicated. Moreover, the diffraction peaks of the irradiated HEAs shift to the left, meaning a slight lattice expansion happens. Meanwhile, the intensity of the diffraction peaks decreases with the increasing of irradiation dose, owing to the increased concentration of defects in the near surface of HEAs.

**Figure 2.** XRD patterns of pristine and He+ ion-irradiated CoCrFeCuNi HEAs at different doses.

Irradiation has been known to enhance diffusion in solids by creating defects, such as vacancies and interstitials. These defects fill the irradiated layer with voids, eventually leading to a decrease in crystallinity. Irradiation-induced defects generally lead to structural damage in the HEAs, such as cluster of small defects and dislocation loops [25]. According to Makinson et al. [26], the peak diffraction intensity of X-rays depends on the concentration of defects in a material. Meanwhile, peak intensity decreases with an increase in defect concentration, and vice versa. In the present study, the defect concentration induced by He+ ion irradiation increases with an increase in irradiation dose of up to 1.0 × <sup>10</sup><sup>18</sup> ions/cm2, leading to a decrease in peak intensity. Moreover, the low solubility of helium in CoCrFeCuNi HEAs leads to its precipitation into bubbles and voids [27], further decreasing the peak intensities.

#### *3.2. Surface Topography*

The SEM images reveal the surface morphology of pristine and irradiated CoCrFeCuNi HEAs at different irradiation doses, as shown in Figure 3a–d. Table 1 shows the components of different regions in the microstructures of pristine and irradiated CoCrFeCuNi alloys estimated by EDX spectroscopy. The analysis was conducted on five pristine and five irradiated regions randomly chosen from the samples. The EDX maps showing the elemental distribution of Cr, Co, Fe, Cu, and Ni on the HEAs surface at different doses, are shown in Figure 4a–d, respectively. The EDX maps show matrix phases enriched with Cr, Fe, Co, and Ni, as well as Cu-rich phases at the phase boundaries with fair amounts of Ni, and no obvious diffusion occurs in the HEA components after irradiation, implying that the composition of HEA is stable after irradiation to some extent. Nevertheless, the distribution of the elements is very homogeneous in both the matrix phases and Cu-rich phases. Moreover, both the matrix phases and Cu-rich phases consist of only one simple phase, which is consistent with previous studies [28]. According to Figure 3b–d, when the irradiation dose increases to 5.0 × <sup>10</sup><sup>17</sup> ions/cm2, the volume swelling of the Cu-rich phases markedly increases in severity. Meanwhile, no significant change is observed in the matrix phases, suggesting the inhomogeneity of radiation resistance on the surface perpendicular to the ion incident direction. The evolution of volume swelling after irradiation is discussed in the subsequent section.

**Figure 3.** SEM images of CoCrFeCuNi HEAs. (**a**) Pristine and (**b**–**d**) He+ ions irradiated at fluences of 2.5 <sup>×</sup> 1017 ions/cm2, 5.0 <sup>×</sup> <sup>10</sup><sup>17</sup> ions/cm2, and 1.0 <sup>×</sup> <sup>10</sup><sup>18</sup> ions/cm2.


**Table 1.** Components of different regions in the microstructure of CoCrFeCuNi HEAs (at. %).

**Figure 4.** Energy-dispersive X-ray spectroscopy (EDX) maps showing elemental distribution of CoCrFeCuNi HEAs. (**a**) Pristine and (**b**–**d**) He<sup>+</sup> ions irradiated at fluences of 2.5 <sup>×</sup> 1017 ions/cm2, 5.0 <sup>×</sup> 1017 ions/cm2, and 1.0 <sup>×</sup> 1018 ions/cm2.

However, at a dose of 1.0 × 1018 ions/cm2, cracking and exfoliation occur, and the subsurface becomes visible after the surface flakes off. Wei et al. [29] reported a similar occurrence on Cu48Zr48Al4 bulk metallic glass composites induced by He+ ion irradiation at the same dose. When high-energy He<sup>+</sup> ions are bombarded at a target material, most ions take electrons from the matrix to form helium atoms, which can cause distortion, elastic rebound, and stress in the lattice [30]. With increasing irradiation fluence, irradiation-induced helium gas pressure develops inside the matrix. At a fluence

of 1.0 × 1018 ions/cm2, excessive pressure of the helium gas was developed inside the irradiated area, which eventually led to cracking and exfoliation [31,32], as shown in Figure 3d.

