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Article

Structure and Properties of NbMoCrTiAl High-Entropy Alloy Coatings Formed by Plasma-Assisted Vacuum Arc Deposition

by
Yurii F. Ivanov
1,
Yuriy Kh. Akhmadeev
1,
Nikita A. Prokopenko
1,*,
Olga V. Krysina
1,
Nikolai N. Koval
1,
Elizaveta A. Petrikova
1,
Oleg S. Tolkachev
1,
Vladimir V. Shugurov
1,
Vladimir V. Uglov
1,2 and
Alexander N. Shmakov
1,3
1
Institute of High Current Electronics SB RAS, 2/3 Akademichesky Ave., 634055 Tomsk, Russia
2
Department of Solid State Physics, Belarusian State University, 4 Nezavisimosti Ave., 220030 Minsk, Belarus
3
Budker Institute of Nuclear Physics SB RAS, 11 Akademika Lavrentieva Ave., 630090 Novosibirsk, Russia
*
Author to whom correspondence should be addressed.
Coatings 2023, 13(7), 1191; https://doi.org/10.3390/coatings13071191
Submission received: 31 May 2023 / Revised: 27 June 2023 / Accepted: 29 June 2023 / Published: 2 July 2023
(This article belongs to the Section Surface Characterization, Deposition and Modification)
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Versions Notes

Abstract

:
The paper analyzes the structure and properties of metal, cermet, and ceramic NbMoCrTiAl high-entropy alloy (HEA) coatings formed on solid substrates by plasma-assisted vacuum arc deposition (from multicomponent gas-metal plasma through Nb, Mo, Cr, and TiAl cathode evaporation in argon and/or a mixture of argon and nitrogen). The analysis shows that all coatings represent a nanocrystalline (3–5 nm) multilayer film. The metal coating has a bcc lattice (a = 0.3146 nm). The ceramic coating has an fcc lattice (an uncertain lattice parameter due to highly smeared diffraction peaks). The coating hardness increases in the order of metal, cermet, and then ceramic, reaching 43 GPa at Young’s modulus equal to 326 GPa. When heated in air, the metal and cermet coatings start to oxidize at 630–640 °C, and the ceramic coating at 770–780 °C.

