**Impact of Chemical Fluctuations on Stacking Fault Energies of CrCoNi and CrMnFeCoNi High Entropy Alloys from First Principles**

#### **Yuji Ikeda 1,2,\*, Fritz Körmann 1,3, Isao Tanaka 2,4,5,6 and Jörg Neugebauer <sup>1</sup>**


Received: 13 August 2018; Accepted: 29 August 2018; Published: 30 August 2018

**Abstract:** Medium and high entropy alloys (MEAs and HEAs) based on 3d transition metals, such as face-centered cubic (fcc) CrCoNi and CrMnFeCoNi alloys, reveal remarkable mechanical properties. The stacking fault energy (SFE) is one of the key ingredients that controls the underlying deformation mechanism and hence the mechanical performance of materials. Previous experiments and simulations have therefore been devoted to determining the SFEs of various MEAs and HEAs. The impact of local chemical environment in the vicinity of the stacking faults is, however, still not fully understood. In this work, we investigate the impact of the compositional fluctuations in the vicinity of stacking faults for two prototype fcc MEAs and HEAs, namely CrCoNi and CrMnFeCoNi by employing first-principles calculations. Depending on the chemical composition close to the stacking fault, the intrinsic SFEs vary in the range of more than 150 mJ/m<sup>2</sup> for both the alloys, which indicates the presence of a strong driving force to promote particular types of chemical segregations towards the intrinsic stacking faults in MEAs and HEAs. Furthermore, the dependence of the intrinsic SFEs on local chemical fluctuations reveals a highly non-linear behavior, resulting in a non-trivial interplay of local chemical fluctuations and SFEs. This sheds new light on the importance of controlling chemical fluctuations via tuning, e.g., the annealing condition to obtain the desired mechanical properties for MEAs and HEAs.

**Keywords:** high-entropy alloy; stacking-fault energy; density functional theory

#### **1. Introduction**

High entropy alloys (HEAs) or complex concentrated alloys (CCAs) based on 3d transition metals have attracted enormous attention recently, in particular due to their outstanding mechanical properties. The equiatomic CrMnFeCoNi HEA, also often termed the Cantor alloy [1], has an excellent combination of strength and ductility [2–5]. Different strategies have been proposed to further improve the mechanical properties, e.g., by tuning the chemical compositions of CrMnFeCo and CrMnFeCoNi towards nonequiatomic alloys [6–10] or by resorting to so-called medium entropy alloys (MEAs), such as CrCoNi alloys [11–13].

A key factor in controlling the underlying deformation mechanism and therewith tuning the mechanical properties is the stacking fault energy (SFE). Low SFEs can induce, e.g., transformationinduced plasticity (TRIP) or twinning-induced plasticity (TWIP) [14–16], and for this reason, SFEs of HEAs and CCAs have been investigated previously in numerous experimental [13,17–19] as well as theoretical studies [20–31]. Interestingly, in a recent experimental work [17], the measured SFEs for equiatomic CrMnFeCoNi revealed large fluctuations, and it was proposed that the SFEs of CrMnFeCoNi may sensitively depend on the local chemical environments in the vicinity of the stacking faults (SFs). Also, in computational works [24,26,29,30], large fluctuations of SFEs have been found for MEAs and HEAs based on 3d transition metals, like CrCoNi and CrMnFeCoNi, which indicates a strong dependence of the SFEs on local chemical fluctuations close to the SFs. A recent experimental work [32] revealed that chemical inhomogeneity in a Cr10Mn30Fe50Co10-based alloy also caused a large deterioration of mechanical properties. Although these results suggest an important role of chemical fluctuations in these alloys, the impact of such fluctuations on the SFEs of MEAs and HEAs has not been intensively investigated yet. In a recent study employing first-principles calculations [24], it was proposed that the SFEs of CrCoNi and CrCoFeNi may depend on the valence electron concentration (VEC) of the elements near the SF. In that study, however, only a very limited number of configurations and local compositions near the SF were evaluated, prohibiting a further quantitative analysis.

In the present study, we comprehensively investigate the impact of compositional fluctuations near the intrinsic SF (ISF) on the intrinsic SFE (ISFE) for the face-centered cubic (fcc) equiatomic CrCoNi and CrMnFeCoNi alloys based on first-principles calculations. The ISFEs were calculated using supercells with and without an ISF, while the chemical disorder was modeled using the coherent potential approximation (CPA) [33–35]. The combination of supercells and the CPA makes it possible to systematically investigate arbitrary local composition ratios in the vicinity of the ISFs and to elucidate the impact of compositional fluctuations for individual elements using relatively small, and thus computationally efficient, supercells.

#### **2. Computational Details**

We investigate the impact of the local chemical environment in the vicinity of ISFs on the ISFEs, as described in Figure 1a. In principle, the impact of the local chemical environment can be investigated by employing large supercells. As this approach turns out to be computationally too demanding for systematically screening a large number of different compositional fluctuations near the ISFE, we resort to an alternative approach. To reduce the computational cost while keeping the key physical ingredients, we combined the CPA with the supercell approach. Specifically, the impact of the local chemical environment near the ISFs on the ISFEs was investigated using six-layer supercells with and without an ISF as shown in Figure 1b, and the compositional fluctuation in the vicinity of the ISF was introduced based on the CPA by modifying the mixing ratios of the chemical elements in the L1 layers close to the ISF (see Figure 1b) from the equiatomic while keeping the equiatomic ratios in the subsequent L2 and L3 layers. The ISF was introduced by tilting the 111 axis of the perfect-fcc simulation cell by 6 112¯/*a*, where *<sup>a</sup>* is the fcc lattice constant [36]. The ISFEs, *<sup>γ</sup>*ISF, were computed with

$$
\gamma\_{\rm ISF} = \frac{E^{\rm fcc} + \rm ISF}{A} \,, \tag{1}
$$

where *E*fcc <sup>+</sup> ISF and *E*fcc are the energies of the simulation cells with and without ISFs, respectively, and *A* denotes the area of the ISFs. As we focused on the impacts of the chemical fluctuations near the ISFs, the computed ISFEs are shown as the differences from the ones without compositional fluctuations. The impacts of lattice vibrations [26,29], magnetic excitations [21], chemical short-range order (SRO) [30], and volume changes [9] on the ISFEs have been investigated previously and were therefore not included in the present study. We also note that local lattice distortions, which cannot be considered in the CPA, possibly affect the ISFEs.

**Figure 1.** (**a**) Schematic of the intrinsic stacking faults (ISF) in a three-component face-centered cubic (fcc) equiatomic disordered alloy. The red line indicates the ISF. The circles represent atoms, which are colored differently according to the chemical elements. Notice that local chemical fluctuations can exist close to the ISF (gray background region) even if, on average, the alloy has an equiatomic composition; (**b**) Projected atomic positions of the simulation cells with and without ISFs. The red background regions represent the six-layer simulation cells, while the dashed boxes indicate the cell shape of the perfect fcc structure. The red lines indicate the ISFs. The circles represent atoms which are colored differently according to the given mixing ratios of the constitutive chemical elements in the CPA to model the compositional fluctuations. The atomic layers are labeled L1, L2, and L3 according to the distance from the ISFs.

SFEs are also often computed based on the axial Ising model [37], in which the SFE is derived from the energy differences among the fcc, hexagonal close-packed (hcp), and sometimes double hcp phases. By construction, however, the axial Ising model cannot capture the impact of local chemical fluctuations close to the SFs, because the perfect fcc and hcp structures do not include the SFs explicitly. The presently employed supercell approach, in contrast, enabled us to explicitly investigate the impacts of local chemical fluctuations close to the SFs.

The electronic structure calculations were performed with the exact-muffin-tin-orbital (EMTO) method [38–42] in combination with the full-charge-density (FCD) method [43,44] within the density functional theory (DFT) framework. The DFT energies were calculated within the generalized gradient approximation (GGA) of the Perdew–Burke–Ernzerhof (PBE) form [45] in the following perturbative manner [46]. The electronic densities were first calculated using the local-density approximation (LDA), and then the total energies were calculated within the GGA-PBE via the FCD method based on the obtained electronic densities. This approach makes the calculations faster and often more stable while keeping the accuracy of the GGA [46], and it has therefore been employed in numerous studies using the EMTO approach [47–56]. The Brillouin zones were sampled by 22 × 22 × 4 *k*-point meshes per six-atom computational unit cell, where the 111 direction in Figure 1b was set to be the *z*-axis. As experiments [57] and first-principles calculations [52,53] have revealed that both CrCoNi and CrMnFeCoNi are paramagnetic (PM) at room temperature, we simulated the magnetic disorder with random magnetic moments by employing the disordered local moment (DLM) model [42,58,59] in combination with the CPA. The lattice constant was fixed to 3.56 Å for CrCoNi and 3.6 Å for CrMnFeCoNi, which are close to the experimental values [1,60–64], and the atomic positions were fixed to keep the rigid-sphere packing.

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

We first discuss the impact of compositional fluctuations in the vicinity of the ISF for CrCoNi with increasing or decreasing individual elemental concentrations. The results are shown in Figure 2a.

**Figure 2.** (**a**,**b**) Differences between the computed ISFEs of CrCoNi from those where no compositional fluctuation exists in the L1 layers. (**a**) Result as a function of the local concentration of the element *M* (*M* = Cr, Co, Ni) in the L1 layers. The other elements were kept equiatomic in the L1 layers. The vertical gray line indicates the point of the ideal solid solution, i.e., without chemical fluctuations near the SF; (**b**) Result as a function of the local composition ratio in the L1 layers. The red circle indicates the equiatomic composition ratio; (**c**) The valence electron concentration (VEC) as a function of the composition ratio.

The ISFE was found to increase monotonically with the increase of local Ni concentration close to the ISFs. This suggests that Ni segregation towards the ISF is thermodynamically limited and that there is a strong driving force to deplete the Ni concentration in the vicinity of the ISFs. This can be intuitively understood by the fact that pure Ni energetically prefers the fcc phase (groundstate of Ni), whereas in the vicinity of the ISF, the stacking order of the close-packed planes is similar to that of the hcp structure. If the layers close to the ISF are fully occupied by Ni, the ISFE increases by more than 100 mJ/m2, and when Ni is fully suppressed from the ISF, the ISFE decreases by more than 50 mJ/m2. This is consistent with previous computational results based on supercell models [24] in which the ISF with the local composition ratio of Cr8Co10Ni14 was 59 mJ/m2 higher than the ISF with the local composition ratio of Cr12Co10Ni10. It should be noted that the results in Ref. [24] were obtained based on non-spin-polarized calculations and another type of supercell approach (slab + vaccum layer), which prohibits a quantitative comparison.

The impact of Cr turned out to be somewhat more complex. The minimum ISFE was found at a Cr concentration of about 0.5, whereas the ISFE increased when the L1 layers were either highly occupied by Cr or mostly free from it. When Cr fully occupied the layers close to the ISF, the ISFE increased by more than 80 mJ/m2. A similar non-linear behavior of the ISFE was also found for the ferromagnetic (FM) fcc Cr*x*Co1−*<sup>x</sup>* binary alloys by first-principles calculations [65]. This suggests that such a complex impact of Cr on the ISFE may be also found in other random 3d transition metal alloys.

When Co fully occupied the L1 layers, the ISFE did not change largely from when no chemical fluctuations existed, which suggests that Co in CrCoNi reveals a small fcc–hcp energy difference. This could be reasoned as follows. The groundstate of pure Co is the FM hcp phase, while with increasing temperature, it experiences both a structural transition to the fcc phase at around 700 K [66] and a magnetic transition at around 1400 K [66]. As found in previous works [67,68], magnetic fluctuations in Co contribute to the stability of the fcc phase at elevated temperatures. For example, in Ref. [68], the fcc–hcp energy difference of pure Co was found to be significantly reduced in the DLM state compared to the FM state. From the above considerations on pure Co, it can be intuited that the small ISFE change with respect to the local Co concentration close to the ISF is caused by the paramagnetic state of CrCoNi.

To further elucidate the impact of the non-linear trend of the compositional fluctuations for the ISFEs, we further explored the full compositional region for the L1 layers, as shown in Figure 2b. Overall, the ISFE varied in the range of approximately 180 mJ/m2, depending on the local composition ratio of the L1 layers. The strongest decrease in the ISFE of more than 60 mJ/m2 was found for nearly equiatomic Cr and Co occupying the L1 layers.

