**Wear and Corrosion Resistance of Al0.5CoCrCuFeNi High-Entropy Alloy Coating Deposited on AZ91D Magnesium Alloy by Laser Cladding**

#### **Kaijin Huang 1,2,3,4,\*, Lin Chen 2, Xin Lin 3, Haisong Huang 4, Shihao Tang <sup>4</sup> and Feilong Du <sup>4</sup>**


Received: 15 November 2018; Accepted: 28 November 2018; Published: 30 November 2018

**Abstract:** In order to improve the wear and corrosion resistance of an AZ91D magnesium alloy substrate, an Al0.5CoCrCuFeNi high-entropy alloy coating was successfully prepared on an AZ91D magnesium alloy surface by laser cladding using mixed elemental powders. Optical microscopy (OM), scanning electron microscopy (SEM), and X-ray diffraction were used to characterize the microstructure of the coating. The wear resistance and corrosion resistance of the coating were evaluated by dry sliding wear and potentiodynamic polarization curve test methods, respectively. The results show that the coating was composed of a simple FCC solid solution phase with a microhardness about 3.7 times higher than that of the AZ91D matrix and even higher than that of the same high-entropy alloy prepared by an arc melting method. The coating had better wear resistance than the AZ91D matrix, and the wear rate was about 2.5 times lower than that of the AZ91D matrix. Moreover, the main wear mechanisms of the coating and the AZ91D matrix were different. The former was abrasive wear and the latter was adhesive wear. The corrosion resistance of the coating was also better than that of the AZ91D matrix because the corrosion potential of the former was more positive and the corrosion current was smaller.

**Keywords:** laser cladding; high-entropy alloy coating; AZ91D magnesium alloy; wear; corrosion

#### **1. Introduction**

High-entropy alloys are a new kind of alloy with excellent properties such as good wear resistance, excellent corrosion resistance, excellent oxidation resistance, low electrical conductivity, low thermal conductivity, and low coefficient of thermal expansion; they were invented by Yeh in 1995 [1–7]. They are composed of five or more elements in the same or an approximately equal molar ratio and have simple BCC and/or FCC solid solution phases. Their microstructure and properties are different from those of traditional alloys such as Fe- and Ni-based alloys or intermetallic compounds such as Ti–Al, Ni–Al, and Fe–Al compounds; this is because of the high entropy effect, severe lattice distortion effect, sluggish diffusion effect, and cocktail effect of the former [5,7].

So far, many methods such as arc melting [8], Tungsten Inert Gas Arc Welding (TIG) [9], Gas Tungsten Arc Welding (GTAW) [10], mechanical alloying [11], DC sputtering [12], thermal spay technology [13], and laser cladding [14–32] have been adopted to prepare different high-entropy alloys or coatings. Among them, the laser cladding method has attracted special attention because of its

advantages of rapid heating, rapid cooling, compact coating, and low dilution rate, etc. [19]. At present, the preparation of high-entropy alloy coatings by laser cladding mainly focuses on the microstructure and properties of different matrix materials such as Ti-6Al-4V [14,15], carbon steel [16–21], stainless steel [22–24], tool steel [25], die steel [26], aluminum [27], magnesium [28–30], and magnesium alloys [31,32], and different high-entropy coatings [14–32]. It should be noted that the study of the preparation of high-entropy alloy coatings on magnesium and magnesium alloys by laser cladding was mainly carried out by a few researchers: Yue [28,29,31], Huang [32], and Meng [30]. The possible reason for this is that good-quality laser-clad coatings on magnesium and its alloys are very difficult to obtain because of its low melting point (922K) and low boiling point (1363K).

It is well known that magnesium and its alloys have been widely used in many fields such as the automotive, communication, and aerospace industries, but their poor wear resistance and corrosion resistance hinder their application in those situations where these properties are required. To solve this problem, a method called laser cladding has been developed to improve the wear and corrosion resistance of magnesium and its alloys, and numerous literature reports have confirmed this result [33,34]. Based on the excellent wear resistance and corrosion resistance of high-entropy alloys [1–7] and on the basis of an author's previous research [32], we decided to continue the preparation of new wear-resistant and corrosion-resistant high-entropy alloy coatings by laser cladding on an AZ91D magnesium alloy substrate, which is widely used. To our best knowledge, no studies have reported an Al0.5CoCrCuFeNi high-entropy alloy coating fabricated by laser cladding on an AZ91D magnesium alloy substrate. The reason for studying the Al0.5CoCrCuFeNi alloy is that it has a simple FCC solid solution phase and exhibits excellent wear resistance due to its large work-hardening capacity [35].

In this paper, mixed powders of Cu, Ni, Al, Co, Cr, and Fe were used to fabricate an Al0.5CoCrCuFeNi high-entropy alloy coating on AZ91D magnesium alloys by laser cladding. The microstructure, wear behavior, and corrosion behavior of the coating are shown in detail.

#### **2. Experimental**

Laser cladding of mixed powders of Cu, Ni, Al, Co, Cr, and Fe was undertaken on an AZ91D magnesium alloy (Guangdong Dongguan Jubao Magnesium Alloy Material Co., Ltd, Dongguan, China). The sample size used in the laser cladding was 100 mm × 50 mm × 10 mm. The powders (Hebei Xingtai Nangong Zhongzhou Alloy Material Co., Ltd, Xingtai, China) had particle sizes of 48–75 μm and purity of 99 wt %. The different element powders were mixed by ball milling (DECO-PBM-V-0.4L, Hunan Changsha Deco Instrument Equipment Co., Ltd, Changsha, China) according to the nominal composition of the Al0.5CoCrCuFeNi high-entropy alloy. The mixed powder from the ball milling was preset on the surface of the AZ91D magnesium alloy using 4 vol % PVA solution. The thickness of the preset layer was about 0.6 mm. The laser cladding experiment was completed under the protection of 99.999% high-purity argon by using a TR050 type CO2 high-power laser. The optimized laser cladding parameters were as follows: laser power was 3000 W, laser scanning speed was 10 mm/s, laser spot diameter was 4 mm, and laser spot overlap rate was 25%.

The phase structure and microstructure of the coating after laser cladding were identified and observed using an X-ray diffractometer (X' Pert PRO), an optical microscope (Axiovert 200MAT), and a scanning microscope (Quanta 400) with an energy spectrum, respectively. The microhardness of the coating cross section was measured using a microhardness tester (MICROMET 3). The loading force was 100 grams and the loading time was 15 s.

The wear resistance of the coating was evaluated by the dry sliding wear method, in which the size of the sample used for testing was 10 mm × 10 mm × 10 mm, and the friction pair used for matching was bearing steel (AISI52100, HV0.1700). Figure 1 shows a schematic illustration of the block-on-ring sliding wear tester. The parameters for the dry sliding wear were as follows: the applied load was 98 N, the dry sliding speed was 0.4187 m/s, the dry sliding wear time was 75 mins, and the dry sliding wear distance was 1884 m. The AZ91D magnesium alloy was selected as the experimental material for the dry sliding wear comparison. For each experimental datapoint, the average value of three experimental results was taken as the final data. Wear weight loss was measured using an electronic analytical balance (Bartorius BS110) with an accuracy of 0.1 mg.

The corrosion resistance of the coating was evaluated by testing the corrosion potential polarization curve of a 3.5 wt % sodium chloride solution. A saturated calomel electrode (SCE) was used as the reference electrode and platinum was used as the counter electrode. The test was performed from −2.5 V to 1.5 V with a scanning speed of 1 mV/s.

The surface morphologies of the worn and corroded specimens were observed using a scanning microscope (Quanta 400) with an energy spectrum.

**Figure 1.** Schematic illustration of the block-on-ring sliding wear tester.

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

#### *3.1. Microstructure and Microhardness*

Figure 2a shows an optical image of the cross section of the coating/AZ91D substrate sample after laser cladding. It can be seen from Figure 2a that there were basically no defects such as microcracks or pores in the coating. The interface between coating and substrate indicates that they were well combined. In addition, small pieces of AZ91D magnesium alloy substrate were also found in the coating (Figure 2a). This is believed to be due to some partially melted Mg being detached from the substrate and becoming trapped within the rapidly solidifying coating. Figure 2b shows the scanning electron microscopy morphology of position A in Figure 2a. It can be seen from Figure 2b that the coating has typical dendrite microstructure characteristics, and the chemical compositions of the dendrite (DR) and inter-dendrite (ID) are given in Table 1. As shown in Table 1, copper segregates significantly in the inter-dendrite region. In other words, local segregation of copper occurred in the Al0.5CoCrCuFeNi high-entropy alloy coating prepared by laser cladding. These results are similar to the microstructure characteristics of Al0.5CoCrCuFeNi high-entropy alloys prepared by arc melting in the literature [35–38]. The difference is that the dendrite microstructure of the laser cladding coating is fine due to the rapid heating, melting, and rapid solidification of the preset layer during laser cladding.

**Table 1.** EDS results of the laser-clad high-entropy alloy coating (atom %).


The copper segregation in the Al0.5CoCrCuFeNi high-entropy alloy can be explained by the mixing enthalpies [39] of different atomic pairs in Table 2. As shown in Table 2, all of the mixing enthalpies of copper with iron, chromium, cobalt, and nickel are positive. This means that copper has a low affinity for these atoms and is easily repelled by them. In the process of laser cladding, different elements in the Al0.5CoCrCuFeNi high-entropy alloy were mixed evenly due to the convection and agitation of the laser cladding pool. However, during the cooling and solidification stage, copper was excluded from the dendrites due to the low affinity of copper atoms with iron, chromium, cobalt, and nickel; thus, copper segregation occurred at the inter-dendrites.

89D/ 8**%**9(/56

**Figure 2.** Morphology of a laser-clad Al0.5CoCrCuFeNi high-entropy alloy coating: (**a**) optical microscope (OM) image; (**b**) SEM image in position A.


