**Microstructure and Mechanical Properties of Precipitate Strengthened High Entropy Alloy Al10Co25Cr8Fe15Ni36Ti6 with Additions of Hafnium and Molybdenum**

#### **Sebastian Haas 1, Anna M. Manzoni 2, Fabian Krieg <sup>1</sup> and Uwe Glatzel 1,\***


#### Academic Editor: Yong Zhang

Received: 28 January 2019; Accepted: 8 February 2019; Published: 12 February 2019

**Abstract:** High entropy or compositionally complex alloys provide opportunities for optimization towards new high-temperature materials. Improvements in the equiatomic alloy Al17Co17Cr17Cu17Fe17Ni17 (at.%) led to the base alloy for this work with the chemical composition Al10Co25Cr8Fe15Ni36Ti6 (at.%). Characterization of the beneficial particle-strengthened microstructure by scanning electron microscopy (SEM) and observation of good mechanical properties at elevated temperatures arose the need of accomplishing further optimization steps. For this purpose, the refractory metals hafnium and molybdenum were added in small amounts (0.5 and 1.0 at.% respectively) because of their well-known positive effects on mechanical properties of Ni-based superalloys. By correlation of microstructural examinations using SEM with tensile tests in the temperature range of room temperature up to 900 ◦C, conclusions could be drawn for further optimization steps.

**Keywords:** HEA; high entropy alloys; compositionally complex alloys; mechanical characterization

#### **1. Introduction**

For centuries, constructions with high mechanical requirements have been built using steel or iron-based alloys, while titanium or aluminum have been used for light-weight constructions. In gas turbines nickel- or cobalt-based superalloys are used because of their extremely good mechanical properties and oxidation resistance over long periods of time at high temperatures. The main properties of these alloys are given by their base element, while the addition of other elements in small amounts leads to fine adjustments of a specific behavior [1].

An unconventional approach of alloy design is based not on a main element, but on a chemical composition that exhibits a large number of elements with none of them dominating. In this group one differentiates between single-phase high entropy alloys (HEA) [2] and multiphase compositionally complex alloys (CCA). Multicomponent alloys were examined first by Yeh [2] and Cantor [3]. Investigating the microstructure of alloys with different numbers of elements in equiatomic mixtures, Cantor found the five-component, equiatomic alloy Co20Cr20Fe20Mn20Ni20 to form a single-phase solid-solution. High configurational entropy, due to the large number of elements and similar concentrations, leads to a decrease in the Gibbs free energy and thus a stabilization of one solid-solution phase if no intermetallic phase with low enthalpy of formation exists [2]. Such a single-phase material was claimed to have outstanding properties, e.g., a good strength-ductility behavior over a wide

temperature range, thermal stability, wear resistance and high resistance against oxidation of the material [4].

Most alloys containing at least five elements in near equiatomic composition, however, do not crystallize as a single-phase solid-solution, but rather form intermetallic compounds as secondary phases with even lower Gibbs free energy. This leads to the formation of compositionally complex alloys (CCA). They cover a wide range of chemical compositions, with a high number of occurring phases and therefore a wide perspective in terms of applications. We restricted ourselves to have at least five elements and no element should dominate the composition. In the domain of high-temperature materials a gap between steels (<650 ◦C) and nickel-based superalloys (>850 ◦C) might be filled by less cost-intensive compositionally complex alloys. The focus should be on mechanical strength, oxidation resistance, processability and, last but not least, the cost factor. We therefore compared our alloy with commercially available alloys. The Ni-Co-Cr-Mo alloy IN 617 shows high-temperature strength and oxidation resistance, while Alloy 800 H has good creep properties and a quite ductile behavior at temperatures below 600 ◦C. The base alloy of our work shows a similar melting range and is optimized with respect to the features mentioned above. Possible industrial applications could be: Chemical industry appliances and especially land-based steam turbines as well as parts for gas turbines in sections with lower temperatures (700–800 ◦C).

The development of CCA for mechanical applications at both room and higher temperature is based on different microstructural approaches (see the review article by Manzoni [5]): Interpenetrating phases lead to a high resistance of the material, and additionally interesting electrical or magnetic properties may appear [6]. One example to strengthen the equiatomic high entropy alloy CoCrFeMnNi is the addition of 2 at.% carbon. This leads to a fcc-microstructure with embedded Cr23C6 particles and higher strength than for the CoCrFeMnNi alloy. The best strength-ductility relation of this carbide-strengthened alloy can be reached by applying the correct thermomechanical treatment [7]. Another example of particle strengthened CCA is the alloy Al12.35Co17.5Cr17.5Fe35.15Ni17.5 (in at.%). This microstructure exhibits a bcc-matrix with coherently embedded particles of B2 structure. The cuboidal nanoscale particles lead to high strength at room and high temperatures [8].

These three mentioned approaches show the diversity in designing CCA for mechanical applications, using strengthening mechanisms based on very different microscopic features. Singh et al. [9] examined the six-component, equiatomic alloy AlCoCrCuFeNi, with a focus on solidification behavior, using different techniques to analyze the microstructure. Even a high cooling rate by splat-quenching leads to a phase decomposition and the cast alloy shows a microstructure with several phases of different crystal structures, unlike expected solid-solution stabilization due to a high configurational entropy. Recognizing the difficulties of avoiding phase separation, a transition was made from developing a HEA to optimizing a CCA.

Much effort was put into optimizing alloys with near equiatomic composition of AlCoCrCuFeNi, investigating changes in composition and finding correct heat treatment parameters [10], resulting in the chemical composition Al10Co25Cr8Fe15Ni36Ti6 (in at.%). This alloy composition is the base alloy for this work, exhibiting L12-ordered, coherently embedded precipitates in a fcc-matrix, comparable to nickel-based superalloys. Additionally the base alloy shows another third phase, appearing in a needle-like shape up to a length of 50 μm. These needles are very rich in aluminum (28 at.%), with a reduction of all other elements, especially Fe, Co and Cr. TEM investigations identified this phase to have a Heusler type structure.

A way to improve the mechanical behavior is directional solidification, where grains are oriented in loading direction. As a material itself has quite good creep properties, for example in the case of superalloys, the grain boundaries play a strong role in weakening the material at higher temperatures. Avoiding grain boundaries inclined at 45◦ by the loading direction, shear stresses along these weaker boundaries can be eliminated and the strength is just defined by the microstructure inside of the longitudinal orientated grains [11].

Nickel-based superalloys provide further ideas to improve the high-temperature properties, due to similarities in microstructure. By the addition of Hf and/or Mo, different approaches for optimization are expected: Molybdenum is supposed to partition to the matrix, to have a solid-solution strengthening effect due to its bigger atomic size and therefore to lead to a distorted lattice, acting as an obstacle for dislocation movement [12]. The positive effect of hafnium is due to a strengthening of grain boundaries: A columnar-grained microstructure for example shows an enhancement of creep ductility and lifetime by the addition of hafnium that strengthens the vicinity of grain boundaries vertical to the load stress [13]. Doherty et al. [14] explain this by Hf participating more likely to the γ -phase Ni3(Al, Hf), leading to a strengthening of it. The grain boundaries are therefore strengthened by the fracture-retarding effect of the interlocking γ -configuration in this areas [15].

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

#### *2.1. Alloy Preparation*

In this work all chemical compositions are given in at.%. The composition of the base alloy Al10Co25Cr8Fe15Ni36Ti6 was changed by adding 0.5 at.% hafnium or 1.0 at.% molybdenum on the expense of aluminum. The added amounts and the reduction of aluminum were identified by simulations using the ThermoCalc software [16] with the database TTNi7 [17]. The base alloy in its original composition is also considered.

All constituents with a purity of 99.99% were cleaned in ethanol using an ultrasonic bath and were then melted in a vacuum induction furnace. The material was distributed randomly in a ceramic crucible in the middle of a water-cooled Cu-coil. After evacuating the chamber twice to a pressure of <sup>5</sup>·10−<sup>4</sup> mbar it was flooded with argon to prevent the evaporation of elements, especially chromium. The ceramic mold was heated up to a temperature of 1400 ◦C by a second coil and a graphite receptor, thus the material remained in liquid state after casting. To achieve directionally solidified grains in the [001]-direction, the Bridgman process was used and the mold was withdrawn through a water cooled baffle with a speed of 3 mm/min. The cast rods, with a diameter of 20 mm and a length of about 110 mm, were homogenized for 20 h at 1220 ◦C in the case of the Mo-containing alloy and at 1140 ◦C in the case of the Hf-containing alloy to avoid eutectic formation determined by differential scanning calorimetry. Subsequent annealing was performed in two different ways for both alloys, 900 ◦C/50 h and 950 ◦C/100 h respectively. After heat treatment the rods cooled down to room temperature in the furnace. To remove the oxide layer the samples were initially sand-blasted and afterwards treated with aqua regia. The rods were cut by electrical discharge machining to obtain samples for microscopic and mechanical characterization.

#### *2.2. Microstructural Observations*

Flat disks were cut from the rods and cut again in the middle to examine the cross and longitudinal section of the microstructure. These surfaces were embedded in a conductive resin, ground, polished with 1 μm diamond slurry and finally polished chemically. The specimens were etched with a solution of 3 g Mo-acid in 100 mL H2O, 100 mL HCl and 100 mL HNO3 to achieve a better phase contrast by dissolving γ - and Heusler type phase. For the examination we used a scanning electron microscope (SEM) Zeiss 1540EsB Cross Beam, operating under an accelerating voltage of 30 kV and using the SE2-detector for imaging. Precipitate size and volume fraction were determined using the classification of the Weka segmentation method [18] in the open source software Fiji [19], based on ImageJ [20,21]. More than 500 particles per state were analyzed.

#### *2.3. Mechanical Tests*

Electrical discharge machining was employed to obtain flat specimens for high temperature tensile tests. The square cross section of the samples was 1.0 × 1.9 mm2 and the gauge length was 8 mm, while the entire length of the sample was 25 mm. The specimens were cut out in such a way that the tensile direction was parallel to the [001]-grain-orientation. Before assembling the specimen, their surfaces were ground and a type-S thermocouple was welded for the regulation of temperature on the lower end of the gauge length. The sample was then attached to the ceramic clamping, the radiant heated furnace was closed and heated up to the desired temperature. Tensile tests were performed with a deformation rate of 0.01 mm/s (corresponding to 1.3 × <sup>10</sup>−<sup>3</sup> 1/s). A load cell and a high-resolution camera were logging about four pairs of values for stress and strain in one second. This lead to engineering stress–strain curves for each test, providing mechanical parameters like ultimate tensile strength (UTS), yield strength (YS) and strain to failure (εf).

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

#### *3.1. Chemical and Microstructural Analysis of Al10Co25Cr8Fe15Ni36Ti6*

The alloy Al10Co25Cr8Fe15Ni36Ti6 exhibited a dendritic microstructure after the casting process. Dendrites could be dissolved by a homogenization heat-treatment at 1220 ◦C for 20 h. Subsequent annealing for 50 h at 900 ◦C lead to a three-phase microstructure: Figure 1a) shows the large (several 10 μm), randomly distributed Heusler type phase, with its characteristic needle-like shape and a volume-fraction of <5%. The γ -microstructure is displayed in Figure 1b) with a higher magnification with L12-ordered cuboidal shaped γ -precipitates and small matrix channels with round secondary γ -particles. The matrix has a face-centered cubic structure [22].

**Figure 1.** SEM-images of (**a**) Heusler type phase with a needle-like shape and a length up to 50 μm and (**b**) cuboidal γ -particles with an edge-length up to 400 nm and matrix channels with small and spherical secondary γ -particles (some 10 nm).

The chemical analyzation of all phases, determined by TEM/EDS, is listed in Table 1 and results in a Co-Fe-Cr rich fcc-matrix and Ni-Al-Ti rich L12-ordered precipitates [10]. The Heusler type phase is very rich in Al, while all other elements are depleted, especially Fe, Co and Cr.

**Table 1.** Chemical composition (in at.%) of all occurring phases in the annealed state (900 ◦C for 50 h) of the alloy Al10Co25Cr8Fe15Ni36Ti6 analyzed by TEM/EDS [10].


#### *3.2. Impact of Bridgman Process on Mechanical Properties*

An increase of tensile strength was expected by the use of the Bridgman process, resulting in directional solidified samples. The base alloy was cast twice, conventionally cast with randomly orientated grains and directionally solidified, as described in Section 2 "Materials and Methods". The heat treatment was equal for both conditions (annealing at 900 ◦C for 50 h). The phase-characteristics (content and size) concerning Heusler type phase and γ -phase were identical for both types of processing techniques and are displayed in Figure 1. Figure 2 shows the grain structures of both states, confirming elongated grains in the direction of load (σ) in the case of directional solidification, while the conventionally cast microstructure exhibits grain boundaries across the orientation of external stress.

**Figure 2.** SEM-image of (**a**) conventionally cast, polycrystalline and (**b**) directionally solidified microstructures of the base alloy Al10Co25Cr8Fe15Ni36Ti6 after annealing at 900 ◦C for 50 h.

Results of tensile tests, carried out over a temperature range from room temperature to 900 ◦C, are displayed in Figure 3. Directional solidification (DS) showed an improvement in two ways as compared to conventional casting (CC). On one hand, DS-samples scattered much less, as standard deviation of particularly ultimate tensile strength (UTS) and also strain to failure (εf) were reduced strongly for each temperature level, see also Table 2. While the deviation is 50 MPa and 8% respectively in CC-state at the most, the maximum deviation in the DS-state is only 16 MPa and 5%. Furthermore, curve progressions for DS-samples seem to be much smoother and more like expected, as the strain to failure increases and the ultimate tensile strength decreases from 600 ◦C to 900 ◦C. All curves at room temperature exhibit high ductility with strain to failure levels about 25% for DS-samples and up to 50% for CC-samples.

**Figure 3.** Stress–strain curves at room temperature up to 900 ◦C for (**a**) conventionally cast samples and (**b**) directionally solidified samples of the base alloy Al10Co25Cr8Fe15Ni36Ti6 in the annealed state (900 ◦C/50 h).


**Table 2.** Ultimate tensile strength (UTS) and strain to failure (εf) at different temperature levels for the conventionally cast (CC) state and the directionally solidified (DS) state.

On the other hand, the most important improvement was the increase of UTS in case of the DS-samples by a factor of 1.5–1.8 in the range of room temperature up to 700 ◦C. Surprisingly, at higher temperatures of 800 and 900 ◦C the improvement by DS processing was only small (a factor of 1.1 and 1.0 respectively for 800 and 900 ◦C).

#### *3.3. Influence of Refractory Elements on Microstructural Characteristics*

The original heat treatment of the base alloy has been investigated and adapted for the alloys containing small amounts of hafnium and molybdenum: The homogenization treatment for the base alloy (1220 ◦C/20 h) was supposed to work for Al9Co25Cr8Fe15Ni36Ti6Mo1 and the Hf-containing alloy as well. In the case of the Mo-alloy the Heusler type phase was completely dissolved after this treatment, while the Hf-alloy exhibited eutectic formations at the grain boundaries and unsolved Heusler type phase in spherical form, attached to the grain boundaries between eutectic regions. In this case the heat treatment temperature needed to be adapted and finally no eutectic formation occurred at 1140 ◦C, but the Heusler type phase still remained unsolved in the homogenized state. These two phenomena are shown in Figure 4.

**Figure 4.** Microstructure of the Hf-containing alloy Al9.5Co25Cr8Fe15Ni36Ti6Hf0.5 after homogenization at (**a**) 1140 ◦C/20 h and (**b**) 1220 ◦C/20 h.

The annealing step 900 ◦C/50 h was taken from the initial heat treatment, where γ´-particles precipitate in the fcc-matrix. The γ -morphology of all alloys after standard treatment is shown in Figure 5 with clear changes: While the base alloy showed cuboidal particles with rounded corners, the Mo-alloy exhibited spherical, and the Hf-alloy showed cubic, sharp-cornered precipitates. These geometries were due to different values of misfit between the γ -phase and the matrix. It was a result of the differences in lattice parameters of both phases [23]. Experiments for quantitative determination of lattice parameters and therefore misfit values have been carried out with synchrotron radiation at photon source BESSY II in Berlin, Germany over a wide temperature range and are in the process of being evaluated.

**Figure 5.** γ -microstructure of the base alloy (**a**) and with the addition of molybdenum (**b**) and hafnium (**c**) after annealing at 900 ◦C/50 h.

