Microstructural Evolution of High-Entropy Intermetallic Compounds during Detonation Spraying

Abstract

This study aims at investigating the feasibility of depositing quality coatings from various high-entropy intermetallic compounds (HEICs) using detonation spraying (DS). Four different HEIC coatings, namely (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, (NbTaVZrHf)Al₃, and (FeNiCoCrMn)(MoCr), were prepared by DS on low alloy steel substrates. The HEIC powders were first prepared by arc melting followed by ball milling and then used as reinforcement particles to deposit HEIC coatings. Elemental segregation was observed for all the as-cast samples. Powders with average particle sizes of about ~25 µm for (NbTaVCrTi)Al₃, ~22 µm for (NbTaVNiFe)Al₃, ~34 µm for (NbTaVZrHf)Al₃, and ~18 µm for (FeNiCoCrMn)(MoCr) were obtained. (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, and (NbTaVZrHf)Al₃ HEICs exhibited a nearly single D0₂₂ (TaAl₃ type) structure, while (FeNiCoCrMn)(MoCr) exhibited a single D8_b (FeCr type) structure. Dense coatings consisted of a lamellar microstructure and sound bonding with the substrate, and low porosity was obtained for all the samples. Crystal structures of the HEIC samples were highly retained during DS, whereas all the samples underwent some degree of oxidation. Microhardness values of 745 HV for (NbTaVCrTi)Al₃, 753 HV for (NbTaVNiFe)Al₃, and 862 HV for (NbTaVZrHf)Al₃ were obtained, which are significantly higher than the microhardness of the substrate (~140 HV). Among all the samples, (FeNiCoCrMn)(MoCr) exhibited the highest microhardness values of about 1047 HV.

1. Introduction

In contrast to conventional alloys, which are typically based on a single principal element and add a limited amount of alloying elements, the proposed alloy design technique in multi-principal element alloys provides a wide chemical window to tailor microstructure and mechanical properties of the materials [1]. These namely high-entropy alloys (HEAs), which are often composed of five or more elements in an equiatomic or near equiatomic ratio and offer a broad chemical composition range to create unique materials, in contrast to the conventional alloying design. Previous studies already demonstrated technologically important properties such as unprecedented toughness [2], high tensile strength [3], good corrosion and oxidation resistance [4], and superior strength-to-weight ratios [5] in HEAs.

Following by the establishment of the concept of high entropy in solid solution alloys, several subclasses of high-entropy materials (HEMs) such as high-entropy ceramics [6], high-entropy glasses [7], and high-entropy intermetallic compounds (HEICs) [8,9] were developed. The latter are simply HEAs, in which the constituent elements occupy distinct lattice sites creating the ordered superlattice of intermetallic compounds. They combine the ordered superlattice of intermetallics and the entropy stabilization effects in a single material, which can potentially create structure–function integrated materials with superior properties compared to disordered HEAs. Recently, several HEICs with interesting properties have been developed, which are reviewed elsewhere [10]. To name a few, ZrTiHfCuNiFe and Fe_0.75Co_0.75Ni_0.75Cu_0.75TiZrHf HEICs with B2-type ordering and high compressive yield strengths of about 2 GPa have been developed [11,12]. It was also reported that the Co₂₅Ni₂₅(HfTiZr)₅₀ HEIC exhibits a high elastic strain of 2% [13]. Moreover, this HEIC shows the Elinvar effect, exhibiting an almost invariant Young’s modulus between room temperature and 627 °C [13]. Al_0.25FeCoNiV is another HEIC with L1₂-type ordering that exhibits an ultimate tensile strength of 1530 MPa and fracture strain of 20% [14]. A high tensile strength of 1611 MPa and a fracture strain of 25% have been also reported for Ni_43.9Co_22.4Fe_8.8Al_10.7Ti_11.7B_2.5 multicomponent intermetallics by tailoring grain–boundary segregation and phase precipitation [15]. Additionally, Xiao et al. [16] designed an L1₂-type Co-Ni-Al-Ti-Ta-Nb-B-based multicomponent intermetallic with a lamellar structure, which demonstrated a yield strength of ~1.0 GPa and a tensile elongation of ~17% at room temperature.