**Table 2.** Values of ΔHij mix (kJ/mol) for atomic pairs of elements [39].

Figure 3 presents the XRD patterns of the Al0.5CoCrCuFeNi high-entropy alloy coating prepared by laser cladding. The analytical result confirmed that the phase was a simple FCC solid solution in the laser-clad coating. This result proved that the Al0.5CoCrCuFeNi high-entropy alloy coating was successfully fabricated by laser cladding on the AZ91D magnesium alloy based on the mixed powders of Cu, Ni, Al, Co, Cr, and Fe of Al0.5CoCrCuFeNi and the optimized laser cladding parameters. This result is consistent with the results reported in the literature [35–38]. Thus, for the laser-clad Al0.5CoCrCuFeNi high-entropy alloy coating, both the dendrites and Cu-rich inter-dendrites were of one simple FCC phase (Figure 1b). The formation of a simple solid solution phase rather than intermetallic compounds is mainly attributed to the significant lowering of free energy by the high enthalpy of mixing [40].

**Figure 3.** XRD patterns of a laser-clad Al0.5CoCrCuFeNi high-entropy alloy coating.

Figure 4 shows the microhardness of the laser-clad Al0.5CoCrCuFeNi high-entropy alloy coating. As can be seen from Figure 4, the microhardness of the coating was HV0.1365—about 3.7 times of the AZ91D matrix (HV0.198) and higher than that of the Al0.5CoCrCuFeNi high-entropy alloy (HV5233) prepared by arc melting [35–37]. The high microhardness of the laser-clad coating was attributed to the solid solution strengthening of different alloy elements and fine grain strengthening of the fine dendritic structure during the cooling process. Obviously, the high microhardness of the coating is more beneficial to increasing the wear resistance of the coating.

**Figure 4.** Microhardness of a laser-clad Al0.5CoCrCuFeNi high-entropy alloy coating.

#### *3.2. Wear Properties*

Figure 5 shows the dry sliding wear weight loss of different samples. According to Figure 5, the wear resistance of the coating was better than that of the AZ91D matrix, and the wear weight loss was about 2.5 times smaller than that of the AZ91D matrix. This is in accordance with their corresponding microhardness values (Figure 4). In other words, the hardness is higher and the wearability is better. Figure 6 shows the surface morphologies of different samples after dry sliding wear. In Figure 6, both of them present obvious groove characteristics. However, the wear mechanisms of the two samples were also significantly different. The coating showed obvious abrasive wear characteristics, and the AZ91D matrix showed obvious adhesive wear characteristics. Some elongated pockmarks and microcracks (Figure 6a) can be found on the worn surface of the AZ91D matrix. Table 3 shows the EDS results of different positions of different worn samples. It can be seen from Table 3 that AZ91D matrix position 1 and coating position 1 are each matrix materials, although each has a small amount of oxidation due to friction heating. Both AZ91D matrix position 2 and coating position 2 are iron oxide particles, but there are a lot of fine particles on the worn surface of the coating, while the worn surface of the AZ91D matrix is much cleaner.

**Figure 5.** Weight loss of different specimens.

896)C-

#### /45/ 8**%**9 @5/



**Table 3.** EDS results of worn surfaces of the AZ91D matrix and the laser-clad specimen (wt %).

The causes of the above phenomena should be related to the hardness of the respective matrix materials. Figure 4 shows that the microhardness of the coating was about HV0.1365, while the microhardness of the AZ91D matrix was only HV0.198. In addition, the coating is composed of Al0.5CoCrCuFeNi high-entropy alloy, and Al0.5CoCrCuFeNi high-entropy alloy has superior work-hardening characteristics [35], but the AZ91D matrix does not have the same excellent work-hardening capability. As a result, the microhardness of the coating continued to increase when dry sliding wear was applied to the bearing steel friction pair and may even approach or exceed the hardness of the bearing steel friction pair; it would then shear off the steel ring material, and iron wear debris would be smeared over the coating surface. With the extension of the dry sliding wear time, the iron wear debris would be oxidized due to heat generated by friction. Therefore, iron oxide abrasive particles will be used as an abrasive cutting the Al0.5CoCrCuFeNi high-entropy alloy coating. In contrast, the AZ91D matrix did not have this process and only the soft AZ91D matrix was transferred to the hard bearing steel surface, so its wear rate (30.38 μg/s) was much higher than that (11.93 μg/s) of the coating.

#### *3.3. Corrosion Properties*

Figure 7 and Table 4 show the potentiodynamic polarization curves and corrosion properties, respectively, of different samples in 3.5 wt % sodium chloride solution. According to the corrosion electrochemical behavior of the two samples and the lack of a passivating region in Figure 7, both the AZ91D matrix and the coating are active dissolved materials. Therefore, using the principle of "the smaller the corrosion current and the higher the corrosion potential, the better the corrosion resistance of the material" to evaluate the corrosion resistance of the active dissolved material, it can be seen that the corrosion resistance of the Al0.5CoCrCuFeNi high-entropy alloy coating prepared by laser cladding was better than that of the AZ91D matrix. This is reflected by the fact that the former had a significantly higher corrosion potential (Ecorr = −0.998 V) than did the latter (Ecorr = −1.46 V), and the former also exhibited a lower corrosion current (icorr = 1.60 × <sup>10</sup>−<sup>4</sup> A/cm2) than did the latter (icorr = 6.20 × <sup>10</sup>−<sup>4</sup> A/cm2).

**Figure 7.** Potentiodynamic polarization curves of different specimens in 3.5 wt % sodium chloride solution.

**Table 4.** Corrosion parameters of different samples in 3.5 wt % sodium chloride solution.


Figure 8 shows the corroded surfaces of two samples. It can be seen from Figure 8a that the magnesium oxide film generated on the AZ91D matrix cannot provide effective protection due to its loose and porous properties and the extremely low equilibrium potential (−2.37 V) of magnesium in the AZ91D matrix (Table 3). The following reactions will occur [41]:

$$\text{Mg(s)} + \text{H}\_2\text{O} \rightarrow \text{Mg(OH)}\_2\text{(s)} + \text{H}\_2\text{(g)}$$

$$\text{Mg(OH)}\_2\text{(s)} + \text{Cl}^- \rightarrow \text{MgCl}\_2 + \text{OH}^-,$$

$$\text{Mg(s)} + \text{Cl}^- \text{(aq)} \rightarrow \text{MgCl}\_2.$$

896)C-/45/ 8**%**9 @5/ **Figure 8.** Corroded surfaces of (**a**) an AZ91D matrix specimen and (**b**) a laser-clad specimen. Therefore, the corrosion of the AZ91D matrix was severe (Table 4 and Figure 8a).

On the other hand, the corrosion attack on the Al0.5CoCrCuFeNi high-entropy alloy coating specimen was much less acute. Figure 8b reveals that the corroded surface has a grainy appearance, indicating that corrosion attacks may mainly occur on dendrite boundaries. It is considered that the presence of Cr, Al, and Ni in the Al0.5CoCrCuFeNi high-entropy alloy coating prepared by laser cladding could form dense chromium oxides, aluminum oxides, and nickel oxides at the surface, and this could be an important factor that contributes to the relatively high corrosion resistance of the Al0.5CoCrCuFeNi high-entropy alloy coating specimen.

It should be pointed out that the composition of the high-entropy Al0.5CoCrCuFeNi selected in this paper is not superior to that of other high-entropy compositions (for example, the one studied in Reference [32]). In this paper, only the high-entropy Al0.5CoCrCuFeNi composition is chosen because of its good corrosion resistance.

As for how to choose the composition of the coating to obtain high corrosion resistance, we think that the composition of the high-entropy alloy should be based mainly on whether the same high-entropy alloy prepared by the arc melting method has excellent corrosion performance. If so, it can be selected as a candidate material for laser cladding to prepare a high-entropy alloy coating; otherwise, it should not be selected.

#### **4. Conclusions**


**Author Contributions:** Investigation—K.H., L.C., H.H., S.T., F.D.; writing—original draft preparation, L.C.; writing—review and editing, K.H.

**Funding:** This research received no external funding.

**Acknowledgments:** This work was supported by the Opening Project (2018CL18) of Material Corrosion and Protection Key Laboratory of Sichuan Province, Sichuan University of Science & Engineering, the Open Research Fund Program (XDKFJJ [2016]06) of Key Laboratory of Advanced Manufacturing Technology, Ministry of Education, Guizhou University, and the Fund (SKLSP201733) of the State Key Laboratory of Solidification Processing in NWPU. The authors are also grateful to the Analytical and Testing Center of Huazhong University of Science and Technology.

**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* **Fabrication of AlCoCrFeNi High-Entropy Alloy Coating on an AISI 304 Substrate via a CoFe2Ni Intermediate Layer**

#### **Wenyuan Cui 1,\*, Sreekar Karnati 1, Xinchang Zhang 1, Elizabeth Burns <sup>2</sup> and Frank Liou <sup>1</sup>**


Received: 30 October 2018; Accepted: 19 December 2018; Published: 20 December 2018

**Abstract:** Through laser metal deposition, attempts were made to coat AlCoCrFeNi, a high-entropy alloy (HEA), on an AISI 304 stainless steel substrate to integrate their properties. However, the direct coating of the AlCoCrFeNi HEA on the AISI 304 substrate was found to be unviable due to cracks at the interface between these two materials. The difference in compositional change was suspected to be the source of the cracks. Therefore, a new transition route was performed by coating an intermediate layer of CoFe2Ni on the AISI 304 substrate. Investigations into the microstructure, phase composition, elemental composition and Vickers hardness were carried out in this study. Consistent metallurgical bonding was observed along both of the interfaces. It was found that the AlCoCrFeNi alloy solidified into a dendritic microstructure. The X-ray diffraction pattern revealed a transition of the crystal structure of the AISI 304 substrate to the AlCoCrFeNi HEA. An intermediate step in hardness was observed between the AISI 304 substrate and the AlCoCrFeNi HEA. The AlCoCrFeNi alloy fabricated was found to have an average hardness of 418 HV, while the CoFe2Ni intermediate layer had an average hardness of 275 HV.