Nickel-based superalloys with extraordinary strength- and creep-properties at high temperatures exhibited more cubic precipitates, a higher γ -volume content (60–70%) and larger γ -particles (up to 500 nm) [24] than the base alloy Al10Co25Cr8Fe15Ni36Ti6 (Vγ = 40%, dγ = 400 nm) after heat treatment at 900 ◦C/50 h. Therefore, several studies with variation of annealing time and temperature (±50 K; +500 h) have been conducted, resulting in an enhancement of both size and volume fraction after an annealing treatment at 950 ◦C for 100 h for the base alloy, already described in [25]. SEM-images of both conditions for the base alloy, as well as for Al9.5Co25Cr8Fe15Ni36Ti6Hf0.5 and Al9Co25Cr8Fe15Ni36Ti6Mo1 are shown in Figure 6, where larger cubic particles in (b), (e) and larger spherical particles in (h) can be detected compared to their original appearance in (a), (d) and (g) after shorter annealing at a lower temperature. Figure 6f shows the spherical Heusler type phase accumulation in the case of the Hf-containing alloy, while the needle-shaped Heusler type phase in the case of the base alloy and the Mo-containing alloy is represented in (c) and (i).

**Figure 6.** SEM-images showing the γ - and Heusler type morphology for three different alloys and two different annealing treatments 900 ◦C/50 h (**a**,**d**,**g**) and 950 ◦C/100 h (**b**,**c**,**e**,**f**,**h**,**i**).

Since volume fraction and precipitate size play an important role for the mechanical behavior, all samples were investigated carefully after both annealing steps using SEM-images. The resulting volume fractions of the γ - and Heusler type phase, as well as the size of γ -particles are listed in Table 3. The size corresponds to the diameter in the case of round particles and to the edge-length in case of cuboidal particles.

**Table 3.** Volume fractions of γ - and Heusler type phase, size and shape of γ -precipitates after the annealing treatments 900 ◦C/50 h and 950 ◦C/100 h.


An increase of γ -size and γ -volume fraction, as well as an increase of Heusler type volume fraction is confirmed by Table 3 in the case of the base alloy for the annealing step 950 ◦C/100 h. While the size of precipitates increased with longer treatment at higher temperatures for the Moand Hf-containing alloys, too, the volume fractions of the Heusler type phase and γ -phase were not enhanced, but decreased about 7–8% in the case of the γ -phase and remained constant in the case of the Heusler type phase.

#### *3.4. High-Temperature Tensile Tests of the Alloys after Annealing at 950* ◦*C for 100 h*

The presented stress–strain diagrams of the base alloy after standard heat-treatment (900 ◦C/50 h) in Figure 3 are completed by Figure 7, where samples were annealed at 950 ◦C for 100 h, resulting in larger γ -precipitates. Samples tested at room temperature showed a brittle fracture behavior, as well as specimens tested at 600 and 700 ◦C. In general, large scattering occurred respective to strain to failure and curve progression. Ultimate tensile strength scattered less within one temperature series.

**Figure 7.** Stress–strain curves for the base alloy Al10Co25Cr8Fe15Ni36Ti6 after annealing at 950 ◦C for 100 h.

Figure 8 shows the stress–strain curves of Al9Co25Cr8Fe15Ni36Ti6Mo1 after the annealing step 950 ◦C/100 h in detail. Scattering of the curve progressions was very high and particularly strain to failure values show large scatter at all temperatures tested. At room temperature and at 600 ◦C the alloy exhibited a ductile behavior, while the ultimate tensile strength at higher temperatures was reached very quickly after short elongations. In these cases, the first cracks appeared rapidly and fracture propagation was fast.

**Figure 8.** Stress–strain curves for the Mo-containing alloy Al9Co25Cr8Fe15Ni36Ti6Mo1 after annealing at 950 ◦C for 100 h.

Stress-strain curves of the hafnium containing alloy are displayed in Figure 9. First of all, various tests at the same temperature were very reproducible with only a little scatter. Tests at room temperature reach the highest ultimate tensile strength with a strain to failure of about 20%. Samples deformed at 600 ◦C exhibited only half the strain to failure. In the range of 600–900 ◦C the samples showed the expected evolution of strength and plasticity, as ultimate tensile strength decreased and strain to failure increased with increasing temperature.

**Figure 9.** Stress–strain curves for the Hf-containing alloy Al9.5Co25Cr8Fe15Ni36Ti6Hf0.5 after annealing at 950 ◦C for 100 h.

#### *3.5. Discussion and Comparison of Mechanical Properties*

In this work, compositionally complex alloys were improved with the goal to increase mechanical properties in the temperature range around 700–800 ◦C. Therefore, discussion will mainly depend on this temperature region.

Figure 10 shows the comparison between different variants of the base alloy respective ultimate tensile strength and strain to failure. The progress of this values from low to high temperatures is indicated by arrows and the significant temperature ranges are marked by differently colored areas. While the polycrystalline samples showed little strength and ductility in the marked red area, both parameters increased using directional solidification (green and blue areas). The two colored areas (green, blue) for directional solidified alloys, with different annealing treatments in contrast, revealed a clear difference: Annealing at 900 ◦C/50 h did in fact lead to smaller precipitates, but these samples exhibited a higher ultimate tensile strength, higher strain to failure and less scatter. The reason for this desirable behavior was due to the volume content of Heusler type phase that was drastically reduced from 9% to 3% after annealing at 900 ◦C/50 h, see Table 3.

**Figure 10.** Evolution of ultimate tensile strength and strain to failure from room temperature (RT) to 900 ◦C for different manufactured and annealed types of the base alloy Al10Co25Cr8Fe15Ni36Ti6.

An overview about all tested tensile samples is shown in Figure 11, including the different types of the base alloy (blue), the Hf- and Mo-containing alloys in red and green respectively, as well as two conventionally used nickel-based alloys that are used in the temperature-range of 680–820◦C, highlighted by the vertical yellow stripe.

**Figure 11.** Ultimate tensile strength for all investigated alloys over the temperature range from room temperature to 900 ◦C. Inconel 617 and Alloy 800 H in a polycrystalline state taken from references [26,27].

Figure 12 shows more detailed views on two types of comparison and allows viewing of the effective aspects of mechanical behavior. Directionally solidified samples were produced to neglect the huge factor of grain-structure and grain-size in the mechanical behavior and to investigate the pure microstructure influence independently.

**Figure 12.** Ultimate tensile strength over the temperature range from room temperature to 900 ◦C for all directionally solidified alloys (**a**) and the conventionally cast, polycrystalline base alloy, compared with the two commercial nickel-based alloys Inconel 617 [26] and Alloy 800 H [27] (**b**).

Thus Figure 12a compares all DS-samples and the base alloy even after two different annealing treatments, leading to the following remarks: Differences in UTS of all tested specimen were getting smaller in the higher temperature range 800–900 ◦C. Between 700 and 800 ◦C, and also at 600 ◦C, there was a clear order: The worst behavior was exhibited by the Mo-containing alloy, the base alloy had a remarkably higher strength and the best mechanical properties could be observed in the case of the Hf-containing alloy. Next to the alloys, annealed under the same conditions, the base alloy after a treatment at 900 ◦C/50 h was observed in the range of the Hf-containing alloy. Since the base and the Mo-containing alloy (950 ◦C/50 h) showed the same characteristics of the Heusler type phase, round γ -particles contributed in a bad way to the mechanical behavior. As the two strongest materials, the base alloy (900 ◦C/50 h) and the Hf-containing alloy (950 ◦C/100 h) exhibited almost identical stress–strain diagrams concerning strength, ductility and reproducibility, the small difference in γ -morphology (sharp corners at the Hf-containing alloy), as well as the shape of Heusler type phase were not a reason for the worse fracture behavior. Consequently, good high-temperature tensile properties occurred in the case of cubic precipitates, assumed that the Heusler type phase content was kept quite low. If the Heusler type phase content rose three times higher in the case of the base alloy, the ultimate tensile strength fell down drastically.

For a comparison with commercially used alloys, the polycrystalline base alloy could be used, see Figure 12b: While the base alloy showed UTS exceeding that of Alloy 800 H over the whole temperature range, it was very similar or only a little worse than the UTS of the alloy Inconel 617 at temperatures up to 700 ◦C. In the important temperature range between 700 and 800 ◦C, UTS of the base alloy was equal in the beginning and even exceeded IN 617 at 800 ◦C by a factor of 1.2.

Figure 13a shows the yield strength of all tested alloys over the temperature range from room temperature to 900 ◦C, where the samples showed a remaining plastic deformation of 0.2%. An interesting and application oriented fact is shown in Figure 13b, where the polycrystalline base alloy was compared to the conventionally nickel-based alloys. Similar to the ultimate tensile strength in Figure 12 the yield strength of the base alloy exceeded Alloy 800 H quite significantly. The more competitive alloy IN 617, however, showed similar values of UTS, except at 800 ◦C, but the yield

strength was not able to reach the levels of the base alloy. In the important temperature range the yield strength of the base alloy overran IN 617 by a factor of about 1.7 at 700 ◦C and 1.5 at 800 ◦C.

**Figure 13.** Yield strength over the temperature range from room temperature to 900 ◦C for all alloys (**a**) and the conventionally cast, polycrystalline base alloy, compared with two commercial nickel-based alloys Inconel 617 [26] and Alloy 800 H [27] (**b**).

#### **4. Conclusions and Outlook**

Summarizing the base alloy Al10Co25Cr8Fe15Ni36Ti6 and the influence of Hf (0.5 at.%) and Mo (1.0 at.%) additions on microstructural and mechanical properties:


The top priority for future work is to reduce or maybe even avoid the Heusler type phase and to clarify the question about its role on mechanical behavior. We can state that a spherical shape and a low content of <3% is desirable for good mechanical properties. Thus it is necessary to gain knowledge about the chemical and thermodynamic stability resulting in an only two-phase microstructure.

**Author Contributions:** Conceptualization, U.G., S.H., A.M.M.; Methodology, S.H., A.M.M., F.K.; Software, F.K., A.M.M.; Validation, S.H., A.M.M.; Formal Analysis, S.H., A.M.M.; Investigation, S.H., A.M.M.; Resources, S.H., A.M.M.; Data Curation, S.H., A.M.M.; Writing-Original Draft Preparation, S.H.; Writing-Review & Editing, S.H., A.M.M., U.G.; Visualization, S.H., A.M.M.; Supervision, U.G.

**Funding:** This research work was funded by German Research foundation (DFG) projects GL 181/50-1 and MA 7004/1-1 and supported by the priority program SPP2006 "Compositionally Complex Alloys—High Entropy Alloys (CCA-HEA)".

**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* **The Effect of Scandium Ternary Intergrain Precipitates in Al-Containing High-Entropy Alloys**

#### **Sephira Riva 1, Shahin Mehraban 1, Nicholas P. Lavery 1, Stefan Schwarzmüller 2, Oliver Oeckler 2, Stephen G. R. Brown <sup>1</sup> and Kirill V. Yusenko 1,3,\***


Received: 4 June 2018; Accepted: 19 June 2018; Published: 22 June 2018

**Abstract:** We investigate the effect of alloying with scandium on microstructure, high-temperature phase stability, electron transport, and mechanical properties of the Al2CoCrFeNi, Al0.5CoCrCuFeNi, and AlCoCrCu0.5FeNi high-entropy alloys. Out of the three model alloys, Al2CoCrFeNi adopts a disordered CsCl structure type. Both of the six-component alloys contain a mixture of body-centered cubic (*bcc*) and face centered cubic (*fcc*) phases. The comparison between in situ high-temperature powder diffraction data and ex situ data from heat-treated samples highlights the presence of a reversible *bcc* to *fcc* transition. The precipitation of a MgZn2-type intermetallic phase along grain boundaries following scandium addition affects all systems differently, but especially enhances the properties of Al2CoCrFeNi. It causes grain refinement; hardness and electrical conductivity increases (up to 20% and 14% respectively) and affects the CsCl-type → *fcc* equilibrium by moving the transformation to sensibly higher temperatures. The maximum dimensionless thermoelectric figure of merit (*ZT*) of 0.014 is reached for Al2CoCrFeNi alloyed with 0.3 wt.% Sc at 650 ◦C.

**Keywords:** high-entropy alloys; in situ X-ray diffraction; grain refinement; thermoelectric properties; scandium effect

#### **1. Introduction**

High-Entropy Alloys (HEAs) are defined according to the atomic percentage of their principal elements, between 5 and 35 at %, or according to their configurational entropy at random state, <sup>Δ</sup>S > 1.5 R (R = 8.314 J·K−1·mol<sup>−</sup>1). Their high compositional complexity draws a hyper-dimensional space whose limits have yet to be explored, for its investigation has mostly focused on equiatomic or near-equiatomic compositions [1,2]. In fact, the study of HEAs is driven by the search of new single-phase systems, whose formation is opposed by the diverse mixing enthalpies, atomic size, and valence electron concentration of the constituent elements. Consequently, single-phase HEAs are rare, whereas systems consisting of multiple solid solutions and ordered intermetallic phases are more common [3].

Following the pioneering publications in 2004 by Yeh [4] and Cantor [5], HEAs have attracted growing research interest due to their outstanding mechanical and thermal properties; including high compression yield, fracture strength, ductility and toughness, as well as extreme corrosion, wear and fatigue resistance. HEAs have thus been proposed as candidates for applications in which high temperature stability is pivotal; e.g., as replacements for conventional binders after liquid-phase or spark-plasma sintering, for liquefied gas storage and for high-temperature thermoelectrics [1,6]. Thermoelectric properties can be tuned with respect to valence electron concentration (VEC) and a dimensionless thermoelectric figure of merit (*ZT*) of 0.012 at 505 ◦C was reached for Al2CoCrFeNi. The phase evolution of HEAs has been extensively studied both ex situ and in situ [7–11].

Following a traditional trend in alloy development, HEAs have seen the addition of selected secondary phases to further tune their mechanical properties. This approach has led to the development of a new class of metal-matrix composites, containing oxides [12], silicon carbide [13] or nanodiamonds [14]. Alloying with elements in low concentrations, on the other hand, has resulted in the precipitation of intermetallic compounds in the matrix phase. However, while intermetallics deeply affect yield strength, hardness, tensile properties, and matrix stabilization—due to the competition between mixing enthalpy of atom pairs and mixing entropy—their proper distribution, size, shape, and volume fraction represent a cause for concern [1]. The synthesis of precipitation-hardened HEAs following the introduction of intermetallics has proven largely unsuccessful when binary compounds are concerned, and no studies have been performed on the formation of stable ternary inclusions as pinning centers in complex multi-principal component alloys [15–17]. This is mostly due to the difficulties in choosing appropriate alloying elements. Their selection should be guided by the following considerations: miscibility for most or all HEA constitutive elements in liquid state, low formation enthalpy for ternary compounds and crystallization in common structure types (e.g., σ-phase, Laves phase).

We recently highlighted the compound-forming ability of scandium and its outstanding effect on the mechanical properties of multicomponent alloys [18]. Scandium forms over three hundred binary and ternary phases with most elements of the periodic table, many of which crystallize in highly symmetrical structures (e.g., space groups *Pm*3*m*, *Fm*3*m*, *P*63*/mmc*). Scandium-based intermetallics (i.e., Al3Sc, *V*- and *W*-phases) are responsible for the enhanced properties of several commercial aluminum alloys and newly developed multicomponent systems [19]. These features make scandium a perfect candidate to achieve precipitation-hardened HEAs. Moreover, its low density makes it ideal in combination with HEA based on 3*d*-elements, which make up to 85% of the known systems [1]. The AlxCoCrCuyFeNi HEA can be considered a model alloy. The nature of the solid solution can be tuned by changing the aluminum and copper content, since Al and Cu act as *bcc*- and *fcc*-stabilizers respectively [20]. Thus, Al2CoCrFeNi and AlCoCrCu0.5FeNi have been widely reported as pure *bcc* phases, while Al0.5CoCrCuFeNi as purely *fcc*-structured [21–23].

We herein report on the effects of 0.3–5 wt.% scandium addition to the microstructure, mechanical and transport properties, thermal stability, and phase evolution upon temperature of the model Al2CoCrFeNi, AlCoCrCu0.5FeNi, and Al0.5CoCrCuFeNi HEAs. We show that scandium forms the same stable ternary intermetallic in all three alloys, but that the compound interacts differently with *bcc*- and *fcc*-structured alloys.

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

The target alloys were prepared using induction melting from pure metallic powders, in a BN crucible in an Ar-filled Customised DAB01 glove-box (Saffron Scientific Equipment Limited, Knaresborough, UK). Complete melting of the samples was achieved above 1300 ◦C. After 5 min at the melting temperature, the sample was cooled down naturally to room temperature. The samples were re-melted three times to assure homogeneity. A brief rationale of the synthesized specimens and their performed analysis is reported in Supplementary Materials Table S1.

Differential scanning calorimetry (DSC) measurements were performed on small pieces of sintered samples (50 mg) placed in an Al2O3 crucible and heated in a Netzsch STA 449 F1 Jupiter. Heating (Selb, Germany) and cooling were performed in flowing Ar gas with a temperature ramp of 10 K·min−<sup>1</sup> from 35 to 1300 ◦C.

