Preparation and Microstructure of Multi-Component High Entropy Alloy Powders Fabricated by Gas Atomization Method

Abstract

As an attractive high-entropy alloy, AlCrCoNiCu high-entropy alloy has excellent corrosion resistance, wear resistance, and anti-bacterial capabilities, and is considered to be a potential substitute material for marine and nuclear industry materials with great potential. One key to further optimizing the performance of high entropy alloy was to prepare high entropy alloy powder materials with uniform composition, good flow-ability, and stable performance. In this work, the AlCrCoNiCu high entropy alloy powder was prepared by the gas atomization method. The results indicated that the powder was spherical in shape, homogeneous in composition, and composed of a face-center cubic (FCC) phase. After adding Fe and Mn elements, FCC and body-center cubic (BCC) phases appeared and the particle size of the powder was mainly located at 10–50 μm. Furthermore, the larger the particle size was, the more obvious the surface roughness was. With the decreasing powder size, its shape became relatively regular, and the surface roughness decreased. This work provided an experimental and theoretical reference for preparing high-performance single-phase and multi-phase high entropy alloy spherical powders.

1. Introduction

High entropy alloys (HEAs) have attracted great attention in past decades as a new class of metallic materials, which have great potential applications in the nuclear industry, aerospace and automotive industry owing to their good ductility and strength, excellent corrosion resistance, and high-temperature oxidation resistance properties [1,2,3]. In recent years, the research of high entropy alloys mainly focused on composition optimization or preparation of high entropy alloy blocks, powders and coatings based on different preparation methods [4,5,6,7]. In general, the properties of high entropy alloys were greatly affected by the solidification rate during the preparation process. Therefore, the preparation method exerted a critical effect on their microstructures, and subsequent performance [8,9,10]. Nowadays, powder metallurgy and surface technology were extensively applied to manufacture high entropy alloy block and coating materials, which could meet the needs of practical projects [11,12,13,14,15]. Nevertheless, both preparation methods required high entropy alloy powder materials with homogeneous composition, uniform dispersion and good flow-ability. As a result, it was recognized that the preparation of high entropy alloy powder with excellent performance had broad market prospects.

The properties of high entropy alloys were greatly affected by their composition. Therefore, composition control was a key way to tune the microstructure, so as to optimize their properties. In fact, high entropy alloy powders with different compositions might be used for laser repair of aluminum alloy cracks due to the mutual diffusion of Al element between AlCrCoNiCu high entropy alloy and aluminum alloy matrix, which thus formed a strong, wear-resistant and corrosion-resistant high entropy alloy protective coating [16]. Meanwhile, the FCC structure contributed to improving the mechanical plasticity of high entropy alloy. Recently, Stepanov et al. [17] found that the CoCrFeNiMn high entropy alloy with FCC structure possessed the equiaxed grains in an as-cast state, and the composition segregation was almost hard to occur during solidification. Moreover, the high entropy alloy with BCC structure usually had high strength, which has great potential in the application of high-strength materials [18]. For those high entropy alloy with the FCC + BCC structure, their strength could even be adequately maintained at temperatures up to 800 °C, which is also accompanied by good plasticity and great work hardening ability. Furthermore, their strength also exhibited an obvious positive temperature coefficient at a high strain rate and over alow temperature range [19]. Therefore, manufacturing high entropy alloy powder or block with FCC and BCC structure could further promote the applications of high entropy alloy. Generally, the high entropy alloys possessed many constituent elements, large differences in atomic radii, and different crystal structures, which made the atomic arrangement deviatefrom the equilibrium position. Thus, the phenomenon of irregular atomic arrangement was more prominent and there exerted a more serious lattice distortion in their lattice structure. When the distortion reached a certain degree, it led to phase transformation, which thus changed the alloy microstructure and even precipitated nanocrystals and amorphous. Based on these results, the addition of trace elements was an important way to regulate the phase transformation in the matrix [20,21,22,23].