1. Introduction

High-entropy alloys (HEAs) and their carbides, nitrides, oxides, etc. [1,2,3,4] feature a unique combination of properties unattainable in most modern materials: high corrosion resistance [5,6], good mechanical characteristics at high [7] and low temperatures [8,9], high wear resistance [10,11], high strength and plasticity [12,13], high hardness (HEAs) and superhardness (ceramic HEAs) [14,15,16,17,18,19], etc. It can thus be expected that such alloys will find a wide range of applications in the industry in the near future [20,21,22,23,24,25,26].
For example, HEAs which are made as light as Al alloys but harder than certain types of metal glass [27] can find application in transport and energy technologies, where high-strength, light alloys are in great demand [28]. HEAs containing refractory Nb, Mo, and Ta are capable of preserving their high hardness even above 1200 °C, which makes them superior to conventional superalloys like Inconel 718 and Haynes 230 [7], and can be employed in gas turbines, jet nozzles, and nuclear reactors [29]. Low-density refractory HEAs can also be used in aerospace technologies for which light materials resistant to high temperature are needed [30]. The outstanding cryogenic properties of HEAs make them an excellent material for manufacturing rocket jackets, pipelines for liquid oxygen or nitrogen transport and storage [31]. As has been shown [21], HEAs can replace Ni-based refractory alloys and become next-generation refractory materials.
Among the promising HEA deposition technologies are plasma or thermal spraying in which HEAs in the form of rods and powders are melted and deposed on the surface of parts [32]. HEAs can be used for electromagnetic noise suppression, especially in electronics [33], and their nitrides as well as HEAs as such for radiation-resistant coating deposition on fuel element shells [34], which extends the application of this type of material in high technologies.
Among the HEA nitride coatings with rather high functional properties obtained to date are (AlCrTaTiZr)N [15], (TiAlCrSiV)N [16], (AlCrMoSiTi)N [17], (AlMoNbSiTaTiVZr)N [18], (AlCrNbSiTiV)N [19], (TiVCrZrY)N [35], (TiVCrZrHf)N [36], (TiHfZrVNb)N [37], and (AlCrTiZrNbY)N [38].
The (AlCrTaTiZr)N and (TiVCrZrY)N coatings [15] were formed by reactive magnetron sputtering without substrate heating and substrate bias. They comprised a single fcc solid solution phase, and their microhardness measured 32 and 17.5 GPa, respectively. The coatings (AlCrTaTiZr)N [39], (AlCrNbSiTiV)N [19], and (AlMoNbSiTaTiVZr)Nx [18] were also deposited by reactive magnetron sputtering but with a substrate bias of −100 V and substrate heating to 373–773 K. They comprised only one fcc solid solution phase, and their microhardness measured 35 GPa, 41 GPa, and 37 GPa, respectively.
The conclusion is that nitride coatings based HEAs, like those based on transition metals, have a single-phase nanostructure. Increasing the substrate temperature and the nitrogen pressure in combination with a negative substrate bias provides the formation of a denser coating structure; the structure transforms from columnar to equiaxed.
The hardness of (TiHfZrVNb)N, formed by the process of separated vacuum arc deposition and reactive magnetron sputtering, increases from 23.4 GPa to 46.5 GPa when increasing the nitrogen concentration from 36% to 51.1% [34]. The structure of arc-deposited (AlCrTiZrNbY)N consists of two phases [38]: a nanocrystalline bcc phase (average crystallite size 15 nm, lattice period 0.342 nm) and a nanocrystalline fcc phase (average crystallite size 7 nm, lattice period 0.437 nm). Among the arc-deposited coatings, TiVZrNbHf nitride coatings formed at a nitrogen pressure of 0.27 Pa shows a very high hardness. Increasing the substrate bias from −50 V to −100 V at a coating texture orientation of [111] increases the coating hardness from 59 to 64 GPa, and at a biaxial orientation of [111] and [110] at −200 V, the hardness increases to 70 GPa; the crystallite size, which in this case measures 57 nm [40].
By now, many methods are available for HEA formation: powder metallurgy methods, hot pressing and casting technologies, etc. [41,42]. With promising prospects for the formation of metal coatings and hard (superhard) ceramic coatings based on HEAs, is the vacuum arc deposition method [43,44,45,46,47,48,49,50,51,52].
As a rule, the elemental composition, phase state, and defect substructure of metals and alloys are studied by scanning electron microscopy, transmission electron microscopy, and X-ray diffraction analysis. Recently, methods that use synchrotron radiation have been actively applied for research in the lattice state of growing HEA coatings.
The paper analyzes the structure and properties of metal, cermet, and ceramic NbMoCrTiAl coatings with the use of synchrotron radiation and other modern materials science methods.
The analysis extends the knowledge of HEA coatings formed by plasma-assisted vacuum arc deposition, including those with a metal–ceramic multilayer structure, which are poorly studied and little addressed in the literature.