#### *3.3. Mechanical Behavior*

A nanoindentation test was conducted to reveal the average microhardness of the HEA surface and thus investigate the irradiation-induced hardening behavior. Figure 5 shows typical (a) depth profiles of nanoindentation hardness and (b) dependence of the Hirr/Hunirr ratio on the indentation depth of pristine and irradiated CoCrFeCuNi HEAs at different irradiation doses. The irradiated region and substrate region are consistent with the dpa profiles (see Figure 1b). Figure 5a shows that the pristine samples decrease in hardness with increasing indent depth at the indentation depth of h > 50 nm. Such depth-dependent hardness behavior is regarded as the indentation size effect (ISE) [33]. By contrast, hardness increases with increasing indent depth at the depth of *h* < 50 nm is considered as reverse ISE, which is usually caused by uncertainty factors from the surface and the geometrical shape of the Berkovich indenter. Therefore, to exclude uncertainty factors from the surface, we ignored the data in regions shallower than 100 nm in this study. Compared with pristine samples, all irradiated samples exhibited different degrees of hardening at *h* > 100 nm. To more clearly characterize the degree of increment, the dependence of the Hirr/Hunirr ratio on the indentation depth of all samples is exhibited in Figure 5b. Notably, the normalized nanoindentation hardness reaches a peak with an increase in indentation depth. The same peak position on all curves is located at a depth of ~146 nm (the dotted transition line). For the shallower region before the transition line, Hirr/Hunirr increases with an increase in indentation depth, which is similar to the results for other metals [25] irradiated with similar ions. However, for the deeper region after the transition line, Hirr/Hunirr decreases with an increase in indentation depth.

**Figure 5.** (**a**) Typical depth profiles of nanoindentation hardness and (**b**) dependence of the Hirr/Hunirr ratio on the indentation depth of pristine and He+ ion-irradiated CoCrFeCuNi HEAs at different irradiation doses.

An irradiation-induced hardened layer forms in the HEAs after He+ ion implantation. When indented into the sample surface, the hemispheric influence zone beneath the indenter can reach 4 to 10 times the indentation depth. As the indenter presses deeper, if the radius of the influence zone is less than the thickness of the hardened layer, the damage grade effect is exerted, and the hardening ratio increases; otherwise the softer substrate effect is exerted, and the hardening ratio decreases [34].

To explain the ISE, Nix and Gao [35] developed a model based on the concept of geometrically necessary dislocations. This model predicts the hardness-depth profile by using the following equation.

$$H = H\_0 \left( 1 + \frac{h^\*}{h} \right)^{1/2},\tag{3}$$

where *H* is the hardness, *h* is the indentation depth, *H*<sup>0</sup> is the hardness at an infinite depth (i.e., macroscopic hardness), and *h*\* is the length that characterizes the depth dependence of hardness for a given material and indenter geometry, which depends on the material and shape of the indenter tip (i.e., statistically stored dislocation density) [36]. To aid the discussion, the hardness data are plotted as *H*<sup>2</sup> versus 1/*h* in Figure 6. *H*<sup>0</sup> is the square root of the intercept for the linear fitting of the hardness data in the near-surface region. The value of *H*<sup>0</sup> excludes the size effect and can be used as a parameter to characterize the hardening effect in HEAs under irradiation.

**Figure 6.** Plot of *H*<sup>2</sup> versus 1/*h* for the pristine and irradiated HEAs at different doses.

In Figure 6, the pristine sample shows good linearity in the range of 100 nm < *h* < 1000 nm. Regardless, the irradiated samples exhibit bilinearity with a shoulder at the depth of about 150 nm consisting of the depth of the transition line, (see Figure 6). Table 2 lists *H*<sup>0</sup> and *h*\* for the pristine samples and irradiated samples calculated by least squares fitting of hardness data in the 100 nm < *h* < 1000 nm range according to Equation (3) [37].


**Table 2.** Calculated *H*<sup>0</sup> and *h*\* based on the Nix–Gao model.

The data in Table 2 shows that the calculated *H*<sup>0</sup> first increases and then decreases with an increase in irradiation dose, which is consistent with the hardness results. The calculated *h*\* decreases systematically, implying a smaller size effect during irradiation, which is consistent with previous studies [37].

#### **4. Discussion**

The schematic in Figure 7 illustrates the swelling effect, as well as the hardening behavior, of Cu-rich phases during He<sup>+</sup> irradiation. With regard to the formation of He bubbles, nuclear

loss-induced collisional damage, such as dislocations, voids, and interfaces, act as sinks for point defects. Dislocations can partly absorb more interstitials than vacancies. A large number of vacancies are consequently generated in the HEAs, and vacancy clusters grow and trap He to form He bubbles. For CoCrFeCuNi HEA, the Cu-rich phases appeared as defect-rich phase boundary regions in the matrix phases, similar to grain boundary, which are also favorable nucleation sites of He bubbles [38]. With the ion dose increasing to 5.0 × <sup>10</sup><sup>17</sup> ions/cm2, the He concentration increasingly rises and reaches near-saturation [39] in the matrix phases. Thus, He bubbles tend to migrate and gather toward the Cu-rich phases, resulting in the increased severity of volume swelling, as shown in Figure 7. We also reported that the irradiation response is inhomogeneous along the ion incident direction because of a perpendicular local shear stress [40]. Cu-rich phases may act as local low-stress regions because of the more loosely packed He atoms. He bubbles may thus preferentially migrate toward the Cu-rich phases to unload the stress, resulting in the increased severity of volume swelling [41,42]. This factor can be potentially influence the volume swelling of Cu-rich phases.