Large variation in the ISFE for CrCoNi was also found in previous computational works using supercell models with and without ISFs, where the computed ISFEs of CrCoNi were distributed in the range of 59 mJ/m2 [24], approximately 230 mJ/m2 [29], or 205 mJ/m<sup>2</sup> [30], or with a standard deviation of approximately 110 mJ/m<sup>2</sup> [26]. Note that the ISFE variation range in Ref. [24] was relatively small compared with that in Refs. [26,29,30]. This might be related to the limited set of considered configurations, or the usage of non-spin-polarized calculations and the suppression of energy fluctuations due to the magnetic degrees of freedom. This is consistent with the finding in Ref. [29], where the fluctuations of the ISFEs in CrMnFeCoNi were much smaller in the nonmagnetic state compared to the spin-polarized one.

As found in Figure 2a,b, the ISFE of CrCoNi varied non-linearly with respect to the local chemical composition close to the ISFs. This became even clearer by considering the VEC dependence on the ISFE, which is often employed to predict the phase stability of HEAs and CCAs [69,70]. The relations between the ISFE and the VEC near the ISFs for CrCoNi and CrFeCoNi have been discussed previously [24]. Figure 2c shows the VEC values for the same composition region as that in Figure 2b. By construction, VEC depends linearly on the chemical concentrations, while the trend does not match that of the actually computed ISFE, except for under Ni-rich conditions. This indicates that the VEC is, in general, not a sufficient quantitative descriptor for estimating the variation of ISFEs.

We next considered the five-component fcc CrMnFeCoNi HEA, the so-called Cantor alloy. The impact of compositional fluctuations in the L1 layers on the computed ISFE are shown in Figure 3a. The ISFE depended on the local concentrations of Cr, Co, and Ni similarly, as found for CrCoNi. The ISFE monotonically increased with an increase in the local Ni concentration in the L1 layers. The ISFE also increased both when Cr nearly fully occupied the L1 layers and when it was nearly fully excluded in the L1 layers, while the ISFE decreased when the local Cr concentration in the L1 layers was around 0.5. Co showed a relatively small impact on the ISFE compared to the other elements also for CrMnFeCoNi with a slight decrease in the ISFE when fully occupying the L1 layers. The impact of Mn on the ISFE is clearly non-linear, as also found for Cr. When the local Mn concentration in the L1 layers was less than 0.4, the ISFE was hardly affected. In contrast, when the local Mn concentration in the L1 layers was larger than 0.4, the ISFE drastically decreased down to approximately 150 mJ/m2 with Mn fully occupying the L1 layers.

**Figure 3.** (**a**,**b**) Differences between the computed intrinsic SFE (ISFEs) of CrMnFeCoNi from those where no compositional fluctuation exists in the L1 layers. (**a**) Result as a function of the local concentration of the element *M* (*M* = Cr, Mn, Fe, Co, Ni) in the L1 layers. The other elements were kept equiatomic in the L1 layers. The vertical gray line indicates the point of the ideal solid solution, i.e., without chemical fluctuations near the SF; (**b**) Result as a function of the local composition ratio in the L1 layers in the pseudoternary region of Cr, Fe0.5Co0.5, and Mn0.5Ni0.5. The red circle indicates the equiatomic composition ratio; (**c**) The valence electron concentration (VEC) as a function of the composition ratio in the same pseudoternary region as that of (**b**).

Recent experiments [63,71] have reported the phase decomposition of the CrMnFeCoNi alloy into body-centered cubic Cr, L10 MnNi, and B2 FeCo after annealing at 450–500 ◦C. To resolve the possible relationship between the precursor of the phase decomposition and the corresponding local chemical fluctuations near the ISFs, the ISFEs were also computed for the local composition ratios in the L1 layers in the pseudoternary region of Cr, Fe0.5Co0.5, and Mn0.5Ni0.5. The results are shown in Figure 3b. The ISFE increased with an increasing amount of Mn0.5Ni0.5, while the lowest ISFE was found when Mn0.5Ni0.5 was fully excluded from the L1 layers and replaced by a mixture of about 0.5 Cr and 0.5 Fe0.5Co0.5. This suggests that ISFs are suppressed in the presence of MnNi clusters. The VEC-derived linear dependencies are shown in Figure3c and reveal once more that the VEC is, in general, not a good predictor of the nonlinear dependence of ISFE on local chemical fluctuations.

#### **4. Conclusions**

We investigated the impact of compositional fluctuations on the ISFEs for the fcc equiatomic CrCoNi and CrMnFeCoNi alloys by employing first-principles calculations by combining the supercell and CPA approaches. For both alloys, the ISFEs were found to vary within a range of more than 150 mJ/m<sup>2</sup> depending on the local chemical environment in the vicinity of the ISFs. The chemical dependencies were shown to be highly non-linear and therewith, strongly deviated from linear Vegard's law-like behavior. Cr caused non-linear behavior in CrCoNi. In the CrMnFeCoNi alloy, the presence of MnNi clusters could suppress ISFs.

The strong SFE dependence on the local chemical compositions in the vicinity of the SFs indicated, on the one hand, that the SFs in HEAs and CCAs may promote particular types of chemical segregations towards the SFs. On the other hand, if chemical fluctuations exist in HEAs and CCAs on a large enough scale, SFs are likely to occur in the local chemical environments with low SFEs. These potential behaviors further complicate the prediction of physical descriptors to determine deformation mechanisms in HEAs and CCAs compared to, e.g., unary metals or ordered alloys. The complexity might be further enhanced in the presence of chemical SRO, which impacts the probabilities of local chemical fluctuations compared to the ideal mixing state. At the same time, the revealed dependence of SFEs on local chemical fluctuations in the vicinity of SFs opens the route towards tuning alloy properties of HEAs and CCAs via controlling chemical fluctuations by tuning, e.g., the annealing conditions in the alloy processing route. Our results encourage further experimental analyses employing, e.g., transmission electron microscopy/energy-dispersive X-ray spectroscopy or atomic probe tomography to explore the role of chemical fluctuations in the vicinity of SFs in more detail.

**Author Contributions:** Y.I. performed the DFT calculations which were analyzed together with F.K. All authors have equally contributed to the overall discussion and preparation of the manuscript as well as read and approved the final manuscript version.

**Funding:** Funding from the Deutsche Forschungsgemeinschaft (SPP 2006), from the Ministry of Education, Culture, Sports, Science and Technology (MEXT), Japan, through Elements Strategy Initiative for Structural Materials (ESISM) of Kyoto University, from the Grant-in-Aid for Scientific Research on Innovative Areas "Nano Informatics" (Grant No. 25106005) from the Japan Society for the Promotion of Science (JSPS), and from NWO/STW (VIDI grant 15707) are gratefully acknowledged.

**Acknowledgments:** Discussions with Blazej Grabowski and Zhiming Li (Max-Planck Institut für Eisenforschung GmbH) are gratefully acknowledged.

**Conflicts of Interest:** The authors declare no conflict of interest. The founding sponsors had no role in the design of the study; in the collection, analyses, or interpretation of data; in the writing of the manuscript, and in the decision to publish the results.

#### **References**


c 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* **Elemental Phase Partitioning in the** *γ***-***γ***" Ni2CoFeCrNb0.15 High Entropy Alloy**

**Bin Han 1,\*, Jie Wei 1, Feng He 1,2, Da Chen 1, Zhi Jun Wang 2, Alice Hu 1,3,4, Wenzhong Zhou 1,3 and Ji Jung Kai 1,3,\***


Received: 28 October 2018; Accepted: 16 November 2018; Published: 28 November 2018

**Abstract:** The partitioning of the alloying elements into the γ" nanoparticles in a Ni2CoFeCrNb0.15 high entropy alloy was studied by the combination of atom probe tomography and first-principles calculations. The atom probe tomography results show that the Co, Fe, and Cr atoms incorporated into the Ni3Nb-type γ" nanoparticles but their partitioning behaviors are significantly different. The Co element is much easier to partition into the γ" nanoparticles than Fe and Cr elements. The first-principles calculations demonstrated that the different partitioning behaviors of Co, Fe and Cr elements into the γ" nanoparticles resulted from the differences of their specific chemical potentials and bonding states in the γ" phase.

**Keywords:** high entropy alloy; gamma double prime nanoparticles; elemental partitioning; atom probe tomography; first-principles calculations

#### **1. Introduction**

Recently, a new class of structural materials, known as high entropy alloys (HEAs), have attracted considerable attention due to their excellent properties and potential applications in the aerospace and energy industries [1–11]. Compared with the conventional alloys, the face-centered cubic (FCC) HEAs exhibit unique properties such as outstanding ductility [7], exceptional fracture toughness [12] as well as excellent corrosion resistance [13]. However, the single-phase FCC HEAs are insufficiently strong, which limits their engineering applications.

The strategy of introducing the dispersed hard D022-structured gamma double prime (γ") or L12-structured gamma prime (γ ) nanoparticles into the FCC matrix (γ phase) has been proved to be one of the most effective approaches to enhance the strength of the FCC HEAs, as it is the case in many superalloys [10,14–16]. It is known that the alloying elements in the γ" or γ phase plays an important role on the stability and the mechanical properties of the nano-precipitated alloys [17–20]. Therefore, it is critical to clarify the partitioning of the alloying elements into the nanoparticles of the FCC HEAs. However, this issue still lacks research because the observation of alloying elements in the nanoparticles which embedded in the FCC matrix is still a challenge. Although the energy dispersive X-ray spectroscopy (EDS) equipped on scan electron microscope (SEM) or on transmission electron microscope (TEM) has been widely used to determine the material composition, it is difficult to distinguish the composition of the nanoparticles from that of the surrounding matrix. Atom probe

tomography (APT), the only technique which can generate the three-dimensional (3D) atom maps of materials in the real space with nearly atomic-scale resolution, has been proved to be a powerful method of characterizing the composition of different kinds of nanoparticles [15,18,21,22]. Recently, we have successfully clarified the partitioning of the alloying elements into the γ nanoparticles in the NiFeCoCrTi0.2 HEA by APT [22]. In this work, the γ" nanoparticles in a Ni2CoFeCrNb0.15 HEA was investigated by APT. Different partitioning behaviors of the alloying elements into the γ" nanoparticles were observed. It was found that the Co element tends to partition into the γ" nanoparticles but Fe and Cr elements are largely depleted from the γ" nanoparticles. The APT results were confirmed by the first-principles calculations from the perspective of the electronic states.

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

An ingot with a composition Ni2CoFeCrNb0.15 was produced by arc melting Fe, Co, Ni, Cr, and Nb metals with high purity (>99.9%) in an argon atmosphere. After repeatedly melted five times, the ingot was then drop-casted into a copper mold to make a slab with a dimension of 5 mm × 10 mm × 50 mm. Afterwards, the slab was solution-treated at 1473 K for 2 h, followed by water quenching. Then, the homogenized slab was cold rolled with a total thickness reduction of 70% and subsequently recrystallized at 1473 K for 4 minutes (min) and water-quenched. At last, aging was performed at 923 K for 40 h and 100 h, respectively, followed by water quenching.

The TEM specimen was prepared by mechanically grinding and followed by ion-milling using a precision ion polishing system (PIPS, Model 695, Gatan, Pleasanton, CA, USA). The TEM (TEM, JEOL 2100F, Tokyo, JAPAN) was operated under 200 keV. Needle specimens for APT analysis were prepared by gallium focused-ion-beam (FIB), with a FIB-SEM dual-beam system (Scios, FEI, Hillsboro, OR, USA), using a conventional lift-out technique [23]. The APT analysis was performed using a local electrode atom probe (LEAP5000 XR, CAMECA, Madison, WI, USA). The samples were run in the voltage mode at a specimen temperature of 50 K, with 200 kHz pulses at a pulse fraction of 20%. An Integrated Visualization and Analysis Software (IVAS, Version 3.8.2) protocol was employed to reconstruct the 3D atomic maps [24].

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

To confirm the formation of D022-structured γ" nanoparticles, the TEM analysis was performed before the APT measurement. Figure 1 shows a bright-field (BF) and a dark-field (DF) TEM images of the sample aged for 40 h. The DF-TEM image recorded from the spot marked with a yellow circle in the inset selected area diffraction pattern (SADP). The nanoparticles can be clearly observed in both the BF-TEM and the DF-TEM images. From the SADP it can be confirmed that the matrix has an FCC structure (γ phase), whilst the nanoparticles have a D022 structure (γ" phase) which is revealed by the additional faint spots. Figure 1c shows the size distribution of the nanoparticles with an average size of 13.5 ± 2.9 nm. The size of the nanoparticles was measured from the length of the nanoparticle along their long axis.