2. Materials and Methods

The materials under study were metal, cermet, and ceramic NbMoCrTiAl coatings deposited on high-purity tungsten, 12Cr18Ni10Ti stainless steel, commercially pure titanium (Grade 2), and a hard WC–8%Co alloy. The gas-metal plasma for HEA coating deposition was produced on the QUINTA setup [53] with several different plasma generators operating at a time: a PINK-P gas plasma generator based on a non-self-sustained arc discharge with thermionic and hollow cathodes and four arc evaporators with niobium (99.8 wt.% Nb), titanium-aluminum (Ti–50 wt.% Al), molybdenum (99.96 wt.% Mo), and chromium (99.99 wt.% Cr) cathodes of diameter 80, 100, 100 mm and of dimensions 400 × 80 mm2, respectively. All plasma generators, except for the evaporator with TiAl, were located on the chamber walls around a table with planetary satellites at equal distances from the chamber center. The evaporator with TiAl was at the top of the chamber perpendicular to the other plasma generators. The experimental arrangement is described in detail elsewhere [54].
The contribution of the ion current of each arc evaporator was assessed by measuring the azimuthal characteristics of the ion current densities from the evaporators and PINK-P generator. The collector was located on a planetary satellite moving along the whole perimeter of the working chamber such that its spacing from all plasma generators was the same. The minimum distance from the generator exit aperture to the collector was 16 cm, and the maximum distance was 52 cm. The results of measurements are presented in Figure 1.
The azimuthal ion current density distribution was used to determine the arc parameters at which the average ion current density to the collector was the same for all evaporators, measuring 2.45 mA/cm2. Such conditions ensured the most uniform ion current density distribution over the chamber during simultaneous operation of the plasma generators, specified temperature level throughout the deposition of coatings, and coating thickness uniformity. The discharge currents of the metal plasma generators for the deposition of HEA coatings are tabulated in Table 1. The discharge current of the PINK-P gas plasma generator was 50 A. The substrates to be coated were fixed in a holder making an angle of 45° with the satellite rotation axis for uniform exposure to all types of plasma. The working gas was argon for the metal coating deposition and a mixture of argon and nitrogen in equal proportions for the ceramic coating deposition; the pressure in all experiments was 0.3 Pa. During the deposition, a negative pulsed bias was applied to the substrates. To optimize the mode of the HEA coating deposition, preliminary experiments were performed on the formation of coatings of different composition with separately and jointly operating plasma generators. The coatings formed in the experiments, in addition to HEA ones, were the following (Table 1): Nb, Mo, Cr, and TiAl, and NbN, MoN, CrN, and (Ti,Al)N, depending on the cathode material and gas mixture (argon and argon with nitrogen, respectively). The deposition time was 30 min. The coating thickness was measured by the Calotest method. The coating growth rates are given in Table 1. Reasoning from the preliminary experiments, metal and ceramic coatings 3 µm thick were formed.
The deposition modes for multilayer cermet coatings were the same as those for metal and ceramic coatings, respectively. The deposition time of each metal and ceramic layer was chosen so that its thickness was 300 nm at a total number of layers of 10. Figure 2 shows a spherical section of such a coating of thickness 3 µm ± 10% after its Calotest. It is clearly seen that the cermet coating has a multilayer structure.
The elemental composition, phase state, and defect substructure of the coatings were examined on a Philips SEM-515 scanning electron microscope with an EDAX ECON IV microanalyzer and on a JEOL JEM-2100F transmission electron microscope. Their phase state and structure were analyzed on an XRD-6000 diffractometer in CuKα radiation (radiation wavelength λ = 1.540598 Å) in the standard Bragg–Brentano mode. The phase composition was analyzed using PDF 4+ databases and POWDER CELL 2.4 full-profile analysis program. Additionally, these were investigated on the VEPP-3 storage ring in synchrotron radiation (INP SB RAS) under normal conditions (tungsten substrate, radiation wavelength λ = 0.15401 nm, diffraction angle 2θ = 20°–80°). The software was MAUD (structural analysis) and Fityk v.1.3.1 (data processing). The heat resistance of the coatings was measured by real-time X-ray diffractometry in synchrotron radiation on the VEPP-3 storage ring (Anton Paar HTK-2000 high-temperature X-ray chamber, OD-3M-350 one-dimensional position-sensitive detector, Fityk v.1.3.1 data processing software). The measurements were performed in air at an operating wavelength of 0.172 nm, temperature of 30–1300 °C, specimen heating rate of 15 °C/min, and exposure time of 1 min per frame. The diffraction angle 2θ was 28°–59°. The hardness of the coatings was measured on a PMT-3 hardness tester (Vickers method, indenter load 0.5 N). Their physicomechanical properties were studied on a TTX-NHT nanohardness tester (acquisition rate 20.0 Hz, maximum linear load 30.0 mN, loading rate 60.00 mN/min, unloading rate 60.00 mN/min, pause 5.0 s, Poisson’s ratio 0.25, Oliver–Pharr analysis), and their tribological properties were studied on a pin-on-disk tribometer and on an oscillating TRIBOtester (TRIBOtechnic, France) using a WC–8%Co ball of diameter 6 mm (load 2 N, wear track radius 2 mm, track length 50 m, ball velocity 25 mm/s). The wear of the coatings was determined from their wear track profilometry.