In spite of the above-mentioned progress, a lack of ductility and a higher cost of HEMs compared to conventional alloys restrict their widespread application. Developing appropriate coating techniques may resolve this dilemma to some extent. In this regard, thermal spray techniques, such as atmospheric plasma spraying and detonation spraying (DS), have already shown promising prospects in fabricating HEM coatings [17,18,19,20]. Both HEA and HEIC coatings have been prepared. For example, Liao et al. [21] deposited CoCrFeNiMn HEA coatings on a 316L stainless steel substrate using the DS technique. The coatings exhibited a dense microstructure with a small amount of flocculent metal oxides and nano-multicomponent metal oxides and demonstrated high bond strength and abrasive wear resistance. Moreover, high-entropy (VCrMoWNi)B-Ni₂B composite coating has been fabricated by plasma spraying, which exhibited good wear resistance owing to the cooperation effect between the hardness of (VCrMoWNi)B and large plastic deformation ability of Ni₂B [22]. Furthermore, it has been also demonstrated that the plasma-sprayed (YYbLuEuEr)₃Al₅O₁₂ high-entropy oxide coating is effective in protecting the nickel-based superalloy [23]. The research on the preparation of coatings with superior multifunctional properties is ongoing [24,25,26,27,28].

Furthermore, the effect of post-deposition treatment processes such as heat treatment, laser melting, and spark plasma sintering on the microstructure and mechanical properties of HEA coatings prepared by DS has been also studied. We et al. [29] reported that annealing at 900 °C decreased the microhardness of CoCrFeNiMn HEA coatings, but at the same time increased their wear resistance due to the formation of strong metallurgical bonding between the splats and the ensuing improvement in the cohesive strength of the splats. Batraev et al. [30] also reported that CoFeNi, CoCuFe_0.8Ni, and Al_0.05CoCu_0.3FeNi coatings with a uniform FCC structure can be obtained by the detonation spraying of metal powder blends and subsequent annealing.

In addition to HEAs, the fabrication of HEIC coatings by detonation spraying has been also demonstrated [31,32]. Our group already fabricated several HEIC coatings such as (CoFeCuNiMn)Al with a single B2-ordered structure [31], (NiCoFeCuMn)Zn₃ with a single γ-brass D82-type structure [32], and (CuFeNiTi)₃Al with a double B2 + FCC phase structures [31]. All coatings exhibited a lamellar microstructure with low porosities and strong bonding with the substrate. Considering the promising characteristics of HEICs, such as exceptional oxidation resistance, high microhardness, and thermal stability, the behavior of HEICs during DS needs to be studied more. Moreover, developing effective coating processes can overcome the problems related to the higher cost of HEICs compared to conventional materials. Therefore, in this study, we prepared several novel HEIC coatings and studied their microstructure and mechanical properties to further unravel the potential of this promising class of materials for practical applications. DS is a thermal spraying technique that is characterized by a fast heating of the particles to 1000–1500 K and accelerating them to the highest velocity of 800–1200 ms⁻¹ by detonation wave. Compared with other thermal spray techniques, DS coatings typically exhibit a uniform and dense microstructure with high hardness and strong bonding to the substrate owing to the sufficiently high velocity of the particles upon hitting the substrate. This technique is, therefore, suitable to deposit quality coatings from various powder materials. To demonstrate the feasibility of DS for preparing HEIC coatings with different compositions, four distinct compositions were utilized as feedstock powder. Particularly, two groups of HEICs (D0₂₂ aluminides and sigma phases) were selected for preparing coatings based on previous studies where high thermal stability, oxidation resistance, and microhardness values were noted for these two groups.