Figure 2 shows the APT results of the sample aged for 40 h. In the three-dimensional (3D) atom map, the γ" nanoparticles are delineated by 50 at.% Ni iso-concentration surfaces in red. It can be observed that the γ" nanoparticles are disk-like. From the sliced atom maps, it can be found that the γ" nanoparticles mainly consist of Ni and Nb. In addition, Co element shows a strong tendency to partition into the γ" nanoparticles, but Fe and Cr elements are largely depleted from the γ" nanoparticles. To clarify the accurate composition of these γ" nanoparticles, the proximity histogram, which is calculated over the iso-concentration surfaces, is plotted (Figure 2b). Therein, the chemical elements are displayed as a function of the distance from the iso-concentration surfaces. The proximity histogram shows that the average concentration of Co is up to 8.2 ± 0.3 at.%, but the average concentrations of Fe and Cr are only 1.3 ± 0.4 and 1.6 ± 0.1 at.% in the γ" nanoparticles, respectively. To investigate the composition stability of γ" nanoparticles, the sample aged for 100 h was also analyzed by APT. The composition of the γ" nanoparticles are summarized in Table 1. It was

found that the composition of γ" nanoparticles in the sample aged for 100 h is almost the same as that in the sample aged for 40 h, which indicates that the composition of γ" nanoparticles have reached the steady state after 40 h aging. The APT results demonstrate that the Co element is much easier to partition into the γ" nanoparticles than Fe and Cr elements. In addition, Co, Fe, and Cr prefer Ni sublattice sites in the Ni3Nb-type γ" nanoparticles because the Ni composition in the γ" nanoparticles are only 65 at.%. It should be noted that a small part of Co, Fe, or Cr atoms may also occupy Nb sublattice sites as the Nb composition in the γ" nanoparticles is about 24%. However, it is difficult to determine which kind of elements occupied Nb sublattice sites only from APT results. Some similar results are also found in the Ni-based superalloys [25–27]. For example, Lawitzki et al. reported that the alloying elements such as Cr occupied both the Ni and Nb sublattice sites of γ" phase in the 718 alloy [25].

**Figure 1.** The BF-TEM (**a**); and the DF-TEM (**b**) images of the sample aged for 40 h. The inset in (**b**) is the SADP along the zone-axis z = [001]. (**c**) The size distribution of the γ" nanoparticles.

**Figure 2.** (**a**) The 3D atom map (62 <sup>×</sup> <sup>64</sup> <sup>×</sup> 80 nm3) and the 4 nm-thick sliced atom maps of Co, Fe, Cr, Ni, and Nb of the 40 h aged sample. In the 3D map, the nanoparticles are delineated by 50 at.% Ni iso-concentration surfaces in red for better illustration. (**b**) The proximity histogram of the iso-concentration surfaces illustrated in the 3D atom map. The alloying elements are shown as a function of the distance from the iso-concentration surface (vertical dashed line).


**Table 1.** Chemical composition of γ" nanoparticles (at.%).

First-principles calculations were performed to confirm the site substitution preferences in the γ" nanoparticles and to investigate the origin of the different partitioning behaviors of the Co, Fe, and Cr elements into the γ" nanoparticles. The calculations employed the plane-wave pseudopotential approximations with the generalized gradient approximations, as implemented in the Vienna ab initio simulation package (VASP) [28]. A plane wave cutoff energy of 500 eV and 9 × 9 × 9 Monkhorst–Pack k-point grids were used in the calculation. A 3D periodic supercell with D022-structured Ni24Nb8 was employed to determine the total energies of the cells. The D022-structured Ni3Nb was fully relaxed, and the lattice parameters were determined to be a = b = 3.643 Å and c = 7.484 Å, which is in good agreement with both the previous reported experimental and theoretical results [27,29].

The formation energies for an element X (X = Co, Fe, and Cr) to substitute a Ni site and a Nb site of the D022-structured Ni3Nb were defined as [30]

$$E\_{X \to Ni} = \left( E\_{Ni\_{23}XNb\_8}^{tot} + \mu\_{Ni} \right) - \left( E\_{Ni\_{24}Nb\_8}^{tot} + \mu\_X \right) \tag{1}$$

$$E\_{X \to Nb} = \left( E\_{Ni\_{24}Nb\gamma X}^{tot} + \mu\_{Nb} \right) - \left( E\_{Ni\_{24}Nb\chi}^{tot} + \mu\_X \right) \tag{2}$$

where *Etot* is the total energy and *μ* is the chemical potential. The chemical potential is defined as the energy per atom of the element in its stable pure phase. Our calculations show that the total energy of *Ni*24*Nb*<sup>8</sup> (*Etot Ni*24*Nb*<sup>8</sup> ) is −222.84 eV and the chemical potentials of Ni and Nb are −5.47 and −10.20 eV, respectively. Table 2 summarized the total energies, the chemical potentials, and the formation energies. The calculation results demonstrate that Co and Fe atoms prefer to occupy the Ni sublattice sites rather than Nb sublattice sites, as the formation energies, *ECo/Fe*→*Ni*, are significantly lower than *ECo/Fe*→*Nb*. However, the formation energies for Cr to occupy Ni and Nb sublattice sites are almost same, which indicates that Cr atoms occupy both the Ni and Nb sublattice sites. In addition, the formation energy of Co that occupies the Ni sublattice site is nearly zero, which is much lower than that of Fe and Cr, indicating that Co is more stable in the D022-structured Ni3Nb than Fe and Cr. Similar results were also reported in the L12-structured Ni3Ti phase [31]. The calculation results confirmed the APT observation that the concentration of Co in the γ" nanoparticles are much higher than that of Fe and Cr.

**Table 2.** The Calculated chemical potentials, total energies and formation energies with the unit of eV.


To further clarify the origin of the formation energy differences of Co, Fe, and Cr in the D022-structured Ni3Nb, the total energy, and the chemical potential of the solute atoms are carefully checked, as the formation energy is determined by these two parts. It is found that the formation energy difference between Fe and Cr is mainly caused by the difference of their chemical potentials because the total energies of *Etot Ni*23*FeNb*<sup>8</sup> and *<sup>E</sup>tot Ni*23*CrNb*<sup>8</sup> are almost the same. However, both the chemical potential and the total energy of Co (*Etot Ni*23*CoNb*<sup>8</sup> ) are higher than that of Fe and Cr, which indicates that the formation energy differences between Co and Fe/Cr not only result from their chemical potential differences, but also come from their total energy differences. As the total energy originates from the charge distribution of the system, we calculated the charge density difference of Ni23CoNb8, Ni23FeNb8, and Ni23CrNb8 systems with reference to Ni24Nb8 system, respectively. From the charge density difference, we can find that the charge accumulation appears between Co and Nb atoms (Figure 3a) but does not appear between Fe and Nb (Figure 3b), or Cr and Nb (Figure 3c). The charge distribution results indicate that the bonding for Co and Nb is much stronger than that of Fe and Cr. The strong Co-Nb bond stabilizes Co atoms in the D022-structured Ni3Nb, which demonstrates that the bonding state of Co plays an important role in lowering the formation energy compared with that of Fe and Cr.

**Figure 3.** The charge density difference on the (001) plane of the Ni23CoNb8 (**a**), Ni23FeNb8 (**b**), and Ni23CrNb8 (**c**) systems with reference to the Ni24Nb8 system. The yellow regions and blue regions correspond to the increased and decreased charge density (0.001 eV/Bohr<sup>−</sup>3), respectively.

#### **4. Conclusions**

In summary, the partitioning of alloying elements into the γ" nanoparticles in a Ni2CoFeCrNb0.15 HEA was studied by APT and first-principles calculations. It was found that the composition of Co in the γ" nanoparticles is up to 8.2 at.% but the composition of Fe and Cr are only 1.3 at.% and 1.6 at.%, respectively. This indicates that the Co element is much easier to partition into the γ" nanoparticles than Fe and Cr elements. In addition, Co and Fe atoms prefer to occupy the Ni sublattice sites but Cr occupies both Ni and Nb sublattice sites. The first-principles calculations demonstrated that the different partitioning behaviors of Co, Fe, and Cr elements into the γ" nanoparticles are attributed to the differences of their specific chemical potentials and the bonding states in the γ" phase. This research paves the way for the composition control of the γ" nanoparticles in the HEAs.

**Author Contributions:** B.H., J.W., and F.H. conceived and designed the experiments; B.H. performed the APT experiment under the supervision of J.J.K.; J.W. performed the first-principles calculations under the supervision of W.Z. and A.H.; F.H. designed and prepared the bulk samples under the supervision of Z.J.W.; and D.C. performed the TEM analysis. All authors discussed the results and approved the final manuscript.

**Funding:** This work was supported by the Hong Kong Research Grant Council for financial support (Grant Nos. CityU 21202517 and 11212915), National Natural Science Foundation of China (Grant Nos. 11605148 and 51771149), and the Collaborative Research Funds (CityU 11205515 and 11209314) from the Research Grant Council, Hong Kong.

**Acknowledgments:** We would like to express our thanks to Junhua Luan for technical support.

**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* **The Effects of Mo and Nb on the Microstructures and Properties of CrFeCoNi(Nb,Mo) Alloys**

#### **Chun-Huei Tsau \* and Meng-Chi Tsai**

Institute of Nanomaterials, Chinese Culture University, Taipei 111, Taiwan; asd99586@yahoo.com.tw **\*** Correspondence: chtsau@staff.pccu.edu.tw

Received: 29 July 2018; Accepted: 27 August 2018; Published: 29 August 2018

**Abstract:** The effects of niobium and molybdenum additions on the microstructures, hardness and corrosion behaviors of CrFeCoNi(Nb,Mo) alloys were investigated. All of the CrFeCoNi(Nb,Mo) alloys displayed dendritic microstructures. The dendrites of CrFeCoNiNb and CrFeCoNiNb0.5Mo0.5 alloys were a hexagonal close packing (HCP) phase and the interdendrites were a eutectic structure of HCP and face-centered cubic (FCC) phases. Additionally, the dendrites of CrFeCoNiMo alloys were a simple cubic (SC) phase and the interdendrites were a eutectic structure of SC and FCC phases. The volume fraction of dendrites and interdendrites in these alloys were calculated. The influences of the volume fraction of dendrite in the alloys on the overall hardness were also discussed. The CrFeCoNiNb alloy had the larger volume fraction of dendrite and thus had the highest hardness among these alloys. The CrFeCoNi(Nb,Mo) alloys also showed better corrosion resistances in 1 M H2SO4 and 1 M NaCl solutions by comparing with commercial 304 stainless steel. The CrFeCoNiNb0.5Mo0.5 alloy possessed the best corrosion resistances in these solutions among the CrFeCoNi(Nb,Mo) alloys.

**Keywords:** CrFeCoNi(Nb,Mo); microstructure; hardness; corrosion; sulfuric acid; sodium chloride

#### **1. Introduction**

High-entropy alloys (HEA) has been announced for more than ten years [1–3]. The concept of HEA provides a new field for alloys design and thus becomes a very important field of materials development. This high-entropy alloy concept is now widely used to develop the high-performance alloys [4,5] and refractory alloys [6,7], also it is also applied in the thin film processes [8–10]. All of these researches are focus on the unique properties of the high-entropy alloys. Corrosion resistance is an important property of high-entropy alloys for structural applications; and many high-entropy alloys possess good corrosion resistances in different solutions are reported, such as FeCoNiCrCu*<sup>x</sup>* high-entropy alloys in 3.5% sodium chloride solution [11], Al7.5Cr22.5Fe35Mn20Ni15 high-entropy alloy in different solutions [12] and Al0.5CoCrFeNi alloy in a 3.5% NaCl solution [13]. That is, the high-entropy alloy concept is used to develop structural alloys with good corrosion resistance.