The transition temperatures highlighted by DSC were used to decide annealing conditions. Samples were heat-treated above their first reversible or irreversible transition temperature with

the following specifics: Al2CoCrFeNi, Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi at 850 ◦C, 12 h; Al2CoCrFeNi + 3 wt.% Sc at 900 ◦C, 12 h; Al0.5CoCrCuFeNi + 3 wt.% Sc and AlCoCrCu0.5FeNi + 3 wt.% Sc at 930 ◦C for 6 h. During annealing, pellets of each sample were sealed in a silica tube under vacuum (<10−<sup>7</sup> Pa) and heated in a furnace. After annealing, the tubes were quenched in ice-cold water.

For microstructure and elemental analysis, all samples were mounted in carbonized resin, ground and polished using MetaDiTM Supreme Polycrystalline Diamond Suspension (1 μm) (Coventry, UK).

The morphology and elemental compositions were analyzed using a Hitachi S-4800 Field Emission scanning-electron microscope (SEM) equipped with energy dispersive X-ray (EDX) analyzer (Tokio, Japan). The average elemental composition was obtained from 2.5 mm maps (Table S2).

The Vickers hardness was measured on a WilsonR VH3100 Automatic Knoop/Vickers Hardness tester (Buehler, Lake Bluff, IL, USA); 25 individual points under a 9.81 N (1 kg) testing load were measured to get statistically significant results.

The density was measured according to Archimedes' principle in water, in the ATTENSION equipment (Biolin Scientific, Stockholm, Sweden). Six measurements were taken for each sample to obtain statistically relevant results

The small punch tests were performed on discs of diameter 12.5 mm and thickness circa 0.8 mm. Measurements were performed with a properly modified Tinius Olsen H25KS Benchtop Tester (Salfords, UK). The setup included a lower die (diameter of 8 mm) and a punch (4 mm diameter). Each experiment was reproduced twice for statistical significance; however, since each specimen had a slightly different thickness, the final results were normalized for the standard 0.5 mm thickness as described in [24].

For powder X-ray diffraction (PXRD), samples were powdered using a Fritsch mini-mill Pulverisette 23 (Idar-Oberschtein, Germany) (steel vial and ball, 10 min at 50 rpm). PXRD data for the annealed powdered samples were collected at ID06B-LVP beam-line at the European Synchrotron Research Facility, ESRF (room temperature, λ = 0.22542 Å) using position sensitive detector. LaB6 (NIST SRM 660c) was used as external standard for calibration.

In situ high-temperature PXRD patterns were collected at the I-11 beam-line at the DIAMOND light source (λ = 0.494984 Å). LaB6 (NIST SRM 660c) was used as external standard for wavelength and sample to detector distance calibration. A wide-angle Mythen-2 Si position sensitive detector. The detector was moved at constant angular speed with 10 s scan time at each temperature and 60 s waiting time to let the temperature stabilize. The powdered alloys were sealed in a 0.5 mm silica glass capillary in vacuum and heated in the capillary furnace from 25 to 1200 ◦C with axial rotation [25]. In all samples, oxidation was detected above 1000 ◦C, which can be due to the reaction of metallic alloy with silica at high temperature, resulting in capillary destruction. In situ low-temperature PXRD profiles were collected at the P02.1 beam-line at the PETRA III synchrotron (λ = 0.207150 Å). LaB6 (NIST SRM 660c) was used as external standard for calibration. A wide-angle position sensitive detector based on Mythen-2 Si strip modules was used. The detector was moved at constant angular speed with 10 s scan time at each temperature and 60 s waiting time to let the temperature stabilize. The powdered alloys were sealed in 0.5 mm silica glass capillaries in vacuum and cooled in nitrogen flow from 300 to 100 K. Temperature was directly measured with a thermocouple during the experiment; error arising from the distance between sensor and capillary was estimated to be below 5%.

Temperature dependent PXRD patterns were analyzed using Powder3D software [26]. Phase composition has been verified using the Powder Diffraction File database [27]. Parametric sequential refinements were performed using the TOPAS 5.0 software [28]. Profile parameters for the Lorentzian function, cell parameters, and phase fractions were refined simultaneously for all phases.

The Seebeck coefficient S (μV·K−1) and electrical conductivity <sup>σ</sup> (kS·cm−1) were measured simultaneously under He atmosphere with a Linseis – Seebeck and Electric Resistivity Unit (LSR-3 1100, Linseis, Selb, Germany) four-point setup with PtRh/Pt and Pt contacts and a continuous reverse of the polarity of the thermocouples (bipolar setup, measurement current: 100 mA) using cuboid samples

(ca. 8 × <sup>2</sup> × 3 mm). Three heating cycles up to 875 ◦C (10 K·min−1, 3 data points per temperature) were performed. Thermal diffusivity was measured up to 875 ◦C (heating/cooling rate 10 K·min−1) under He atmosphere with a Linseis LFA1000 (Selb, Germany) apparatus. Simultaneous heat loss and finite pulse corrections were applied using Dusza's model [29]. Values were averaged from five measurement points at each temperature. For calculation of thermal conductivity κ, they were multiplied with the Dulong−Petit heat capacity Cp and the density as derived by the weight and the volume determined by Archimedes' principle. The single values of each sample are given in Table S3. According to experimental Cp values of materials with similar compositions (e.g., 0.60 J·g−1·K−<sup>1</sup> for AlCoCrFeNi at 25 ◦C) [30], the room temperature heat capacity of these materials is about 20% higher than the Dulong–Petit value of 0.49 J·g−1·K−1; this probably adds this uncertainty to the values of <sup>κ</sup> and thus *ZT*.

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

Al2CoCrFeNi, Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi were synthesized via induction melting in the atomic compositions reported in the Table S1. While all three systems are widely reported as single-phase [1], refinements of PXRD data obtained with synchrotron radiation highlight the presence of a secondary *fcc* phase in the *bcc*-structured AlCoCrCu0.5FeNi alloy (Figure 1c) and of a very small secondary *bcc* phase in the mainly *fcc*-structured Al0.5CoCrCuFeNi alloy (Figure 1b). Unlike Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi, Al2CoCrFeNi appears to be a solid solution forming a disordered CsCl structure-type. (Figure 1a). This is consistent with a previously reported investigation on CoCrCuFeNi-based systems—which displayed a phase separation due to the positive mixing enthalpy between copper and other elements [31]—and with the Hume–Rotary classification maps in ref. [32].

**Figure 1.** Powder X-ray diffraction (PXRD) Rietveld refinements performed from the DIAMOND light source (I11, λ = 0.494984 Å) of (**a**) Al2CoCrFeNi. The section between 12 and 16 degrees 2θ is enlarged to show the symmetric shape of the first reflection. (**b**) Al0.5CoCrCuFeNi. The section between 12 and 15 degrees 2θ is enlarged to show the asymmetry of the first reflection, which can only be fitted by taking into account a small amount of *bcc* phase. (**c**) AlCoCrCu0.5FeNi. The section between 13 and 15 degrees 2θ is enlarged to show the presence of an *fcc* phase.

The presence of the (100) diffraction line in the *bcc*-structured PXRD pattern indicates an ordered superstructure, generally attributed to Al and Ni ordering [22,33,34]. Lattice parameters of the as-cast alloys are the following: *aB*<sup>2</sup> = 2.877(2) Å for Al2CoCrFeNi; *afcc* = 3.601(1) Å for Al0.5CoCrCuFeNi and *abcc* = 2.891(3) Å for AlCoCrCu0.5FeNi. All values are consistent with the literature within the experimental error [20,22,35–37]. Small differences between the reported results arise from the high sensitivity of the HEA to synthetic pathway and to minor compositional variations. Therefore, nominally equivalent starting materials can in turn display profoundly different crystal structures, microstructures, and phase transitions.

The microstructures and elemental distributions of as-cast Al2CoCrFeNi, Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi are reported by means of EDX element mapping in Figures S1–S3, respectively. Two phases are present in the as-cast AlCoCrCu0.5FeNi alloy, whose microstructure is characterized by a matrix and a circular secondary phase of darker color. Elemental distributions appear completely homogeneous (Figure S3). Annealing results in the growth of a darker secondary phase and the slight segregation of Al from the rest of the elements. The as-cast Al0.5CoCrCuFeNi HEA consists of dendritic-like and interdendritic-like regions, the first being richer in Co, Cr and Fe; and the latter Cu-rich (Figure S2). Both the dendrite-like and interdendrite-like matrix have been previously linked to a simple *fcc* phase, and copper segregation has been explained through its high mixing enthalpy with cobalt, chromium, iron, and nickel [20]. The microstructure changes drastically after annealing, as the two phases cannot be easily differentiated. Nevertheless, element segregation persists, with Cu and Al separating from Co, Cr, Fe, and Ni. The as cast Al2CoCrFeNi microstructure is dominated by large (∼100 μm) unstructured grains which, unlike previous studies, show no trace of non-equiaxed dendrites [21,37]. With respect to other elements, chromium segregation is clearly visible, though it can be reduced by annealing (see Figure S1). On the other hand, annealing causes the formation of a homogeneously dispersed secondary phase, which appears as black dots.

The microstructure of all alloys is strongly affected by even a 3 wt.% addition of scandium. In the case of Al2CoCrFeNi (Figure 2a, Figure S4), the large grains of homogeneous compositions are refined and scandium segregates in the inter-grain volume. The scandium-based intermetallic forming in the inter-granular region appears to contain all elements in the same relative fractions as the main phase—with the notable exception of chromium (Table S1). In Al0.5CoCrCuFeN (Figure 2b), scandium is dispersed more homogeneously and aids the formation of a globular microstructure. Nevertheless, most of it segregates in the inter-granular region, with copper- and nickel-rich areas (Figure S5). Lastly, following scandium addition the original columnar cellular microstructure of the AlCoCrCu0.5FeNi HEA turns into equiaxed non-dendritic-like grains (Figure 2c). The chemical inhomogeneity due to scandium segregation appears surprising, considering the strongly negative mixing enthalpy of the metal with iron, nickel, cobalt and aluminum [38–40]. The driving force of the segregation is thus the formation of the secondary phase.

**Figure 2.** Overview of the effect of scandium addition on the microstructure of the three systems (scanning electron microscope-back scatter detector (SEM-BSE) images). **Left**: Microstructure of the as-cast (**a**) Al2CoCrFeNi, (**b**) Al0.5CoCrCuFeNi and (**c**) AlCoCrCu0.5FeNi High-Entropy Alloys (HEAs) before and after 3 wt.% scandium addition. **Right**: scandium distribution of the areas displayed in (**b**) according to energy dispersive X-ray (EDX) maps.

The secondary phase is the same in all systems and can be indexed as a ternary intermetallic analogous of MgZn2-type (Figure 3). Compounds of this structure type have been reported for scandium with several metals, in compositions such as AlCuSc, AlCoSc, Al1.06Cr0.94Sc, AlFeSc, and AlNiSc [18]. As confirmed by elemental composition maps (Figures S4–S6), the ternary phase is a highly disordered structure containing all five elements and scandium.

**Figure 3. Left**. PXRD profile (DIAMOND = 0.22542 Å) of the as-cast disordered CsCl-structured Al2CoCrFeNi alloy before (blue) and after (orange) a 3 wt.% Sc addition. The PXRD profile of the scandium phase is indexed with black hexagons. **Right**. Crystal structure of the hexagonal MgZn2-type intermetallic along the c-axis, depicted as its AlCuSc analogue (Al in green, Cu in orange and Sc in blue; a, b and c are cell axis).

Knowledge about scandium-containing ternary compounds is still fragmentary. Out of all the cited MgZn2-type intermetallic phases, only AlCuSc has been thoroughly investigated, due to its effect on the mechanical properties of Al-based alloys, as part of the so-called *W*-phase. In particular, it was shown that the microhardness of the *W*-phase is much higher than the one of the Al3Sc phase (5150–5170 MPa against 3900–4300 MPa) [18].

The formation of a very hard phase in the HEA matrix affects its mechanical properties. As shown in Figure 4, single phase *bcc* alloys are harder than duplex-structured and *fcc* alloys. This is hardly surprising, considering the stronger interatomic forces involved in the *bcc* vs. *fcc* packing of alloys. On the other hand, increasing scandium content in Al0.5CoCrCuFeNi does not affect hardness. Indeed, it is even detrimental to the hardness of the AlCoCrCu0.5FeNi alloy and is not accompanied by an increase in ductility (as shown by the disk punch tests presented in Figure S7). Only in the originally hard Al2CoCrFeNi HEA the formation of the intermetallic results in an increase in hardness. The addition of 0.5 wt.% Sc causes a 20% hardness enhancement, as well as visible grain refinement. Disc punch tests performed on the HEA with 0, 0.5 and 2 wt.% scandium additions show a decisive increment in brittleness, proportional to the concentration of scandium (Figure S8).

Differential scanning calorimetry (DSC) was performed on alloys from room temperature to <sup>1300</sup> ◦C with a 10 K·min−<sup>1</sup> heating rate. The second cycle of heating and cooling, which is less influenced by effects of the synthetic route on the specimens—i.e., internal stress-strain, magnetic ordering—is reported in Figure 5.

**Figure 4.** Vickers hardness values for Al2CoCrFeNi (red), Al0.5CoCrCuFeNi (green) and AlCoCrCu0.5FeNi (blue) HEAs with 0, 0.5, 2 and 3 wt.% Sc additions. Values are an average of 25 indentations at 1 HV. SEM images of the microstructures of all alloys are shown in circles of 30 μm diameter.

**Figure 5.** Differential scanning calorimetry (DSC) of (**a**) Al2CoCrFeNi, (**b**) Al0.5CoCrCuFeNi and (**c**) AlCoCrCu0.5FeNi with (dotted line) and without (solid line) 3 wt.% scandium. The second heating/cooling cycle is reported for each specimen, in red and blue respectively.

The DSC profile of Al2CoCrFeNi (Figure 5a) has a sigmoid-like deviation between 600 and 700 ◦C. In a theoretical work, Gao associates this feature with the transition *bcc*<sup>1</sup> *+ bcc*<sup>2</sup> *+* CsCl → *bcc+*CsCl, but this is incompatible with the crystal structure of the Al2CoCrFeNi alloy as per Figure 1a. The divergence between the results reported here and Gao's interpretation might arise from the profound differences in the crystal structures of the nominally equivalent starting material [41]. To unequivocally identify the cause of the transition, a more detailed investigation of the effect of temperature on phase stability is needed. The corresponding Sc-containing sample displays a sharp reversible peak at 1150 ◦C, probably corresponding to the melting and crystallization of the intermetallic phase, preceded by irreversible peaks. These two endothermic peaks, located at 906 ◦C and 966 ◦C, might correspond to phase transitions occurring in the scandium-phase. The sigmoid-like deviation clearly visible in the pristine alloy is hardly distinguishable from the background line in the Sc-containing specimen but is located at higher temperature (between 700 and 800 ◦C).

Al0.5CoCrCuFeNi shows a slight reversible peak centered at 760 ◦C, as well as two reversible peaks above 1150 ◦C (Figure 5b). In previously reported DSC curves, the two endothermic peaks of Al0.5CoCrCuFeNi at 1140 and 1270 ◦C have been linked to the melting of interdendritic-like and dendritic-like material, respectively. The results of Jones et al. highlight a third reversible peak, appearing at 850 ◦C and related to the dissolution and recrystallization of the L12 phase, which might form during prolonged heat treatment below 850 ◦C and is dependent from the sample cooling rate [42,43]. Its absence is indicative of the purity of the as-cast *fcc* sample, which is confirmed by high-resolution PXRD data (Figure 1b). The scandium containing specimen displays three irreversible signals upon heating (the first two, endothermic, at 921 and 1000 ◦C; the latter, exothermic, at 1073 ◦C). Upon cooling, the Sc-containing specimen behaves very similarly to its corresponding pristine alloy. The irreversible transitions occurring in the sample might thus relate solely to the intermetallic, and have little impact on the matrix.

Finally, AlCoCrCu0.5FeNi displays a reversible transition at circa 624 ◦C, which is maintained in the Sc-containing sample (Figure 5c). Irreversible phenomena occur in the pristine alloy above 1150 ◦C, as in the previous sample. The scandium-containing alloy presents two irreversible endothermic peaks (at 906 and 966 ◦C) upon heating and two exothermic peaks upon cooling (at 1102 and 1206 ◦C), which are too far from the heating ones to be considered part of the same phenomenon.