3. Results

Table 2 presents the elemental composition of the metal, cermet, and ceramic HEA coatings. It is readily seen that in terms of the relative content of metal atoms, the coatings fit the definition of high-entropy materials.
As can be seen in Figure 3, the X-ray diffraction spectra of the three types of coating deposited on WC–8%Co and W have a much broader profile with respect to the reference lines of WC (Figure 3a) and W (Figure 3b). This suggests that the coatings can be both X-ray amorphous and nanocrystalline. The X-ray diffraction spectrum of the metal coating (Figure 3a, spectrum 1), unlike the others, reveals several peaks which allow one to determine the type and parameters of its lattice. Our analysis shows that the coating has a bcc lattice with a = 0.3166 nm. Note that Mo, Nb, and Cr contained in the HEA coating have this type of lattice.
The X-ray diffraction spectra measured via synchrotron radiation reveal broad peaks from HEA phases (Figure 3b). Judging from spectra 1, 3 in Figure 3b, the coatings have a structure close to that of Mo (bcc lattice); its unit cell parameter measures a = 0.3146 nm. The metal NbMoCrTiAl coating reveals peak asymmetry toward larger angles and likely because its metal peaks are superimposed. For the cermet coating (NbMoCrTiAl–(NbMoCrTiAl)N), even broader peaks of irregular shapes are observed, which can also be explained by the superposition of the peaks of its individual elements. The size of the coherent scattering regions is 5–6 nm for the metal coating and 2–3 nm for the cermet one. The X-ray diffraction spectrum of the ceramic coating has a single broad peak (Figure 3b, spectrum 2), which makes it impossible to determine the type and parameter of its lattice.
Our analysis by transmission and scanning transmission electron microscopy shows that the metal, ceramic, and cermet HEA coatings represent a multilayer film (Figure 4).
According to the data from the X-ray spectrum analysis and X-ray mapping, the layers differ in elemental composition, which is exemplified for the metal and ceramic HEA coatings in Figure 5 and Figure 6. As it follows from our analysis, the metal and ceramic coatings (Figure 4) are rich in Ti and Al in their bright layers and in Nb, Mo, and Cr in their dark layers; the nitrogen distribution over the ceramic coating volume is uniform.
It can be supposed that the main cause of the nonuniform element distribution is in the evaporation of four cathodes located on the chamber walls around the table with planetary satellites and from the location of the TiAl cathode at the top of the chamber perpendicular to the other cathodes.
Figure 7 shows the structure of the metal and ceramic coatings imaged by transmission electron microscopy.
The data in Figure 7 suggest that both coatings have a columnar structure (Figure 7c,d) formed by crystallites of size 3–5 nm (Figure 7d). The ratio of the diffraction ring radii (Figure 7e,f) attests that the metal and the ceramic HEA coating has a bcc and an fcc lattice, respectively.
The cermet HEA coating, as expected, consists of alternating nanocrystalline metal and ceramic layers with a bcc (Figure 8c,d) and an fcc lattice, respectively (Figure 8a,b).
It should be noted that the column height in the metal and ceramic coatings is much greater than the thickness of the individual layers (Figure 7) and can reach the total coating thickness, which adversely affects the crack resistance of the coatings. Such a columnar structure throughout the thickness of ceramic HAE coatings is described elsewhere [55]. In the cermet coating with alternating metal and ceramic layers, the column height is limited to the thickness of a single layer (Figure 8), which greatly increases its crack resistance compared to the metal and ceramic coatings. Besides, such cermet multilayer coatings have higher hardness, compared to metals, and higher fracture toughness, compared to single-phase ceramics [56].
The results of heat resistance tests with the use of synchrotron radiation (VEPP-3 storage ring) show that on heating in air, the ceramic coating starts to oxidize at 770–780 °C. The phase state of the heated coating is impossible to determine due to its thermal expansion and changes in the interplanar spacings of its reflections. At 1070–1080 °C, the X-ray patterns of the coating reveal new reflections presumably from oxides of the alloy elements. Further increasing the temperature shifts the reflections toward larger angles due to the deformation of the specimen and failure of the optical scheme. At ~1180 °C, all reflections previously present on the X-ray patterns disappear. On cooling to ~800 °C, intense reflections of the specimen holder are observed, suggesting that the stress arising in the coating deforms the specimen so that its surface comes out of the incident synchrotron radiation. Therefore, it is impossible to determine the final coating composition. Figure 9 presents the total set of X-ray patterns as projections of their intensity onto the diffraction angle–temperature plane for the ceramic HEA coating on a W substrate during heating in air from room temperature to 1300 °C.
The metal and cermet HEA coatings heated in air demonstrate almost the same behavior. At 630–640 °C, the coatings start to oxidize. At 1070–1080 °C, new reflections appear, presumably from oxides of their elements. At 1260–1270 °C, the HEA reflections disappear completely. Further, the coatings start to fail, and their final state is impossible to determine.
The mechanical and tribological parameters of the HEA coatings are presented in Table 3. The data presented allow the following conclusions. First, the metal HEA coating has a rather high hardness; the proportional summation of the hardness of its elements gives an HV = 1.1 GPa. Second, the hardness of the ceramic HEA coating is at the level of superhardness. In some cases [57,58], superhardness can be displayed by ceramic HEA coatings. Third, the wear resistance of the ceramic and cermet coatings is more than 20 times higher than that of the metal coatings. Such wear resistance is reported for the ceramic HEA and amorphous metal HEA coatings as well as for multilayer cermet coatings [58,59].