This work focused on the preparation and microstructure observations of high entropy alloy spherical powder fabricated by the gas atomization method where the phase structure and content were regulated by adding strengthening elements Fe and Mn. The results indicated that the obtained single-phase AlCrCoNiCu spherical powder and multi-phase AlCrCoNiCuFe (Mn) spherical powder had a uniform composition, high performance, and good powder flow-ability. Moreover, the evolution mechanisms of microstructure and properties of powders with different compositions and particle sizes were further determined by using multiple characterization methods. The corresponding results revealed the relationship between the surface roughness of the sphere and the particle size, which thus provided an experimental and theoretical basis for further promoting the application of high entropy alloys.

2. Experimental

The high entropy alloy ingots were produced by vacuum arc melting according to the designed composition ratio with high-purity Al, Co, Cu, Ni, Cr, Fe, and Mn blocks materials (the specific composition was shown in

Table 1. Chemical compositions of 5 kinds of high entropy alloy ingots produced by vacuum arc melting. (at. %).

Powders	Al	Co	Cu	Ni	Cr	Fe	Mn
AlCrCoNiCu	20	20	20	20	20	-	-
1#	15.00	6.00	15.00	33.00	15.00	15.00	-
2#	6.00	9.00	34.00	22.00	7.00	-	22.00
3#	16.00	8.00	29.00	21.00	6.00	-	20.00
4#	15.00	6.00	15.00	34.00	15.00	15.00	-

). Before alloy melting, the vacuum chamber was filled with high-purity Ar gas (99.999% purity) with a pressure of 0.05 MPa as the protective gas. During the melting process, the alloy ingot was repeatedly turned over for at least three times, and electromagnetic stirring was added in the last two times to ensure uniform composition (Figure 1a,b). In order to explore the influence of Fe and Mn elements on the microstructure of high entropy alloy spherical powder, five kinds of high entropy alloy spherical powders were prepared by using plasma rotary electrode atomization pulverizing equipment (SL-ZFDW-04, Xi’an, China). Powder preparation was carried out in Ar or He protective atmosphere, and the oxygen increment of powder preparation was controlled under 100 ppm (Figure 1c,d). The alloy ingot was first melted into solution, followed by forming the small alloy droplets by gas impact, which can be cold condensed into powder under a vacuum environment. In order to obtain alloy powder with stable particle size distribution, the specific process parameters were determined as: current 700 A, voltage 60 V, rotational speed 23,000 rpm, and feed speed 2.0 mm/s. The microstructure and micro-area composition of the surface for the high entropy alloy spherical powder were characterized by optical microscope (OM), scanning electron microscopy (SEM), and energy dispersive spectroscopy (EDS). The microstructure evolution of the high entropy alloy spherical powder under different components and preparation technology was analyzed. The phase composition of the high entropy alloy powder was performed by X-ray diffraction (XRD). At the same time, the phase was verified by SEM and Transmission electron microscopy (TEM). It was noteworthy that the microstructure of the alloy powder was observed and analyzed at the nanoscale.

3. Results and Discussion

Under low-scale SEM observation as shown in Figure 2a, it can be found that the surface of AlCrCoNiCu high entropy alloy spherical powder was not smooth. There existed some grain structures or lattice structures. After the surface was magnified, it was obvious that small particles were observed on its surface as depicted in Figure 2b. Its fine particle structure is fine grains, and the average size of these grains is about 1 μm. This grain structure is mainly due to the centrifugal force of the high-entropy alloy electrode rotating at high speed when the plasma rotating electrode is atomized, the liquid is thrown out and crushed into fine droplets, and the rapid condensation process is formed. It can be seen from Figure 2b that there are still certain gaps between the grains, and the formation of these gaps is mainly caused by the volume shrinkage of the liquid phase during the rapid solidification process [24]. From the element distribution of a single sphere in Figure 2c, the five components of high entropy alloy spherical powder (Al, Cr, Co, Ni and Cu) were uniformly distributed in the microspheres, producing a homogeneous solid solution structure [25]. The contents of the five elements were listed in

Table 2. The real element content of five kinds of high entropy alloy spherical powder with determined by the SEM-ESD analysis (at.%).