Reversible phenomena thus occur in all systems above 600 ◦C; whereas irreversible peaks appear in the scandium-containing alloys at ca. 900 ◦C. To investigate the nature of these transitions, the pristine alloys were annealed above their average first reversible transition temperature, and the scandium-containing specimens at the temperatures of their first irreversible transition. Annealing time was shortened for the samples at the highest temperature (930 ◦C) in order not to lose aluminum. Therefore, Al2CoCrFeNi, Al0.5CoCrCuFeNi, and AlCoCrCu0.5FeNi were annealed at 850 ◦C for 12 h. Al2CoCrFeNi + 3 wt.% Sc was treated at 900 ◦C for 12 h, while Al0.5CoCrCuFeNi + 3 wt.% Sc and AlCoCrCu0.5FeNi + 3 wt.% Sc at 930 ◦C for 6 h. The corresponding element distribution and microstructures are presented in the following pages. Figure 6 reports the microstructure and element distribution of the annealed Al2CoCrFeNi + 3 wt.% Sc alloy. With respect to Figure S1b (the annealed pristine alloy), microstructure is refined and the intermetallic scandium phase grows in a dendritic-like structure. The secondary phase which that have appeared in the pristine alloy (in the form of black dots) is not displayed by the matrix. A comparison with the scandium-containing alloy prior to annealing (Figure S4) shows a more homogeneous distribution of chromium, even though the metal still visibly segregates along the grain boundaries.

**Figure 6.** SEM images of the microstructure of annealed Al2CoCrFeNi (**a**), Al0.5CoCrCuFeNi (**b**) and AlCoCrCu0.5FeNi (**c**) before and after 3 wt.% Sc addition. Blank alloys were annealed at 850 ◦C for 12h; Al2CoCrFeNi + 3 wt.% Sc at 900 ◦C, 12 h; Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi at 930 ◦C, 6h. EDX maps show the scandium element distribution of the areas depicted on the right side (green).

The heat treatment of the pristine Al0.5CoCrCuFeNi (Figure 6b) causes the disruption of the original dendritic- and interdentritic-like microstructure. However, the microstructure of the scandiumcontaining alloy after annealing is quite similar to the as-cast pristine alloy: a columnar cellular structure displaying strong element segregation (Figure S2b). The light phase consists mostly of Cu and Sc, and within it regions rich in Al and Ni. The darker areas are rich in Co, Cr, and Fe. It is of interest to note that this element segregation is different from the one prior annealing (Figure S5), as Ni is depleted from the matrix.

The annealed AlCoCrCu0.5FeNi + 3 wt.% Sc has a complex microstructure (Figure 6c). Acicular particles grow in an unstructured matrix, the bigger of those being easily removed from the specimen during polishing (and leaving behind elongated valleys). Elemental segregation occurs to form several diverse regions: A Co-rich matrix, areas rich in Al and Ni, and Cu-Sc precipitates. Iron clearly segregates from Al and Ni, but does not follow the trend of other elements. Finally, Cr coalesces in circular well-shaped areas. Combining this information, we can note that the brittle phase is, as expected, rich in Cu and Sc. The annealed scandium-containing alloy is very different not only from the pristine alloy, but also from the original as cast specimen. Indeed, the annealed pristine alloy shows no clustering of Cr, Fe, and Cu (Figure S3); and the scandium-containing as cast alloy clearly highlights the coexistence of Co, Cr, and Fe (Figure S6).

To further investigate the nature of the transitions occurring in our HEA systems and how they are affected by the intermetallic phases, we performed in situ PXRD of disordered CsCl-type Al2CoCrFeNi and the *fcc*-structured Al0.5CoCrFeNi with and without 3 wt.% scandium. Quantitative data are available only for the main phases, whose reflections had enough intensity to be detected and refined. The development of scandium-rich phases could not be followed directly.

Several works have highlighted that Al0.5CoCrCuFeNi consists of a mixture of two *fcc* phases of similar cell parameters, corresponding to dendritic and interdendritic phase [44–46]. Chen reported a single *fcc* structure below 500 ◦C, and a duplex *fcc/bcc* structure above 600 ◦C [47]. According to Jones, on the other hand, two *fcc* phases of similar cell parameters co-exist and are stable up until their melting temperatures (1150 ◦C for the Cu based solid solution and 1350 ◦C for the multi-component phase) [44]. As shown in Figure 1b, our system consists of an *fcc* phase (*afcc* = 3.601(1) Å) and a small (<10%) amount of *bcc* phase (*abcc* = 2.867(4)). The second *fcc2* phase (*a f cc*<sup>2</sup> = 3.609(2) Å) reported in Figure 7a can be considered a minor admixture in an exsolution equilibrium with the *bcc* phase. Heating does not affect the primary phase: reflections become sharper, but their relative intensity remains similar. The addition of scandium (Figure 7b) has no impact on the *fcc* matrix, but influences the *bcc* → *fcc*<sup>2</sup> equilibrium by shifting it towards the formation of *bcc*.

**Figure 7. Left**: High-temperature behavior (I11 at DIAMOND, λ = 0.494984 Å) of Al0.5CoCrCuFeNi with (**b**) or without (**a**) scandium according to in situ PXRD data. **Right**: Weight percentage of the major phases in each sample: *fcc* (blue), *bcc* (red), *fcc*<sup>1</sup> (green), others (pink). The yellow areas mark the start of oxidation.

The thermal behavior of the CsCl-type Al2CoCrFeNi sample is displayed in Figure 8a. We observe the decomposition of 40 wt.% of the CsCl-type phase above 620 ◦C and the exsolution of an *fcc* phase (*a f cc*<sup>1</sup> = 3.635(8) Å). At 750 ◦C a second *fcc*<sup>2</sup> phase of lattice parameter *a f cc*<sup>2</sup> = 3.631(9) Å develops from *fcc*<sup>1</sup> and grows almost linearly with temperature. The continuous presence of the (100) diffraction line in the PXRD pattern of the alloy suggests that the Al-rich sub-structure responsible for the CsCl-type crystal structure might not be involved in the exsolution process. Results for the Sc-containing Al2CoCrFeNi alloy are presented in Figure 8b. The exsolution of the *fcc* phase from the CsCl-type one follows a different pathway, with *fcc*<sup>1</sup> and *fcc*<sup>2</sup> forming at the same time. The transformation occurs at higher temperatures: a 3 wt.% scandium addition is enough to stabilize the main CsCl-type phase for about 150 ◦C.

**Figure 8. Left**: High-temperature behavior (I11 at DIAMOND, λ = 0.494984 Å) of CsCl-type Al2CoCrFeNi with (**b**) or without (**a**) scandium according to in situ PXRD data. **Right**: Weight percentage of the major phases in each sample: *fcc* (blue), *bcc* (red), *fcc*<sup>1</sup> (green), others (pink). The yellow areas mark the start of oxidation.

The thermal expansion coefficient of Al2CoCrFeNi and Al2CoCrFeNi + 3 wt.% Sc can be evaluated by in situ high-temperature PXRD measurements by fitting the corresponding dataset to:

$$\mathbf{a}(\mathbf{T}) = \boldsymbol{\alpha}\_{\mathrm{T}\_{\mathrm{o}}} \cdot \exp\left[\boldsymbol{\alpha} \cdot (\mathbf{T} - \mathbf{T}\_{\mathrm{o}}) + \frac{\beta}{2} \cdot \left(\mathbf{T}^{2} - \mathbf{T}\_{\mathrm{o}}^{2}\right)\right] \tag{1}$$

where αTo is the cell parameter α at the reference temperature (To). Figure 9 reports the variation of the lattice parameter αCsCl upon heating and cooling. A satisfactory fitting can be obtained for Al2CoCrFeNi for <sup>α</sup> and <sup>β</sup> equalling 3.6(2) × <sup>10</sup>−6·K−<sup>1</sup> and 1.69(3) × <sup>10</sup>−8·K<sup>−</sup>2, respectively (Figure 9a); and for Al2CoCrFeNi + 3 wt.% Sc for <sup>α</sup> and <sup>β</sup> equaling 4.2(1) × <sup>10</sup>−6·K−<sup>1</sup> and 1.62(2) × <sup>10</sup>−8·K−2, respectively (Figure 9b). As highlighted in Figure 9c, the scandium-containing CsCl-type phase has slightly larger cell parameters than the regular alloy and follows the same trend with temperature. This trend is consistent to the one of pure iron in the investigated temperature range, but differs strongly from the behavior of other cubic metals which constitute the alloy (in particular, from chromium and nickel) [48–51].

**Figure 9.** Thermal expansion data fitted with Equation (1) for (**a**) Al2CoCrFeNi and (**b**) Al2CoCrFeNi + 3 wt.% Sc. In red*:* data collected upon heating from 100–400 K (PETRAIII, λ = 0.207150 Å) and from 300–1100 K (DIAMOND, λ = 0.494984 Å); in blue: data collected upon cooling from 300–100 K (PETRAIII, λ = 0.207150 Å). Low temperature data are highlighted in the inset. (**c**) Thermal expansion curves for Al2CoCrFeNi and Al2CoCrFeNi + 3 wt.% Sc with respect to their constitutive cubic metals [48–51].

The addition of scandium to the Al-containing HEAs results in the precipitation of a ternary intermetallic of MgZn2-type along the grain boundaries. Microstructural data as well as high-temperature PXRD suggest an extraordinary stability of Sc-based precipitates, which form in all alloys containing aluminum and first raw transition metals. The matrix coherence following the formation of the secondary phase is maintained and the Vickers hardness increases up to 20% due to precipitation hardening (for 0.5 wt.% Sc addition to Al2CoCrFeNi). The intermetallic phase has high thermal stability and affects the *bcc* → *fcc* exsolution equilibrium by stabilizing the body-centered cubic phase with respect to the face-centered cubic one. This effect is likely related to the segregation of part of the *fcc*-stabilizing elements (i.e., Ni and Co) in the ternary compound. Out of *bcc*-stabilizers, aluminum segregates in the MgZn2-type phase, but chromium has much lower affinity for the secondary phase and coalesces in the matrix (as highlighted by EDX maps). Lattice thermal expansion data of the CsCl-type alloy with or without scandium show that both systems follow the same trend upon heating. However, the cubic phase in the scandium-containing HEA has slightly larger cell parameters: this is a further indication of the compositional difference between the matrix of the two systems.

While the thermoelectric properties of AlxCoCrFeNi (0 ≤ *x* ≤ 3) [6,30,52] are discussed in detail in literature, the effect of Sc alloying on the transport properties of Al2CoCrFeNi is herein reported for the first time (Figure 10). The absolute Seebeck coefficient of *n*-type materials decreases slightly with increasing Sc content. The electrical conductivity for both compositions increases by about 1000 S·cm−<sup>1</sup> compared to pure Al2CoCrFeNi [6,52]. In accordance with the higher electrical conductivity, the thermal conductivity shows similar, also slightly increased values compared to ref. [6]; whereas the thermal conductivity for pure Al2CrCoFeNi from ref. [52] is extraordinarily high.

Power factor and *ZT* value increase with temperature until a peak *ZT* value of 0.012 at 700 ◦C (for 5 wt.% Sc addition) and 0.014 at 650 ◦C (for 0.3 wt.% Sc) is reached. The subsequent decrease of *ZT* with temperature is mainly attributed to the decrease of the absolute value of *S* due to excitation of minority charge carriers. According to PXRD patterns of the samples after thermoelectric measurements, no degradation was observed (Figure S12). In addition, the samples show a good cyclability within three subsequent heating and cooling cycles (Figures S13 and S14).

**Figure 10.** Thermoelectric properties of Al2CrCoFeNi + 5 wt.% Sc (blue) and Al2CrCoFeNi + 0.3 wt.% Sc (orange) averaged over heating and cooling cycles (without first heating): electrical conductivity σ (**top**, **left**), Seebeck coefficient S (**top**, **right**), thermal conductivity κ (**middle**, **left**), power factor PF (**middle**, **right**) and thermoelectric figure of merit *ZT* (**bottom**, **left**). For comparison, literature data for Al2CrCoFeNi (brown [6] and pink [52] dashed line) were added.

#### **4. Conclusions**

Ever since HEAs have been proposed as promising materials for high-temperature applications, their phase stability, and the effect of selected intermetallic compounds on their performances upon heating have seldom been studied. We report a thorough investigation on the effect of small scandium additions to the microstructure, phase stability, and mechanical and thermoelectric properties of the model HEAs Al2CoCrFeNi, Al0.5CoCrCuFeNi and AlCoCrCu0.5FeNi. High-temperature in situ PXRD investigations highlights the existence of secondary phases in these systems, traditionally though as consisting of a single solid solution. The addition of scandium to the studied HEAs causes the precipitation of a MgZn2-type intermetallic phase. This hexagonal compound leads to chemical inhomogeneity along grain boundaries, causing grain refinement and a powerful increase in hardness (for Al2CoCrFeNi, a 0.5 wt.% addition of scandium enhances hardness by 20%). Upon heating, *bcc* and CsCl-type phases turn into one or more *fcc* phases, whereas *fcc* phases appears very stable. Scandium acts as stabilizer of the *bcc* phase, by affecting the exsolution equilibrium *bcc* → *fcc*. Regarding thermoelectric properties, the alloying of scandium in Al2CoCrFeNi increases the electrical conductivity by 14%. The maximum *ZT* value of 0.014 reached for 0.3 w% Sc addition at 650 ◦C is in the same order of magnitude as for other HEAs, but too low to make these materials promising for current thermoelectric applications. Still, HEAs are a young material class and optimized HEA will have the advantage of extreme thermal stability.

The results reported here open new possibilities in the design and further improvement of the properties of multicomponent alloys. They could be used as the basis for designing new stable and metastable multicomponent single-phase alloys with improved mechanical properties. The effect of minor elements additions to HEA should be thoroughly investigated to tackle specific technological needs.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/1099-4300/20/7/488/s1.

**Author Contributions:** Conceptualization, K.V.Y.; Investigation, K.V.Y., S.R., S.S., S.M.; Resources, N.P.L.; Writing-Original Draft Preparation, S.R.; Writing-Review & Editing, S.R., S.S., O.O., K.V.Y., Supervision, O.O., S.G.R.B.

**Funding:** This research was funded by the National research network of Wales (Ser Cymru, project NRN046) and by the European Space Agency (contract number 4000111643/NL/PA).

**Acknowledgments:** The Authors gratefully acknowledge the financial support provided by the Welsh Government and Higher Education Funding Council for Wales through the Sêr Cymru National Research Network in Advanced Engineering and Materials and by the European Space Agency. The authors thank the Materials Advanced Characterization Centre (MACH1) at Swansea University, the DIAMOND light source (Oxfordshire, UK) and the PETRA III synchrotron (Hamburg, Germany) for providing measurement time and technical support. A PhD scholarship from the "Studientstiftung des deutschen Volkes" for S.S. is gratefully acknowledged.

**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 Cold Rolling on the Phase Transformation Kinetics of an Al0.5CoCrFeNi High-Entropy Alloy**

#### **Jun Wang \*, Haoxue Yang, Tong Guo, Jiaxiang Wang, William Yi Wang and Jinshan Li**

State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi'an 710072, China; hxyang@mail.nwpu.edu.cn (H.Y.); believegt@163.com (T.G.); 2013301034@mail.nwpu.edu.cn (J.W.); wywang@nwpu.edu.cn (W.Y.W.); ljsh@nwpu.edu.cn (J.L.)

**\*** Correspondence: nwpuwj@nwpu.edu.cn; Tel.: +86-29-88460568; Fax: +86-29-88460294

Received: 25 October 2018; Accepted: 29 November 2018; Published: 30 November 2018

**Abstract:** The solid state phase transformation kinetics of as-cast and cold rolling deformed Al0.5CoCrFeNi high-entropy alloys have been investigated by the thermal expansion method. The phase transformed volume fractions are determined from the thermal expansion curve using the lever rule method, and the deformed sample exhibits a much higher transformation rate. Two kinetic parameters, activation energy (*E*) and kinetic exponent (*n*) are determined using Kissinger– Akahira–Sunose (KAS) and Johnson–Mehl–Avrami (JMA) method, respectively. Results show that a pre-deformed sample shows a much lower activation energy and higher kinetic exponent compared with the as-cast sample, which are interpreted based on the deformation induced defects that can promote the nucleation and growth process during phase transformation.

**Keywords:** high-entropy alloy; phase transformation; kinetics; deformation; thermal expansion

#### **1. Introduction**

High-entropy alloys (HEAs) are a new class of alloys designed based on a unique alloy concept that have multi alloy components (normally five or more principal elements) with equal or near-equal atomic composition [1,2]. HEAs are mainly solid solution based alloys, which own very attractive properties, like high strength and hardness, high fracture toughness, excellent corrosion resistance and unique physical properties, making them the potential engineering materials in many industry areas [3–7].

Thermal induced phase transformations are one of the main transitions inside metallic materials which can be used to control the final phases of the material, and thus to tune the properties. In order to know the details of the phase transformation processes, one of the useful methods is to study the phase transformation kinetics, which can give a lot of useful information like phase transformation speed and transition mode [8,9]. However, until now, there have only been very few papers referring to the kinetic analysis of high-entropy alloys.