**Keywords:** high-entropy alloy; laser metal deposition; elemental powder; graded material

#### **1. Introduction**

As a novel metallic alloy system, high-entropy alloys (HEAs) have received considerable attention in the past decade. The name HEA indicates that the mixing of the principal elements in the alloy leads to a substantial change in entropy. This change in entropy promotes the formation of a simple solid solution instead of complex compounds. One of the extensively studied HEAs is equiatomic AlCoCrFeNi, which shows high hardness, good wear behavior but low tensile ductility [1–7]. As-cast AlCoCrFeNi alloy showed a tensile elongation of 1.0%, while post-heat treatment, the elongation was increased to 11.7% [4]. Wang et al. studied the compressive properties of AlCrFeCoNi HEA prepared by vacuum arc melting. They found that this alloy showed large strain hardening and compressive strength up to 2004 MPa with a 32.7% compressive plasticity [6]. Munitz et al. reported the impact of heat treatment of AlCoCrFeNi HEA, in which the BCC (Body-centered cubic) matrix transformation occurred between 650 and 975 ◦C. This transformation led to a substantial increase in microhardness [5]. Further modification of this alloy system through the addition of titanium, leading to AlCoCrFeNiTi*x* (*x* = molar ratios), was found to be promising for wear protection [1]. Further, AlCoCrFeNi HEA solidified with dendritic and interdendritic microstructures due to elemental segregation. Dendritic segregation regions were found to be Al- and Ni-rich, while interdendritic areas were Fe- and Cr-rich, and the distribution of Co was uniform. Body-centered cubic (BCC) Fe and Cr precipitates, and B2

(ordered BCC) Al- and Ni-rich matrices were observed in previous studies [2,4,6,8,9]. Most of these studies are based on material fabricated through processes such as casting and arc melting. Unlike these early studies, laser metal deposition (LMD) was implemented in this study.

LMD is capable of fabricating freeform three-dimensional metallic components [10–12] and has been used to fabricate several HEAs [12–16]. Chen et al. fabricated AlxCoFeNiCu1-*<sup>x</sup>* (*x* = 0.25, 0.5 and 0.75 atom %, respectively) HEAs using elemental powders on the AISI 304 substrate. They reported an increase in hardness with an increase in aluminum content [16]. He et al. used laser cladding to produce FeCoCrNiAlTi*x* (*x* = 0, 0.25, 0.5, 0.75 and 1 atom %, respectively) coating on Q253 steel through the use of elemental powders. Addition of titanium was observed to improve the hardness and wear resistance of the HEA [15]. Similarly, FeCoCrAlCu HEA coating by laser cladding demonstrated good wear resistance under a dry sliding condition [17].

In this paper, the feasibility of coating an AlCoCrFeNi HEA on an AISI 304 stainless steel substrate was investigated. Sole LMD fabrication of AlCoCrFeNi HEA components is very costly due to the need for high-purity (i.e., 99.9%) raw powders of elements such as Co, Cr and Ni. AISI 304 stainless steel, on the other hand, is a low-cost structural material. However, AISI 304 is a soft material with low wear resistance. It is widely used in industrial facilities, transportation equipment and architectural applications. Therefore, by coating AlCoCrFeNi HEA on AISI 304, it can enhance the hardness of AISI 304 structures. This combination of materials could facilitate fabrication of components for applications that require both hardness and wear resistance.

However, direct coating of AlCoCrFeNi HEA on AISI 304 is difficult due to the change in chemistry, thermal expansion and residual stress of the dissimilar materials. For example, the measured coefficient of thermal expansion (CTE, 10−6/K) for AlCoCrFeNi HEA was 9.03 (293–303 K), 12.47 (368–378 K) and 13.54 (423–773 K) [18]. However, the CTE values of AISI 304 were 14.7 (293 K), 16.3 (400 K), 19.5 (700 K) and 20.2 (800 K) [19]. Harihar et al. observed crack formation at the bottom of an AlCoCrFeNi deposit when deposited on an AISI 304 substrate. Due to the brittleness of the deposited material, the deposit broke off from the AISI 304 substrate easily [12]. An extensive network of cracks occurred when a TiVCrAlSi HEA was cladded on a Ti-6Al-4V substrate. This was attributed to the difference between the thermal expansion coefficients and residual stresses associated with the high cooling rate in laser cladding [20].

Therefore, to facilitate the dissimilar material bond, an intermediate layer was necessary and could accommodate the residual stresses and variation in chemistry change [10,21,22]. Intermediate layers of Fe/Cr/V were used between AISI 316 stainless steeland Ti-6Al-4V to facilitate a similar material bond [10]. Currently, there are few studies available identifying the viable intermediate layer between AlCoCrFeNi HEA and AISI 304. In this study, an attempt was made to coat the equiatomic AlCoCrFeNi HEA on the AISI 304 substrate using LMD. The objective was to obtain a strong bond between the two materials. We first demonstrated the issues with direct-coating the HEA onto the substrate. Then we proposed a candidate intermediate material and proved its viability.

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

Elemental powders of gas-atomized aluminum (Al), chromium (Cr), cobalt (Co), nickel (Ni) and iron (Fe) from Atlantic Equipment Engineers Inc. were used as precursor materials. These powders, weighed in required ratios, were mixed using a Turbula mixer (Glen Mills Inc., Clifton, NJ, USA) for 1 h to obtain homogeneous blends. Commercially procured AISI 304 bar stock (dimensions: 2.75 inch × 2 inch × 0.25 inch) was used as the substrate material for the deposition. The particle size distribution of the elemental powders stated by the producer is as tabulated in Table 1. Elemental analysis of the elemental powders is listed in Table 2. Elemental compositions (atom %) of the as-blended CoFe2Ni intermediate layer and AlCoCrFeNi alloy are given in Table 3.


**Table 1.** Particle size distribution of the precursor elemental powders.

**Table 2.** Elemental analysis (atom %) of elemental powders as provided by the manufacturer.


**Table 3.** Nominal compositions (atom %) of CoFe2Ni and AlCoCrFeNi alloy powder blends.


The laser deposition process was performed in an LMD system whose schematic representation is as seen in Figure 1a. The heat source was a 1 kW continuous-wave YAG fiber laser (IPG Photonics, Oxford, MA, USA) with a 2 mm beam diameter. The powders were fed using a vibration X2 powder feed system procured from Powder Motion Labs. The powder was introduced into the melt pool through an alumina tube. A computer numerical control (CNC) table was used to facilitate the movement during the deposition. Argon gas was used to ensure an inert atmosphere and act as a carrier gas to deliver the powder mixture to the melt pool.

In the current setup, the 2 mm spot size is insufficient to attain a large capture efficiency of the powder. This is due to the scatter of the powder flow out of the powder feed tube. This scatter was suspected to vary with individual precursor powder. Therefore, in order to obtain as-deposited compositions that are close to as-blended compositions, the capture efficiency during the deposition process needed to be increased. A trochoidal toolpath (shown in Figure 1b) was designed to create a large enough melt pool to improve capture efficiency during deposition. This toolpath was inspired by "weave"-style toolpaths that are commonly used in welding.

The AISI 304 substrates were cleaned with acetone to remove the impurities such as dirt and oil from the surface. A preheating scan was conducted by running the laser across the substrate surface. To ensure a successful start, the power of the initial five layers of the deposition was carried out at 750 W and 8.5% (3.36 g/min) powder feed rate. The remainder of the deposit was run at a power level of 550 W and 8.5% (3.36 g/min) powder feed rate. The thickness of each layer is 1 mm.

**Figure 1.** Schematic of the experimental setup, (**a**) laser metal deposition (LMD) system and (**b**) the trochoidal tool path.

After laser deposition, vertical transverse sections of the specimens were cut using a wire electric discharge machine (Hansvedt Industries Inc., Rantoul, IL, USA) and mounted in Bakelite for polishing and etching. The metallographic specimens were first ground using 240, 400, 600 and 800 grit silicon carbide papers and then polished using 15 μm, 9 μm and 3 μm diamond suspensions. The final step of polishing involved 0.05 μm colloidal silica suspension. To reveal the microstructure, the electrolytic etching was carried out in the nitric acid solution (70 mL nitric acid, 30 mL distilled water) at 5 V for 5 seconds. Scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS) and electron backscatter diffraction (EBSD) were performed on Helios Nanolab 600 SEM (Thermo Fisher Scientific, Waltham, MA, USA). The SEM image was acquired by an Everhart-Thornley detector. The EDS element was analyzed by the factory standardizations provided in the Aztec software. The EBSD step size was selected to be 2.5 μm. EBSD data acquisition and analysis were conducted using Aztec and Channel 5 software, respectively. Grain size was measured by the line intercept method, and the misorientation angle was 10◦. Optical microscopy images were collected using a Hirox optical microscope. X-ray diffraction patterns were collected using Philips X'pert MRD using Cu anode. The Vickers hardness was measured using a Struers Duramin hardness tester (Struers Inc., Cleveland, OH, USA) at a 9.8 N load and a 10 s load duration. The reported hardness results were the average of three indentations.

#### **3. Results and Discussions**

#### *3.1. Direct Coating of AlCoCrFeNi HEA on AISI 304 Substrate*

The direct LMD of the AlCoCrFeNi HEA on the AISI 304 substrate will be discussed first. Figure 2a shows a portion of the vertical transverse section of the HEA deposit near the AISI 304 substrate. An area close to the crack zone, as marked in the dashed-line box, is shown in Figure 2b with high magnification. A network of cracks, mostly transverse and horizontal in orientation, were found to be prevalent. Cracking occurred at the bottom of this HEA deposit. This could be attributed to the mismatch between the thermal expansion coefficients. The CTE of this HEA was reported to be 9.03 (10<sup>−</sup>6/K, 293–303 K) while the value of AISI 304 was 14.7 (10−6/K, 293 K) [18,19].

**Figure 2.** (**a**) Optical microscopy image of the vertical transverse cross-section of direct AlCoCrFeNi HEA coating on AISI 304 substrate, (**b**) a high-magnification view of the dashed–line-boxed area in (**a**).

The elemental composition distribution along the interface between the HEA deposit and the AISI 304 substrate is shown in Figure 3. At the bottom of the melted metal, the composition mixing was significant during the laser deposition process (see Figure 3). The bottom of the deposit had high susceptibility of cracking in the transverse cross-section, as seen in Figure 2.