CrFeCoNi alloy has a very good corrosion resistance property is reported in our previous study [14]. It has a granular FCC structure and some HCP precipitates. However, the hardness of CrFeCoNi alloy is too low (HV144) to limit its structural application. Molybdenum has a benefit on the corrosion resistance of stainless steels is well known [15,16]. Niobium also can improve the corrosion resistance of FeCuNbSiB amorphous alloys [17]. Therefore, this study adds Mo and/or Nb into CrFeCoNi alloy and tests their properties. The microstructures revolution, hardness and polarization behaviors of corrosion of the CrFeCoNi(Nb,Mo) alloys in H2SO4 and NaCl solutions are all tested to evaluate their commercial application.

#### **2. Experimental**

The CrFeCoNiNb, CrFeCoNiNb0.5Mo0.5 and CrFeCoNiMo alloys were prepared by arc melting using appropriate amounts of the elements with purities above 99.9%. The alloys were made under a partial pressure of argon atmosphere (400 torrs). The bottoms were remelted at least 4 times to ensure homogeneity. Table 1 lists the chemical compositions of the alloys, the maximum deviation of each element in the alloys was less than 1 atomic percent. The microstructural evolution of the as-cast alloys was observed using a field emission scanning electron microscope with an energy dispersive spectrometer (SEM/EDS, JEOL JSM-6335, JEOL Ltd., Tokyo, Japan), which was operated at 15 kV. The structures were characterized by X-ray diffraction (XRD) using a Rigaku ME510-FM2 (Rigaku Ltd., Tokyo, Japan) with Cu-K (with a wavelength of 1.5406 Å) radiation operated at 30 kV at a scanning rate of 0.04 degree/s. The microstructures and lattice images of the alloys were obtained using a high-resolution transmission electron microscope (HREM, JEOL JEM-3000F, JEOL Ltd., Tokyo, Japan), which was operated at 300 kV. The corresponding diffraction patterns (DP) were obtained from the high-resolution lattice images by fast Fourier transformation (FFT) in Gatan digital micrograph software. The hardness of the alloys was measured using both a Mitutoyo Akashi MVK-G1500 microhardness tester (Mitutoyo Co., Kanagawa, Japan) under a load of 10 gf and a Matsuzawa Seiki MV1 Vicker's hardness tester (Matsuzawa Co., Akita, Japan) under a load of 30 kgf.

**Table 1.** The average chemical compositions of the as-cast CrFeCoNi(Nb,Mo) alloys analyzed by SEM/EDS.


Polarization curves of the as-cast alloys were obtained in a potentiostat/galvanostat (Autolab PGSTAT302N, Metrohm Autolab B.V., Utrecht, The Netherlands) using a three-electrode system at a scanning rate of 1 mV/s. The CrFeCoNi(Nb,Mo) alloys for polarization testing were mounted in epoxy resin and the exposed surface area of each was fixed at 19.64 mm<sup>2</sup> (with a diameter of 5 mm). The reference electrode was a saturated silver chloride electrode (Ag/AgCl) and the counter electrode was a smooth Pt sheet. All the potentials that are below a saturated silver chloride electrode (SSE), whose reduction potential is 222 mV higher than that of the standard hydrogen electrode (SHE) at 25 ◦C [18]. The specimens whose polarization curves were obtained were all mechanically wet-polished using 1200 SiC grit paper. Test solutions with a concentration of 1 M were prepared from reagent-grade sulfuric acid (H2SO4) and sodium chloride (NaCl) that were dissolved in distilled water. To eliminate any effect of dissolved oxygen, the solutions were deaerated by bubbling nitrogen gas through them before and during the polarization experiments. The polarization test started after the specimen, counter electrode and reference electrode were placed in the bubbling solution for 900 s.

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

The microstructures of as-cast CrFeCoNi(Nb,Mo) alloys are displayed in Figure 1. They indicated that the Nb and Mo additions could change the granular microstructures of CrFeCoNi alloy to a dendritic microstructure of CeFeCoNi(Nb,Mo) alloys; and their interdendrites all showed a eutectic structure. Also, some precipitates were only observed in the as-cast CrFeCoNiNb alloy, shown in Figure 1a; the precipitates were not found in the other two alloys under the as-cast state. The chemical compositions of the phases in these CrFeCoNi(Nb,Mo) alloys are listed in Table 2. The deviation of each element-content in the phases were less than 1 atomic percent. The dendrites of these three alloys had higher niobium and/or molybdenum contents. Additionally, the precipitates in CrFeCoNiNb alloy had similar compositions with the HCP dendrites but the lattice constants of these two phases were quite different.

**Figure 1.** The SEM micrographs of as-cast (**a**) CrFeCoNiNb; (**b**) CrFeCoNiNb0.5Mo0.5; and (**c**) CrFeCoNiMo alloys.

Figure 2 shows the XRD patterns of the as-cast CrFeCoNi(Nb,Mo) alloys. The lattice constants of the phases are also marked in the figure. It indicated that every alloy had two major phases. The major phases in the as-cast CrFeCoNiNb and CrFeCoNiNb0.5Mo0.5 alloys were the FCC and HCP phases, the lattice constants of both the FCC and HCP phases in these two alloys were also very close. No peak of the precipitates was found in the as-cast CrFeCoNiNb alloy because its volume fraction was too small. However, the phases in the as-cast CrFeCoNiMo alloy were an FCC phase and a simple cubic

(SC) phase, the SC phase (isometric cubic) had a large lattice constant of 8.398 Å which indicated that a unit cell had many atoms.


**Table 2.** The average chemical compositions of the phases in the CrFeCoNi(Nb,Mo) alloys analyzed by SEM/EDS.

**Figure 2.** X-ray diffraction (XRD) patterns of the as-cast CrFeCoNi(Nb,Mo) alloys.

The transmission electron microscopy (TEM) bright field (BF) images of the dendrite, the matrix of the interdendrite and the precipitate of as-cast CrFeCoNiNb alloy are shown in Figure 3a–c, respectively. Figure 3a displays the image of the dendrite and inserts are the corresponding lattice image and the FFT DP, which were taken from the zone axis of [0110]; and the FFT DP indicates that the dendrite was a single HCP phase. Figure 3b is the TEM image of the matrix of interdendrite in the alloy, the inserts are the corresponding lattice image and FFT DP, which were taken from the zone zxis of [011]; and this FFT DP indicated that it was an FCC structure. Figure 3c shows the TEM image of the precipitate and its corresponding lattice image and FFT DP, which were taken from the zone axis of

[0111]. The precipitates showed a HCP structure and its lattice constants of *a*- and *c*-axes were 2.88 and 4.70 Å, respectively.

**Figure 3.** Transmission electron microscopy (TEM) (bright field) BF images of the phases in as-cast CrFeCoNiNb alloy: (**a**) an image of dendrite, inserts are the corresponding lattice image and fast Fourier transformation (FFT) diffraction patterns (DP) taken from the zone axis of [0110] which shows a hexagonal close packing (HCP) structure; (**b**) an image of interdendrite, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [011] which shows an FCC structure; and (**c**) an image of precipitate, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [0111] which shows a HCP structure.

Figure 4a shows the TEM BF images of the dendrite in CrFeCoNiNb0.5Mo0.5 alloy, inserts are the corresponding lattice image and FFT DP, which were taken from the zone axis of [2423]; and this FFT DP indicated that this phase was a HCP structure. Figure 4b show the TEM BF image of the interdendrite in CrFeCoNiNb0.5Mo0.5 alloy, inserts are the corresponding lattice image and FFT DP, which were taken from the zone axis of [011]; and this FFT DP indicated an FCC structure.

**Figure 4.** TEM BF images of the phases in as-cast CrFeCoNiNb0.5Mo0.5 alloy: (**a**) an image of dendrite, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [2423] which shows a HCP structure; (**b**) an image of interdendrite, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [011] which shows an FCC structure.

Figure 5a shows a TEM BF image of the dendrite of as-cast CrFeCoNiMo alloy, inserts are the corresponding lattice image and FFT DP, which were taken from the zone axis of [001]. The image indicated that the dendrite was a single phase. However, the lattice image indicated that the unit cell of this phase had a large lattice constant of 8.398 Å. Therefore, the diffraction spots of the FFT DP were very close which indicated that this phase had a large lattice constant. Additionally, the lattice points from the lattice image showed only 1-fold symmetry and thus the unit cell had a SC structure (i.e., an isometric cubic structure) with a large lattice constant. This also meant the SC structure had a complex structure, this complex structure needs further investigation. Figure 5b displays the image of the eutectic interdendrite in the as-cast CrFeCoNiMo alloy. Both of the corresponding lattice image and FFT DP taken from the zone axis of [112] confirmed that the matrix of interdendrite was an FCC structure.

**Figure 5.** TEM BF images of the phases in as-cast CrFeCoNiMo alloy: (**a**) an image of dendrite, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [001] which shows a SC structure; (**b**) an image of interdendrite, inserts are the corresponding lattice image and FFT DP taken from the zone axis of [112] which shows an FCC structure.

The influence of niobium and molybdenum additions on the volume fractions of the dendrites and interdendrites of these alloys were different. Table 3 lists the volume fraction of the dendrites of as-cast CrFeCoNi(Nb,Mo) alloys. To determine of the volume fraction of dendrites was by drawing arbitrary lines in photos and measuring the intercept lengths of intercepted dendrites. From this, the volume fraction of the dendrites was calculated by the equation [19]:

$$V\_d = L\_d = \frac{\sum L\_a}{L\_T} \tag{1}$$

where *Vd* is the volume fraction of the dendrites, *Ld* is the linear fraction of the dendrites, *La* is the intercepted length of each dendrite and *LT* is the total length. The results indicated that the CrFeCoNiNb alloy had more volume fraction of dendrites; and molybdenum addition would decrease the volume fraction of dendrites.


**Table 3.** The volume fraction of the dendrites in the as-cast CrFeCoNi(Nb,Mo) alloys.

The overall hardness of as-cast CrFeCoNi(Nb,Mo) alloys and the microhardness of the dendrites and interdendrites of the alloys are list in Table 4. The hardness of as-cast CrFeCoNi alloy only had HV 144. However, the hardness increased sharply after additions of niobium and/or molybdenum. Because of niobium and molybdenum had larger atomic radiuses. The atomic radiuses of Cr, Fe, Co, Ni, Nb and Mo are 0.128, 0.124, 0.125, 0.125, 0.143 and 0.140 nm, respectively [20]. Therefore, the hardness increased significantly because of the larger lattice distortion and forming the different phases after additions of niobium and molybdenum. The interdendrites of the alloys had almost the same hardness (about HV 400). The dendrites of the alloys were harder than the interdendrites (about HV 700). The overall hardness of an alloy was contributed by the volume fractions of the dendrite (the hard part) and the interdendrite (the soft part) in this alloy. Therefore, decreasing the volume fraction of the hard part, that is, the dendrites, would result in decreasing the overall hardness of the alloy. Both the dendrites and interdendrites of CrFeCoNiNb0.5Mo0.5 alloy had the highest hardness but the overall hardness of this alloy was lowest among these alloys. This was contributed by the lowest volume fraction of dendrite (the hard part) in the CrFeCoNiNb0.5Mo0.5 alloy. On the contrary, the CrFeCoNiNb alloy had the highest volume fraction of dendrite and thus had the highest overall hardness among these alloys.



The polarization behaviors of the as-cast CrFeCoNi(Nb,Mo) alloys and 304 stainless steel in 1M deaerated H2SO4 solution at 30 ◦C are shown in Figure 6. The polarization data were also compared with those of commercial 304 stainless steel (304SS) whose composition was by weight 71.61% Fe, 18.11% Cr, 8.24% Ni, 1.12% Mn, 0.75% Si, 0.05% Co, 0.02% Mo, 0.05% C, 0.03% P and 0.02% S. The important data of these polarization curves are listed in Table 5. The corrosion potential (*E*corr) of CrFeCoNi(Nb,Mo) alloys were very close and all nobler than that of 304SS. Table 6 lists the standard electrode potential of selected elements [21]. The standard electrode potential of niobium is lower than that of molybdenum, which means that the niobium is more active than molybdenum. Therefore, the CrFeCoNiMo alloy had the highest *E*corr and the CrFeCoNiNb alloy had the lowest *E*corr among these three alloys. The polarization curve below *E*corr was the cathodic polarization curve; and the curve above *E*corr was the anodic polarization curve. The corrosion current densities (*i*corr) of CrFeCoNi(Nb,Mo) alloys were less or equal to the *i*corr of 304SS. The polarization curve of 304 stainless steel displayed a large anodic peak; and the anodic peaks of CrFeCoNi(Nb,Mo) alloys were significantly less than that of 304SS and thus CrFeCoNi(Nb,Mo) alloys had lower passivation potential (*E*pp) and anodic critical current density (*i*crit). This meant that the CrFeCoNi(Nb,Mo) alloys were easy to enter passivation regions and form passive films during corrosion in H2SO4 solution by comparing with 304 stainless steel. The large anodic peak of 304 stainless steel in H2SO4 solution was caused by formation of iron hydroxide and higher oxides of iron and chromium [22]. The lowest current densities of the passivation regions (*i*pass) of these alloys were around 10–20 A/cm2. Additionally, the main

passivation region of CrFeCoNi(Nb,Mo) alloys were broader than that of 304SS. All of these alloys had a similar breakdown potential (*E*b) of about 1 V (SSE).