The as-cast HEAs always exhibit coarse grains especially for big ingots. In order to obtain fine equiaxed grains, methods like cold rolling or forging are incorporated during the treating process [10,11]. Whether this kind of treatment can affect the subsequent phase transformation process is still unclear for HEAs.

Al*x*CoCrFeNi (0 ≤ *x* ≤ 2) HEAs are one of the alloy systems that intrigue the research interest due to many aspects. First, this alloy system owns very good mechanical [12–15] and physical properties [16–18]. Second, with the increasing Al content, the main phase of the alloy moves from pure face-centered-cubic (FCC) to FCC + pure body-centered-cubic (BCC), and pure BCC phase, making the mechanical and physical properties adjustable [16–25]. Thus, the microstructure, phases and properties of this alloy system have been intensively investigated. In this study, a binary phase

(FCC + BCC) Al0.5CoCrFeNi HEA with a balanced strength and plasticity are chosen to investigate the effect of cold rolling on the subsequent solid state phase transformation kinetics.

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

The ingots with a nominal composition of Al0.5CoCrFeNi were prepared by melting pure elements (at least 99.95 wt.%) in a vacuum induction-melting furnace. The furnace is firstly heated to 500 ◦C and held for 2 h to remove the water vapor, and vacuumed to below 10 Pa. Then pure argon is backfilled to expel the rest of air until the vacuum goes back to standard atmospheric pressure. This process is repeated three times in order to gain an oxygen-free environment as much as possible. Afterwards, melting and casting is performed with the protection of high purity argon and the alloy is heated at about 1550 ◦C for 15 min. Approximately 15 kg ingot is produced by casting the melt into a steel crucible with a height of 180 mm, upper inner diameter of 140 mm and bottom inner diameter of 130 mm.

Samples with the size of 8 × <sup>20</sup> × 70 mm3 are taken from the center of the ingot. Then, the sample is cold rolled to a thickness reduction of 20%. Then, samples were machined to *φ* 6 × 25 mm for thermal expansion measurement both from as-cast and cold rolled plates. The thermal expansion curves were tested using a Netzsch® DIL-402C dilatometer (Selb, Germany) under the protection of argon with constant heating rate of 4 K/min, 6 K/min, 8 K/min and 10 K/min.

The microstructure and phases were characterized by scanning electron microscope (SEM, TESCAN MIRA3 XMU (Brno, Czech Republic), and the working distance and the energy of beam used are 15 mm and 20 kV, respectively) and X-ray diffractometer (XRD, DX 2700 (Dandong, China), and the voltage and current used during measurement are 40 kV and 30 mA, respectively), respectively. Uniaxial tensile tests are carried out with a MTS SANS CMT5105 (Shenzhen, China) mechanical tester at the strain rate of 10−<sup>3</sup> s−<sup>1</sup> and the specimens are prepared along the rolling direction (RD) direction.

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

#### *3.1. Microstructure*

The microstructure and XRD patterns of Al0.5CoCrFeNi high-entropy alloys at as-cast and 20% cold rolled (CR) condition are shown in Figures 1 and 2. Except for the grain after cold rolling showing a very small amount elongation along the rolling direction, the as-cast and cold rolled sample show quite similar microstructure (Figure 1). Both the samples are mainly FCC phase together with a small amount of BCC phase (Figure 2), indicating 20% cold rolling has no significant effect on the microstructure and phase transition of the Al0.5CoCrFeNi high-entropy alloy.

**Figure 1.** The microstructure of the Al0.5CoCrFeNi high-entropy alloys at (**a**) as-cast and (**b**) 20% cold rolled condition. The arrow represents the rolling direction.

**Figure 2.** The XRD patterns of Al0.5CoCrFeNi high-entropy alloy at as-cast and 20% cold rolled condition.

#### *3.2. Transformed Volume Fraction*

Figure 3a shows the thermal expansion curves of Al0.5CoCrFeNi HEA at as-cast and 20% CR condition measured at the heating rate of 10 K/min. There is a peak around 1100 K for both curves indicating a phase transition exist, which is proved to be an FCC-BCC phase transition [8]. The first step for phase transition kinetic analysis is to determine the phase transformed volume fraction (*f*) as a function of temperature or time. Figure 3b shows how to calculate the phase transformed volume as a function of temperature with lever rule method using thermal expansion curves shown in Figure 3a. The lines AD and BE can be obtained by extending the linear expansion section of the thermal expansion curve. Then, the point C is the intersection of AC and thermal expansion curve (*f* = 0 when C is at point D and *f* = 100% when C is point E). Finally, the theoretical volume fraction of phase transition can be calculated using the simple lever rule method assuming the volume changes of phase transition is in proportion to the length variations, which is expressed in Equation (1) [26]:

$$f = \frac{|AC|}{|AB|},\tag{1}$$

where *f* is the volume fraction, and *|AC|* and *|AB|* are the length of the dashed line shown in Figure 3b, indicating the instant thermal expansion length and overall length, respectively.

**Figure 3.** (**a**) thermal expansion curves of Al0.5CoCrFeNi high-entropy alloy measured at as-cast and 20% cold rolled condition at heating rate of 10 K/min; (**b**) an example of as-cast sample shows how to calculate the phase transformed volume fraction as function of temperature.

Figure 4 shows the transformed volume fraction as a function of temperature at different heating rates for the as-cast (Figure 4a) and 20% CR (Figure 4b) samples. All of the *f–T* curves show the similar "S" type trend, indicating that the transformation is controlled by a typical nucleation-growth mechanism [27]. Compared with the as-cast sample, it can be seen from Figure 4 that the phase transition of cold rolled sample starts at a much higher temperature (1033 K and 1055 K at 4 K/min for as-cast and cold rolled samples, respectively) but ends at a much lower temperature (1195 K and 1151 K at 4 K/min for as-cast and cold rolled samples, respectively), indicating a much higher transformation rate when the sample is deformed.

**Figure 4.** The phase transformed volume fraction as a function of temperature curves of (**a**) as-cast and (**b**) 20% cold rolled Al0.5CoCrFeNi high-entropy alloy.

#### *3.3. Activation Energy*

Activation energy (*E*) and kinetic exponent (*n*) are the two most important kinetic parameters. Activation energy is the parameter that can assess phase transition energy barrier. The Kissinger– Akahira–Sunose (KAS) method [28] and Friedman method [29] can be used for the determination of activation energies. Here, we use KAS method, which has the following form [28]:

$$\ln\left(\frac{T^2}{\varrho}\right) = -\mathbb{C} + \frac{E}{RT'} \tag{2}$$

where *T* is temperature at certain transformed volume fraction, *φ* is the heating rate, *E* is the activation energy, *C* is a constant and *R* is the molar gas constant. Assuming the activation energy is the same at the same transformed fraction at different heating rates, then, by linear regression between ln*(T2/φ)* and *1/T* using the data taken from Figure 4, the activation energy can be directly determined by the slope of the curve.

The determined activation energy *E* for as cast and 20% CR Al0.5CoCrFeNi HEA at four different transformed volume fractions, *f* = 0.2, 0.4, 0.6 and 0.8 are shown in Figure 5. The activation energies of as-cast and 20% CR Al0.5CoCrFeNi HEA show the same decreasing trend with the increasing transformed volume fraction, *f*, however, the value are quite different. For the as-cast sample, the activation energies are 262 kJ/mol and 126 kJ/mol when transformed volume fraction is 0.2 and 0.8, respectively. The mean value is 181 kJ/mol—while, for the 20% CR sample, *E* are 223 kJ/mol and 110 kJ/mol when transformed volume fraction is 0.2 and 0.8, respectively. The average activation energy is 159 kJ/mol, which is 10% lower than the as-cast sample, indicating that the phase transformation process of the deformed sample is much easier than the as-cast sample.

**Figure 5.** The activation energies at different phase transformed volume fractions for as-cast and cold rolled Al0.5CoCrFeNi high-entropy alloys calculated using the KAS method.

#### *3.4. Kinetic Exponent*

The kinetic exponent, also called the Avrami exponent, is a key kinetic parameter that can directly obtain the phase transition mode. The kinetic exponent is generally determined based on the well-known Johnson–Mehl–Avrami (JMA) model [30], which has the following form:

$$f = 1 - \exp(-\mathcal{K}t^n),\tag{3}$$

where *n* is the Avrami exponent, *K* is a rate constant and *t* is the time. In case of phase transformations that are carried out at constant heating rate, the time can be expressed as [31]:

$$t - t\_0 = \frac{T - T\_0}{\mathcal{Q}},\tag{4}$$

where *T*<sup>0</sup> and *t* are the starting temperature and time of phase transition. Then, the kinetic exponent can be obtained by some mathematical treatment of Equations (3) and (4), which can be expressed as:

$$m = -\frac{\ln(-\ln(1-f))}{\frac{E}{RT}}.\tag{5}$$

Based on Equation (5), we can determine the kinetic expoent *n* from the slope of ln(−ln(1 − *f*)) vs. 1/*T* curve.

For a real phase transformation, *n* will change with the transformed volume fraction. A simple derivative equation of Equation (5) can determine the Avrami exponent as variables, which has been used in many previous published papers [31–33]. In this case, *n* is called local Avrami exponent, which can be described as:

$$m = -\frac{\partial \ln(-\ln(1-f))}{\partial \frac{E}{RT}}.\tag{6}$$

By taken the activation energy at a different transformed volume fraction using Equation (2), Avrami exponent *n* can be calculated by the differential of the curve of *E*/*RT* − ln[− ln(1 − f)]. Figure 6 represents the kinetic exponent at different transformed volume fractions at the heating rate 10 K/min for the as-cast and CR sample. According to Figure 6, two important pieces of information can be drawn. First, both of the as-cast and deformed samples show three typical stages of transformation: (I) rapid decrease of the kinetic exponent when *f* < 5%, indicating that the nucleation rate of the sample is decreasing at the beginning; (II) slow decrease of the kinetic exponent from 4 to 1.5 during 9% < *f* < 95%, indicating the decreasing nucleation rate, and the growth mode of the nuclei changes from interface controlled to diffusion controlled; (III) strong increase of the kinetic exponent when *f >* 95%, an indication that some new inhomogeneous nucleation occurs inside the sample. Second the kinetic exponent, *n*, is higher for the sample pre-deformed than the as-cast sample. This means that the

nucleation rate of the deformed sample is much higher and the growth of the nuclei is much easier than the as-cast sample.

**Figure 6.** Variation of *n* with phase transformed fraction for as-cast and deformed Al0.5CoCrFeNi high-entropy alloy at the heating rate of 10 K/min.

#### *3.5. Discussion*

For the deformed sample like cold rolling, normally, there will be an anisotropic effect due to the stress state during deformation are different at different directions. For the cold rolled sample, the rolling direction (RD) and the transverse direction (TD) always owns different textures, stress and properties. For the above results, all of the analysis is based on the sample that is taken from the RD direction. Figure 7 shows the thermal expansion curves that are both taken from the RD and TD direction of a 20% CR sample. According to the figure, it can be seen that the rolling direction has quite a limited effect on the thermal expansion behavior. The reason may be due to the fact that a 20% reduction is not large enough to generate a big difference, and it can also be found from Figure 1 that the CR sample has a similar microstructure as the as-cast sample. Large differences can be made evident at much larger CR reduction thickness, and the grain will be severely deformed and elongated along the rolling direction [12]. In this case, the thermal expansion behavior could be different for different directions. Thus, it can be concluded that there is no typical anisotropic effect in the 20% CR sample, and, during the solid state phase transformation process, the nucleation and growth process can be treated as homogeneous.

**Figure 7.** The thermal expansion curves of the samples taken from the rolling direction (RD) and transverse direction (TD) for 20% cold rolled Al0.5CoCrFeNi high-entropy alloy measured at the heating rate of 10 K/min.

There is not much difference for the microstructure (Figure 1) and phases (Figure 2) between the as-cast and deformed sample; however, according to the *f* –*T* curve, activation energy and kinetic exponent, there is a large difference. The microstructure of Al*x*CoCrFeNi (0 ≤ *x* ≤ 2) HEAs has been intensively investigated in different heat treatment conditions [17–25]. Different from the thermo-mechanical processing that can affect the transformation pathways in Al0.3CoCrFeNi [34], 20% cold rolling cannot alter the phase transition product [12]. Thus, the microstructure and phases exhibit no difference as shown in Figures 1 and 2. However, the deformed sample shows a much faster transformation rate, much lower activation energy and much larger kinetic exponents, an indication of a large different for the transformation kinetics. Figure 8 exhibits the tensile stress–strain curves of Al0.5CoCrFeNi HEAs at as-cast and 20% CR conditions. There exists a large difference: the deformed sample shows a much larger stress but lower plasticity. The yield stress increased from 402 MPa (as-cast sample) to 755 MPa (CR sample), while the strain decreased from 33.7% (as-cast sample) to 6.2% (CR sample), indicating a strong strain-hardening effect of the cold rolling. In the present study, due to the small reduction rate of the CR, there is not much difference in the macroscopic microstructure; however, 20% CR is enough to generate a large crystal lattice distortion, and many crystal defects, like vacancy, dislocation and substructure inside the grain. Then, the deformed sample stored a large quantity of energy, which is beneficial for the phase transition during the continuous heating. Thus, because of all the above factors that are generated by cold rolling, the deformed sample shows a much higher transformation rate since all the defects and stored energy inside the sample can promote the nucleation rate and grain growth speed.

**Figure 8.** The engineering stress–strain curves of Al0.5CoCrFeNi high-entropy alloys at as-cast and 20% CR conditions.

#### **4. Conclusions**

The effects of cold rolling on the phase transformation kinetics of an Al0.5CoCrFeNi high-entropy alloy are studied by the thermal expansion method. The transformed volume fraction–temperature curves of as-cast and 20% cold rolling Al0.5CoCrFeNi high-entropy alloys are determined by the lever rule method using thermal expansion curves. Using the KAS method, the mean activation energy for as-cast and deformed HEAs calculated are 181 kJ/mol and 159 kJ/mol, respectively. The kinetic exponent is determined using the local Avrami exponent method based on the JMA equation. Both as-cast and deformed samples show the three stages of kinetic exponent variations, and the deformed alloy shows a much larger value compared with the as-cast alloy, indicating a much higher nucleation and growth rate. The reason for the kinetic exponent variations are thought to be related with the deformation induced strain and defects inside the sample that can generate a much higher energy that favors the nucleation and growth process.

**Author Contributions:** Conceptualization, J.W, and J.S.L.; Investigation, J.W., H.X.Y., T.G., J.X.W.; Writing—original draft, J.W., H.X.Y., T.G.; Writing—review & editing, J.W., W.Y.W., J.S.L.

**Funding:** This research was funded by the Natural Science Foundation of China, grant number 51571161 and 51774240 and the fund of State Key Laboratory of Solidification Processing in NWPU, grant number 121-TZ-2015.

**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* **Al-Ti-Containing Lightweight High-Entropy Alloys for Intermediate Temperature Applications**

#### **Minju Kang, Ka Ram Lim, Jong Woo Won, Kwang Seok Lee and Young Sang Na \***

Korea Institute of Materials Science, 797 Changwondae-ro, Seongsan-gu, Changwon, Gyeongnam 642-831, Korea; minju214@gmail.com (M.K.); krlim@kims.re.kr (K.R.L.); jwwon@kims.re.kr (J.W.W.); ksl1784@kims.re.kr (K.S.L.)

**\*** Correspondence: nys1664@kims.re.kr; Tel.: +82-55-280-3377

Received: 19 April 2018; Accepted: 8 May 2018; Published: 9 May 2018

**Abstract:** In this study, new high-entropy alloys (HEAs), which contain lightweight elements, namely Al and Ti, have been designed for intermediate temperature applications. Cr, Mo, and V were selected as the elements for the Al-Ti-containing HEAs by elemental screening using their binary phase diagrams. AlCrMoTi and AlCrMoTiV HEAs are confirmed as solid solutions with minor ordered B2 phases and have superb specific hardness when compared to that of commercial alloys. The present work demonstrates the desirable possibility for substitution of traditional materials that are applied at intermediate temperature to Al-Ti-containing lightweight HEAs.

**Keywords:** high-entropy alloys; alloys design; lightweight alloys

#### **1. Introduction**

Recently, high-entropy alloys (HEAs) have attracted considerable attention because of their extraordinary properties [1–10], and numerous HEAs have been reported with various compositions [1–10]. Three major HEAs are CoCrFeMnNi alloy which has potential applications in cryogenic environments [5,6], refractory VNbMoTaW alloy for high-temperature structural applications [7,8], and AlCoCrFeNi alloy which maintains high strength up to intermediate temperatures [9,10]. AlCoCrFeNi HEA is relatively lightweight and has excellent specific strength around intermediate temperatures. It is a possible alternative for Ti alloys or wrought superalloys such as Inconel 718 [9,10]. Additional weight reduction can improve the competitiveness of the lightweight HEAs; therefore, we attempted to develop new HEAs that contain lightweight elements, namely Al and Ti.