2. Materials and Methods

2.1. Synthesis of HEIC Powders

(NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, (NbTaVZrHf)Al₃, and (FeNiCoCrMn)(MoCr) HEICs were fabricated by arc melting under argon atmosphere. The powder mixtures were first prepared by weighting and mixing high-purity (>99.9 wt.%, Alfa Acer Company, Karlsruhe, Germany) elemental powders in an agate mortar. The mixed powders were then placed in a copper mold equipped with a water cooling system and melted by a 5SA arc melting furnace (TORRES LLC, Ekaterinburg, Russia) five times to ensure chemical homogeneity. The samples were flipped over between the melting passes.

To obtain HEIC powders, the solidified bottoms were crushed, ground, and subsequently ball milled for 1 h using a high energy ball mill (XQM-0.2S, Changsha Tianchuang Powder Technology Co., Ltd., Changsha, China). The powders were loaded into a Zirconia jar filled with Zirconia balls with a powder-to-ball ratio of 5:1, and the milling was carried out at a rotation speed of 600 rpm. The cups were sealed under an Ar atmosphere to protect the powders from oxidation. The final powders for DS were obtained by passing the milled powders from a 200-mesh sieve.

2.2. DS Coating

A commercial detonation spraying gun CCDS2000 (Siberian Technologies of Protective Coatings LLC, Novosibirsk, Russia) was employed for the DS process (Figure 1). The spraying distance was 270 mm, the feed rate was ~0.4 g/shot, and the shot frequency was 3 shots/second. The explosive gas was a mixture of C₃H₈ + C₂H₂ + O₂ (Chelyabtekhgaz, Chelyabinsk, Russia) along with N2. These parameters were determined by a trial-and-error approach. A low-alloy steel with a composition of (wt.%) 0.34% C, 0.38% Si, 0.68% Mn, 1.47% Cr, 1.53% Ni, 0.25% Mo, 0.017% S, <0.035% P, and bal. Fe was used as the substrate.

2.3. Characterization

The crystal structure of the powders and DS coatings was characterized by X-ray diffraction (XRD, Rigaku Ultima IV, Rigaku, Tokyo, Japan) using Cu–Kα radiation. The XRD patterns were collected in the 2θ range of 5 to 95 degrees at a scan rate of 2 deg/min. For DS coatings, the top surface was polished for XRD measurements and the cross-section was polished for microstructural observations. First, the surface was polished to a mirror state by a polishing–grinding machine (EcoMet/AutoMet 250, Buehler, Leinfelden-Echterdingen, Germany). For polishing, the samples were mounted inside a bakelite resin and were only exposed to the desired surface. To obtain a mirror finish, P1500 abrasive paper (~9 μm), Al₂O₃ slurry (~3 μm), and polishing cloth were utilized. The microstructure of the samples was analyzed by an scanning electron microscope (SEM, Jeol JSM7001F, Tokyo, Japan) equipped with energy-dispersive X-ray spectroscopy (EDS; Oxford INCA X-max 80, Oxford Instruments, Abingdon, UK) for chemical analysis. For each measurement, EDS was collected from 5 points, and the average values were reported. The level of porosity in the coatings was determined by Thixomet Pro 5.3 software. Porosity was estimated as the ratio of the total area occupied by pores to the entire area of the micrographs obtained at a magnification of 100×. An FM-800 microhardness tester (Future-Tech Corp, Kawasaki, Japan) was employed to measure the microhardness changes along the coating/substrate interfaces. The measurements were conducted with a load of 50 g and a holding time of 10 s on the polished surfaces. For each measurement, the microhardness was recorded at five locations, and the average value was reported.