The variation in Vickers hardness across the HEA–AISI 304 direct coating is presented in Figure 4. The average Vickers hardness of the HEA deposits was 412 HV, while that of the substrate was 161 HV. Since the coefficients of thermal expansion are mismatched between HEA and the substrate, residual stresses were developed during the laser deposition process. The AISI 304 substrate had a high elongation rate from 28% to 50% in the temperature range of 300–500 ◦C [23]. However, the tensile elongations of the AlCoCrFeNi HEA were 1% (as-cast condition) and 11.7% (after heat treatment) [4]. A difference in ductility exists between the substrate and the HEA. Having an intermediate material to bridge these differences was deemed necessary.

**Figure 3.** Elemental composition distribution along the interface between the AISI 304 substrate and the HEA deposit.

**Figure 4.** Vickers hardness profile of the direct coating of the AlCoCrFeNi alloy on AISI 304.

#### *3.2. A New Transition Route*

A blend of Fe, Co and Ni powders was selected as the candidate intermediate material. Since they are among the constituents of the AlCoCrFeNi HEA, no special procurement was needed. A Fe–Co–Ni ternary phase diagram at 1073 K compiled from experimental data is shown in Figure 5 [24]. Fe, Ni and Co have excellent mutual solubility, and no brittle intermetallic phases are expected. From the phase diagram, an atomic composition ratio of Fe, Ni and Co of 50%, 25% and 25%, respectively, was chosen. The selected ratio is expected to bridge the material composition gap between the AlCoCrFeNi HEA and AISI 304. This new transition route, AISI 304 substrate → CoFe2Ni intermediate layer → AlCoCrFeNi HEA, was then carried out and characterized.

**Figure 5.** Ternary alloy phase diagram of Fe–Co–Ni at 1073 K [24].

#### *3.3. AlCoCrFeNi HEA–AISI 304 with an Intermdeiate Layer*

#### 3.3.1. Microstructure

The CoFe2Ni intermediate layer was coated on the AISI 304 substrate using premixed elemental powder. Then, the AlCoCrFeNi HEA was coated on the intermediate layer by LMD. The intermediate layer composition was theorized to avoid the formation of intermetallic compounds and bridge the large gap in strength differences. Figure 6a,b shows the optical images of etched surfaces of transverse sections of these deposits. Unlike the HEA–AISI 304 direct coating, no apparent cracks were observed, which indicated an improvement in bonding. However, issues of microporosity persisted. A dendrite microstructure was observed along the interface between the intermediate layer and the HEA.

**Figure 6.** The optical microstructure of (**a**) the CoFe2Ni intermediate layer and the AISI 304 substrate and (**b**) the AlCoCrFeNi alloy deposit and the CoFe2Ni intermediate layer.

(**a**) (**b**)

A high-magnification secondary electron image of the AlCoCrFeNi HEA deposit is shown in Figure 7, where a two-phase dendritic microstructure was observed. The area fraction of the dendritic microstructure was ~52%, while the interdendritic area fraction was ~48%. The interdendritic region is named A, and the dendritic region is named B. The mean elemental compositions of A and B (average from three arbitrary points) were analyzed by EDS, and the results are listed in Table 4. It is shown that the atomic percentages of Al and Ni were ~29% in A and ~41% in B. The percentages of Fe and Cr were ~54 atom % in A and 43 atom % in B. These results indicate that Fe and Cr were rich in A, while Al and Ni were rich in B. The composition of Co did not show evident differences between A and B. The mixing enthalpies between Fe–Cr, Fe–Ni, Fe–Co, Fe–Al, Cr–Ni, Cr–Co, Cr–Al, Ni–Co, Ni–Al and Co–Al were −1, −2, −1, −11, −7, −4, −10, 0, −22 and −19 kJ/mol, respectively [6,25]. The mixing enthalpy of Al and Ni was higher than other pairs, which indicated that Al and Ni tended to form atomic pairs and segregate. Similar results have been reported for the AlCoCrFeNi HEA, with this microstructure being attributed to the spinodal decomposition [2,4–6,9].

**Figure 7.** Secondary electron image of the AlCoCrFeNi HEA microstructure at a magnification of 10000.

**Table 4.** Elemental compositions analyzed by energy dispersive X-ray spectroscopy (EDS)of the AlCoCrFeNi HEA shown in Figure 7.


XRD was used to identify the crystal structures of the intermediate layers and the HEA. A transition of the crystal structure was observed from the AISI 304 substrate to the AlCoCrFeNi alloy. The XRD patterns of the AISI 304 substrate, the CoFe2Ni intermediate layer and the AlCoCrFeNi alloy are shown in Figure 8. The present phases and the corresponding crystallographic information are summarized in Table 5. The peak patterns of FCC were observed in the CoFe2Ni intermediate layer, while BCC peak patterns were detected in the AlCoCrFeNi alloy. Löbel et al. found BCC and B2 (ordered BCC) phases in AlCoCrFeNiTi*<sup>x</sup>* (*x* = 0) when fabricated via arc melting [1]. A similar result was reported by Shiratori et al., when casting was employed to produce an AlCoCrFeNi HEA [26]. Due to the same basic lattice structure and lattice parameters, the B2 ordered structure is very hard to detect from XRD, as the peak patterns of B2 and BCC are the same [2,9]. However, the evidence of the existence of the B2 phase was found from the EDS analysis above. Previously, an AlCoCrFeNi HEA was reported to also contain the FCC crystal structure with preheating or post-heat treatment [5,13,26]. The FCC structure was not found in this work, which could be because the high cooling rate during LMD inhibited the formation of the FCC crystal structure [5,13,26].

**Figure 8.** XRD pattern of the AISI 304 substrate, the CoFe2Ni intermediate layer and the AlCoCrFeNi HEA.



The evolution in chemistry from the intermediate layer to the substrate was characterized by an EDS line scan first. The quantitative results are shown in Figure 9a. The EDS measured results of the AISI 304 substrate (Cr: ~18–19 atom %, Fe: ~70–72 atom %, Ni: ~9–10 atom % in Figure 9a) did not vary from the nominal AISI 304 elemental compositions. Mn (~1–2 atom %) was detected in the AISI 304 substrate by EDS but is not shown in Figure 9. The percentages of Co (~17–22 atom %) and Ni (~21–23 atom %) reduced, while the Fe (~54–56 atom %) content increased from the intermediate layer to the AISI 304 substrate. A small amount of Cr (~3–5 atom %) was present in the intermediate layer, because the substrate was mixed with the intermediate layer. The composition distribution from the HEA to the intermediate layer is shown in Figure 9b. The constituents of the AlCoCrFeNi HEA

were detected by EDS (Al: ~16–17 atom %, Co: 19–20 atom %, Cr: ~17 atom %, Fe: ~25 atom %, Ni: ~20–21 atom %). The difference between the as-blended (20 atom %) and as-deposited aluminum (~16–17 atom %) percentages is suspected to be a consequence of inconsistency in capture efficiencies of the constituent powders, and evaporation due to differences in melting point. Al and Cr were present in the intermediate layer as seen in Figure 9b, and their total content was ~4–5 atom %.

**Figure 9.** Elemental composition distribution along the boundary, (**a**) CoFe2Ni intermediate layer and AISI 304 substrate and (**b**) AlCoCrFeNi HEA and CoFe2Ni intermediate layer.

#### 3.3.2. EBSD

Figure 10a shows the inverse pole figure (IPF) map obtained from the bottom of the HEA section of the specimen. The measured area was approximately 3.4 mm × 1.2 mm of the cross-section parallel to the build direction (BD), which spanned from the left to the right of the specimen. The difference in color indicates the different crystallographic orientations. From Figure 10a, the overall constitution can be classified into two zones—the edge zone (1 and 3) and the middle zone (2). In areas 1 and 3, the grains were observed to be elongated along the build direction (see 1 and 3 in Figure 10a). The distributions of the intercept lengths (using 100 horizontal lines) in different areas are depicted Figure 10b. The median linear intercept for areas 1 and 3 was 72.5 μm, while it was 127.5 μm for area 2. From the linear intercept distribution of area 2, 25% of the intercept values were greater than 300 μm, whereas only 14% of the intercept values were above 300 μm for areas 1 and 3. This grain morphology is likely to be a consequence of deposition toolpath and variation in cooling rate at edges and in the middle [27,28]. Figure 10c,d show the {100}, {110} and {111} pole figures of different areas, which give the distribution of the pole density along the build direction. The pole figure of the areas 1 and 3 (Figure 10c) suggests that the orientations of the grains were close to the <100> direction. However,

the grains were random in orientation and did not appear with obvious texture in area 2 (Figure 10d). Further study is necessary to investigate the impact of this toolpath on the grain morphology.

**Figure 10.** (**a**) Inverse pole figure (IPF)IPF map of the bottom of the HEA section in the specimen; the measured region was approximately 3.4 mm × 1.2 mm, from the left to the right side in the cross-section parallel to the build direction (BD); (**b**) distribution of the intercept length of grains with the bin size of 10 μm; (**c**) pole figure of areas 1 and 3; and (**d**) pole figure of area 2 in (**a**).

#### 3.3.3. Vickers Hardness Analysis

Figure 11 gives the Vickers hardness distribution of the AlCoCrFeNi HEA deposited on the AISI 304 substrate with the CoFe2Ni intermediate layer. The Vickers hardness of the CoFe2Ni intermediate layer was around the 275 HV, which could be attributed to the solid solution strengthening. Table 6 lists the Vickers hardnesses of the AlCoCrFeNi HEA, annealed AISI 304, aged Inconel 625, and annealed duplex steel SAF 2205 [29–31]. The average Vickers hardness of the HEA deposit was in the range of 418 HV, because of the second-phase strengthening [4].

According to the XRD results, the AISI 304 substrate and the CoFe2Ni intermediate layer had an FCC structure, while the AlCoCrFeNi HEA had a BCC structure. The transition from FCC to BCC structure is also expected to enhance the hardness. The high hardness is expected to correlate with good performance in strength and wear resistance [1,16].

**Figure 11.** Vickers hardness profile of the AISI 304 substrate—AlCoCrFeNi HEA with the CoFe2Ni intermediate layer.



#### **4. Conclusions**

An AlCoCrFeNi HEA was coated on an AISI 304 substrate by laser metal deposition (LMD) technology. The coating on the substrate without and with the intermediate layer was characterized and discussed. The main conclusions are as follows:


**Author Contributions:** For this research article, W.C. designed and performed the experiments, analyzed data and wrote the manuscript; S.K. assisted in EDS analysis and manuscript review; X.Z. assisted in SEM analysis; E.B. contributed in EBSD analysis, and F.L. guided the research project.