**Figure 7.** Schematic of radiation response in CoCrFeCuNi HEAs.

At ion doses ranging from 2.5 × 1017 ions/cm2 to 5.0 × 1017 ions/cm2, HEAs exhibit marked hardening behavior, which can be attributed to obstacles to dislocation glide, proposed by Orowan and Seeger [43,44]. In general, irradiation-induced defect clusters and/or dislocation loops, as well as He bubbles, pin dislocation lines and impede dislocation glide, causing HEAs to harden [45]. However, as previously mentioned, as the irradiation dose increases, numerous He bubbles migrate and gather toward the Cu-rich phases to form cavities, which may induce swelling and softening of the Cu-rich phases. While the irradiation-induced dislocations continue to increase in the matrix phases, the overall result for the HEA ultimately reflects the hardening effect.

When the ion dose increases to 1.0 × 1018 ions/cm2, the hardness of the HEA slightly decreases. To avoid surface irregular interference, we selected the flatter subsurface for hardness testing. Consistent with the SEM micrograph in Figure 3d, at the highest dose, the uncovering of the subsurface after exfoliation corresponds to a new irradiation plane, except that it is only affected by diminished irradiation effects, and therefore the hardness decreases. In addition, recovery [46–48] might occur during irradiation. The thermal-spike effect during long-term irradiation can lead to annealing of defects and annihilation of helium bubbles, thus reducing the hardening effect. We hypothesize that the

weakened irradiation effect on the uncovering subsurface, combined with recovery, causes softening of the HEA.

#### **5. Conclusions**

CoCrFeCuNi HEAs prepared by arc melting were irradiated with a 100 keV He<sup>+</sup> ion beam at fluences of 2.5 × <sup>10</sup><sup>17</sup> ions/cm2, 5.0 × 1017 ions/cm2, and 1.0 × <sup>10</sup><sup>18</sup> ions/cm<sup>2</sup> at room temperature. XRD results proved that all diffraction peak intensities along the (111), (200), and (220) lattice planes decreased regularly, indicating that the irradiated layer was filled with voids, leading to a decrease in crystallinity. When the irradiation dose increased to 5.0 × <sup>10</sup><sup>17</sup> ions/cm2, the irradiation-induced swelling became increasingly severe in the Cu-rich phases, whereas no significant change occurred in the matrix phases. This difference suggested that the degree of swelling varied between the two phases under the same irradiation condition. This finding contributed to the migration and gathering of the He bubble-induced noncompact structure and the lateral inhomogeneity of radiation damage. At a higher dose, 1.0 × <sup>10</sup><sup>18</sup> ions/cm2, cracking and exfoliation occurred revealing the subsurface after the surface flaked off mainly from the excessive pressure of the He gas. Nanoindentaion was performed to indicate that hardening induced by He+ ion irradiation occurred in both irradiated samples. The depth-dependent hardness behavior was explained by the Nix-Gao model. At ion doses of 2.5 × <sup>10</sup><sup>17</sup> ions/cm2 and 5.0 × <sup>10</sup><sup>17</sup> ions/cm2, the HEAs exhibited an obvious hardening behavior, which can be explained by dislocation-dominated hardening effect. When the ion dose reached 1.0 × <sup>10</sup><sup>18</sup> ions/cm2, the hardening effect decreased probably because of the weakened irradiation effect on the uncovered subsurface, combined with the long-term thermal-spike effect-induced recovery.

**Author Contributions:** Conceptualization, K.Z. and Y.W.; software, W.T.; validation, B.W., K.Z. and Y.W.; formal Analysis, Y.W.; investigation, Y.W. and Y.L.; resources, B.W.; data curation, Y.W. and Y.F.; writing-original draft preparation, Y.W.; writing-review & editing, Y.W. and K.Z.; visualization, Y.W.; supervision, B.W. and K.Z.; project administration, B.W.

**Funding:** This research was funded by the National Natural Science Foundation of China (Grant No. 51401028, No. 51271193, No. 11402277, No. 11790292) and the Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDB22040303).

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


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