3. Results and Discussion

3.1. Microstructure of HEICs

The microstructure of the as-cast HEIC samples was studied on SEM in back-scattered electron (BSE) mode to distinguish elemental segregation based on their atomic number. The BSE-SEM images and the corresponding EDS maps are shown in Figure 2. A dendritic microstructure with clear elemental segregation can be observed for all the as-cast HEICs. For (NbTaVCrTi)Al₃, three distinct regions, bright contrast, gray contrast, and dark gray contrast, can be distinguished in the BSE-SEM micrograph. EDS maps indicate that the bright contrast region is enriched in Nb and Ta heavy elements, the dark gray region is Al- and Cr-rich, and the gray region is rich in Ti and V. The same elemental segregation can be also observed in the microstructure of (NbTaVNiFe)Al₃. Na and Ta were segregated into bright contrast regions, Fe and Ni were mainly partitioned into gray contrast regions, and V was segregated around Nb-Ta (bright contrast regions), while Al indicated a relatively uniform distribution with minor segregation into dark gray regions. For (NbTaVZrHf)Al₃, Nb, Ta, and Zr were segregated together in the dendritic arms (bright gray regions), V, and to a lower extent Al, were segregated into inter-dendritic regions (gray regions), and Hf indicated a relatively uniform distribution between dendritic and inter-dendritic areas. Furthermore, the microstructure of the (FeNiCoCrMn)(MoCr) HEIC also indicated the segregation of Mo and Ni/Mn (minor), while other elements exhibited a nearly uniform distribution. The chemical composition of these distinct microstructural regions is shown in

Table 1. The chemical composition of different microstructure features is indicated in Figure 2 for the as-cast HEICs.

		EDS Chemical Composition (at. %)
HEICs		Al	Nb	Ta	V	Ti	Cr	Zr	Hf	Ni	Fe	Mn	Co	Mo
(NbTaVCrTi)Al3	D	73.98 ± 1.2	13.63 ± 0.6	7.58 ± 0.2	1.11 ± 0.03	3.10 ± 0.05	0.59 ± 0.04
	ID1	83.18 ± 1.4	0.08 ± 0.01	0.11 ± 0.01	1.54 ± 0.01	0.26 ± 0.03	14.83 ± 0.3
	ID2	74.05 ± 1.1	0.93 ± 0.02	1.63 ± 0.04	6.86 ± 0.1	5.75 ± 0.2	10.78 ± 0.4
(NbTaVNiFe)Al3	D	73.94 ± 1.2	12.05 ± 0.4	10.95 ± 0.5	2.01 ± 0.01					0.35 ± 0.01	0.69 ± 0.04
	ID1	75.00 ± 0.9	0.09 ± 0.01	0.09 ± 0.01	1.01 ± 0.01					5.53 ± 0.1	18.27 ± 0.5
	ID2	91.96 ± 1.0	0.17 ± 0.04	0.12 ± 0.03	0.27 ± 0.02					0.17 ± 0.03	0.11 ± 0.02	O = 7.20 ± 1.5
(NbTaVZrHf)Al3	D	75.13 ± 0.8	5.36 ± 0.2	6.37 ± 0.3	2.77 ± 0.06			5.26 ± 0.2	5.11 ± 0.2
	ID1	73.81 ± 1.2	2.44 ± 0.1	2.99 ± 0.1	12.09 ± 0.5			4.10 ± 0.1	4.57 ± 0.1
	ID2	81.58 ± 0.9	0.45 ± 0.04	1.20 ± 0.02	7.17 ± 0.3			1.71 ± 0.05	2.18 ± 0.1			O = 5.72 ± 1.3
(FeNiCoCrMn)(MoCr)	D						34.65 ± 0.9			14.01 ± 0.5	11.09 ± 0.4	9.51 ± 0.2	11.70 ± 0.3	19.04 ± 0.6
	ID						30.35 ± 1.0			18.20 ± 0.7	10.08 ± 0.1	13.31 ± 0.2	13.14 ± 0.4	14.92 ± 0.5

.