**Acknowledgments:** The authors gratefully acknowledge the financial support from NSF (National Science Foundation) grants CMMI-1547042 and CMMI-1625736. The support from the Intelligent Systems Center (ISC) and Materials Research Center (MRC) for the help in sample preparation and materials characterization is also appreciated.

**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* **Processing and Characterization of Refractory Quaternary and Quinary High-Entropy Carbide Composite**

#### **Hanzhu Zhang and Farid Akhtar \***

Division of Materials Science, Luleå University of Technology, 971 87 Luleå, Sweden; hanzhu.zhang@ltu.se **\*** Correspondence: farid.akhtar@ltu.se

Received: 10 April 2019; Accepted: 28 April 2019; Published: 6 May 2019

**Abstract:** Quaternary high-entropy ceramic (HEC) composite was synthesized from HfC, Mo2C, TaC, and TiC in pulsed current processing. A high-entropy solid solution that contained all principal elements along with a minor amount of a Ta-rich phase was observed in the microstructure. The high entropy phase and Ta-rich phase displayed a face-centered cubic (FCC) crystal structure with similar lattice parameters, suggesting that TaC acted as a solvent carbide during phase evolution. The addition of B4C to the quaternary carbide system induced the formation of two high-entropy solid solutions with different elemental compositions. With the increase in the number of principal elements, on the addition of B4C, the crystal structure of the HEC phase transformed from FCC to a hexagonal structure. The study on the effect of starting particle sizes on the phase composition and properties of the HEC composites showed that reducing the size of solute carbide components HfC, Mo2C, and TiC could effectively promote the interdiffusion process, resulting in a higher fraction of a hexagonal structured HEC phase in the material. On the other hand, tuning the particle size of solvent carbide, TaC, showed a negligible effect on the composition of the final product. However, reducing the TaC size from −325 mesh down to <1 μm resulted in an improvement of the nanohardness of the HEC composite from 21 GPa to 23 GPa. These findings suggested the possibility of forming a high-entropy ceramic phase despite the vast difference in the precursor crystal structures, provided a clearer understanding of the phase transformation process which could be applied for the designing of HEC materials.

**Keywords:** high-entropy ceramic; solid-state diffusion; microstructure; phase evolution; hardness

#### **1. Introduction**

High-entropy ceramics (HECs), as a new class of ceramic materials, are developed from the concept of high-entropy alloys (HEAs). In HEAs, multiple principal elemental metals are incorporated to form a single phase alloy or multiphase composites [1,2]. With a similar design concept, HECs consist of multiple principal ceramic compounds such as metallic oxides, nitrides or carbides. In recent work, it has been found that the high entropy effect applies in the HECs system consisting of multiple components [3,4]. Solid solution phases that lack long-range order can be formed attributed to the minimization of Gibbs energy. Christina M. Rost et al. [5] fabricated an entropy stabilized oxide that showed a single-phase face-centered cubic (FCC) structure. Single phase solid solutions have been reported in systems like high entropy borides [6] and high entropy carbides [7,8]. Compared with conventional ceramic materials, single phase HECs show superior mechanical properties contributed by the strain strengthening effect from lattice distortion. For example, the high-entropy carbide fabricated by E. Castle et al. [7] exhibited a single phase FCC structure and enhanced hardness (36.1 GPa) compared to all the component carbides.

Due to the high melting point of ceramic materials, the solid-state pulsed current processing (PCP) is the economically preferable processing route as it offers relatively low sintering temperatures and short sintering times [9]. PCP utilizes pulsed current, directional pressure, and high vacuum during sintering and offers the advantage of providing homogenously densified materials [10,11]. Due to the strong covalent bonding and complex crystal structures of refractory ceramics, the synthesis of high entropy ceramics has been performed with ceramic precursors containing one nonmetal element (for example carbide, nitride or boride) and a similar crystal structure [6,7,12,13] using PCP. The selection approach of the ceramic precursors minimizes the geometrical difference among the starting ceramic precursor materials and achieves a single-phase high-entropy phase during PCP [14].

Our previous work on the high-entropy ceramic composite shows the possibility of forming a single-phase hexagonal structure from carbide precursors with a vast difference in the crystal structures (FCC, hexagonal and rhombohedral) and two nonmetal atoms (carbon and boron), attributing to the independent diffusion of metal and nonmetal atoms during the PCP processing [14]. However, the phase formation rules in a high-entropy ceramic system are still elusive and how B4C influences the diffusion process remains unclear. In order to develop a better understanding of the formation mechanism of the high-entropy ceramic phase, a systematic study was established and conducted in this work. Quaternary HEC composite was processed from four refractory carbides, HfC, Mo2C, TaC, and TaC. The phase evolution on B4C addition was studied by investigating the microstructure and phase composition of the five-component carbides. Moreover, due to the importance of starting particle size in solid-state diffusion [15], HEC composites were fabricated from precursors with different particle sizes. The effects of tuning the particles size of the solvent and the solute carbides on the microstructure, phase composition, and mechanical properties are discussed.

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

HfC (<1.25 μm, Sigma Aldrich, Darmstadt, Germany) and (−325 mesh, American Elements, Los Angeles, LA, USA), Mo2C (−325 mesh, Alfa Aesar, Haverhill, MA, USA) and (2.6 μm, Nanografi, Ankara, Çankaya), TaC (−325 mesh, Alfa Aesar, Haverhill, MA, USA) and (1μm, US Research Nanomaterials, Houston, TA, USA), TiC (5 μm, H.C. Starck, Munich, Germany) and (2 μm, Alfa Aesar, Haverhill, MA, USA) and B4C (1–7 μm, Alfa Aesar, Haverhill, MA, USA) were utilized to synthesize high-entropy ceramic composites. Starting materials with different particle sizes were mixed following the designed recipe, with a molar ratio of 2:1:2:2:2. The powder mixture was homogenized in a ball milling machine for two hours, using 4 mm stainless steel balls as a milling media and a powder to ball mass ratio of 1:5. The homogenized powder mixture was filled into a graphite die with a diameter of 10 mm and prepressed before sintering. The sintering was conducted in SPS-530ET (Dr. Sinter Spark Plasma Sintering System, Fuji electronic industrial Co., Ltd., Tsurugashima, Japan) in a glovebox. The samples were heated to 1800 ◦C with a heating rate of 100 ◦C/min and then held at 1800 ◦C for 5 min under vacuum. A uniaxial pressure of 60 MPa was applied during the process.

Before the microstructure characterization and phase identification, the samples were cold mounted in epoxy and polished following standard metallurgical sample preparation procedures. The microstructure was observed using a Scanning Electron Microscope JSM-IT300 (JEOL, Tokyo, Japan) operating at an acceleration voltage of 15 kV. The elemental composition was analyzed using energy dispersive spectrometer (EDS) mounted on JSM-IT300 that was calibrated with Cobalt. X-ray diffraction (XRD) was conducted using an X-ray diffractometer (Empyrean, PANalytical, Malvern, UK) with Cu-Kα radiation (wavelength 0.154 nm). The scanning was performed from 5 to 120◦ (2 Theta), with a step size of 0.02◦. The XRD data was investigated using the software PANalytical X'Pert HighscorePlus with the PDF-4 database. The nanohardness of sintered HEC composites was determined using an MTS NanoIndenter XP with a Berkovich diamond indenter. The measurements were performed with a maximum load of 100 mN at ambient temperature in air. The force-displacement curves of the indenter were recorded. Due to different phase distribution and grain sizes in the microstructure of the PCP samples, the nanohardness values for sample 4-HEC and 5-HEC were obtained by performing at least

5 measurements at specific regions based on the optical microscope mounted on MTS NanoIndenter XP, while the hardness values for sample HEC(+) and HEC(fine) were obtained by performing matrixes indentations and calculating the average value from indentations that were not located in porosity.

#### **3. Results and Discussions**

#### *3.1. E*ff*ect of B4C Addition into (HfMoTaTi)C*

The microstructure of the quaternary HEC composite (4-HEC) sintered from HfC, Mo2C, TaC, and TiC shows two distinct phases, as shown in Figure 1. The bright phase with a size of 1–5 μm is dispersed uniformly in the dark matrix phase in Figure 1b. Based on the energy-dispersive X-ray spectroscopy (EDS) mapping analysis in Figure 1g, the bright phase is rich in Ta whilst the dark phase contains all constitutional elements Hf, Mo, Ta, Ti and C. Because of the sensitivity of backscattered electron detector (BED) to the atomic number, phases with different densities appear with different contrast in the BED microstructure. The constituent carbides show a vast difference in the densities, from 14.62 g/cm<sup>3</sup> for TaC, 12.2 g/cm3 for HfC, 9.18 g/cm<sup>3</sup> for Mo2C to 4.93 g/cm<sup>3</sup> for TiC, therefore the consistent contrast of the dark region in the Figure 1b implies that it represents the quaternary ceramic phase containing all constitutional elements Hf, Mo, Ta, Ti, and C. The average atomic ratio of each phase obtained by performing EDS point analysis on several point locations are listed in Table 1. As EDS lacks the accuracy of quantitative analysis of light elements [16], the atomic ratio of the four metals is normalized to Ta. The results show that the bright phase contains Ta and C as the dominating elements and a minor amount of Mo and trace amount of Hf and Ti, while the dark phase contains all four metal elements with a relatively lower amount of Ta. According to the X-ray diffraction (XRD) data in Figure 1a, the PCP 4-HEC composite consists of two face-centered cubic (FCC) crystal structures with similar lattice parameters (a1 = 0.4429 nm, a2 = 0.4399 nm), as marked in the inset in Figure 1a. The BED microstructure, EDS analysis and XRD data, Figure 1 and Table 1, suggest the formation of high-entropy FCC solid solution containing all constitutional elements. In the previous work on high-entropy ceramic B4(HfMo2TaTi)C [14], it has been reported that TaC has the lowest metal vacancy formation energy among the precursor carbides, thus it acts as the solvent FCC lattice during the formation of the high-entropy phase [7], i.e., constituent atoms except Ta diffuse into the vacancies in the TaC lattice to form the multicomponent solid solution. Therefore, the Ta-rich phase and the high-entropy phase can be regarded as intermediates in the phase transformation from constitutional carbides to the hexagonal high-entropy ceramic phase. The multicomponent interdiffusion induces TaC lattice distortion, which in this case results in the reduction in the initial lattice parameter of TaC, 0.4460 nm (ICDD reference pattern of TaC: No. 03-065-0282). Based on the quantitative results shown in Table 1, the high-entropy phase with higher content of Hf, Mo, and Ti metal atoms experience intense atomic diffusion compared to the Ta-rich phase. Hence, in the quaternary high-entropy phase, the atomic position exchange between Ta and other metal atoms with similar or smaller atomic radii results in a smaller lattice parameter (0.4399 nm) than the TaC-rich phase (0.4429 nm). Furthermore, the high-entropy phase shows a hardness of 28.4 GPa, which is 23.5% higher than that of the Ta-rich phase (23 GPa) as shown in Figure 2, due to the lattice distortion induced strain strengthening effect. This result is in line with the previous reports on high-entropy materials [6,7].