**Figure 6.** Polarization curves of CrFeCoNi(Nb,Mo) alloys and 304 stainless steel in 1 M deaerated H2SO4 solution at 30 ◦C.

**Table 5.** Polarization data of CrFeCoNi(Nb,Mo) alloys and 304 stainless steel in 1 M deaerated H2SO4 solution at 30 ◦C.


**Table 6.** Standard electrode potential at 25 ◦C [21].


The micrographs of CrFeCoNi(Nb,Mo) alloys after the polarization test in 1 M deaerated H2SO4 solution at 30 ◦C are shown in Figure 7. Both of the dendrites and interdendrites of CrFeCoNiNb alloy were significantly corroded after test, as shown in Figure 7a; but the FCC phase (the matrix of interdendrite) was severely corroded than the HCP phase. On the contrary, only the FCC phase (the matrix of interdendrite) of CrFeCoNiNb0.5Mo0.5 alloy was slightly corroded, the HCP phase almost maintained its original shape, as shown in Figure 7b. This also proved that the CrFeCoNiNb0.5Mo0.5 alloy had the minimum *i*corr among these CrFeCoNi(Nb,Mo) alloys, as listed in Table 5. The micrograph of CrFeCoNiMo alloy also displayed a severely corroded surface after polarization test, as shown in Figure 7c. Also, the FCC phase (the matrix of interdendrite) was more corroded than the SC phase. Therefore, in the local cells of the CrFeCoNi(Nb,Mo) alloys, the FCC phase in the interdendrites behaved as an anode and another phase (e.g., the HCP phase of CrFeCoNiNb and CrFeCoNiNb0.5Mo0.5 alloys and the SC phase of CrFeCoNiMo alloy) behaved as a cathode.

**Figure 7.** SEM micrographs of the alloys after polarization test in 1 M deaerated H2SO4 solution at 30 ◦C, (**a**) CrFeCoNiNb alloy; (**b**) CrFeCoNiNb0.5Mo0.5 alloy; and (**c**) CrFeCoNiMo alloy.

The polarization curves of the as-cast CrFeCoNi(Nb,Mo) alloys in 1 M deaerated NaCl solution at 30 ◦C are shown in Figure 8. The values of *E*corr and *i*corr of these alloys are listed in Table 7. All of these data are also compared with commercial 304 stainless steel. The *i*corr of these four alloys were also close. In addition, the *E*corr of the CrFeCoNi(Nb,Mo) alloys were very close and much nobler than 304 stainless steel. However, the passivation regions of CrFeCoNi(Nb,Mo) alloys were much broader than that of 304 stainless steel. Both of the CrFeCoNiMo and CrFeCoNiNb0.5Mo0.5 alloys had significantly anodic peaks and passivation regions; but CrFeCoNiNb and 304 stainless steel did not display anodic peaks. Adding Mo can reportedly increase the corrosion resistance of the alloy in a solution that contains chloride ions because molybdenum can increase the stability of the passivation films of steels [15,23]. The polarization curves of CrFeCoNi(Nb,Mo) alloys indicated that the increasing of Mo-content resulted in forming anodic peak and passivation regions. In addition, the cathodic limiting current densities (*i*L) were observed in the polarization curves of CrFeCoNiNb and CrFeCoNiMo alloys. The cathodic limiting current density (*i*L) related to the maximum reaction rate, which was limited by the diffusion rate of hydroxyl ions (OH−) in solution [18].

**Figure 8.** Polarization curves of CrFeCoNi(Nb,Mo) alloys and 304 stainless steel in 1 M deaerated NaCl solution at 30 ◦C.

**Table 7.** Polarization data (*i*corr and *E*corr) of CrFeCoNi(Nb,Mo) alloys and 304 stainless steel in 1 M deaerated NaCl solution at 30 ◦C. ǂ


The micrographs of the as-cast CrFeCoNi(Nb,Mo) alloys after polarization test in 1 M deaerated NaCl solution at 30 ◦C are shown in Figure 9. Similar to the results of these alloys tested in 1 M deaerated H2SO4 solution at 30 ◦C, the major corroded areas of these alloys were the matrixes of the interdendrites (FCC phase) of CrFeCoNi(Nb,Mo) alloys. On the contrary, almost no corrosion occurred on the dendrites of these CrFeCoNi(Nb,Mo) alloys. Furthermore, no deep-type pitting was observed indicated that these alloys had a good corrosion resistance in NaCl solution. This also proved that molybdenum could improve the localized corrosion resistance.

**Figure 9.** *Cont*.

**Figure 9.** SEM micrographs of the alloys after polarization test in 1 M deaerated NaCl solution at 30 ◦C, (**a**) CrFeCoNiNb alloy; (**b**) CrFeCoNiNb0.5Mo0.5 alloy; and (**c**) CrFeCoNiMo alloy.

#### **4. Conclusions**

All of the CrFeCoNi(Nb,Mo) alloys displayed dendritic microstructures. The major two phases of CrFeCoNiNb and CrFeCoNiNb0.5Mo0.5 alloys were the HCP and FCC phases, where the dendrites were a single HCP phase. The major two phases of CrFeCoNiMo alloys were the SC and FCC phases, where the dendrites were a single SC phase. All of the interdendrites in the CrFeCoNi(Nb,Mo) alloys were eutectic structures.

The microhardness and overall hardness of CrFeCoNi(Nb,Mo) alloys increased by comparing with CrFeCoNi alloy because the elements of niobium and molybdenum had larger atomic radiuses. The microstructures significantly influenced the overall hardness of these alloys. The highest overall hardness of CrFeCoNiNb alloy was caused by its larger volume fraction of the dendrites. On the contrary, the lowest overall hardness of CrFeCoNiNb0.5Mo0.5 alloy was caused by its less volume fraction of the dendrites.

The corrosion resistances of CrFeCoNi(Nb,Mo) alloys in 1M deaerated H2SO4 and NaCl solutions were better than commercial 304 stainless steel. Additionally, the CrFeCoNiNb0.5Mo0.5 alloy had the best corrosion resistances in these solutions form the polarization curves and the micrographs after corrosion test. In these CrFeCoNi(Nb,Mo) alloys, the FCC phase behaved as an anode of the local cell in the alloy and was thus severely corroded than another phase in 1M deaerated H2SO4 and NaCl solutions.

**Author Contributions:** C.-H.T. conceived and designed the experiments; M.-C.T. performed the experiments; M.-C.T. and C.-H.T. analyzed the data; C.-H.T. contributed reagents/materials/analysis tools; C.-H.T. wrote the paper. Both authors have read and approved the final manuscript.

**Acknowledgments:** We are grateful to the Ministry of Science and Technology of Republic of China for its financial support under the project MOST 106-2221-E-034-008.

**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 Molybdenum Additives on Corrosion Behavior of (CoCrFeNi)100***−x***Mo***x* **High-Entropy Alloys**

#### **Wenrui Wang 1,\*, Jieqian Wang 1, Honggang Yi 1, Wu Qi <sup>1</sup> and Qing Peng <sup>2</sup>**


Received: 6 October 2018; Accepted: 26 November 2018; Published: 28 November 2018

**Abstract:** The present work investigates the influence of micro-alloyed Mo on the corrosion behavior of (CoCrFeNi)100−*x*Mo*<sup>x</sup>* high-entropy alloys. All of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys exhibit a single face-centered cubic (FCC) solid solution. However, the (CoCrFeNi)97Mo3 alloy exhibits an ordered sigma (*σ*) phase enriched in Cr and Mo. With the increase of *x* (the Mo content) from 1 to 3, the hardness of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys increases from 124.8 to 133.6 Vickers hardness (HV), and the compressive yield strength increases from 113.6 MPa to 141.1 MPa, without fracture under about a 60% compressive strain. The potentiodynamic polarization curve in a 3.5% NaCl solution indicates that the addition of Mo has a beneficial effect on the corrosion resistance to some certain extent, opposed to the *σ* phase. Furthermore, the alloys tend to form a passivation film in the 0.5 M H2SO4 solution in order to inhibit the progress of the corrosion reaction as the Mo content increases.

**Keywords:** (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys; high entropy alloy; microstructure; mechanical properties; corrosion behavior

#### **1. Introduction**

Traditional alloys only have one major element as a matrix [1]. With the increase of the amount of alloying elements and the concentration of minor elements, the alloy forms a fragile intermetallic phase, which not only increases the difficulty of the microstructure, but also may result in a reduction in the mechanical performance [2–4]. To overcome these difficulties, high entropy alloys (HEAs) are invented with extensive research interests; HEAs usually compose of five or more kinds of major elements, with the concentration of each principal element being between 5 and 35 at % [5–7]. HEAs tend to generate a face-centered cubic (FCC), body-centered cubic (BCC), or hexagonal closed-packed (HCP) multicomponent solid solution phase [8,9]. Some HEAs have been confirmed to achieve a series of excellent properties, such as high strength, high hardness, and glorious corrosion resistance [10–12].

The CoCrFeNi alloy has received extensive attention for its outstanding corrosion resistance, ductility, and structure stability [13,14]. However, because of the poor mechanical strength, the application of the CoCrFeNi alloy in engineering has been limited. It has been shown that an ordered σ strengthening phase can be formed by a certain amount of Mo additives, resulting in precipitation strengthening [15–17]. However, the excessive addition of Mo causes a large amount of the coarse *σ* strengthening phase, which may lead to a rapid increase of the alloy brittleness. Furthermore, because of the low electronic potential of Mo, the excessive content of Mo may reduce the corrosion resistance [18–20]. Referring to the chemical composition of austenitic stainless steel, the content of Mo in stainless steel is generally less than 3 wt %. Therefore, we also supply a small

amount of Mo in addition to the CoCrFeNi alloy, in order to study the corrosion resistance of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloy for the development of an HEA system, with a good performance of both strength and corrosion resistance.

In this paper, the as-cast (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* =1, 2, and 3 wt %) HEAs have been prepared mainly by vacuum arc melting. The excellent corrosion behavior was investigated by the electrochemical experiments.

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

#### *2.1. Samples Fabrication*

Elements Co, Cr, Fe, Ni, and Mo with purities of over 99.9 wt % were prepared as raw materials previously. The as-cast (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3, represented by Mo1, Mo2, and Mo3, respectively) were prepared by vacuum arc melting and were fabricated under the WK-∏ vacuum arc melting furnace at least five times in the crucible, so as to ensure the chemical homogeneity. The size of the ingot was approximately Φ 35 × 10 mm, and the ingot was annealed for 3 h under 500 ◦C and was cooled in the air so as to release the residual stress caused by rapid cooling during casting.

### *2.2. Microstructure of the (CoCrFeNi)100*−*xMox Alloys*

The crystalline phases of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys were identified by X-ray diffraction (XRD), using an Ultima IV X-ray diffractometer with Cu K*α* radiation. The X-ray diffractometer has an operating voltage of 30 kV and an operating current of 20 mA with the diffraction angle (2*θ*) from 20 to 90◦, at a scanning rate of 4◦/min.

Etching the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys with aqua regia (HNO3: HCl = 1:3, volume fraction) and using FE-SEM JEOL JEM-7600F scanning-electron microscopy (SEM) (JEOL Ltd., Tokyo, Japan) equipped with an energy dispersive spectrometer (EDS) allowed for an analysis of the microstructure and composition.