4. Conclusions

In sum, we have analyzed the structure and properties of metal, cermet, and ceramic NbMoCrTiAl high-entropy alloy (HEA) coatings formed on solid substrates by the process of plasma-assisted vacuum arc deposition (from multicomponent gas-metal plasma through Nb, Mo, Cr, and TiAl cathode evaporation in argon and a mixture of argon and nitrogen). The analysis shows that all coatings represent a nanocrystalline (3–5 nm) film with a multilayer structure; the multilayer structure owes to the experimental technique used. The metal coating has a bcc lattice (a = 0.3146 nm). The ceramic coating has an fcc lattice; its parameter is uncertain due to highly smeared diffraction peaks. The hardness and Young’s modulus of the coatings increase in the order of metal, cermet, and then ceramic, reaching 43 and 326 GPa, respectively. The wear resistance of the ceramic and cermet HEA coatings is more than 20 times higher than that of the metal one. On heating in air, the metal and cermet coatings start to oxidize at 630–640 °C, and the ceramic coating at 770–780 °C. At 1260–1270 °C, the coatings fail completely.

Author Contributions

Conceptualization, Y.F.I.; methodology Y.F.I.; validation, O.V.K. and V.V.U.; investigation E.A.P., N.A.P., A.N.S. and O.S.T.; resources V.V.S.; data curation N.N.K.; writing—original draft preparation, Y.F.I. and O.V.K.; writing—review and editing, Y.F.I., N.A.P. and O.V.K.; visualization, E.A.P., N.A.P., A.N.S. and O.S.T.; funding acquisition, Y.K.A. All authors have read and agreed to the published version of the manuscript.

Funding

This work was carried out with the financial support of the Russian Federation represented by the Ministry of Science and Higher Education (project No 075-15-2021-1348) within the framework of event No. 3.1.4, 3.1.5, 3.1.12 and 3.1.13.

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Not applicable.

Acknowledgments

TEM investigations was carried out using the equipment of the CSU NMNT TPU, supported by the RF MES project No. 075-15-2021-710.

Conflicts of Interest

The authors declare no conflict of interest.