Most HEAs are developed by a trial-and-error approach based on the effects such as mixing enthalpy and valence electron concentration (VEC) [11,12]. This conventional method is not effective when developing HEAs with a new combination of elements among numerous possibilities. The CALPHAD (Computer Coupling of Phase Diagrams and Thermochemistry) approach may be the best way for designing new HEAs because of its capability in predicting the phase stability [13–16]. However, the application of CALPHAD to the design of new HEAs is challenging because of the lack of a reliable thermodynamic database to cover the entire composition range [13–15]. F. Zhang et al. [17] reported a new approach to design new multi-component FCC HEAs by binary phase diagrams. The FCC single phase formation in the CoCrFeMnNi HEA was predicted using this approach [17]. This method is effective at finding "matching elements" that form a single solid solution and is suitable for designing novel HEAs.

In this work, we sought to design lightweight HEAs, which contain lightweight elements, namely Al and Ti, by using binary phase diagrams. Seven HEAs containing Al and Ti were designed and their mechanical properties were compared with those of commercial alloys.

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

The selection of candidate elements was conducted based on their binary phase diagrams. The details regarding the design of the Al-Ti-containing lightweight HEAs are described in Section 3.1.

The aforementioned Al-Ti-containing lightweight HEAs were fabricated by vacuum plasma arc melting (PAM) with high-purity elements. The HEA button ingots were re-melted 4 to 5 times in a melting furnace for the homogenization. The HEA plates were fabricated by vacuum induction melting using a graphite mold. The microstructures of the HEAs were analyzed by optical microscopy (OM), FE-SEM (model: SU-6600, HITACHI), and transmission electron microscopy (TEM, model: Tecnai F20, FEI). The TEM samples were prepared using focused ion beam (FIB). The crystal structure of the material was examined by X-ray diffraction (XRD) measurements on the as-casted material using a MXP21VAHF diffractometer with a CuKα radiation source (model: D/Max-2500VL/PC, RIGAKU). The Vickers hardness tests were carried out using a conventional indenter with a load of 2.94 N for 15 s. A minimum 10 tests were carried out on specimen.

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

#### *3.1. Selection of Candidate Elements Based on Their Binary Phase Diagrams*

To design lightweight HEAs, we selected Al and Ti as the basic elements owing to their low density. The key idea of designing new HEAs using their binary phase diagrams is finding elements that mix. Therefore, the solubility of Al and Ti should be investigated first. Figure 1 shows the Al-Ti phase diagram [18]. It is very complex, and numerous intermediate ordered phases exist. However, β-Ti (BCC phase) and α-Ti (HCP phase) appear in the Ti-rich region as a solid solution. These solid solution regions suggest the possibility of the formation of single solid solution. Therefore, the lightweight elements, Al and Ti, could be the basic elements for lightweight HEAs.

**Figure 1.** Al-Ti phase diagram [18]. It has solid solution region within a certain temperature and composition range that suggest the possibility of the formation of single solid solution.

The design of the Al-Ti-containing HEAs consisted of three steps. In the first step, the candidate elements were selected from various Ti-X binary phase diagrams wherein "X" represents elements that form homogeneous solid solution with Ti within a certain temperature range [18]. In the second step, a second series of candidate elements were selected from various Y-Al binary phase diagrams, wherein "Y" represents elements that show adequate solubility in Al within a certain composition and

temperature range. The final step was the selection of the final candidate elements from the X-Y binary phase diagrams.

From the Ti-X binary phase diagrams, we found 8 candidate elements, Cr, Hf, Mo, Nb, Ta, V, W, and Zr, in which each had a solid solution region in a certain range [18]. Figure 2 shows the Ti-X (X = Cr, Hf, Mo, Nb, Ta, V, W, or Zr) binary phase diagrams. All elements had a large area of solid solution within a certain temperature range, and all the homogeneous solid solution phases had BCC crystalline structure, indicated in green (Figure 2).

**Figure 2.** Ti-X phase diagrams where "X" is (**a**) Cr; (**b**) Hf; (**c**) Mo; (**d**) Nb; (**e**) Ta; (**f**) V; (**g**) W; and (**h**) Zr [18].

The next step consisted of finding a second series of candidate elements. Their binary phase diagrams are shown in the supplementary data (Figure S1). 6 candidate elements, Cr, Hf, Mo, Nb, V, and Zr, were chosen from the Y-Al binary phase diagrams. These elements are not form solid solution in all composition ranges, although the solid solutions contain Y's own crystalline structure (BCC) at a certain temperature and composition range. Among the 6 candidates, Nb and Zr were removed because they form solid solution in very restricted range.

From the first and second steps, Cr, Hf, Mo, and V were selected. To select the final candidate elements, the binary phase diagrams between these elements were investigated. Hf-Cr, Hf-Mo, and Hf-V showed complex phase diagrams, which could have possibly formed some intermediate ordered phases (supplementary data, Figure S2). Therefore, Cr, Mo, and V were selected as the elements for the Al-Ti-containing HEAs.

#### *3.2. Microstructure of the Al-Ti-Containing HEAs*

Ternary, quaternary, and quinary HEAs were designed by adding the selected elements to Al-Ti. Seven HEAs, numbered #1 to #7, were fabricated and their compositions are detailed in Table 1. The XRD profiles of the Al-Ti-containing HEAs are shown in Figure 3. The AlCrMoTi (#4) and AlCrMoTiV (#7) had a single BCC structure, and ordered BCC peaks appeared for the other HEAs.

**Table 1.** Valence electron concentration (VEC), atomic size difference (δ), and enthalpy of mixing (ΔHmix) of designed HEAs.


**Figure 3.** The XRD patterns of Al-Ti-containing HEAs.

The microstructures of AlCrMoTi (#4) and AlCrMoTiV (#7) which had a single BCC structure, are shown in Figure 4a,b. Both showed a dendritic morphology, and the dendritic growth was suppressed by the addition of V. The TEM analysis was performed for AlCrMoTiV (#7), and the results are shown in Figure 4c,d. Figure 4c is bright field image of AlCrMoTiV (#7) and Figure 4d is the diffraction pattern taken from [001]B2/BCC zone axis of Figure 4c. Although the XRD results demonstrate that it is a single BCC structure, diffraction pattern reveals (100) superlattice reflections marked with red-dotted circle that indicate an ordered B2 phase. Nano-scale B2 phases were observed; however, it was not detected by the XRD owing to their small fraction and size. Because AlCrMoTiV (#7) contained minor ordered precipitates, it could be defined as an ordered solid solution [19].

**Figure 4.** Microstructures of (**a**) AlCrMoTi and (**b**) AlCrMoTiV. Both have dendritic morphology. (**c**) bright field image and (**d**) diffraction pattern of AlCrMoTiV.

One of the physical properties of HEAs is the high-entropy effect [20]. The solid solution is stabilized because of the high configurational entropy [20,21], and this effect increases with an increased number of elements [20]. This trend can be observed in our work. As the number of elements increased from ternary to quinary, the phases progressively became simpler. The high-entropy effect is dominant at high temperatures according to G = H − TS, where G is Gibbs free energy, H is enthalpy, T is temperature, and S is entropy [20,21]. Thus, this HEA can exist as a solid solution at elevated temperatures, even though a small fraction of the B2 phase existed in an as-cast state at room temperature.

#### *3.3. Application of Empirical Parameters*

The empirical parameters for the HEAs originated from the classic Hume-Rothery rules [22,23]. Guo et al. [24] reported that −22 ≤ ΔHmix ≤ 7 kJ/mol, and δ ≤ 8.5% are required for the sole simple phases (i.e., FCC, BCC, and their mixtures, including both ordered/disordered cases). Zhang et al. [19] proposed −15 ≤ ΔHmix ≤ 5 kJ/mol and δ ≤ 6.5%, whereas Yang et al. [12] proposed δ ≤ 6.6%. In the case of VEC, 8 < VEC and VEC < 6.87 were suggested for the single FCC and BCC structures, respectively [20]. These parameters were statistically determined; therefore, there are differences and exceptions depending on the work [12,19,22,25].

The empirical parameters, valence electron concentration (VEC), atomic size difference (δ), and enthalpy of mixing (ΔHmix) of the designed HEAs were calculated and are shown in Table 1 and Figure 5 [22]. The VEC and δ values satisfied the existing criteria. However, the ΔHmix values were positioned relatively below the alloys which form solid solutions (Figure 5b). All the HEAs satisfy Guo's criterion although only AlCrMoTi (#4), AlMoTiV (#6), and AlCrMoTiV (#7) satisfy Zhang's criterion. The microstructural analysis demonstrates that AlCrMoTi (#4) and AlCrMoTiV (#7) are ordered solid solutions; the empirical parameters of HEAs in present work agree with previous research and support the idea that using the binary phase diagrams can be a solution to screen the proper candidate elements for the design of novel HEAs.

**Figure 5.** (**a**) VEC and (**b**) δ − ΔHmix plot of Al-Ti-containing HEAs. Yellow colored regions highlight Zhang's criterion [19] for single phase to form.

#### *3.4. Specific Hardness of the Al-Ti-Containing HEAs*

It has been reported that HEAs have a severe lattice distortion owing to the atomic size differences of elements, which induces solid solution strengthening [25]. Figure 6 shows the relationship between the hardness and δ. The hardness varied across the δ values. Moreover, Figure 6b clearly shows the hardness variation with the addition of Cr, Mo, and V. The hardness increased remarkably with the addition of Cr. The atomic radius of Cr is 1.25 Å which is quite different than the others [24]; this causes further severe lattice distortion.

The hardness of the AlCrMoTi (#4) and AlCrMoTiV (#7) solid solutions were compared with that of other HEAs and commercial alloys and is summarized in Table 2 [26–31]. The hardness of the AlCrMoTi (#4) and AlCrMoTiV (#7) were 606 ± 11 and 556 ± 25 HV, respectively. These values are 17% and 7% higher than that of the AlCoCrFeNi HEA, which is representative of HEAs for intermediate temperature applications [9]. The theoretical densities are also lower than that of the AlCoCrFeNi HEA, so the specific hardnesses were 30% and 19% higher than that of the AlCoCrFeNi HEA. Furthermore, the specific hardness of the HEAs in the present work are considerably higher than that of competitive conventional alloy systems. The AlCrMoTi (#4) showed a 29% improved specific hardness when compared with that of Ti-6Al-4V alloy, which is an alloy that is commonly used for intermediate temperature applications. It is expected that a great improvement in the hardness can accomplish the weight lightening through the gage reduction of machineries despite their higher density when compared to that of the Ti-6Al-4V alloy. Further research on the fine-tuning of the elemental composition for lighter and cheaper alloys as well as a detailed microstructural and high-temperature property analysis could open a new path toward lightweight HEAs for intermediate temperature applications.

**Figure 6.** (**a**) Relationship between the hardness and atomic size difference; (**b**) hardness variation with the Cr, Mo, and V addition to ternary HEAs. The hardness increases remarkably with the Cr addition.


**Table 2.** Comparison of specific hardness of HEAs with other alloys.

#### **4. Conclusions**

In summary, HEAs which contained Al and Ti were designed based on their binary phase diagrams. This approach is powerful for screening candidate elements for novel HEAs. The candidate elements that were selected formed a solid solution within a certain temperature and composition range. The high-entropy effect is enhanced with an increased number of elements; therefore, the AlCrMoTi and AlCrMoTiV HEAs are verified to be solid solutions with a minor ordered B2 phase. These HEAs have a superb specific hardness when compared to that of Ti-6Al-4V and Inconel 718 alloys and show promise as future substitutes for Ti alloys for intermediate temperature structural applications. Furthermore, fine-tuning of the elemental composition of HEAs can lead to the development of novel light-weight HEAs.

**Supplementary Materials:** The following are available online at http://www.mdpi.com/1099-4300/20/5/355/s1, Figure S1: Y-Al phase diagrams where "Y" is (a) Cr, (b) Hf, (c) Mo, (d) Nb, (e) V, and (f) Zr [18]. Figure S2: (a) Hf-Cr, (b) Cr-Mo, (c) Cr-V, (d) Hf-Mo, (e) Hf-V, and (f) Mo-V binary phase diagrams [18]. Cr, Mo, and V were selected which have simple phase diagrams.

**Author Contributions:** M.K., K.R.L., J.W.W. and Y.S.N. conceived and designed the experiments; M.K. and K.R.L. performed the experiments and analyzed the data; K.S.L. and Y.S.N. contributed reagents/materials/analysis tools; M.K. wrote the paper.

**Acknowledgments:** This study was supported financially by Fundamental Research Program (PNK5610) of the Korean Institute of Materials Science (KIMS), and by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (MSIP) of Korea (NRF-2016M3D1A1023534).

**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* **High Strength and Deformation Mechanisms of Al0.3CoCrFeNi High-Entropy Alloy Thin Films Fabricated by Magnetron Sputtering**

**Wei-Bing Liao 1, Hongti Zhang 2,3, Zhi-Yuan Liu 4, Pei-Feng Li 5, Jian-Jun Huang 1, Chun-Yan Yu 1,\* and Yang Lu 2,6**


Received: 23 January 2019; Accepted: 2 February 2019; Published: 4 February 2019

**Abstract:** Recently, high-entropy alloy thin films (HEATFs) with nanocrystalline structures and high hardness were developed by magnetron sputtering technique and have exciting potential to make small structure devices and precision instruments with sizes ranging from nanometers to micrometers. However, the strength and deformation mechanisms are still unclear. In this work, nanocrystalline Al0.3CoCrFeNi HEATFs with a thickness of ~4 μm were prepared. The microstructures of the thin films were comprehensively characterized, and the mechanical properties were systematically studied. It was found that the thin film was smooth, with a roughness of less than 5 nm. The chemical composition of the high entropy alloy thin film was homogeneous with a main single face-centered cubic (FCC) structure. Furthermore, it was observed that the hardness and the yield strength of the high-entropy alloy thin film was about three times that of the bulk samples, and the plastic deformation was inhomogeneous. Our results could provide an in-depth understanding of the mechanics and deformation mechanism for future design of nanocrystalline HEATFs with desired properties.

**Keywords:** high-entropy alloys; thin films; hardness; deformation behaviors; nanocrystalline

#### **1. Introduction**

It is well known that among all the alloy composition design systems, high-entropy alloys (HEAs) are a brand-new concept based on novel multi-component system composition designs. They contain at least four or five principal metal components and simply form a single face-centered cubic (FCC), body-centered cubic (BCC) or hexagonal close-packed (HCP) phase [1–6]. This novel concept is an important breakthrough of the past 25 years [7,8], as it is completely different from the traditional alloy design concepts in which one major component was selected, and other minor components were added to improve their related physical and chemical performances. It is worth mentioning that HEAs not only have simple phase structures, but also possess many excellent mechanical and physical properties, such as high tensile strength [9–11], good ductility at ambient and cryogenic temperatures [12,13], superior wear and fatigue resistance [14], and strong radiation tolerance [15,16]. These unique features qualify HEAs as potential engineering materials to meet the demanding requirements for complex and harsh environment applications, particularly in the turbine, aerospace, and nuclear industries [17–21]. However, the chemical composition of HEAs contains multiple elements which would naturally raise the cost for industrial application, and limit HEAs extensive development. As a consequence, to reduce the cost for future industrial applications and take full advantage of the above excellent comprehensive properties, HEA thin films (HEATFs) can be efficiently prepared and simultaneously coated on the surface of industrial components, especially for those complex geometry components. In these cases, the HEATFs will play an important role. The initial research of HEATFs is associated with the recent rapid development of HEAs and the high throughput preparation idea [22–26]. As the geometric size and microstructures of the thin films are different from the three-dimensional bulk samples, their performances under loading and service conditions could be completely different [27,28]. So far, HEATFs were verified to have remarkable effects on the hardness [24]. A series work on HEATFs was done not only on the high throughput fabrication but also on the mechanical properties, including the hardness and corrosion properties [29–36]. The previous work has greatly promoted the industrial application of HEATFs. Unfortunately, the related deformation behaviors have not been clearly revealed until now. To facilitate the use of HEATFs and provide a continuous coating technique, the deformation behaviors and reliability of HEATFs merit further investigation. Therefore, in this study we prepared the Al0.3CoCrFeNi HEATFs with a main simple FCC structure by magnetron sputtering, and fabricated nano-scaled pillars on the surface of the thin film by focus ion beams (FIBs), then utilized in situ scanning electron microscopy (SEM) compression to study the deformation behaviors of the HEATFs.