Previous studies suggested that the fundamental characteristics of constituent elements such as the melting point, atomic size, electronegativity, and the enthalpy of the mixing of binary pairs are possible factors for determining elemental segregation in HEICs [33]. Obviously, elements with the highest melting point segregate from melt during the first stages of solidification, forming dendritic arms, while elements with a low melting point mainly remain in melt until the last stages of solidification, enriching inter-dendritic regions upon solidification. However, this trend may be disrupted by the atomic size of the elements and the chemical affinity and enthalpy of the mixing of binary elemental pairs. The segregation of Nb and Ta in all HEIC samples can be attributed to their higher melting points compared to other constituent elements. It seems that segregation in (NbTaVCrTi)Al₃ is highly dominated by the melting point of the elements. Moreover, the segregation of Ti near Nb-Ta dendritic arms can be attributed to its positive enthalpy of mixing with Ta (1 kJ/mol) and Nb (2 kJ/mol) [34]. Similarly, the partial segregation of Al with Ni in (NbTaVNiFe)Al₃ may be ascribed to the more negative enthalpy of the mixing of Al with Ni (−22.3 kJ/mol) compared to other elements [34]. Apparently, the same arguments can explain the segregation of V in the inter-dendritic regions (NbTaVZrHf)Al₃. Finally, the more homogenous microstructure of (FeNiCoCrMn)(MoCr) HEIC is attributed to the proper combination of elements occupying distinct sublattices in the crystal structure. However, the minor segregation between Mo and Ni/Mn may be ascribed to the positive enthalpy of mixing of Mo with Mn (4.9 kJ/mol) and the more negative enthalpy of the mixing of Mn and Ni (−8.2 kJ/mol) [34] compared to other binary pairs in this system.

The secondary electron SEM images of the as-milled HEIC powders are shown in Figure 3. Particles with irregular morphologies, which is a characteristic of ball-milled powders, can be observed for all the samples. The average particle size was estimated by ImageJ (National Institutes of Health and the Laboratory) considering a total number of about 100 particles for each sample. The obtained average particle size was ~25 µm for (NbTaVCrTi)Al₃, ~22 µm for (NbTaVNiFe)Al₃, ~34 µm for (NbTaVZrHf)Al₃, and ~18 µm for (FeNiCoCrMn)(MoCr) HEIC powders. While some fractions of fine particles can be observed for all samples, the majority of the particles fall in a narrow range around the reported average particle size. (FeNiCoCrMn)(MoCr) exhibited the smallest particle size among all the ball-milled powders in this study, which may be attributed to its higher microhardness compared to other HEICs (see below). Some levels of agglomeration can be also detected for all the powders. It should be also noted that ball milling for only 1 h did not induce any changes in the microstructural segregation observed in the as-cast samples.

Figure 4 shows the XRD results of the HEICs in the as-cast and DS-deposited states. The as-cast (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, and (NbTaVZrHf)Al₃ exhibited a D0₂₂ HEIC phase (TaAl₃-based structure, space group I4/mmm, COD database code 4312726) with a minor amount of Al₂O₃ (COD database code 9009675) and TaO (COD database code 1538733) oxides. Considering the binary TaAl₃ structure, it can be depicted that five different elements have been successfully replaced by Ta in the D0₂₂ lattice of TaAl₃. On the other hand, the crystal structure of (FeNiCoCrMn)(MoCr) could be indexed with a single D8_b-FeCr-type (σ-phase, space group P42/mnm, COD database code 9016031) structure.

3.2. DS Coatings

The XRD patterns of the DS coatings (Figure 4) show that (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, and (NbTaVZrHf)Al₃ HEICs highly retain their D0₂₂ structure, while the amount of oxide phases significantly increases after the DS process. As explained below, these oxide phases are mainly formed during the flight of the particles toward the substrate, shattered on the surface upon hitting the substrate, and then distributed within the microstructure of the coatings. XRD indicated Al₂O₃ and TaO as the main oxide phases, reflecting the higher oxidation of these elements by an explosion wave. (FeNiCoCrMn)(MoCr) coating also retains its crystal structure after DS; however, some level of oxidation is also clear for this sample. The oxidation of powders typically occurs during thermal spraying techniques [35]. Moreover, a clear peak broadening can be observed for all the samples. XRD peak broadening may be due to the lattice strain and/or decrease in the crystallite size. Both strain broadening and crystallite size-related broadening are possible phenomena during the DS process. It should be also noted that no other intermetallic phases could be detected, indicating good thermal stability for all the samples.