Powders	Al	Co	Cu	Ni	Cr	Fe	Mn	C
AlCrCoNiCu	22.37	17.84	21.26	21.26	17.26	-	-	-
1#	8.55	0.08	18.63	44.47	13.59	14.66	-	-
2#	0.63	8.61	35.92	22.92	6.63	-	24.98	-
3#	1.55	5.90	30.38	17.61	5.64	-	21.87	17.06
4#	12.78	8.61	16.66	28.48	10.75	10.50	-	20.81

. It was seen that the ratios of Al, Cu, Ni elements were relatively higher, while the content of Cr and Co was relatively lower compared with the settings. The reason for this was mainly attributed to the high melting point of Cr and Co elements, which thus resulted in significant element segregation during the electron beam smelting process.

It can be seen from Figure 3b that the four kinds of powders added with Fe/Mn were also mainly FCC structures. However, the 1# specimen consisted of FCC and BCC phases, and the phase of 2# powder was a single FCC phase. In terms of phase constitution shown in Figure 3b, the strongest diffraction peak (111)_FCC intensity of FCC solid solution phase of 2# powder was lower than that of 1# powder, which was mainly attributed to the high cooling rate of small particles. To some extent, this inhibited the diffusion of atoms and increased the solid solubility of solute atoms. Finally, the transformation process of the alloy phase structure from BCC to FCC was accelerated. Therefore, the FCC solid solution phase content of 2# powder was more than that of 1#. Since the Al element of 3# powder increased, the (110)_FCC diffraction peak intensity of 3# powder was significantly higher than that of 2# powder and simultaneously indicated more BCC content. The main difference between 4# powder and 1 # powder was that the contents of Mn and Fe were different. The increasing Ni made the intensity of (110)_FCC diffraction peak becoming higher.

Figure 4 shows the overall surface SEM morphology of 1# and 2# powders and the partially enlarged view of larger particles. It can be seen from Figure 4a1,b1 that the particles of the two kinds of powders are relatively uniform, basically spherical, with good spherical regularity, indicating that they have good powder flowability [26]. There are satellite spheres on the surface of powder particles with larger particle sizes, and the larger the particle size, the rougher the surface and the more prone to defects. The fine particles bonded to the surface of the particles with larger particle sizes are mainly caused by the different solidification speeds of the powder. The difference in solidification speed between large and small powders may be due to the difference in volume. When the droplets have solidified, the large droplets may still be in a liquid or semi-solid state, and these small particles in the solidified state may be embedded on the surface of large particles when they collide with unsolidified large particles (Figure 4a2,b2). Another reason may be that there are more defects on the surface of larger powders, which can provide more nucleation sites for the subsequent formation of satellite spheres.

The parameters of four kinds of high entropy alloy powders prepared by the gas atomization method were verified by experiments and the corresponding particle size distributions were listed in

Table 3. Powder size distribution of four kinds of high entropy alloy powders prepared by gas atomization method.

Powders	Wt. %
	<10 μm	10~50μm	>50μm
1#	18.68	78.02	3.3
2#	4.29	84.28	11.43
3#	9.79	84.76	5.05
4#	22.06	76.29	1.65

. Figure 4c indicated the relationship between the cumulative mass distribution of the corresponding powder and the particle size. One can find that the particle size distribution range of the four kinds of powders was relatively concentrated. All particle sizes were less than 80mm and the powder size was mainly distributed in the range of 10~50 μm. Generally speaking, if the particle size was less than 100 mm, it could meet the requirements of coaxial powder feeding for laser additive manufacturing/remanufacturing [27]. Thus, the powder yield was high. As we all know, AlCrCoNiCu high entropy alloy spherical powder was mainly composed of FCC [28].