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

The target with a composition of Al0.3CoCrFeNi was prepared by metallurgy with high-purity (>99.99%) raw metal materials of aluminum, cobalt, chromium, iron, and nickel. The size of the target is φ76.2 × 3.175 mm. The Al0.3CoCrFeNi HEATFs were deposited on silicon wafer substrates by magnetron sputtering. Before putting the target in the vacuum chamber, it was cleaned by argon ion bombardment for about 2 min to remove the oxide or contaminants on the surface. To ensure a uniform deposition a rotation speed of the silicon wafer substrate was set at 2 rpm. The surface roughness of the as-deposited HEATFs was determined by white light interferometry (WLI) using Wyko NT9300 Surface Profiler (Veeco Instruments, Plainview, NY, USA), while the surface morphology and detail nanostructures were characterized by scanning electron microscopy (SEM) and atomic force microscopy (AFM) (Bruker Dimension IconTM, Billerica, MA, USA) with ScanAsyst (Bruker Dimension IconTM, Billerica, MA, USA) at room temperature. To investigate the phase structure of the as-deposited HEATFs, high-energy synchrotron radiation X-ray in transmission mode at 11-ID-C of Advanced Photon Source (APS) was used. The X-ray beam wavelength was 0.117418 Å. The detail microstructures of the as-deposited HEATFs were observed by high-resolution transmission electron microscopy (HRTEM) using a JEOL JEM-2100F instrument (JEOL, Akishima, Tokyo, Japan) operated at 200 kV. The chemical composition was analyzed by the energy dispersive X-ray spectrometer (EDS) equipped in the transmission electron microscopy (TEM). Nanoindentation experiments were performed using a Hysitron TI750 nanoindenter (Hysitron, Inc., Minneapolis, MN, USA) with a Berkovich tip. To avoid any potential effects of the substrate on the experiment, the indentation depth was kept to be less than 10% of the whole thickness of the HEATFs. Micropillars were fabricated out of the Al0.3CoCrFeNi HEAHFs by using a FEI Scios focused ion beam (FIB) (USA) (Thermo Scientific™, Hillsboro, OR, USA) at 30 kV/10pA as the final etching condition. The height of the nanopillars was kept to be less than the thickness of the Al0.3CoCrFeNi HEATFs. The in situ SEM compression tests were conducted at room temperature using a PI 85 PicoIndenter (Hysitron Inc.) with a flat punch diamond tip inside a FEI Quanta 450 FEG (USA) (Thermo Scientific™, Hillsboro, OR, USA), under

displacement-control mode and at a strain rate of around 5×10-3 <sup>s</sup>-1. Raw load-displacement data were used to calculate the engineering stress and strain.

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

Figure 1a,b show the two-dimensional (2D) and the three-dimensional (3D) surface roughness and profiles of the as-deposited Al0.3CoCrFeNi HEATFs, respectively, and Figure 1c,d show the X-profile and Y-profile of the corresponding positions selected on the HEATFs as marked in Figure 1a. It can be clearly seen that there are fine undulating nanostructures on the surface of the Al0.3CoCrFeNi HEATFs prepared by the magnetron sputtering deposition technique; however, the entire surface is very flat and smooth, with a roughness Ra of less than 3.5 nm.

**Figure 1.** Surface profiles of the Al0.3CoCrFeNi high-entropy alloy thin films (HEATFs) characterized by white light interferometry (WLI) technique. (**a**) 2D surface profiles, (**b**) 3D surface profiles, (**c**,**d**): the profiles of the *x*-axis and *y*-axis as marked in (a) respectively.

The SEM surface morphology and the specific fine nanostructures of the as-deposited Al0.3CoCrFeNi HEATFs with a magnification of 50,000 times are shown in Figure 2a. It demonstrates that these fine nanostructures are well-knit and compact. The thickness of the HEATFs is about 4 μm, as shown in Figure 2b. To characterize the feature of the HEATFs in more detail, AFM scanning experiments were further conducted. Figure 2c,d show the 2D and 3D AFM images of the surface feature of the Al0.3CoCrFeNi HEATFs. Uniform nanostructures are clearly observed, and the heights of these undulating nanostructures were less than 5 nm, which is well consistent with that typically observed by the surface profile. All these experimental data verified that there were a lot of fine nanostructures on the surface of the Al0.3CoCrFeNi HEATFs, and the surface was very smooth as a whole, with a roughness Ra of less than 5 nm.

Figure 3a shows the TEM images and the corresponding EDS analysis of the Al0.3CoCrFeNi HEATFs. It can be seen that there are a lot of nanocrystalline structures in the HEATFs, with a grain size order of ~10 nm. The elemental distribution of the as-deposited HEATFs is homogenous. The bright and dark places shown in the TEM-EDS images are ascribed to the uneven sample thickness. Quantitative analysis by EDS confirms that the chemical composition is nearly the same as the composition of the sputtering target, as shown in Table 1.

**Figure 2.** Surface morphologies and microstructures of the Al0.3CoCrFeNi HEATFs characterized in detail by scanning electron microscopy (SEM) and atomic force microscopy (AFM). (**a**) The SEM image of the surface morphologies, (**b**) the cross-section the HEATFs deposited on the silicon substrate, (**c**) the AFM image of the surface structure, and (**d**) 3D surface structures of the HEATFs.

**Figure 3.** Element distribution mapping for the Al0.3CoCrFeNi HEATFs by TEM-EDS. The top TEM image shows the region analyzed. (**a**) Explanations for subfigure a; (**b**) explanations for subfigure b; (**c**) explanations for subfigure c; (**d**) explanations for subfigure d; (**e**) explanations for subfigure e; (**f**) explanations for subfigure f.


**Table 1.** The chemical composition of the as-deposited Al0.3CoCrFeNi HEATFs compared with that of the sputtering target.

To obtain the phase structural information of the Al0.3CoCrFeNi HEATFs, high-energy synchrotron radiation X-ray studies were undertaken. Figure 4a shows the X-ray line profiles of the HEATFs. The (111), (200), (220), and (311) phase peaks were observed and identified to be a typical FCC crystalline structure, whilst a small peak appeared before the (111) peak, which means that a minor ordered BCC NiAl type phase structure was in the HEATFs [37]. The corresponding diffraction patterns are exhibited in Figure 4b. The weak continuous rings certify that there are tiny polycrystalline structures in the HEATFs. Interestingly, the diffraction rings of the HEATFs were discrete, with obvious intensity differences, indicating that there were strong textures in the HEATFs. This event could be ascribed to the preferred growth of the thin film induced by the silicon substrate. It should be noted that a rigorous diffraction-intensity-distribution calculation of the solid solution phases responsible for certain orientations in the HEATFs is worthy of a focused topic. However, it is beyond the scope of this work. The synchrotron X-ray experimental results provide cogent evidence that the magnetron sputtering technique is an effect way to prepare the HEATFs with a simple phase structure. Moreover, it could also lead to a wide research range of HEAs, studying the corresponding properties from mesoto nanometer regimes.

**Figure 4.** Phase structures of the HEATFs by high-energy synchrotron radiation X-ray. (**a**) synchrotron X-ray line profiles and (**b**) typical diffraction pattern.

Figure 5a shows the nanoindentation properties of the as-deposited HEATFs. Since a series of 4 × 4 matrix array indentation points were tested in sequence, the average values of elastic modulus and hardness were accurately calculated and were identified to be about 186.01 GPa and 11.09 GPa, respectively. It should be noted that the hardness of the HEATFs is about three times higher than that of the as-cast bulk Al0.3CoCrFeNi HEA sample, but the elastic modulus is nearly the same [38,39]. This enhanced hardness can be ascribed to the nanocrystalline strengthening mechanism which induced the hardening by a large number of grain boundaries observed in Figure 3a. Figure 5b is the typical nanoindentation load-depth curve of the HEATFs. It can be seen that as the loading force increases, the depth of the indenter pressed into the film gradually increases. After unloading, an irreversible depth was retained, indicating a plastic deformation has occurred on the surface of the HEATFs. Figure 5c exhibits the SEM image of the impression mark. The indentation profiles are self-similar. A remarkable pile-up (marked with red arrows) around the indentation can be clearly observed, suggesting that high localized plastic deformation occurred during nanoindentation.

**Figure 5.** Nanoindentation properties of the as-deposited Al0.3CoCrFeNi HEATFs. (**a**) Elastic modulus and hardness of the HEATFs at different indentation points, (**b**) typical nano-indentation load–depth curve, and (**c**) typical SEM image of the impression mark.

To further characterize the mechanical properties of the Al0.3CoCrFeNi HEATFs, a nanopillar sample with a diameter of 738 nm was fabricated from the HEATFs, and in situ SEM compression tests were conducted on the nanopillar sample, as shown in Figure 6. The entire compression deformation process of the nanopillar can be divided into the following stages. Initially, the nanopillar was deformed elastically, and no significant trace appeared on the surface of the nanopillar, as shown in Figure 6a. Secondly, with the increase of the compressive stress, a large localized metal flow and plastic deformation occurred at the top part of the nanopillar, as marked with a red arrow in Figure 6b. It indicates that the deformation of the Al0.3CoCrFeNi HEA nanopillar was inhomogeneous. After that, it can be observed that a slip was generated at the top part of the nanopillar, which was marked with a red arrow in Figure 6c. The occurrence of the slip is not only related to the plastic deformation, but also has an impact on the work hardening and the serration behavior of the HEAs [40]. The corresponding compression engineering stress–strain curve of the Al0.3CoCrFeNi HEA nanopillar is shown in Figure 6d. The yield strength of the nanopillar is about 1024 MPa, which is also about three times that of the bulk sample [38]. It is consistent with the above experimental results obtained by nanoindentation, as the strength is directly proportion with the hardness. After yielding the nanopillar exhibits work-hardening up to an ultrahigh strength. The compressive strength and corresponding strain were ~2075 MPa and ~11.39%, respectively. Following this, softening dominates until final fracture at a strain of ~12.14%. In general, the compression results further confirmed that the yield strength of the Al0.3CoCrFeNi HEATFs is about three times that of the bulk samples, and the plastic deformation is inhomogeneous.

**Figure 6.** Compression properties for the Al0.3CoCrFeNi pillar with a diameter of 738 nm prepared from the HEATFs. (**a**) Elastic deformation stage, (**b**) localized plastic deformation occurred at the top of the pillar, (**c**) a slip generated at the top part of the pillar, (**d**) a typical compression engineering stress–strain curve of the pillar.

#### **4. Conclusions**

In conclusion, the Al0.3CoCrFeNi HEATFs prepared by the magnetron sputtering technique were smooth, with a surface roughness Ra of less than ~5 nm. The chemical composition was homogeneous and amounts of nanocrystallines with a main single FCC phase structure formed in the HEATFs. The hardness of the HEATFs was ~11.09 GPa, and the yield strength of the nanopillar prepared from the HEATFs was ~1024 MPa. Both the hardness and the yield strength were about three times that of the bulk samples. Simultaneously, it was found that the plastic deformation of the HEATFs was inhomogeneous and localized. The present study could provide useful insights in the design and application of HEATFs for functional micro- and nano-devices.

**Author Contributions:** Formal analysis, W.-B.L. and Z.-Y.L.; Funding acquisition, W.-B.L., C.-Y.Y. and Y.L.; Investigation, W.-B.L., H.Z., P.-F.L., J.-J.H. and C.-Y.Y.; Methodology, W.-B.L. and Y.L.; Supervision, Y.L.; Writing—Original Draft, W.-B.L.; Writing—Review & Editing, W.-B.L.

**Funding:** This work was supported by the National Natural Science Foundation of China (Grant No. 51801128, 51701125), the Youth Innovation Talent Project of Guangdong Province (Grant No. 2017KQNCX175), the Natural Science Foundation of Shenzhen University (Grant Nos. 2017069, 201553, 827-000180), the Teaching Reform Research Project of Shenzhen University (Grant No. JG2018090), and the Shenzhen Basic Research Project JCYJ20170302142339007. Y. Lu gratefully thanks the funding support from Shenzhen Science and Technology Innovation Committee under the grant JCYJ20170413141157573.

**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* **Effects of Nitrogen Content on the Structure and Mechanical Properties of (Al0.5CrFeNiTi0.25)Nx High-Entropy Films by Reactive Sputtering**

#### **Yong Zhang 1,\*, Xue-Hui Yan 1, Wei-Bing Liao <sup>2</sup> and Kun Zhao <sup>3</sup>**


Received: 9 July 2018; Accepted: 19 August 2018; Published: 21 August 2018

**Abstract:** In this study, (Al0.5CrFeNiTi0.25)Nx high-entropy films are prepared by a reactive direct current (DC) magnetron sputtering at different N2 flow rates on silicon wafers. It is found that the structure of (Al0.5CrFeNiTi0.25)Nx high-entropy films is amorphous, with x = 0. It transforms from amorphous to a face-centered-cubic (FCC) structure with the increase of nitrogen content, while the bulk Al0.5CrFeNiTi0.25 counterpart prepared by casting features a body-centered-cubic (BCC) phase structure. The phase formation can be explained by the atomic size difference (δ). Lacking nitrogen, δ is approximately 6.4% for the five metal elements, which is relatively large and might form a BCC or ordered-BCC structure, while the metallic elements in this alloy system all have a trend to form nitrides like TiN, CrN, AlN, and FeN. Therefore, nitride components are becoming very similar in size and structure and solve each other easily, thus, an FCC (Al-Cr-Fe-Ni-Ti)N solid solution forms. The calculated value of δ is approximately 23% for this multicomponent nitride solid solution. The (Al0.5CrFeNiTi0.25)Nx films achieve a pronounced hardness and a Young's modulus of 21.45 GPa and 253.8 GPa, respectively, which is obviously much higher than that of the as-cast Al0.5CrFeNiTi0.25 bulk alloys.

**Keywords:** high-entropy films; phase structures; hardness; solid-solution; interstitial phase

#### **1. Introduction**

High-entropy films (HEFs) are a brand-new type of alloy film, which has been developed recently based on the design concept of high-entropy alloys (HEAs) [1]. HEFs can be defined as multiple-component films with high-entropy mixing. Generally, the HEAs are composed of multi-principal-elements (at least five elements with five at a % ≤ each element content ≤ 35 at %) [2,3]. HEAs feature higher mixing entropy than traditional alloys and tend to form disorder face-centered cubic (FCC) and/or body-centered cubic (BCC) phase structures rather than ordered intermetallic compounds [4–6]. Due to severe lattice distortion and solid solution strengthening, attributable to the multi-components, the HEAs show many excellent mechanical properties, such as high-strength, good ductility, and high-wear and corrosion resistances [1]. Based on a similar scientific concept, HEFs have been designed and investigated gradually and show a great potential for application in the coating industry [7–12]. To date, many excellent properties of HEFs have been discovered and studied, such as high-wear resistance [13,14], high-corrosion resistance [15–17], diffusion-barriers effects [18–20], solar-thermal-conversion effects [10,21], plastic-deformation characteristics [22,23], thermal stabilities [7,24,25], and soft magnetic properties [26]. However, there are few theories that can explain the mechanism of the phase formation of the HEFs.

The phase formation and mechanical properties of (Al0.5CrFeNiTi0.25)Nx high-entropy thin films with different N2 flow rates are studied in this paper. The phase structures and mechanical properties of AlxCrFeNiTi0.25 (x: molar ratio, x = 0, 0.25, 0.5, 0.75, and 1.0) bulks have previously been studied systematically [27]. Herein, the optimal composition of AlxCrFeNiTi0.25 (x = 0.5) alloy is selected as a magnetron-sputtering target to explore the phase formation mechanism and mechanical properties of high-entropy thin films. Significantly, the phase structures of high-entropy thin films transform from amorphous to FCC with the increase of nitrogen content, while the bulk Al0.5CrFeNiTi0.25 alloy attains BCC phase structure. Phase formation mechanism is explored from both theoretical calculations and experiments. Concerning high entropy thin films, the ability for solid solution structure formation is first discussed from the atom radius difference (δ). This study can help to provide new insights into understanding the phase formation mechanism of multi-component alloy thin film solids with small atoms such as nitrogen.

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

The (Al0.5CrFeNiTi0.25)Nx films were deposited on p-type Si (100) wafers by a direct current (DC) magnetron sputtering using non-equal atomic ratio Al0.5CrFeNiTi0.25 targets of Φ 60 mm in diameter and 2 mm in thickness. The alloy target was prepared by arc-melting and the smelting was repeated at least five times to ensure uniform mixing of components. Prior to deposition, the Si substrates were cleaned sequentially and rinsed by acetone, alcohol, and distilled water in an ultrasonic bath. Pre-sputtering was an effective way to remove oxide or contaminants on the surface of the target. When the base pressure held at 2.0 × <sup>10</sup>−<sup>4</sup> Pa, high purity argon was injected into the vacuum chamber and the target was cleaned by argon ion bombardment for 15 min at a power of 100 W. The deposition of the (Al0.5CrFeNiTi0.25)Nx films were carried out in an Ar + N2 mixed atmosphere under a DC power of 100 W with a working pressure of 0.5 Pa. The schematic diagram of reactive sputtering is shown in Figure 1. The metal atoms escaped the surface of the target due to the bombardment of high energy particles. Due to the nitrogen atmosphere, different metal atoms reacted with N-atoms and were deposited on the substrates in the end. During the deposition, the total flow rate of Ar + N2 was maintained at 30 standard cubic centimeters per minute (sccm) and the ratio was N2/(Ar + N2) and RN2 was controlled at 0, 10%, 20%, 30%, 40%, and 50%, respectively. The work distance between the substrate and the target was 60 mm and the deposition time was maintained at 60 min. No external heating or bias was used on the substrate during the deposition process.