The coating/substrate cross-section was characterized by BSE-SEM for bonding quality and possible chemical reactions at the interfaces. As shown in Figure 5, all DS coatings indicate a sound bonding with a low alloy steel substrate, where the coating thickness varies from 100 to 150 µm. Moreover, no significant cracks, pores, or other defects can be observed at the coatings/substrate interfaces and within the coatings. However, some dark particles can be detected at the interfaces, especially within the substrate side. Some dark regions can be also detected within the coatings, especially for (NbTaVCrTi)Al₃ and (NbTaVNiFe)Al₃ coatings. On the other hand, in (NbTaVZrHf)Al₃ and (FeNiCoCrMn)(MoCr), neither interfaces nor the coatings indicate significant dark particles. The EDS maps and EDS point chemical analysis (

Table 2. The chemical composition of different microstructure features is indicated in Figure 6 for the DS coatings.

		EDS Chemical Composition (at. %)
HEICs	Point	Al	Nb	Ta	V	Ti	Cr	Zr	Hf	Ni	Fe	Mn	Co	Mo	O	Por. (%)
(NbTaVCrTi)Al3	1	75.70 ± 1.4	12.42 ± 0.7	6.06 ± 0.1	1.52 ± 0.03	3.51 ± 0.06	0.80 ± 0.04									0.4
	2	55.22 ± 1.0	0.12 ± 0.03	0.16 ± 0.01	0.05 ± 0.01	1.60 ± 0.04	0.59 ± 0.03								42.26 ± 2.7
	3	37.42 ± 1.1	1.92 ± 0.06	1.84 ± 0.02	1.72 ± 0.04	3.29 ± 0.07	2.04 ± 0.1								51.76 ± 2.8
	4	63.84 ± 1.3	5.47 ± 0.4	6.16 ± 0.3	5.11 ± 0.2	7.36 ± 0.3	7.84 ± 0.7								4.22 ± 1.0
(NbTaVNiFe)Al3	1	27.29 ± 1.1	11.51 ± 0.4	12.66 ± 0.3	11.29 ± 0.7					19.67 ± 1.5	17.60 ± 1.3					0.32
	2	49.03 ± 1.2	0.22 ± 0.01	0.19 ± 0.02	0.27 ± 0.02					0.33 ± 0.01	0.37 ± 0.07				49.59 ± 3.1
	3	77.23 ± 1.5	7.05 ± 0.3	4.76 ± 0.2	2.49 ± 0.1					4.10 ± 0.3	4.36 ± 0.2
	4	20.03 ± 0.9	7.57 ± 0.1	20.38 ± 1.1	6.39 ± 0.1					15.96 ± 1.0	7.89 ± 0.5				21.79 ± 2.9
(NbTaVZrHf)Al3	1	49.54 ± 1.0	0.47 ± 0.03	0.49 ± 0.01	0.84 ± 0.04			0.67 ± 0.05	0.89 ± 0.04						47.10 ± 2.7	0.30
	2	29.32 ± 1.3	19.24 ± 0.5	17.12 ± 0.7	10.62 ± 0.8			5.91 ± 0.2	4.47 ± 0.2						13.32 ± 1.8
	3	69.36 ± 1.7	5.58 ± 0.1	4.73 ± 0.2	4.50 ± 0.1			4.36 ± 0.3	4.12 ± 0.1						7.36 ± 1.0
(FeNiCoCrMn)(MoCr)	1						8.21 ± 0.6			34.44 ± 1.4	18.71 ± 1.3	7.81 ± 0.6	24.11 ± 1.4	4.28 ± 0.7	2.44 ± 0.7	0.18
	2						35.76 ± 1.3			0.67 ± 0.06	4.27 ± 0.2	3.35 ± 0.3	1.3 ± 0.05	54.65 ± 1.76
	3						18.64 ± 1.1			5.09 ± 0.1	5.62 ± 0.3	2.57 ± 0.1	4.82 ± 0.2	9.77 ± 0.7	53.49 ± 2.5
	4						1.67 ± 0.02			7.40 ± 0.2	2.26 ± 0.1	0.3 ± 0.05	6.72 ± 0.3	53.65 ± 1.2	28.01 ± 2.0