**Figure 1.** XRD patterns and backscattered electron microstructure of PCP 4-HEC (**a**,**b**) and 5-HEC composite (**d**,**e**), and the EDS mapping analysis (**g**); (**c**) shows the volume change of the material during sintering; (**f**) is the EDS qualitative spectra on different phases of 5-HEC.



**Figure 2.** Elastic recovery parameter (ERP), nanohardness and the Young's modulus of the PCP HECs. The nanohardness values refer to the FCC solid solution in 4-HEC, 5-HEC, and average hardness properties for HEC(+) and HEC(Fine).

The curve of *Z*-axis displacement as a function of temperature was recorded during the PCP. As shown in Figure 1c, the decline of *Z*-axis refers to the thermal expansion of the material, while the up-climbing region corresponds to the shrinkage of the bulk volume. The reduction of the volume typically refers to the occurrence of sintering phenomenon where the powder material becomes compacted and forms a densified solid mass. Since the sintering temperature of powder material is normally 2/3–3/4 of the melting point [17], the theoretical sintering temperature for current quaternary refractory carbide mixture should be above 1800 ◦C. Figure 1c shows that the shrinkage of the four-component carbide system 4-HEC takes place from 1000 ◦C to 1600 ◦C, while the same phenomenon for the 5-HEC composite was postponed to a higher temperature range 1370 ◦C–1690 ◦C. This indicates that the addition of B4C to the carbide system hinders the solid-state atomic diffusion required for sintering in the multicomponent carbides, leading to a delay of the formation of the high-entropy ceramic phase. A detailed investigation on the sintering behavior of ceramic precursors to form high-entropy ceramic composites will be reported later, elsewhere.

The addition of B4C to the precursor carbides resulted in the formation of multiple phases during PCP. According to the backscattered electron microstructure in Figure 1e, the PCP 5-HEC composite exhibits three distinct phases. Similar to the 4-HEC, the brightest phase in 5-HEC is rich in Ta, which is coordinated with the fact that Ta has the highest atomic number among the constituent elements, therefore, the Ta-rich phase appears as the brightest phase in the BED microstructure. Two high-entropy solid solutions with different elemental compositions were formed in the five-component system, 5-HEC. The atomic ratio of metal atoms Hf, Mo, Ta and Ti in the gray (HEC1) and dark phase (HEC2) in Figure 1e are shown in Table 1, with HEC2 showing higher content of the solute metal atoms (Hf, Mo, and Ti) in the structure. The higher content of Ti in HEC2 agrees more with the darker contrast of HEC2 than HEC1 in the microstructure, as Ti has the smallest atomic number among the constitutional metal elements. According to the bond dislocation enthalpy (BDE) of transition metal carbides at 298 K [18], Ti-C has the lowest BDE of 423 ± 30 KJ/mol among the solute carbides, while Mo-C2 and Hf-C have a BDE of 500 and 540 ± 25 KJ/mol, respectively, and the covalent atomic radii vary as Ti < Mo < Hf. These factors might contribute to preferable diffusion of Ti over Mo and Hf during the formation of the multicomponent solid solution, resulting in the metal content ratio in HEC2 phase as Ti > Mo > Hf in Table 1. A pronounced boron diffraction peak in the EDS pattern was revealed at the HEC2 phase (Figure 1f), suggesting that HEC2 experienced a more intensive diffusion of B atoms than HEC1. Additionally, the diffractions peaks of B4C were not detected in the XRD diffractogram of the 5-HEC composite in Figure 1d, suggesting the participation of B4C in the formation of high-entropy solid solutions. However, the diffusion priority of Ti, Mo and Hf was not observed in the high-entropy

phase in the 4-HEC composite and HEC1 phase in 5-HEC, suggesting that the incorporation of B4C to the transition metal carbides might have promoted the diffusion process of the metal atoms towards a more energetically favorable state.

The XRD data in Figure 1d shows that the 5-HEC composite contains both FCC and hexagonal structured phases. Similar with the 4-HEC, the FCC pattern is generated from diffraction of two FCC crystal structures with similar lattice parameters, including the presence of an FCC Ta-rich phase. Based on the aforementioned discussions, the HEC1 phase that contains a lower content of the solute atoms should have a closely-matched crystal structure and lattice constant with the solvent TaC lattice due to the reduced atomic position change in the host lattice corresponding to an FCC structure, while the HEC2 phase corresponds to the hexagonal structure. The HEC1 phase shows a lattice parameter of 0.4499 nm, which is slightly greater than that of TaC (0.4460 nm). Assuming that the HEC1 is formed from only transition metal carbides, the lattice parameter should decline as observed in the FCC high-entropy solid solutions in 4-HEC composite (Figure 1a). It is known that the metal-boron bond length is longer than metal-carbon bond length, for example, Ta-C = 2.22 Å [19] and Ta-B = 2.41 Å [20], the expansion of the lattice can be contributed by the addition of B atoms in the formation of an HEC2 solid solution. The formation of high-entropy solid solutions HEC1 and HEC2 that contain all constitutional elements confirms the possibility of processing high-entropy ceramics from a precursor system containing more than one nonmetal atoms and with different crystal structures. The formation of hexagonal structure is likely to be attributed to more B atoms diffusing in the FCC lattice, which induces more severe lattice distortion and consequently leads to the crystal structure change from FCC to a hexagonal structure. In the 5-HEC composite, the FCC structured HEC1 solid solution shows a nanohardness of 27.4 GPa and Young's modulus of 505.8 GPa, which is close to the FCC solid solution phase in 4-HEC (28.4 GPa and 495.2 GPa for nanohardness and Young's modulus, respectively).

#### *3.2. E*ff*ect of Di*ff*erent Starting Particle Sizes*

It is well known that the particle size of the precursors is an essential parameter influencing the atomic diffusion and phase evolution during solid-state sintering in PCP [15]. Fine particles promote the solid-state atomic diffusion by reducing the diffusion distance and promote the kinetics of phase transformation [21,22]. To investigate the effect of particle size on the formation of high-entropy ceramics, the same carbide systems with different particle sizes are sintered in PCP (as listed in Table 2). 5-HEC that utilized relative larger particle sizes of solute metal carbides (HfC, Mo2C, and TiC) is discussed in the previous section and is denoted as HEC(++) in the following discussion. HEC(fine) and HEC(+) contains precursor carbides with the finest particle sizes and a larger particle size of the solvent carbide (TaC), respectively.


**Table 2.** Precursors with different particle sizes are utilized to study the effect on the phase evolution.

The microstructure of the PCP HEC(+) and HEC(fine) show the presence of two phases in Figure 3. According to the compositional mapping analysis in Figure 3d, all four metal elements are distributed in bright and dark phases. The spot-shaped mapping for C is attributed to porosity in the samples, which possibly caused the diamond polishing agents being introduced during the sample preparation procedure. For both PCP HEC(+) and HEC(fine) sample, the bright phase is rich in Ta and Hf while the dark phase contains a greater amount of Mo and Ti. The quantitative analysis results of the selected areas in Table 3 show that these two phases are high-entropy solid solutions with different elemental

compositions. The average atomic ratios are normalized to Ta. The bright phase is rich in Ta, while the dark phase has a higher content of other metals (Ti > Mo > Hf ≈ Ta). Based on the discussion about BED microstructure of different transition metal carbides, the dark region with a higher solute metal content refers to more equilibrium composition and a higher extent of phase transformation towards the high-entropy solid solution.

**Figure 3.** Microstructure (**a**,**b**) and X-ray diffraction phase identification (**c**) and EDS mapping analysis (**d**) of PCP HEC(+) and HEC(fine).

**Table 3.** The quantitative analysis of the metal atom contents of different phases in HEC(+) and HEC(fine).


The XRD patterns of HEC(+) and HEC(fine) show high similarity (Figure 3c). Similar to the diffractogram of HEC(++), the PCP HEC(+) and HEC(fine) composite reveal diffraction patterns from two FCC and one hexagonal structure. Since the elemental mapping from each region with a constant contrast shows a homogenous distribution of constituent elements, a possible reason why the third phase is not distinguishable in the microstructure is that the two FCC phases have indistinguishable contrast in BED images, which suggests that these phases have a similar atomic composition. As the bright phase with less content of foreign atoms (except Ta), it refers to the component with less lattice distortion, therefore it is extrapolated to retain a FCC crystal structure, whilst the dark phase exhibits a hexagonal structure.