**Figure 1.** The schematic diagram of reactive sputtering.

The crystal structures of (Al0.5CrFeNiTi0.25)Nx films were analyzed by a glancing-incidence (1◦) X-ray diffractometer (XRD, BRUKERD8 Discover, Bruker, Karlsruhe, Germany) using the Cu Kα radiation at a scanning rate of 4◦/min. The scanning step was 0.02◦ with a scanning range of 20◦–90◦. The morphology studies and thickness measurements were carried out using field-emission scanning electron microscopy (SEM, Auriga Field Emission Scanning Electron Microscope, Carl Zeiss, Jena, Germany) equipped with an Energy Dispersive X-ray Spectrometer (EDX) operated at 10 kV. The surface roughness of the coatings were measured by an atomic force microscope (AFM, Veeco DI-3100, Bruker, Beijing, China). The hardness and modulus of the as-deposited films were tested at five points for each sample with a nano-indenter using a Berkovich triangular pyramid indenter. The distance between each indentation was 50 μm. The Poisson's ratio and Elastic modulus of the indenter tip were 0.07 and 1.141 × <sup>10</sup><sup>6</sup> MPa, respectively. Micrographs of indentations were tested by a laser scanning confocal microscope (LSCM, OLS-4100, Olympus, Tokyo, Japan).

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

#### *3.1. Phase Structures*

#### 3.1.1. Phase Formation Mechanism Analysis

The XRD patterns of (Al0.5CrFeNiTi0.25)Nx films deposited at different RN2 are evaluated in Figure 2. An amorphous structure was observed with x = 0. As the N2 flow rate increased, the phase structures of (Al0.5CrFeNiTi0.25)Nx films showed a tendency to crystallize. Using XRD analysis, thin films displayed a simple FCC structure and the grains were nanoscale at about 20 nm. The bulk Al0.5CrFeNiTi0.25 alloys counterpart featured a BCC structure [27]. Due to the high-cooling rate, films were far from the equilibrium that could be achieved by bulk alloys, which led to the difference in structure even if they were the same system. Many similar results have been reported for other system film alloys [16,28–31]. While the metallic elements in this alloy system, which all had trends to form nitrides like TiN, CrN, AlN, and FeN, had results that showed only a simple FCC structure in an XRD pattern, rather than a various nitrides phase. It can be inferred that the solid solution occurred between nitrides in high-entropy thin films. Mutual dissolution between carbides and nitrides has been reported in many studies. Figure 3 shows the UC-ZrC0.81 pseudo-binary system that showed a miscibility relationship and the substitutional solid solution structure was formed with the carbides as solid solution units [32]. The nitrides also obtained similar results, such as Ti-Al-N, Cr-Al-N, Ti-Cr-N and Ti-Zr-N, Ti-Al-Si-N nitride films [33–37].

**Figure 2.** X-ray patterns of the (Al0.5CrFeNiTi0.25)Nx films deposited at different RN2.

**Figure 3.** The UC-ZrC0.81 pseudo-binary phase diagram. Reproduced with the permission from Reference [34].

A previous study showed that FCC phases were found to have lower δ, whereas BCC phases showed higher δ [38]. Regarding (Al0.5CrFeNiTi0.25)Nx with x = 0, δ was approximately 6.4%, which was relatively large, and might form a BCC or ordered BCC structure. Metallic atoms solved each other and formed a BCC solid solution structure (Figure 4a). Accompanying the increase in nitrogen, nitrides formed in (Al0.5CrFeNiTi0.25)Nx and were similar in size and structure. The δ between nitrides was relatively small, thus an FCC structure might have formed. When RN2 was equal to 10% and 20%, a low concentration of nitrogen existed in the lattice. However, small amounts of nitrogen cannot facilitate lattice reconstruction. Thus, it still maintained an amorphous structure, however, with a higher nitrogen content, the formation of nitrides was promoted that efficiently improved the order of the lattice. The nitride components were becoming more similar in size and structure and solved each other easily, thus, an FCC (Al-Cr-Fe-Ni-Ti)Nx solid solution formed, which is shown in Figure 4b,c.

**Figure 4.** Schematic of lattice structure. (**a**) Al0.5CrFeNiTi0.25 bulk; (**b**) (Al0.5CrFeNiTi0.25)Nx films; and (**c**) schematic diagram of high-entropy thin film formation.

Figure 5 displays the distribution of component elements of (Al0.5CrFeNiTi0.25)Nx thin films deposited at RN2 = 50% through EDX, including surface, and line and point scanning. As presented in each element map in Figure 5a, the distribution of component elements was uniform, which indicates a segregate-free characteristic. The results of line and point scanning also reflected the same element distribution characteristics. The internal picture in Figure 5b is the composite surface-scanning of this region and segregate-free characteristics can be observed. Viewing the comparative analysis of the energy spectrum of Point A and Point B in Figure 5c, the distribution of elements was stable and had a slight fluctuation. Moreover, the lattice constant of the FCC solid solution structure was calculated by the Prague formula. The FCC solid solution structure with a lattice constant about 4.093 Å was completely different with nitrides such as TiN, CrN, and so on, which also confirmed the occurrence of a solid solution.

**Figure 5.** Energy dispersive X-ray spectroscopy of (Al0.5CrFeNiTi0.25)Nx high-entropy thin films deposited at RN2 = 50%: (**a**) face-scanning; (**b**) line-scanning; the insert picture is the composite surface scanning of this region; and (**c**) point-scanning.

#### 3.1.2. Solid Solution Formation Ability

Regarding high-entropy alloys, the solid-solution formation ability was estimated by the value of Ω and δ [39], which were defined as follows:

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

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

where *Tm* was the melting temperature of the n-elements alloy, Δ*Smix* was the mixing-entropy of an n-element system, Δ*Hmix* was the mixing-enthalpy of an n-element system, *ci* was the atomic percentage of the component, *ri* was the atomic radius, and *<sup>r</sup>* <sup>=</sup> *<sup>n</sup>* ∑ *i*=1 *ciri* was the average atomic radius. Reviewing the previous study, it was concluded that the region where solid-solution structures were formed was in the range of 1.1 to 229.8 for Ω and 0.8% to 6.6% for δ. The radius of component elements in (Al0.5CrFeNiTi0.25)Nx are shown in Table 1, and the specific Ω and δ values of (Al0.5CrFeNiTi0.25)Nx at different N2 flow rates are shown in Table 2.

Concerning (Al-Cr-Fe-Ni-Ti)Nx, the δ was described as a function of nitrogen content, as shown in Figure 6a. The analytical curve was calculated based on the atomic ratio of the target, which was the theoretical value. While different deposition yields of component elements can cause fluctuations in the atomic ratio, the experimental values were well matched with theoretical values, as shown by the star points in Figure 6a. To clearly distinguish different phase regions of film alloys in this study, the nitrogen content was also described as a function of δ in Figure 6b. Regarding Al0.5CrFeNiTi0.25, δ was about 6.4% located in the SSS zone in Figure 6b, and the BCC solid solution phase structure was stable. When increasing the nitrogen content, the FCC structure formed in the (Al0.5CrFeNiTi0.25)Nx system. The calculated value of δ was about 23% for that FCC structure. When the nitrogen content was low, the amorphous structure was stable and located in the middle region.

**Figure 6.** (**a**) Analytical curve of δ as a function of nitrogen content in (Al0.5CrFeNiTi0.25)Nx and (**b**) analytical curve of nitrogen content as a function of δ in (Al0.5CrFeNiTi0.25)Nx films (12% according to Inoue Principle, 15% and 59% according to the Hume–Rothery Rule).

**Table 1.** The radius of component elelments in (Al0.5CrFeNiTi0.25)Nx films.



**Table 2.** The nitrogen content, Ω value, δ value, and the roughness of films at different RN2.

#### *3.2. Deposition Rates*

Figure 7 presents the deposition rate of (Al0.5CrFeNiTi0.25)Nx as a function of RN2 and the exact value of thickness is shown in the internal table. Following the addition of nitrogen, the deposition rate of films was significantly reduced by about 30%. During the deposition, the total flow rate of Ar + N2 was maintained at 30 sccm. Thus, the density of argon decreased gradually with the increase of N2 flow rate and resulted in a lower efficiency of argon-ion bombardment of the target. Under higher RN2, the deposition rate continued to slow and maintain a lower sputtering yield. Added to the effect caused by low argon gas density, the ceramicization of the target surface was also an important factor. This was mainly due to the N-containing layer that formed on the target surface, which reduced the conductivity of the target and resulted in a low sputtering yield. This phenomena was "target poisoning" [40,41].

**Figure 7.** Deposition rates of (Al0.5CrFeNiTi0.25)Nx films at different RN2.

#### *3.3. Surface Morphologies*

The AFM images and SEM micrographs of the (Al0.5CrFeNiTi0.25)Nx films at different RN2 are shown in Figure 9. Initially, a very dense and smooth surface with low surface roughness was obtained. Looking at SEM micrographs, it was observed that a small amount of nanostructure precipitation occurred as the RN2 increased to 10%. When the N2 flow rate increased, the number and particle size of the nanostructures were increased significantly. Viewing the AFM images, many needle nanostructures were also observed in the high-entropy thin films. Additionally, the roughness of films was measured by AFM. The value of Ra gradually became larger and the exact value of surface roughness at different RN2 is shown in Table 2. The heights of the needle structure in all the HEFs were less than 60 nm, which was consistent with the results since the as-deposited HEFs were smooth and homogeneous, as shown in SEM micrographs.

The surface nanostructures displayed by AFM images were different from each other and the size of the needle structures became larger with the increase of RN2. These nano-scaled needle structures were related to the structure zone models, which had a great impact on the microstructural evolution of thin film. Initially, as RN2 was low, the growth of the HEFs effect by nitrogen ions was limited. The nucleation barrier was generally expected to be small, which contributed to the formation of randomly oriented islands for films deposited on the silicon substrates without substrate heating. Accompanying the increase of RN2, a strong driving force was obtained, which facilitated the surface atom diffusion and grain boundary motion. To achieve a stable state, the overall surface and interface energy were minimized through forming the new islands. Islands with lower energy consumed the others. Thus, the roughness of the HEFs increased gradually with the increase of RN2.

#### *3.4. Mechanical Properties*

Figure 8a displays the hardness and Young's modulus of (Al0.5CrFeNiTi0.25)Nx films as a function of the N2 flow rate. Through the increase of the nitrogen content, the hardness and Young's modulus of the films showed a significant upward trend. The high-entropy thin films showed the highest hardness and Young's modulus at 21.78 GPa and 253.8 GPa, respectively. Compared with the bulk Al0.5CrFeNiTi0.25 alloys (the value of hardness was about 6 GPa), the hardness of the (Al0.5CrFeNiTi0.25)Nx HEFs was improved significantly. Considering the load-depth curves of Figure 8b, the manifest and precise change of the hardness and Young's modulus of the HEFs was obtained. When the probe depressed into the same depth, a larger force was required with the increase of the RN2. This was attributed to the fact that solid solution significantly increased the hardness. It is well known that solid solution strengthening increases the yield strength of the material by increasing the stress τ to move dislocations [42]:

$$
\Delta \tau = G b \epsilon^{3/2} \sqrt{c}
$$

where *c* was the concentration of the solute atoms, *G* was the shear modulus, *b* was the magnitude of the Burger's vector, and *ε* is the lattice strain due to the solute. While increasing the N2 flow rates, the parameters *c* and *ε* were increased correspondingly, which resulted in a higher *τ* value. Thus, the effect of solid-solution strengthening was enhanced gradually with the increase of RN2. When the RN2 = 40% and 50%, the nitrogen content in films increased slowly. The difference in nitrogen content was only 2.66%, as shown in Table 2. The mechanical properties of the HEFs became stable.

**Figure 8.** Mechanical properties of (Al0.5CrFeNiTi0.25)Nx films deposited at different RN2. (**a**) Hardness and modulus and (**b**) load-depth curve.

#### **4. Discussion**

It is well known that the solid solutions can be divided into three types: Substitutional solid solutions; interstitial solid solutions; and vacancy solid solutions. Hume–Rothery rules are a set of basic rules that describe the formation of substitutional and interstitial solid solutions. Regarding a substitutional solid solution, the atomic radius of the solute and solvent atoms must differ by no more than 15%, while solute atoms should have a radius no larger than 59% of the radius of the solvent atoms for an interstitial solid solution [43]. However, the non-metallic elements, such as H, B, C, and N, having a very small atomic radius can form compounds with metal elements. When the ratio of a non-metal atom radius to a metal atom radius (Rx/RM) is less than 0.59, the compounds have a simple structure such as M4X, M2X, MX, MX2, which is called an interstitial phase. When Rx/RM > 0.59, the compounds have a complex structure such as Fe3C, thus, there is a interstitial compound [44]. Regarding (Al0.5CrFeNiTi0.25)Nx, the metallic elements all satisfy the formation rule for an interstitial phase and different nitrides with simple structures can be formed. The metallic elements will tend to be ionic and the ionic radius might well reflect the actual particle. Since most of the elements in this alloy system are transitional metals, their ionic radii are very close to each other. Nitride components are becoming more similar in size and structure and solve each other easily. Thus, the FCC structured phase is formed in (Al0.5CrFeNiTi0.25)Nx.

The correlation between the composition, processing, microstructures, and the properties of high-entropy thin films can be summarized systematically and is shown in Figure 10. First, the films are prepared by a magnetron sputtering technique and different N2 flow rates are selected to adjust the composition as well as to change the microstructures. This finds that the nitrogen content in the thin films increases gradually with the increase of the nitrogen content and the structure of films transform from amorphous to FCC. Second, it is further found that the mechanical properties (hardness/Young's modulus) of the high-entropy crystalline films are much better than the high-entropy amorphous films. All told, there is an increasing trend of hardness with higher nitrogen for the high-entropy thin films. Third, due to the difference in the preparation process, the phase structure and the properties of the films are different from the bulk samples, although they are in the same composition system. This study can help enrich the cognition between composition processing, microstructures, and the properties for high-entropy materials, especially for high-entropy thin films.

**Figure 9.** AFM images and SEM micrographs of (Al0.5CrFeNiTi0.25)Nx films at different RN2.

**Figure 10.** Relationship between composition, processing, properties, and the structures of high-entropy films.

#### **5. Conclusions**

The (Al0.5CrFeNiTi0.25)Nx high-entropy thin films were deposited on silicon wafers through magnetron sputtering without substrate bias and heating. It was found that the phase structures of the HEFs transformed from an amorphous to an FCC structure with the increase in nitrogen content—the formation of a nitride solid solution was the main reason. Using the XRD analysis, the grain size was nanoscale at about 20 nm. The formation ability of the solid solution phase for the Al-Cr-Fe-Ni-Ti-N system film alloys was discussed regarding δ. When the value of δ was higher than 23%, the solid solution structures were stable in (Al0.5CrFeNiTi0.25)Nx thin films. The (Al0.5CrFeNiTi0.25)Nx HEFs deposited at RN2 = 40% and 50% yielded a maximum hardness and modulus of 21.78 GPa and 253.8 GPa, respectively, which is much higher than the as-cast Al0.5CrFeNiTi0.25 bulk alloys. The enhancement in hardness was mainly attributed to solid-solution strengthening and the lattice distortion. Additionally, smaller grain-size was a beneficial factor for increasing hardness.

**Author Contributions:** X.-H.Y. finished the preparation of high-entropy thin films. X.-H.Y. and W.-B.L. together carried out analyzing and writing. K.Z. provided equipment. Y.Z. offered the theoretical guidance.

**Funding:** Y. Z. would like to thank the financial supports from the National Science Foundation of China (NSFC, Granted Nos. 51471025 and 51671020). W.-B.L. gratefully acknowledges the financial support from the Natural Science Foundation of Shenzhen University (Grant No. 2017069) and the Shenzhen Basic Research Project JCYJ20170302142339007, and the Youth Innovation Talent Project of Guangdong Province (Grant No. 2017KQNCX175), and the National Natural Science Foundation of China (Grant No. 51801128).

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

#### **References**


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