#### *2.3. Mechanical Properties*

The microhardness was measured using a Wolpert-401MVD Vickers hardness tester (WOLPERT Co., Norwood, MA, USA) with loads of 500 g and a duration of 10 s. The measurements were performed at 10 different locations on each sample, and the average value of the 10 measurements was calculated. Compressive tests were carried out on the Φ 3 × 6 mm samples, using the universal testing machine (CMT 4305) (MTS Co., Eden Prairie, MN, USA) with a strain rate of 10−<sup>3</sup> s<sup>−</sup>1.

#### *2.4. Electrochemical Corrosion Test*

The electrochemical experiments were performed on (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys, using the Potentiostat Workstation Versa STAT MC (PARSTAT 4000, AMETEK Co., Princeton, NJ, USA). A three-electrode electrochemical cell using a saturated calomel electrode (SCE) as a reference electrode, a platinum plate as an auxiliary electrode, and a sample as a working electrode were tested. The electrochemical experiments of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys were conducted in a 3.5% NaCl and 0.5 M H2SO4 solution at room temperature, respectively. The potentiodynamic polarization measurements were taken at a scan rate of 1 mV/s from a potential scanning range of −0.5 V to 1.1 V.

#### *2.5. Corroded Microstructure*

The electrochemically tested alloys were cleaned using an ultrasonic cleaner, and then dried in nitrogen. SEM and EDS were used to study the morphology of the corrosion surface of the high-entropy alloys.

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

#### *3.1. Microstructure of the (CoCrFeNi)100*−*xMox Alloys*

Figure 1 represents the XRD pattern of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, 3) alloys. All of the samples of alloys have a single FCC solid solution structure, which was confirmed by the predecessors [21]. As the Mo content increases, the peak intensity changes, but the FCC phase is kept. As the Mo content becomes 3 wt % (Mo3 structure), the small peak on the left of the matrix FCC phase in the XRD patterns is identified as Cr and Mo rich *σ* phase, which agrees with the previous remarks [22].

**Figure 1.** XRD (X-ray diffraction) patterns of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, 3) high entropy alloys.

The SEM images of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys presented in Figure 2 shows that the alloys are composed of typical dendritic structures. Region A is interdendrites and region B is dendrites. Figure 2c presents the SEM image of the Mo3 alloy, and region D is the grain boundary. Figure 2d is a part of the SEM of the Mo3 alloy. The EDS results of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys are shown in Table 1. According to the EDS, the dendrite is the Co and Fe rich phase, and the interdendrite is the Cr and Mo rich phase. Combined with the XRD and EDS results, the Mo3 alloy exhibits Cr and Mo rich σ phase in the interdendrite. When the content of Mo is 1 and 2 wt %, there is no formation of a precipitate phase in the alloy because of the high entropy effect. However, HEAs undergo spinodal decomposition inside the crystal grains during cooling, leading to the formation of microstructures with the same structure but different compositions.

**Figure 2.** SEM (scanning-electron microscopy) images of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* high-entropy alloys: (**a**) *x* = 1; (**b**) *x* = 2; (**c**) *x* = 3; (**d**) 10 × magnification.


**Table 1.** Element concentration determined using the energy dispersive spectrometer (EDS) of the three samples of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys (at %).

#### *3.2. Mechanical Properties*

Figure 3 shows the Vickers hardness (HV) of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys as a function of the Mo content. The alloy hardness increases from 124.83 to 133.60 HV, with the Mo content increasing from 1 to 3 wt %. When the content of the Mo element is 3 wt %, the presence of the σ phase results in a remarkable increase in the hardness.

**Figure 3.** Vickers hardness of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* high-entropy alloys as a function of the Mo content.

Besides the hardness, we have examined the stress–strain relationships. The compressive stress–strain curves and the inner longitudinal-section SEM images of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys are shown in Figure 4. After yielding, the strength of the alloys increases continuously. All of the three samples do not break under about a 60% compressive strain, indicating that the alloys possess a good ductility, flexibility, and fracture strain. As shown in Figure 4b–d, the deformation of the Mo3 subgrain boundaries is more prominent in the angle of 45◦, probably due to the resolved shear stress.

Table 2 lists the mechanical properties of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys. The Mo1 and Mo2 alloys exhibit a similar behavior. The yield stress of the Mo3 alloy increases significantly because of the second-phase hardening by the σ phase, as reported in the literature [23,24].

**Figure 4.** (**a**) Compressive stress–strain curves and the inner longitudinal-section SEM images of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys after compression deformation: (**b**) *x* = 1; (**c**) *x* = 2; (**d**) *x* = 3.



#### *3.3. Environmental Effect on Corrosion Behavior*

#### 3.3.1. Corrosion Behavior in Chloride-Containing Solutions

Figure 5 lists the polarization curve of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in a 3.5% NaCl solution. The corrosion potential of the alloys gradually shifts to more positive potentials with a decreasing Mo content.

Table 3 presents the electrochemical parameters of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in a 3.5% NaCl solution. The corrosion current densities (*icorr*) of the Mo1, Mo2, and Mo3 alloys were 0.4, 0.24, and 6.6 <sup>μ</sup>A/cm2, and the corrosion potentials (*Ecorr*) were −199, −277, and −493 mV, respectively. The breakdown potential (*Eb*) gradually shifts to more positive potentials. *Epit* is a primary passivation potential. Δ*E* is the passive region width, defined as the difference between the *Eb* and *Epit*. The *icorr* value of the Mo3 alloy is an order of magnitude higher than the other two. The corrosion resistance of Mo3 was dropped. Therefore, the presence of the Cr and Mo rich *σ* phase in the Mo3 alloy leads to the diminution of the corrosion resistance.

**Figure 5.** Polarization curves of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in a 3.5% NaCl solution. **Table 3.** Electrochemical parameters of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys in a 3.5% NaCl solution.


Figure 6 presents the microstructure of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys after potentiodynamic polarization in a 3.5% NaCl solution. DR is dendrite and IR is interdendrite. According to the SEM images, the majority of types of corrosion were mainly the pitting corrosion. The Mo3 alloy is more susceptible to pitting corrosion, which is consistent with the polarization curve results. Region A is a Cr-rich phase and region B is a Mo-rich phase. There is an element of segregation that causes the corrosion to occur. The interdendritic phase of Mo3 is the Cr- and Mo-rich phase, and the dendrite is a Cr- and Mo-depleted phase. Therefore, galvanic corrosion occurred at the junctions of dendrites. The XRD and SEM result show that the Mo and Cr rich *σ* phase appeared inside the interdendrite, the corrosion occurred at the interfaces around *σ* phase, as shown in Figure 6c,d. The results of the EDS of the alloy after corrosion are summarized in Table 4.

**Figure 6.** SEM images of (CoCrFeNi)100−*x*Mo*<sup>x</sup>* after potentiodynamic polarization in a 3.5% NaCl solution: (**a**) *x* = 1; (**b**) *x* = 2; (**c**) *x* = 3; (**d**) the interdendrite morphology of the Mo3 alloy.


**Table 4.** EDS results for the Mo3 alloy after the potentiodynamic polarization in a 3.5% NaCl solution.

#### 3.3.2. Corrosion Behavior in Acid Solutions

Figure 7 shows the polarization curve of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in 0.5 M H2SO4. Table 5 presents the electrochemical parameters of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in 0.5 M H2SO4. The *ipp* is the lunt current density. It can be seen from the electrochemical parameters that the *icorr* was 34.1, 28.0, and 15.4 μA/cm2, respectively. The *Ecorr* shifted to more positive potentials, and the *icorr* value dropped as the Mo content increased. This suggests that the alloys tend to form a passivation film to inhibit the progress of the corrosion reaction.

**Figure 7.** Polarization curves of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys in a 0.5 M H2SO4 solution

**Table 5.** Electrochemical parameters of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys in a 0.5 M H2SO4 solution.


Figure 8 shows the SEM microstructure of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys after potentiodynamic polarization in a 0.5 M H2SO4 solution, combined with EDS, because the low potential of Mo is enriched and the Cr2O3 is insufficient in region A. Consequently, region A is more susceptible to corrosion in the H2SO4 solution. When the content of Mo is 3 wt %, as Figure 8c indicates, the results of the EDS analysis show that the concentration of element Cr in region A is reduced, and the content of Mo in region B is significantly higher than that in region A, the effect of Mo was to form Mo (VI) oxyhydroxide or molybdate (MoO4 <sup>2</sup>−), decreasing the rate of dissolution in active zones. Hence, the corrosion of the Mo3 alloy is concentrated in region A.

**Figure 8.** SEM images of (CoCrFeNi)100−*x*Mo*<sup>x</sup>* after potentiodynamic polarization in a 0.5 M H2SO4 solution: (**a**) *x* = 1; (**b**) *x* = 2; (**c**) *x* = 3; (**d**) partial view of the Mo3 alloy.

Figure 9a shows the comparison of the corrosion behavior between the HEAs and the conventional corrosion resistant alloys in a 3.5% NaCl solution, compared with those of the conventional corrosion resistant alloys [25,26]. The HEAs are located in the upper part of Figure 9a, the *Ecorr* of the HEAs are more positive than those of the Mn alloys, Ni alloys, and some of the Ti alloys. On the other hand, the *icorr* of the HEAs are much lower than some of Mn alloys and are comparable with the Ti alloys, which indicates that the corrosion resistance of the HEAs is comparable or even better than those of the conventional alloys. However, the partial *icorr* of the HEAs is higher than that of the total, because the presence of the σ phase is catastrophic for HEAs. Figure 9b presents the comparison of the corrosion behavior between the HEAs and the conventional corrosion resistant alloys in the 0.5 M H2SO4 solution. Compared with the conventional alloys [27,28], the *Ecorr* of the HEAs are much more positive than those of the Ti alloys and Ni alloys. The *icor*<sup>r</sup> of the HEAs are much lower than the Ti alloys, Ni alloys, and some of the Cu alloys. As a general trend, the corrosion resistance of the HEAs in the 0.5 M H2SO4 solution is better than those of the conventional alloys.

**Figure 9.** Comparison of the *icorr* and *Ecorr* between high entropy alloys (HEAs) and conventional alloys: (**a**) in a 0.5 M NaCl solution; (**b**) in a 0.5 M H2SO4 solution.

#### **4. Conclusions**

The (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys have been synthesized. Their microstructures, mechanical properties, and corrosion behaviors have been experimentally investigated. The microstructures of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* (*x* = 1, 2, and 3) alloys belong to a single FCC structure. The increase in Mo promotes the formation of the Cr- and Mo-rich σ phase. The hardness and compressive yield strength increase obviously with an increase of the Mo content from 1 to 3 wt %. Regarding the potentiodynamic polarization curves of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys in a 3.5% NaCl solution, the

curves of the Mo1 and Mo2 alloys indicated that the increase of the Mo content increases the corrosion resistance of the chloride environment to some extent. However, the Cr and Mo rich *σ* phase is present at the grain boundary of the Mo3 alloy, resulting in a decreasing in corrosion resistance in the 3.5% NaCl solution. Furthermore, the potentiodynamic polarization curves of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys in the 0.5 M H2SO4 solution yielded an extensive passive region, and as the content of the Mo increased, the corrosion current density gradually decreased. Therefore, in an acidic solution, the addition of Mo has a positive effect on the corrosion resistance of the (CoCrFeNi)100−*x*Mo*<sup>x</sup>* alloys.

**Author Contributions:** Conceptualization, W.W.; formal analysis, J.W.; methodology, J.W. and W.Q.; resources, W.W.; writing (original draft), J.W.; writing (review and editing), J.W., H.Y. and Q.P.

**Funding:** The authors are grateful for the financial support provided by the Fundamental Research Funds for the Central Universities (FRF-TP-16-044A1 and FRF-GF-17-B18) and the National Natural Science Foundation of China (21703007).

**Acknowledgments:** The authors would like to thank the project team assistance with funding. Many thanks to the tutor for his thoughtful and thorough guidance.