) indicated that oxide particles are rich in Al and Ta, which were also determined by XRD analysis to be Al and Ta oxides for (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, and (NbTaVZrHf)Al₃. For (FeNiCoCrMn)(MoCr), a few discrete oxide particles observed at the interface in Figure 5 could not be detected by XRD, probably due to their low content and non-homogenous distribution. In spite of this minor oxidation, all the coatings exhibited dense lamellar microstructures.

It is known that during thermal spraying techniques, the particles experience three different states of fully molting, semi-molting, and non-melting [36]. Detonation waves can heat the particles to about 1000–1500 K in a short span of time and create thermally softened particles moving toward the substrate at a speed of about 800–1200 ms⁻¹ [21]. However, not all the particles are fully melted during DS because the size of the HEIC particles is not uniform and the heat input of the explosion wave is short and non-uniform. Therefore, typically particles with two morphologies of closely packed semi-molten and well-flattened fully molten can be observed on the surface of coatings. Figure 6 shows the microstructural features of the HEIC coatings. For all DS coatings, the microstructure is dominated by a large number of lamellar structures, and it can be observed that many oxide particles are embedded. It should be noted that all DS coatings demonstrated remarkable density and compactness with a low level of porosity (Table 2).

Meanwhile, several distinct regions can be detected in the microstructure of the HEICs deposited by DS. For (NbTaVCrTi)Al₃, dark, bright, gray, and light gray regions can be observed. The bright, gray, and light gray regions are similar to those observed in the as-cast sample, and the dark regions are oxide phases formed during the DS process. There are a lot of bright flocculent structures and small fragments within the microstructure of the (NbTaVCrTi)Al₃ HEIC coating. These are dendritic arms that were shattered upon hitting the substrate at high speeds and redistributed within the microstructure as discontinuous layers and randomly scattered fine fragments. It can be depicted that there was a smaller amount of fully molten particles than semi-molten particles in the HEIC coating, and these fully molten particles experienced more severe oxidation than semi-molten particles owing to accelerated diffusion at higher temperatures. At the same time, the thin oxide layers formed on the surface of particles are also shattered upon hitting the substrate and distributed within the microstructure as fine dark particles. The same microstructure features can be also observed for the (NbTaVNiFe)Al₃ and (NbTaVZrHf)Al₃ coatings, but the content of dark oxide regions highly decreases. The (FeNiCoCrMn)(MoCr) coating also indicated similar microstructure features to that of the as-cast sample in addition to the dark oxide regions. The dark oxide region is rich in Cr and Mo (Table 2), indicating preferential oxidation of these elements during DS. Moreover, the content of manganese decreases after DS, while other elements remain in nearly the same ratios, which may be attributed to its lower evaporation temperature compared to other elements [21]. Additionally, for all HEIC coatings, EDS maps indicated a more uniform distribution of the elements compared to the as-cast samples, which can be ascribed to microstructural refinement and elemental diffusion during DS. Furthermore, oxygen enrichment could be observed in nearly all areas (Table 2).