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Coatings 13 01191 g001 550
Figure 1. Azimuthal distribution of ion current density vs. arc current in argon (a) and mixture of argon and nitrogen in equal proportions (b).
Figure 1. Azimuthal distribution of ion current density vs. arc current in argon (a) and mixture of argon and nitrogen in equal proportions (b).
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Figure 2. Characteristic spherical section of cermet coating on WC–8%Co alloy after the Calotest.
Figure 2. Characteristic spherical section of cermet coating on WC–8%Co alloy after the Calotest.
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Coatings 13 01191 g003 550
Figure 3. Fragments of X-ray diffraction spectra of metal (1), ceramic (2), and cermet HEA coatings (3) deposited on WC-8%Co (a) and W (b). Measurements on XRD-6000 diffractometer (a) and on VEPP-3 storage ring in synchrotron radiation (b).
Figure 3. Fragments of X-ray diffraction spectra of metal (1), ceramic (2), and cermet HEA coatings (3) deposited on WC-8%Co (a) and W (b). Measurements on XRD-6000 diffractometer (a) and on VEPP-3 storage ring in synchrotron radiation (b).
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Figure 4. Structure of metal (a) and ceramic HEA coating (b). Scanning transmission electron microscopy.
Figure 4. Structure of metal (a) and ceramic HEA coating (b). Scanning transmission electron microscopy.
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Figure 5. Metal HEA coating in characteristic X rays from Cr (a), Ti (b), Mo (c), Nb (d), Al (e).
Figure 5. Metal HEA coating in characteristic X rays from Cr (a), Ti (b), Mo (c), Nb (d), Al (e).
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Figure 6. Ceramic HEA coating in characteristic X rays from Cr (a), Ti (b), Mo (c), Nb (d), Al (e), N (f).
Figure 6. Ceramic HEA coating in characteristic X rays from Cr (a), Ti (b), Mo (c), Nb (d), Al (e), N (f).
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Figure 7. Structure of metal (a,d,f) and ceramic HEA coatings (b,c,e) in bright field (a,b) and dark field (c,d), and electron diffraction patterns (e,f) with arrows for reflections in dark field.
Figure 7. Structure of metal (a,d,f) and ceramic HEA coatings (b,c,e) in bright field (a,b) and dark field (c,d), and electron diffraction patterns (e,f) with arrows for reflections in dark field.
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Figure 8. Structure of cermet HEA coating in bright field (a,c) and respective electron diffraction patterns (b,d) taken from its red-circled regions in (a,c).
Figure 8. Structure of cermet HEA coating in bright field (a,c) and respective electron diffraction patterns (b,d) taken from its red-circled regions in (a,c).
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Figure 9. Total set of X-ray patterns as projections of their intensity onto a diffraction angle–temperature plane for ceramic HEA coating on a W substrate during heating in air from room temperature to 1300 °C.
Figure 9. Total set of X-ray patterns as projections of their intensity onto a diffraction angle–temperature plane for ceramic HEA coating on a W substrate during heating in air from room temperature to 1300 °C.
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Table
Table 1. Discharge currents of metal plasma generators and coating growth rates.
Table 1. Discharge currents of metal plasma generators and coating growth rates.
Evaporated CathodeDischarge Current, ACoating Composition, TypeCoating Growth Rate, µm/h ± 10%.
Cr150Cr, metal2.2
CrN, ceramic2.4
TiAl180TiAl, metal2.4
(Ti,Al)N, ceramic5.2
Mo90Mo, metal1.1
MoN, ceramic1.4
Nb120Nb, metal0.9
NbN, ceramic2
Cr, TiAl, Mo, and Nb togetherDischarge currents are indicated aboveNbMoCrTiAl, metal6.3
(Nb,Mo,Cr,Ti,Al)N, ceramic6
NbMoCrTiAl/(Nb,Mo,Cr,Ti,Al)N, cermet6.2
Table
Table 2. Elemental composition of HEA coatings according to X-ray spectrum analysis.
Table 2. Elemental composition of HEA coatings according to X-ray spectrum analysis.
HEA Coating Type Elements, at %
AlTiCrNbMoN
Metal16.017.419.319.328.00.0
Cermet9.717.211.217.721.123.1
Ceramic13.96.89.36.87.855.4
Table
Table 3. Mechanical and tribological characteristics of HEA coatings.
Table 3. Mechanical and tribological characteristics of HEA coatings.
Coating TypeHV, GPaE, GPak, 10−5, mm3N−1m−1μ
Metal15113150.64
Cermet271760.640.20
Ceramic433260.650.47
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MDPI and ACS Style

Ivanov, Y.F.; Akhmadeev, Y.K.; Prokopenko, N.A.; Krysina, O.V.; Koval, N.N.; Petrikova, E.A.; Tolkachev, O.S.; Shugurov, V.V.; Uglov, V.V.; Shmakov, A.N. Structure and Properties of NbMoCrTiAl High-Entropy Alloy Coatings Formed by Plasma-Assisted Vacuum Arc Deposition. Coatings 2023, 13, 1191. https://doi.org/10.3390/coatings13071191

AMA Style

Ivanov YF, Akhmadeev YK, Prokopenko NA, Krysina OV, Koval NN, Petrikova EA, Tolkachev OS, Shugurov VV, Uglov VV, Shmakov AN. Structure and Properties of NbMoCrTiAl High-Entropy Alloy Coatings Formed by Plasma-Assisted Vacuum Arc Deposition. Coatings. 2023; 13(7):1191. https://doi.org/10.3390/coatings13071191

Chicago/Turabian Style

Ivanov, Yurii F., Yuriy Kh. Akhmadeev, Nikita A. Prokopenko, Olga V. Krysina, Nikolai N. Koval, Elizaveta A. Petrikova, Oleg S. Tolkachev, Vladimir V. Shugurov, Vladimir V. Uglov, and Alexander N. Shmakov. 2023. "Structure and Properties of NbMoCrTiAl High-Entropy Alloy Coatings Formed by Plasma-Assisted Vacuum Arc Deposition" Coatings 13, no. 7: 1191. https://doi.org/10.3390/coatings13071191

APA Style

Ivanov, Y. F., Akhmadeev, Y. K., Prokopenko, N. A., Krysina, O. V., Koval, N. N., Petrikova, E. A., Tolkachev, O. S., Shugurov, V. V., Uglov, V. V., & Shmakov, A. N. (2023). Structure and Properties of NbMoCrTiAl High-Entropy Alloy Coatings Formed by Plasma-Assisted Vacuum Arc Deposition. Coatings, 13(7), 1191. https://doi.org/10.3390/coatings13071191

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