The surface morphology of 1# and 2# powders with the corresponding size range less than 10 μm, 10–50 μm, and lager than 50 μm were shown in Figure 5. The surface of the powder with a particle size range large than 50 μm was rough and uneven but had poor smoothness. Moreover, a large number of satellite-like powders were clearly observed. With the reduction of the powder size, there existed only a few satellite-like powders found on the surface of two kinds of powders for the range of 10–50 μm. At the size range of powders with less than 10 μm, the surface of the two kinds of powders was smooth with fine cellular crystals, while no satellite-like powder was observed. In general, with the decreasing powder size, the surface morphology of the 1# and 2# powders changed from coarse petal-like dendrites to fine cellular crystals. Further high magnification SEM observation on the surface of 3# and 4# powders was shown in Figure 5. The alloy powders were mainly spherical or quasi-spherical and the surface of the powders was relatively rough, which also accompanied by many small particles and droplet coating characteristics. Such phenomena could also be seen in other studies [29,30]. For individual powders with large particle sizes, the sphere collapsed and cracked, and meanwhile, degree of sphericity was very low. It was worth noting that the surface characteristics of powder were largely related to their size. The larger the particle size was, the more significant the surface roughness was. With the decrease in powder size, its shape becomes relatively regular and the surface roughness also significantly decreased. Nevertheless, some small powders were spherical in shape with smooth surfaces and few droplet coating features. Furthermore, there existed a large number of satellite-like powders in the alloy powder, which were formed in the process of gas atomization. The molten droplets collide formed a whole with the solidified powder. Therefore, the solidified alloy powder was completely or partially covered. It was worth noting that the larger the particle size was, the more obvious the coating characteristics were. The shape of the alloy powder was mainly affected by the initial shape of the molten droplet, the surface tension of the liquid, the pressure of the atomized gas, and the self-weight of the molten droplet. When the molten droplet was very small, the self-gravity and the atomized gas had a negligible influence on the shape of the powder. Since the surface tension of the molten droplet was the dominant factor, the powder particles indicated a regular spherical. The larger the particle diameter was, the larger the droplet volume was, and thus the lower the cooling rate was, which led to the more obvious solidification shrinkage and the rougher surface. As shown in Figure 5, 3# and 4# samples have uneven surfaces and many voids on the surface after magnification, which might be attributed to the high content of Fe and Ni elements and the high melting point. Such a case made the rapid solidification process shrink significantly. Not only the surface was rough, but also many voids occurred. As shown in Figure 5, the larger the powder size, the rougher its surface. One reason is that the large powder has a more complex gradient and surface area. During the solidification process, the surface is more likely to produce micro-elements and tissue bias. The other reason is that the large powder surface can provide more formal nuclear points. The temperature gradient inside and outside the powder affects the growth of grains. In fact, during powder preparation, gas impingement parameters affect powder formation and size distribution. In addition, the different components of the powder will also affect the uniformity of the elements of the powder. For example, the melting point of Ni element is higher, the condensation speed is slower, and Ni element segregation occurs during the condensation process.

As shown in Figure 6, 1# powder particles’ interior was mainly uniform gray when the particle size was less than 10 μm. For the particle size of 10–50 μm, the gray part became less and some light-colored areas began to appear. This color difference might be ascribed to the different phase contents. Nevertheless, when the particle size was greater than 50 μm, there existed many light-colored areas in the powder section and even dark cavities appeared. By contrast, one can see from Figure 6 that for 2# powder particles, no obvious change occurred in its interior whereas the characteristics of satellite-like powders were obviously observed as the particle size increased. With particle size greater than 50 μm, there also existed a few dark cavities. The element contents of four high entropy alloy spherical powders were listed in Table 3. One can find that the ratio of Al, Cu, Cr, and Co elements was relatively low, while the content of Ni was relatively high compared to the settings. The main reason for this was attributed to the fact that the alloy ingot was impacted by gas so as to form small alloy droplets and then was condensed into powder after when it was melted into alloy solution. During the condensation process, server composition segregation occurred owing to the high melting point of Ni element and the slow condensation speed. Figure 7 presented a further EDS surface scanning of 3# and 4# powders in order to determine the elemental composition of their surfaces. From Table 3, it was obvious that the C element interfered with the results to a certain extent, which thus resulted in the lower contents of all elements than the design. Under the condition of gas atomization, the rapid solidification process was involved in the large size of Al atoms, the high diffusion activation energy in the solid state, and the small crystallization latent heat. As a result, these led to the formation of non-equilibrium solidification of divorced eutectic structure. Assuming that the solid phase did not diffuse and the liquid phase was uniformly mixed during the solidification process, the average alloy element content of the solid phase might be lower than the nominal value of the alloy. Moreover, the solute elements were also pushed to the front of the liquid phase near the solid-liquid interface.