**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 Zr Addition on the Microstructure and Mechanical Properties of CoCrFeNiMn High-Entropy Alloy Synthesized by Spark Plasma Sintering**

#### **Hongling Zhang 1, Lei Zhang 1, Xinyu Liu 2, Qiang Chen <sup>1</sup> and Yi Xu 1,\***


Received: 12 September 2018; Accepted: 16 October 2018; Published: 23 October 2018

**Abstract:** As a classic high-entropy alloy system, CoCrFeNiMn is widely investigated. In the present work, we used ZrH2 powders and atomized CoCrFeNiMn powders as raw materials to prepare CoCrFeNiMnZr*<sup>x</sup>* (*x* = 0, 0.2, 0.5, 0.8, 1.0) alloys by mechanical alloying (MA), followed by spark plasma sintering (SPS). During the MA process, a small amount of Zr (*x* ≤ 0.5) can be completely dissolved into CoCrFeNiMn matrix, when the Zr content is above 0.5, the ZrH2 is excessive. After SPS, CoCrFeNiMn alloy is still as single face-centered cubic (FCC) solid solution, and CoCrFeNiMnZr*<sup>x</sup>* (*x* ≥ 0.2) alloys have two distinct microstructural domains, one is a single FCC phase without Zr, the other is a Zr-rich microstructure composed of FCC phase, B2 phase, Zr2Ni7, and σ phase. The multi-phase microstructures can be attributed to the large lattice strain and negative enthalpy of mixing, caused by the addition of Zr. It is worth noting that two types of nanoprecipitates (body-centered cubic (BCC) phase and Zr2Ni7) are precipitated in the Zr-rich region. These can significantly increase the yield strength of the alloys.

**Keywords:** high-entropy alloy; mechanical alloying; spark plasma sintering; nanoprecipitates; mechanical properties

#### **1. Introduction**

During the past few years, high-entropy alloys (HEAs), a type of multi-principal-element alloy, have drawn widespread attention form worldwide material scientists [1–3]. Compared to conventional alloys, which are composed of one or two major elements, HEAs usually contain 5–13 principal elements, and the content of each component is between 5 and 35 at.% [4]. Due to the high mixing entropy in the multi-component alloy systems, HEAs generally form a simple face-centered cubic (FCC) or body-centered cubic (BCC) structure solid solution and exhibit many novel properties, for example, superb mechanical properties [5,6], outstanding resistances to wear [7], oxidation [8] and corrosion [9], as well as a good temperature stability [10]. Therefore, the concept of HEAs provides a new approach to design alloys with excellent properties to meet different environmental requirements.

Casting [11] and powder metallurgy [12] are two common methods for preparing HEAs. For casting, there are two kinds of defects which are the vaporization of low melting-point elements and the elements segregation in the ingot [13]. For powder metallurgy, the HEAs can be prepared by MA and then consolidated. The HEAs powders prepared by MA are uniform and fine, and even form nanocrystalline. Hence a rapid sintering is needed to maintain this microstructure [14]. SPS is a new field-assisted sintering technique which can effectively suppress the grain coarsening by rapid heating and efficient densification in a few minutes [15–17]. Compared with MA, gas atomization

can produce more homogeneous powders. The preparation of HEAs powders by gas atomization is a rapid solidification process [18], in which grain growth and element segregation are inhibited, and the high cooling rate even leads to the formation of an amorphous phase [14].

In the previous studies, two classical HEAs, CoCrFeNiMn with a single FCC phase and AlCoCrFeNi with a single BCC phase, have been extensively studied. Among these studies, there are lots of experiments that have explored the influence of alloying elements such as Mo [5], Nb [6], Ti [19], Si [20], and Zr [21] on the microstructure and properties of AlCoCrFeNi. The addition of these alloying elements leads to the formation of intermetallic compounds and significantly improves the yield strength of AlCoCrFeNi. For example, after the addition of Zr, Laves phase appears in the BCC-phase matrix, the yield strength of AlCoCrFeNi increases from 1320 MPa to 1560 MPa, and the plastic strain increases from 22.5% to 29.5%. It is feasible to change the phase composition and improve the mechanical properties of AlCoCrFeNi by adding an alloying element. However, there are little similar data on the CoCrFeNiMn alloy [22,23]. In the present work, the CoCrFeNiMn alloy was selected as the matrix and the effect of Zr addition was investigated. There are several reasons for choosing Zr as an alloying element. Firstly, both Zr and Ti are located in the same family in the periodic table of elements, so they have some similar physico-chemical properties. In previous research, the addition of Ti can result in precipitates and improve the superplasticity of CoCrFeNiMn [23]. Secondly, the large atomic size difference and negative enthalpies of mixing [24] between Zr and other constituent elements can produce a strong lattice distortion and even change the phase composition, which may improve the mechanical properties of the CoCrFeMnNi.

In this paper, we focused on the novel high-entropy alloy system of CoCrFeNiMnZr*<sup>x</sup>* (*x* = 0, 0.2, 0.5, 0.8, 1.0). First of all, CoCrFeNiMn metallic powders were prepared by gas atomization, and then, Zr derived from the in-situ dehydrogenation of ZrH2 powders was solid-dissolved into CoCrFeNiMn by MA. The CoCrFeNiMnZr*x* alloy powders were subsequently sintered by SPS. The alloying behavior, microstructures, and mechanical properties of both powders and sintered alloys were investigated.

#### **2. Experimental**

Equimolar ratio CoCrFeNiMn powders were prepared by gas atomization with high purity Ar. The atomization pressure was 4 MPa. Then, using CoCrFeNiMn powders and ZrH2 powders as raw materials, the CoCrFeNiMnZr*<sup>x</sup>* (*x* values in molar ratio, *x* = 0, 0.2, 0.5, 0.8 and 1.0, denoted by Zr0, Zr0.2, Zr0.5, Zr0.8, and Zr1.0, respectively) alloys in nominal compositions were prepared by high-energy ball milling (referred as MA) and SPS. The reaction (ZrH2 = Zr(s) + H2(g)) occurs during MA and SPS. In this experiment, the mixed powders and zirconia balls (5 mm and 10 mm in diameter with the mass ration of 1:1) were put into stainless-steel vials at a mass ratio of 10:1, then mounted on a planetary ball miller (QM-S3P4, Nanjing NanDa Instrument Plant, Nanjing, China) and milled at a rate of 300 rpm under the protection of argon gas. All the mixed powders were milled for 30 h and stopped for 10 min every 20 min to prevent the powders from overheating. Beforehand, in order to study the alloying behavior, mixed Zr0.5 alloy powders were milled for 50 h, and a small amount of powder was taken every 10 h for X-ray diffraction tests. Finally, the gas-atomized CoCrFeNiMn powders without ball milling (Zr0w) and the alloy powders with 30 h of ball milling were sintered by SPS (Dr. Sinter-3.20 MKII, SCM, Japan) at 900 ◦C for 8 min under a uniaxial pressure of 40 MPa with a vacuum of 1 × <sup>10</sup>–3 Pa. The sintered columnar samples have a diameter of 30 mm and a height of 10 mm.

The oxygen content of the CoCrFeNiMn powder was analyzed by the fusion method on an O/N analyzer (736 series, LECO, Saint Joseph, MI, USA). X-ray diffraction with Cu Kα radiation (XRD, XPertPowder, PANalytical, Almelo, Netherlands) was used to analyze the phase composition of the alloy powder and bulk at a speed of 4◦/min, and the 2 thetas ranging from 20◦ to 100◦. The PDF-2 2004 database was used for the phase assignment. The morphology and chemical composition were characterized by scanning electron microscopy (SEM, Quanta, FEI, Hillsboro, TX, USA) equipped with energy dispersive spectrometry (EDS, Inca X-Max, Oxford instruments, Oxford, UK) operating

at 20 kV. Cuboidal specimens with a side length of 10 mm and a height of 5 mm were cut from sintered samples and then ground and polished for SEM and EDS analyses. The crystal structure was identified by transmission electron microscopy (TEM, JEM2100F, JEOL, Tokyo, Japan) operating at 200 kV. The disc-shaped TEM samples with a diameter of 3 mm were electropolished (TenuPol-5, Struers, Ballerup, Denmark) in an electrolyte composed of 20 vol.% perchloric acid and 80 vol.% methanol at −15 ◦C, and the applied voltage and current were 60 mV and 80 mA, respectively. Archimedes' method was used to measure the density of samples. Cylindrical specimens (7 mm in height and 3.5 mm in diameter) were used for the compressive tests (810 series, MTS, Minneapolis, MN, USA) at room temperature. Samples of each composition were tested for 3 times.

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

#### *3.1. Phase and Microstructure of the Gas Atomized HEA Powders*

The chemical analysis of the atomized HEA powder is shown in Table 1. As can be seen, the composition of the CoCrFeNiMn powder prepared by gas atomization was consistent with the nominal composition. The low oxygen content of the powder means that very little oxidation occurred during the atomization. The XRD pattern indicates the formation of the CoCrFeNiMn HEA powders with a single FCC phase (Figure 1a). The morphology and microstructure of the gas-atomized powders are depicted in Figure 1b,c. The atomized powders are spherical or quasi-spherical in diameter, ranging from 5 to 150 μm, and a small amount of fine powders adhere to the surface of large powders to form satellite structures. During the process of gas atomization, solidified fine droplets can easily adhere to the surface of large droplets that remain molten. The zoom-in image indicates that the microstructure consists of cellular and dendritic structures at a submicron scale.

**Table 1.** The chemical compositions of the gas atomized CoCrFeNiMn high-entropy alloy (HEA) powder.

**Figure 1.** The X-ray diffraction (XRD) pattern (**a**), morphology (**b**) and microstructure (**c**) of the gas-atomized Zr0 HEA powders.

#### *3.2. Microstructure and Phase Evolution during MA*

To investigate the alloying behavior in the MA process, X-ray diffraction (XRD) was performed on the milled Zr0.5 powders at 10 h intervals until the milling time reached 50 h (Figure 2). It can be seen that the XRD pattern of primary blending powders includes diffraction peaks of CoCrFeNiMn alloy with FCC structure and ZrH2 (PDF 073-2076). After 10 h of ball milling, the diffraction peaks of ZrH2 could still be detected, but their intensity was significantly reduced. This indicates that ZrH2 decomposed and the Zr dissolved into the matrix. After 20 h of ball milling, some diffraction peaks of ZrH2 disappeared due to the alloying. As the milling time increasing to 30 h, all the diffraction peaks of ZrH2 became absent and the XRD pattern became similar to that of Figure 1, which suggests that the Zr was fully dissolved into the FCC matrix to form a supersaturated solid solution. Additionally, as the ball milling time was extended, the diffraction peaks appeared to broaden. Throughout the milling process, the lattice strain increase and grain refinement are the main reasons for the peak broadening and intensity reduction [25]. On the other hand, ZrO2 diffraction peaks were observed after 40 h of ball milling, and the intensity of the ZrO2 diffraction peaks increase with the extending of ball milling time. The raw materials are usually contaminated by milling media in the process of ball milling [26,27].

**Figure 2.** The XRD patterns of the Zr0.5 alloy powders with different milling times.

Table 2 shows the average crystal size and lattice strain of the Zr0.5 alloy powders with different milling times that have been calculated using the Williamson–Hall equation method. As the milling time was extended from 10 h to 50 h, the crystal size decreased from 14.8 nm to 8.1 nm, and the lattice strain increased from 0.237% to 0.746%. Thus, the high lattice strain and the formation of nanocrystalline are the main reasons for the mentioned diffraction peaks broadening and the diffraction intensity reduction.


**Table 2.** The crystal size and lattice strain of the Zr0.5 HEA powders with different milling times.

Figure 3 shows the morphologies of the Zr0.5 alloy powders that have undergone different milling times. The initial powder is a mixture of spherical atomized CoCrFeNiMn powders and irregularly shaped ZrH2 powders, and the size of the ZrH2 powders are smaller than the CoCrFeNiMn powders (Figure 3a). The difference in size between the two powders will affect the alloying process. After 10 h of ball milling, the relatively large-sized CoCrFeNiMn powders became elliptical and fine sheet

powders were deposited on their surfaces. This indicates that the alloy powders were deformed, broken, and welded during ball milling and, at the same time, the fine Zr (or ZrH2) powders and broken CoCrFeNiMn powders were agglomerated on the surface of the large powder particles. As the ball milling time extended, the CoCrFeNiMn powders were wrapped by Zr-rich alloy powders and the outer layer of the CoCrFeNiMn powders was alloyed with Zr under the impact of the high-energy small balls. Meanwhile, the cycle of crushing and agglomeration of the outer layer continued, which causes the powders to gradually refine and promotes the alloying and diffusion among the different alloy elements [28,29]. After 30 h of ball milling, the particle size hardly changed, which means that the crushing and agglomeration reached a dynamic balance in the MA process.