For the solid-state phase transformations, the starting particle size has been reported to have a strong influence on the reaction kinetics by tuning the contact area between the solid particles [23–25]. Therefore, it was expected that PCP HEC(fine) with finest starting particle size should have a more promoted phase transformation than HEC(+) and HEC(++) which were fabricated with the same sintering route and sintering conditions. Comparing HEC(fine) with HEC(+), the results show high similarity in the microstructures and phase composition (Figure 3). Both composites consist of a

FCC and a hexagonal crystal structure, with high-entropy solid solutions. With HfC having the lowest formation enthalpy (−1.826 eV) [14] and highest BDE (9.18 g/cm3) among the precursors [18], the promotion of the solid-state diffusion can be observed by the increased Hf content in the FCC phase in HEC(fine) than HEC(+) (Hf/Ta = 1 and 0.4, respectively) according to the EDS quantitative results in Table 3. On the other hand, HEC(++) has a complicated phase composition compared to HEC(fine). The Ta-rich phase, representing the least diffusion, was observed in HEC(++) (Figure 1e). A lower fraction of the hexagonal structured high-entropy solid solution in HEC(++) composite can be concluded from the microstructure and lower intensity of X-ray diffraction peaks (for example at 2θ = 44.2◦) in Figure 3c. Therefore, the degree of diffusion is assumed to be in the order of HEC(++) < HEC(+) ≈ HEC(fine). The results suggest that the particle size of solvent carbide TaC is not as essential as the solute carbides in terms of tailoring the phase composition in the multicomponent carbide system. Due to the porosity and small grain size, the nanoindentation testing of HEC(+) and HEC(fine) composites were obtained from 30 indentations, therefore the data represents the overall mechanical properties of the bulk materials instead of each individual high-entropy phase. Although the PCP HEC(fine) and HEC(+) show the same phase composition and similar microstructure, the HEC(fine) composite sintered from precursors with finest grain size shows improved nanoindentation hardness of 23.1 GPa compared to the PCP HEC(+) composite (21.4 GPa). The experimental hardness value is close to the theoretical hardness calculated from the rule of mixture (23.2 GPa). The enhancement of the hardness caused by utilizing small starting particle size during sintering has been reported before [26]. Moreover, Young's modulus of HEC(fine) improves from 380.7 GPa for HEC(+) to 401 GPa, suggesting stronger atomic bonding in the HEC(fine) composite.

#### **4. Conclusions**

A high-entropy ceramic (HEC) composite was synthesized from HfC, Mo2C, TaC, and TiC by pulsed current processing (PCP). The PCP 4-HEC composite contained a high-entropy phase and a Ta-rich phase. Both constituents showed a face-centered cubic (FCC) structure and corresponded to the different extent of phase transformation towards a high-entropy phase. By introducing B4C in the HEC composite, high-entropy solid solutions that contained all principal elements (including B) were formed. The high-entropy phase with less intense interdiffusion remained FCC structure as the solvent carbide, TaC, while the one with more complete phase transformation experienced crystal structure change from FCC to a hexagonal structure. The results showed the feasibility of synthesizing HEC materials from the multi-principal ceramic system with different crystal structures.

In order to investigate the effect of different particle sizes of solvent and solute components on the phase composition and properties, HEC composites fabricated from precursors with different size of solvent (TaC) and solute carbides (with face-centered cubic (FCC) structure) were PCP consolidated and characterized. The particle size of the solvent carbides was found to be more essential for the interdiffusion process and final phase compositions than that of the solute carbide. HEC composites processed from the fine particle size of solvent carbides HfC, Mo2C, and TiC showed a higher content of the hexagonal structured HEC phase. On the other hand, reducing the TaC particle size to nanoscale showed negligible influence on the phase composition, but resulted in enhancement of the microhardness from 21.4 GPa to 23.1 GPa.

**Author Contributions:** Conceptualization, F.A.; methodology, F.A. and H.Z.; formal analysis, H.Z.; investigation, H.Z.; writing—original draft preparation, H.Z.; writing—review and editing, F.A.; supervision, F.A.; project administration, F.A.

**Funding:** This work was supported by the Swedish Foundation for Strategic Research (SSF) for Infrastructure Fellowship, grant number RIF14-0083.

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

#### **References**


© 2019 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* **Microstructure and Mechanical Properties of Particulate Reinforced NbMoCrTiAl High Entropy Based Composite**

**Tianchen Li 1, Bin Liu 1,\*, Yong Liu 1, Wenmin Guo 2,\*, Ao Fu 1, Liangsheng Li 1, Nie Yan <sup>3</sup> and Qihong Fang <sup>4</sup>**


Received: 13 June 2018; Accepted: 6 July 2018; Published: 10 July 2018

**Abstract:** A novel metal matrix composite based on the NbMoCrTiAl high entropy alloy (HEA) was designed by the in-situ formation method. The microstructure, phase evolution, and compression mechanical properties at room temperature of the composite are investigated in detail. The results confirmed that the composite was primarily composed of body-centered cubic solid solution with a small amount of titanium carbides and alumina. With the presence of approximately 7.0 vol. % Al2O3 and 32.2 vol. % TiC reinforced particles, the compressive fracture strength of the composite (1542 MPa) was increased by approximately 50% compared with that of the as-cast NbMoCrTiAl HEA. In consideration of the superior oxidation resistance, the P/M NbMoCrTiAl high entropy alloy composite could be considered as a promising high temperature structural material.

**Keywords:** high entropy alloys; metal matrix composites; mechanical properties

#### **1. Introduction**

Recently the concept of high-entropy alloys (HEAs), consisting of four or more principle metallic elements in equiatomic or near-equiatomic ratios, has attracted considerable interest owing to its tendency to form solid solution with outstanding mechanical and functional properties [1–6]. By introducing refractory elements including group IV (Ti, Zr, and Hf), V (V, Nb, and Ta), and VI (Cr, Mo, and W) in HEAs, the developed refractory HEAs (RHEAs) possess high melting temperature, outstanding strength and hardness, high thermal stability, and softening resistance at elevated temperatures, which opens up the possibilities of the alloy developments satisfying structural demands at high temperature.

Since the first WMoTaNb RHEA was reported in 2010, many literatures have been focused on the preparation methods and mechanical properties of RHEAs [7,8]. For example, the BCC WMoTaNb and WMoTaNbV HEAs maintained yield strength of 405 and 477 MPa at 1600 ◦C respectively, which were much higher than that of Inconel 718 and Haynes 230 [9]. However, most of these RHEAs are generally characterized with high density and poor high temperature oxidation resistance. Such drawbacks are the current bottleneck for utilizing RHEAs as structural materials. In order to improve the oxidation resistance properties of the RHEAs, it was reported that the Ti, Al, Cr, and Si elements are essential for the formation of protective oxide scales. Compared with the traditional alloys, it was widely reported that the RHEAs possessed low diffusivity as well as a wider range of composition design [10]. In addition, it was also reported that the density of RHEAs could be significantly reduced by the additions of light elements, such as Ti [11,12]. Therefore, on the basis of design principles above, a new RHEA with equiatomic composition of NbMoCrTiAl was designed by B. Gorr et al. [13,14]. The as-cast alloy with low density (7.58 g/cm3) exhibits excellent oxidation resistance. The declining oxidation rate at 1100 ◦C was mainly due to the formation of Al, Cr-rich oxide layers. However, the compressive strength of the as-cast NbMoCrTiAl alloy was reported to be as low as 1010 MPa and its high temperature strength is insufficient.

As reported in the literature, most of the RHEAs were prepared by the arc melting process. These as-cast alloys generally exhibited a dendritic structure with severe composition segregation [11–13]. Powder metallurgy (P/M), which is considered a high efficiency technique, is widely used to synthesize non-equilibrium materials. The materials characterized with fine grains are without chemical segregation and evaporation of alloying elements caused by arc-melting [15,16]. As reported, P/M WMoTaNbV shows fine grains of micrometer-scale and ultra-high compressive yield strength of 2612 MPa [17]. Therefore, it is an effective way to prepare high performance RHEAs.

Another prospective candidate for the high strength structural materials was the metal matrix composite (MMCs). The high performance was mainly attributed to the interaction between the particles and dislocations, called dispersion strengthening effect and second-phase strengthening effect. As the NbMoCrTiAl matrix is selected, the crucial challenge is to select the suitable strengthening phase and numbers of papers can be used as references [18–21]. The reinforcement particles include carbides (TiC, SiC), oxides (Al2O3, Y2O3), nitrides, (Si3N4, BN), and borides (TiB2, LaB6). There are mainly two ways to prepare carbides and oxides strengthened composites. One simple measure is to add them directly, another way is to form in-situ carbides or oxides by introducing C and O elements, which can effectively reduce material defects like pores and cracks. It was reported that the TiC and Al2O3 are also commonly used as strengthening phases because of their intrinsic high strength characteristics [22,23].

In this work, in order to further increase the strength of NbMoCrTiAl RHEA, the P/M composite with dispersed carbides and oxides was developed through ball milling and spark plasma sintering (SPS) method. The phase constitutions, microstructures, and mechanical properties of the NbMoCrTiAl HEA composite were investigated in detail.

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

Powders of Nb, Mo, Cr, Ti, and Al with purity higher than 99.5 wt. % and particle size of ≤45 μm were used as starting materials. The elemental powder was mixed in equiatomic composition and then processed by high energy planetary ball milling for 39 h at 300 rpm in a pure argon atmosphere (YXQM-2L, MITR, Changsha, China). Tungsten carbide vials and balls were utilized as the milling media with a ball-to-powder ratio of 10:1. The empirical addition amount of stearic acid was 3.5 wt. %. The stearic acid was introduced as process control agent (PCA) to prevent cold welding as well as particle agglomeration, and also acted as additive for generating particulate reinforced HEA composite. Subsequently, the as-milled powders were consolidated by SPS (D25/3, FCT, Munich, Germany) under 30 MPa axial pressure at 1700 ◦C with the heating rate of 100 K/min. During the sintering process, the sample was held at the sintering temperature for 30 min. Finally the sample cooled to room temperature in furnace.

The samples after sintering were analyzed by X-ray diffraction (XRD, D/MAX-2250, Rigaku, Tokyo, Japan) with a Cu Kα radiation at 40 kV and 200 mA. The microstructure of the composite was characterized by field emission scanning electron microscope (FESEM, Quanta FEG250, FEI, Hillsboro, OR, USA). The approximate volume fraction and average grain size of different phases was evaluated by at least ten images using Image Pro software. The phase composition and element distribution were investigated with electron probe microanalysis (EPMA, EPMA-1600, JEOL, Tokyo, Japan). Microhardness of the sample was measured by a Vickers microhardness tester (MicroMet-5104, Buehler, Lake Bluff, IL, USA) with an indenter load of 100 g and the holding time was 15 s. Cylindrical specimens of compression tests (4 mm in diameter and 6 mm in height) were cut and machined from the sintered bulk. Room temperature compression tests were carried out on an Instron-3369 (INSTRON, Norwood, MA, USA) universal testing machines at strain rate of 10−<sup>3</sup> s<sup>−</sup>1.