The microhardness of the coatings was measured along the cross-section of the coatings perpendicular to the coating/substrate interface (Figure 7). The substrate (low alloy steel in annealed condition) exhibits a relatively low microhardness value of about 130 HV. All of the DS coatings demonstrated significantly higher microhardness values compared to the substrate. It can be seen in Figure 7 that the microhardness of (NbTaVCrTi)Al₃ changes from 680 to 790 HV, with an average microhardness of 745 HV. (NbTaVNiFe)Al₃ indicated microhardness values that changed from a minimum of 719 HV near the substrate to a maximum of 834 HV in the outer regions, with an average value of 753 HV. The obtained microhardness values for (NbTaVNiFe)Al₃ are very close to those of (NbTaVCrTi)Al₃. On the other hand, (NbTaVZrHf)Al₃ demonstrated higher microhardness values compared to (NbTaVNiFe)Al₃ and (NbTaVCrTi)Al₃, which fluctuated from 829 HV to 894 HV, with an average value of 862 HV. Finally, the microhardness values of (FeNiCoCrMn)(MoCr) lie above all other coatings, indicating a narrow range of 999 HV to 1094 HV, with an average value of 1047 HV. These measured microhardness values, especially (FeNiCoCrMn)(MoCr) and (NbTaVZrHf)Al₃, are significantly higher than Al_0.25CoCrFeNiSi_0.6 (~550 HV) [35] and CoCrFeNiMn (~520 HV) [29] HEA coatings, as well as higher than the Ni-7 wt.% FeNiCrV-TiNb composite coating [33] fabricated by DS.

Typically, fluctuations in the microhardness values measured along the cross-section of the coatings/substrate are observed, which can be ascribed to the segregation, oxidation, and the existence of distinct microstructure features in the DS coatings. Here, the microhardness graph indicates differences in microhardness values between four materials. Moreover, all the coatings exhibited narrow fluctuations in the microhardness values in spite of the clear segregation and the presence of distinct microstructural features in the coatings. This may be ascribed to the use of fine powder feedstocks and microstructure refining during the DS process. A hardening area (~20 µm) can be also observed near the coating/substrate interface on the substrate side of all the samples. This is due to the work of hardening the hitting particles on the substrate surface. Furthermore, no softening area near the coating/substrate was observed in the samples. This region is typically observed in laser cladding or other coating techniques, which involve melting and intermixing of the coating and substrate at the interface. Finally, it can be understood that the microhardness of HEICs is mainly determined by crystal structure, atomic bonding, and microstructural features, rather than solely chemical composition. This can be understood from the higher microhardness value of (FeNiCoCrMn)(MoCr) compared to other aluminides in this study.

HEICs typically provide superior thermal stability and mechanical properties, especially at high temperatures. However, the lack of ductility makes it challenging to fabricate bulk components from HEICs. Therefore, fabricating coatings can provide an alternative approach to utilizing the technologically important properties of HEICs. HEIC coatings are particularly interesting candidates for protecting components working at high temperatures, such as jet engines and power plants. However, the interactions between coatings/substrates and the performance of HEIC coatings at high temperatures need to be studied further.

4. Conclusions

The microstructural evolution and microhardness of (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, (NbTaVZrHf)Al₃, and (FeNiCoCrMn)(MoCr) HEICs during detonation spraying was studied. The following conclusions can be drawn:

• The as-cast (NbTaVCrTi)Al₃, (NbTaVNiFe)Al₃, and (NbTaVZrHf)Al₃ indicated clear segregation with a nearly single D0₂₂ (TaAl₃ type) structure, while (FeNiCoCrMn)(MoCr) exhibited a lower degree of segregation and a single D8_b (FeCr type) structure;
• In all HEICs, Nb and Ta segregated together in the dendritic arms, which could be attributed to their higher melting points compared to other constituent elements. The segregation of other elements could be also described by the corresponding enthalpy of mixing of binary pairs;
• All the HEIC coatings fabricated by detonation spraying retained their crystal structures and exhibited a dense lamellar microstructure consisting of different features and segregated regions with oxide particles embedded. All the coatings indicated sound bonding with the substrate and a low level of porosity;
• (FeNiCoCrMn)(MoCr) exhibited the highest microhardness values of 1047 HV among all the HEIC coatings. The microhardness values were 745 HV for (NbTaVCrTi)Al₃, 753 HV for (NbTaVNiFe)Al₃, and 862 HV for (NbTaVZrHf)Al₃.