**Figure 3.** The scanning electron microscopy (SEM) images of the Zr0.5 alloy powders with different milling times: (**a**) 0 h; (**b**) 10 h; (**c**) 20 h; (**d**) 30 h; (**e**) 40 h; (**f**) 50 h.

Figure 4a displays the XRD patterns of HEA powders with different contents of Zr after ball milling of 30 h. Compared with Figure 2, no extra peaks can be found in the XRD patterns of the Zr0, Zr0.2, and Zr0.5 alloys, which suggests that the ZrH2 was completely decomposed and the Zr was effectively dissolved into the CoCrFeNiMn HEA matrix when the Zr content is not higher than 0.5. Nevertheless, the diffraction peak of ZrH2 can be found in the XRD patterns when the Zr content is up to 0.8 and 1.0, and the intensity of the peak increases with the increase of the ZrH2 content.

**Figure 4.** The XRD patterns of HEAs: (**a**) CoCrFeNiMnZr*x* (*x* = 0, 0.2, 0.5, 0.8, 1.0) alloy powders after milling for 30 h, (**b**) sintered CoCrFeNiMnZr*x* (*x* = 0w, 0, 0.2, 0.5, 0.8, 1.0) HEAs.

#### *3.3. Phase Evolution and Microstructure after SPS*

XRD patterns of the sintered CoCrFeNiMnZr*<sup>x</sup>* (*x* = 0w, 0, 0.2, 0.5, 0.8, 1.0) alloys are shown in Figure 4b. Obviously, the ZrH2 diffraction peaks in the Zr0.8 and Zr1.0 powders disappear in the sintered samples. This suggests that the excessive ZrH2 was decomposed and that the Zr remainder was fully alloyed during SPS. The FCC phase is still the dominant phase, meanwhile, some weak diffraction peaks which are identified as Zr2Ni7 (PDF 071-0543), σ phase, and ordered BCC (B2) phase appear. Compared with the XRD patterns of the powders after ball milling, the broad diffraction peaks became narrow due to internal energy releases during the SPS process.

The Zr0w and Zr0 alloys have the same diffraction peaks and consist of an FCC phase, indicating that the MA and SPS processes do not alter the FCC phase of the alloys. The diffraction peaks of the Zr2Ni7, the B2 phase, and the σ phase can be found in the Zr0.2, Zr0.5, Zr0.8, and Zr1.0 alloys. The σ phase was widely reported in previous research [30,31] and it can be identified as the NiCoCr (PDF 021-1271) σ phase, which has a tetragonal structure (a = 8.85 Å and c = 4.59 Å) in the present experiment. The B2 phase has been often found in HEAs containing CoCrFeNi [4,21,32]. These indicate that during the SPS process, some of the alloy atoms diffuse under thermal activation conditions to form ordered solid solutions and intermetallic compounds.

With the increase of the Zr content, the diffraction intensity of the B2 phase decreases while the diffraction intensity of the Zr2Ni7 compound increases. This indicates that when the Zr content is low, the free energy is reduced by the formation of the ordered phase. However, as the Zr content increases, the formation of the Zr2Ni7 is more effective to reduce the free energy, and the formation of the B2 phase is suppressed. On the other hand, the diffraction peaks of the σ phase almost disappear in the Zr1.0 alloy, which also means that the formation of the Zr2Ni7 suppresses the formation of the σ phase. These can be attributed to the lattice strain and a lager negative enthalpy of mixing, caused by the addition of Zr.

Figure 5 displays the SEM (operating at back-scattered electron (BSE) mode) images of the CoCrFeNiMnZr*<sup>x</sup>* alloys synthesized by MA and SPS. As can be seen from Figure 5a,b, both the Zr0w and Zr0 alloys exhibit a single phase feature, which is consistent with the XRD results. After the addition of Zr, there are two distinct microstructural domains: the bright and gray regions (Figure 5c–f).

**Figure 5.** The SEM (operating at back-scattered electron (BSE) mode) images of the CoCrFeNiMnZr*x* alloys synthesized by mechanical alloying (MA) and spark plasma sintering (SPS). (**a**) *x* = 0w; (**b**) *x* = 0; (**c**) *x* = 0.2; (**d**) *x* = 0.5; (**e**) *x* = 0.8; (**f**) *x* = 1.0. Any Zr-rich particles were wrapped by a plastic face-centered cubic (FCC) phase and indicated by yellow arrows.

EDS was performed to investigate the two regions using the Zr0.2 alloy and the results are shown in Figure 6 and Table 3. On the one hand, the bright region contains a higher amount of Zr than the nominal composition whereas Zr is barely found in the gray regions. On the other hand, the content of Co, Cr, Fe, Ni, and Mn are relatively lower in the bright region, and they are almost equal in the atomic ratio in both regions. Obviously, some CoCrFeNiMn HEA particles are not completely broken by the high-energy ball milling due to their high toughness [14,33], and Zr is not uniformly distributed in the matrix. The gray regions can be identified as a pure CoCrFeNiMn FCC phase and the bright region consists of FCC phase, B2 phase, Zr2Ni7, and σ phase. For the gray region, the shapes present are ovals and strips, which were inherited from the MA process. Furthermore, the hard and brittle Zr-rich particles were wrapped by a plastic FCC phase during the ball milling process, a few small bright microstructures are indicated by yellow arrows (Figure 5). Table 4 shows the density of the sintered CoCrFeNiMnZr*x* alloys. It can be seen that the addition of a relatively low-density Zr element reduces the density of the alloy system (the density of Co, Cr, Fe, Ni, Mn, and Zr are 8.9 g·cm−3, 7.19 g·cm−3, 7.87 g·cm<sup>−</sup>3, 8.9 g·cm<sup>−</sup>3, 7.44 g·cm<sup>−</sup>3, and 6.49 g·cm<sup>−</sup>3, respectively).

**Figure 6.** The energy dispersive spectrometry (EDS) mappings of the Zr0.2 alloy synthesized by MA and SPS.





Figure 7a shows the grain structure of Zr0 and the selected area electron diffraction (SAED) pattern is related to the red circle marking area. It indicates that the sintered Zr0 alloy has a micron grain size and a single FCC structure which is consistent with XRD results. Meanwhile, some twins are found. After the Zr addition, there are two obviously different areas distinguished by grain size (Figure 7b). The SAED pattern of red circle marking area demonstrates that the micro-crystal has an FCC structure. EDS results (not shown here) revealed that these nanocrystals are Zr-rich while the micro-crystals do not contain Zr. This confirms the SEM results: the Zr and the outer layer of the atomized powder are alloyed during the ball milling process, and as time passes, the fine Zr-rich powders agglomerated on the outer layer is continuously refined. Furthermore, many spherical nanoprecipitates are found in the nanocrystal area (Figure 7c). Figure 7d–g show the high-resolution TEM (HRTEM) images with corresponding fast Fourier transformations (FFT) of nanoprecipitates. These indicate that there are two

kinds of nanoprecipitates which are BCC structure precipitates and Zr2Ni7 precipitates. In the previous studies [3,34], nanoprecipitates have been considered to be beneficial to the mechanical properties of the material due to the ability of pinning dislocations.

**Figure 7.** The transmission electron microscopy (TEM) images of sintered (**a**) Zr0 and (**b**) Zr0.2 HEAs. The insets in (**a**,**b**) are the selected area electron diffraction (SAED) patterns of the red circle marking area. The dashed line in (**b**) distinguishes the two regions with different grain sizes: the micro-crystals on the left side of the dashed line and the nanocrystals on the right side of the dashed line; (**c**) is the magnifying image of the yellow circle marking area in (**b**); (**d**,**e**) are the high-resolution TEM (HRTEM) of nanoprecipitation; (**f**,**g**) are the fast Fourier transformations (FFT) of (**d**,**e**), respectively.

#### *3.4. Mechanical Properties*

Figure 8 shows the compressive stress-strain curves of the CoCrFeNiMnZr*<sup>x</sup>* alloys at room temperature. The values of yield strength (YS) σ*y*, compressive strength (CS) σmax, and plastic strain limit (PS) ε*<sup>p</sup>* of alloys are listed in Table 5. The maximum yield strength reaches 820 MPa when the Zr content is up to 0.8, but the compressive strength of the Zr*<sup>x</sup>* (*x* = 0.2, 0.5, 0.8, 1.0) alloys are lower than that of alloys without Zr. Because of the addition of Zr, the yield strength evidently improved while the ductility declined simultaneously. Zr0 and Zr0w alloys, exhibit excellent plasticity, and both of the two alloys are not broken in the pressure range of the compressor. The yield strength of Zr0 is higher than that of Zr0w, which suggests that MA can improve the strength of the alloy by refining the grain [14]. Furthermore, the yield strength of the CoCrFeNiMnZr*x* (*x* = 0.2, 0.5, 0.8, 1.0) alloys are significantly higher than that of the Zr0 alloy, but the plasticity strongly reduced at the same time. This can be attributed to solid solution strengthening and second phase strengthening caused by the addition of Zr, as well as fine grain strengthening obtained by mechanical ball milling.

**Figure 8.** The compressive engineering stress–strain curves of the CoCrFeNiMnZr*x* alloys.

**Table 5.** The mechanical properties of the CoCrFeNiMnZr*x* alloys.


The fractographs of the CoCrFeNiMnZr*<sup>x</sup>* (*x* = 0.2, 0.5, 0.8, 1.0) alloys after compressive tests are depicted in Figure 9. The fracture surface of the Zr0.2 alloy presents a step-like pattern with quasi-cleavage characteristics. A small radiation zone appears on the right side of Figure 9a, indicating the localized ductile fracture. Complete cleavage fracture patterns can be seen in Figure 9b–d, which thus deteriorates the alloy plasticity. For the Zr0.5 alloy, the fracture surface consists of many small cleavage planes, with both inter-granular fractures and trans-granular fractures having occurred. Large bright and flat cleavage planes appear in the fractures of Zr0.5 and Zr1.0 alloys, which is in accord with the totally brittle fracture of the alloys. The above results indicate that the CoCrFeNiMnZr*<sup>x</sup>* (*x* = 0.2, 0.5, 0.8, 1.0) alloys exhibit brittle fractures, which can be attributed to a large lattice distortion and the brittle intermetallic compounds present in the Zr-rich region.

**Figure 9.** SEM micrographs of the fracture surface of CoCrFeNiMnZr*x* alloys. (**a**) *x* = 0.2; (**b**) *x* = 0.5; (**c**) *x* = 0.8; (**d**) *x* = 1.0.

#### **4. Conclusions**

The alloying behavior, microstructures, and mechanical properties of both powders and bulk CoCrFeNiMnZr*x* alloys were investigated. The main conclusions of the present work are given below:

(1) Metastable single FCC phase CoCrFeNiMnZr*x* alloy powders were prepared after 30 h of ball milling, which means that MA could enhance the solid solubility of large size elements in the matrix. However, toughness powder is difficult to break during ball milling, which would affect the alloying process and lead to the element segregation. Toughness powders can retain their original properties during ball milling, which means that MA can be used to prepare HEAs composites.

(2) The phase composition of HEAs can be changed by adjusting the content of the alloying element. After SPS, CoCrFeNiMn maintains a single FCC phase. However, two distinct microstructural domains, one composed of a micro-sized CoCrFeNiMn alloy without Zr and the other one composed of Zr-rich multi-phase microstructures consisting of a nano-sized FCC phase, B2 phase, Zr2Ni7, and σ phase, appear in the CoCrFeNiMnZr*<sup>x</sup>* (*x* ≥ 0.2) alloys. When the elements with different physicochemical properties are dissolved into the HEA matrix, the local structural stability will be destroyed, and the local atomic arrangement will change from disorder to order.

(3) The addition of Zr significantly increases the yield strength of the CoCrFeNiMn alloy. The Zr0.8 alloy exhibits the highest yield strength up to 820 MPa which is almost twice as much as the Zr0 alloy (420 MPa). The high yield strength makes it a promising candidate for precision equipment manufacturing.

**Author Contributions:** Conceptualization, Y.X.; data curation, L.Z.; formal analysis, H.Z.; investigation, H.Z. and L.Z.; methodology, Y.X; project administration, Y.X; resources, Y.X; writing—original draft, H.Z. and L.Z.; writing—review & editing, H.Z., X.L., Q.C. and Y.X.

**Acknowledgments:** The authors thank the National Key Project of Research and Development Program of China (2016YFB1100202) for financial support.

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

#### **References**


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