Figure 8a–d displayed the TEM and corresponding selected electron diffraction (SAED) calibration results of 4# powder. From the electron diffraction pattern, one can find that the dendrite region was an FCC solid solution phase rich in Co, Ni, and Cu, which was consistent with the results of the XRD examination. According to the results of energy spectrum surface scanning in Figure 8e, the Ni element was almost uniformly distributed in the sample and the distribution areas of Fe and Cr were consistent, which thus producing the dendrite-like morphology. The same phenomenon was also observed in reference [31,32,33]. Al element was distributed among dendrites and Cu element was mainly distributed at the interface around the Al element. Fe, Cr, and Ni elements as the main component were located at the dendrite region. Since the mixing enthalpy of Al and Ni elements was the lowest, they were enriched in inter-dendritic. The mixing enthalpy of Cu, Fe, and Cr elements was relatively higher than that of Al and Ni elements. Therefore, it was expelled from the dendrite region and was distributed near the Al-Ni region. In addition, it can be seen from Figure 9 that a small number of dislocation lines were found in the dendrite grains of the alloy (shown by the red arrows), which might be resulted from the stress concentration generated during gas atomization. Another reason for this was that the Al element intensified the lattice distortion effect of the high entropy alloy, thereby reducing the systematic stacking fault energy. For face-centered cubic (FCC) structure metals, no obvious twins were generated during tensile deformation at room temperature due to their sufficient slip system. The main slip plane of alloy dislocation movement was generally a lattice plane and the slip direction was <110> lattice direction.

Figure 10 shows the tensile engineering stress-strain curves of four samples which prepared by laser cladding on the ZL114A aluminum alloys using four powders with different elemental compositions. According to our previous research results, the ultimate strength of ZL114A aluminum alloy reached 305 MPa. As shown in Figure 10, it is evident that the performance of all samples has been improved and the ultimate tensile strength of the 2# sample increased by 30MPa compared with the base material (ZL114A aluminum alloy). The performance improvement of samples is mainly due to the addition of Mn and Fe elements. Solid solution strengthening of Mn and Fe elements and rapid cooling of laser cladding cause fine-grain strengthening to improve its strength. On the basis of multi-principal high-entropy alloys with a simple FCC solid solution phase structure, the addition of large-radius atoms Fe and Mn can solid-solution strengthen the alloy. Moreover, high-entropy alloys have many elements, large differences in atomic radii, and different crystal structures, which make the arrangement of atoms in the alloy deviate from the equilibrium position, and the phenomenon of irregular arrangement of atoms is more prominent, so there is serious lattice distortion in the structure. When the distortion reaches to a certain extent, it will lead to phase transformation of the alloy and change the microstructure of the alloy. Therefore, adding an appropriate proportion of Fe and Mn will promote the formation of the BCC phase, thereby improving the performance of the material. More detail about the tensile mechanical parameters were presented in

Table 4. The yield strength, ultimate strength, and elongation of four samples which prepared by laser cladding using four powders.

Samples	Yield Strength(MPa)	Ultimate Strength (MPa)	Elongation(%)
1#	288.5	325.5	6.4
2#	305.7	336.2	9.2
3#	291.3	328.0	7.4
4#	279.3	318.0	6.1

. The strength improvement of all samples indicated that the AlCoCuNiCr-based high-entropy alloys can be considered to be applied to structural damage additive remanufacturing forming materials of commonly used aluminum-based metal parts.

4. Conclusions

The multi-component high entropy alloy powder materials with uniform composition and good powder flow-ability were successfully prepared by the gas atomization method. Moreover, the microstructure evolution mechanism of powders with different compositions and particle sizes was revealed by various characterization methods. The results were drawn as follows:

(1)

The spherical powder of AlCrCoNiCu high entropy alloy was obtained, which was composed of FCC phase with uniform composition. On the basis of AlCrCoNiCu, dual-phase high entropy alloy spherical powder was prepared by adding trace Fe, Mn and other strengthening elements.

(2)

The powders possessed good fluidity and the particle size was mainly 10–50 μm. The characterization results showed that the surface roughness of powder particles was largely related to their size. The larger the particle size was, the more significant the surface roughness was.

(3)

With the decrease in powder size, its shape becomes relatively regular, and the surface unevenness decreased. Some particles with larger particle sizes were bonded with fine particles mainly due to the different cooling speeds of the powder. The larger the particle size was, the more obvious the coating characteristics were.