#### **3. Results**

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

The general microstructure of the HEA composite is shown in Figure 1. The specimens are highly densified by the reaction sintering process without visible pores. The composite presented three obviously identifiable contrasts, namely light, grey, and black regions (marked as A, B, and C). The volume fraction of the light phase, grey phase, and black phase are 60.8%, 32.2%, and 7.0%, respectively. Since the content of carbon and oxygen cannot be measured accurately by EDS, EPMA experiment was employed to analyze the chemical composition in these areas (Table 1). The results indicate that the region A consists of only 10.66 at. % Ti and 16.79 at. % Al, while Ti and Al atoms are enriched in the region B and C, respectively. Noticeably, the region C in Figure 1 was enriched in Al and O, indicating the formation of Al2O3.

**Figure 1.** Microstructure of the NbMoCrTiAl composite after spark plasma sintering at 1700 ◦C for 30 min under a pressure of 30 MPa.

**Table 1.** Chemical composition of high entropy alloy (HEA) composite determined by electron probe



The XRD patterns of the RHEA composite were illustrated in Figure 2. Two types of characteristic diffraction peaks could be identified, indicating the present material is primarily composed of BCC crystalline structure (a = 0.315 nm) with diffraction peaks at 2θ = 43.66◦, 50.82◦ 74.67◦ accompanied with a certain amount of TiC phase (a = 0.435 nm) as the reinforcing phase. Al2O3 cannot be detected by X-ray analysis, which was probably ascribed to the limit detection of 5%. In consideration of the EDS analysis results, it could be concluded that the light, grey, and black phase were identified as BCC HEA phase, TiC phase, and Al2O3 phase, respectively. The average particle size of TiC and Al2O3 were measured to be approximately 4.60 μm and 1.57 μm, respectively.

**Figure 2.** XRD pattern of the NbMoCrTiAl composite.

In order to further investigate the element distribution, the element mapping of the RHEA composite was shown in Figure 3. It was clear that the Mo and Cr elements are enriched almost entirely in the light region (BCC phase). On the contrary, the Ti and C elements are enriched almost entirely in the grey region (TiC phases) with certain amount of Nb, while O and Al are enriched in the black region (Al2O3 phase). No apparent debonding between dispersed particles and the HEA matrix could be found, which suggested that the secondary phase grains and the HEA grains possessed good bonding.

**Figure 3.** The elemental mapping of the composite obtained by electron probe microanalysis (EPMA).

#### *3.2. Mechanical Properties*

The room temperature engineering stress-strain curves of the Nb20Mo20Cr20Ti20Al20 composite was given in Figure 4. Despite no ductility, the fracture strength of the composite was 1542 MPa, increasing by more than a half than that of the as-cast alloy (1010 MPa) [13]. The fracture strength was also much higher than that of the traditional WMoTaNb and WMoTaNbV RHEAs [24], which have yield strengths of 1246 MPa and 1058 MPa, respectively, as shown in Figure 4. As a result of the plasticity and oxidation resistance of the NbMoCrTiAl matrix at high temperature, it was concluded that the P/M NbMoCrTiAl HEA composite had relatively superior mechanical properties and could be considered as a potential high temperature structural material [25].

**Figure 4.** The room temperature engineering stress-strain curves of the composite and typical RHEAs reported in the literatures [13,24].

Microstructure features of the fracture surface of the composite deformed at room temperature are shown in Figure 5. Cleavage fracture is featured by river pattern, which could be clearly seen in Figure 5. This means that the NbMoCrTiAl HEA composite shows typical cleavage fracture mode without ductility. The analysis result is in accordance with the deformation curves above.

**Figure 5.** Microstructure features of the fracture surface of the composite deformed at room temperature (**a**) at low magnifications and (**b**) high magnifications.

#### **4. Discussion**

#### *4.1. Phase Formation*

Although a lack of HEA phase diagrams limit the availability of thermodynamic data, some thermodynamic properties, such as atomic size, melting point, mixing entropy, mixing enthalpy, and valence electron concentration, can be used as criteria to predict the phase formation in the HEA system. According to the empirical rules, BCC solid solutions were generally formed in HEA systems under the following certain conditions. (1) The electron concentration (VEC) was lower than 6.87 [26]; (2) a new parameter combining effects of entropy and enthalpy (Ω) was proposed to be more than 1 and atomic size differences (δ) was less than 6.6% [27]. δ, Ω, and VEC are calculated to be 5.4, 2.92 and 4.80 respectively, so BCC solid solution could be anticipated, which tallied well with the research results produced by H. Chen et al. [13]. Guo et al. [28] further proposed that all ductile alloys have a VEC ≤ 4.4, while all brittle alloys have a VEC ≥ 4.6. Therefore, the NbMoCrTiAl RHEA should be without ductility and it is indeed a type of brittle RHEA.

As mentioned above, BCC phase, TiC phase, and Al2O3 phase can be observed in Figure 1. The actual composition of BCC phase is Nb20.3Mo31.0Cr19.3Ti10.7Al16.7, ignoring the contributions

made by interstitial elements of carbon and oxygen. According to the actual composition of the BCC high entropy phase, all the parameters calculated for the composition presented in this paper are summarized (δ = 5.6, Ω = 3.46, VEC = 4.96). According to the data, the HEA phase should be of BCC structure, and Figure 2 verifies accuracy of the prediction. As a result, we could confirm that such non-equiatomic phase composition is able to enhance the solid solution especially at high temperature [29].

Combining high energy ball-milling of powder precursors with the following SPS compaction, the HEA composite was successfully synthesized in this work. The addition of stearic acid can be used as an effective way for the preparation of composite materials. During the process of high energy ball-milling, PCA is used to not only prevent cold-welding, but also act as the source of the elements carbon and oxygen. Table 2 shows the values of Hmix (kJ/mol) calculated by Miedema's model for atomic pairs between elements with Nb, Mo, Cr, Ti, Al, and C. As shown in Figure 2, this certain concentration of Nb incorporated during grain growth agrees with the large solubility of solute Nb-carbide in solvent TiC mainly because of their largest negative value of Hmix and identical crystal structure. While Mo, Cr, and Al could not dissolve in the TiC due to their relatively weak cohesive force [30].

**Table 2.** Mixing enthalpies (kJ/mol) of unlike atomic pair [30].


#### *4.2. Strengthening Mechanism*

The composite created a kind of material with highly improved compressive strength, increasing from 1010 MPa to 1542 MPa. This remarkable increase in mechanical properties can be attributed to the following reasons. The most important one is the effect of second phase strengthening. It can be found that introduction of brittle ceramic reinforcements would significantly increase mechanical properties of the material. As listed in Table 3, the yield strength of (FeCrNiCo)Al0.7Cu0.5 base HEA reached 630 MPa accompanied by a high strain of 42.7%. With the addition of 10 vol. % TiC, the yield strength increased to 1290 MPa, while plastic strain dropped to 29.2% [31]. FeCoCrNiMn high entropy alloy matrix nanocomposite with addition of SiC was prepared by hot isostatic pressing, which result in the value of yield strength increased to 1600 MPa [18]. In addition, the interphase boundary between the in-situ reinforced particles and the solid solution matrix is free of cracks and pores, resulting in the more possibilities for the design of advanced high strength structural components.

**Table 3.** Room temperature mechanical properties of HEAs and their composites.


Another main factor is grain refinement strengthening mechanism. The grain size of the as-homogenized NbMoCrTiAl phase is about 100 μm, while that of the composite is much lower. This is because the PM process is an effective approach to refine grains compared to arc-melting [16]. In addition, the hard phases evenly distributed at the grain boundary, acting as the barrier of grain

growth. Praveen et al. [32] reported that the grain growth of the FeCoCrNi composite could be obviously suppressed by two-phase mixture (FCC-HEA and carbide) microstructure.

The third reason is the influence of solution strengthening. It is true that the solution element could have remarkable strengthening effect. In this study, the BCC phase contains 0.69 at. % O and 1.28 at. % C, which may cause significant increase in strength because these two interstitial atoms have a substantially smaller size than any of the metal atoms in the HEA. These interstitial atoms will produce a substantial strain field, with which gliding dislocations will interact. There have been studies on the influence of interstitial elements, such as element C, O, and N. Wang et al. [33,34] reported that the addition of 1.1 at. % carbon to a novel single-phase FCC Fe40.4Ni11.3Mn34.8Al7.5Cr6 HEA not only markedly increased the yield strength from 159 MPa to 355 MPa, but also led to a 25 % increase in the elongation to fracture. An increase in yield strength was also observed in the C-dissolved FeCoCrNi fabricated by selective laser melting [35]. With regard to the carbide, the lattice constant of TiC (0.435 nm) is slightly larger than pure TiC (0.433 nm). The different crystal parameter may result from the solution of atoms with large radius, such as element Nb (r = 0.148 nm) in this composite. Therefore, we may reasonably conclude that solution-strengthening effect resulted in the improved mechanical properties.

#### **5. Conclusions**


**Author Contributions:** Y.L., B.L. and W.G. conceived and designed the experiments T.L. and A.F. prepared the NbMoCrTiAl HEA composite and performed the microstructural characterization of the composite under the supervision of Y.L. and B.L.; the mechanical testing under the supervision of B.L. and W.G. All authors discussed the results and approved the final manuscript.

**Funding:** This research was funded by the National Natural Science Foundation of China (51771232), the National Key Research and Development Plan of China (2016YFB0700302), the Science and Technology Planning Project of Hunan Province of China (2015SK1002-1), Natural Science Foundation of Hunan Province (2018JJ3477), the State Key Lab of Powder Metallurgy and Project of Innovation, and Entrepreneur Team Introduced by Guangdong Province, China (201301G0105337290).

**Conflicts of Interest:** The authors declare no conflicts 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/).
