First-Principles Calculations and Experimental Study of H + -Irradiated Zr/Nb Nanoscale Multilayer System

: Nanoscale multilayer coating (NMC) based on Zr/Nb layers (100/100 nm) before and after H+ irradiation was investigated by combining experimental techniques with ﬁrst-principles calculations. Detailed studies of structural and phase state and defect structure of Zr/Nb NMC were performed using methods of transmission electron microscopy, X-ray structural analysis, glow discharge optical emission spectrometry, and the Doppler broadening spectroscopy using variable energy positron beam. The ﬁrst-principles calculations of binding energies for hydrogen in metal Zr/Nb layers was carried out by the pseudopotential method within the density functional theory framework. First-principles calculations and experimental data indicate the presence of macro-and microstrains predominantly in the zirconium layers of Zr/Nb NMC. The main feature of the studied Zr/Nb NMC is the predominant hydrogen localization in Zr layers near the interfaces. The annihilation of positrons is shown to occur mainly in the Zr layers in the vicinity of the interface.


Introduction
Nanoscale multilayer coatings (NMC) are increasingly used in power engineering, electronics, mechanical engineering, optics, biotechnology, and other industries [1][2][3][4][5][6]. It is caused by the possibility of significant modification of physical-mechanical properties and corrosion resistance of structural materials. The layers' thickness and composition variations can achieve higher strength, hardness, and form nanocomposite with a wide range of functional purposes [7][8][9][10][11]. FCC/BCC metals are widely approved as radiation-resistant NMC due to the high density of interfacial dislocations [7]. The recombination of vacancies and interstitial atoms takes place during irradiation in the incoherent interfaces [12,13]. Some examples of such systems are Cu/Nb [10], V/Ag [9], Cu/W [8], Zr/Nb [14][15][16][17], etc. For example, Cu/Nb nanolayers show very high stability under helium ion irradiation [18]. The Cu/V system might have a self-healing effect on radiation defects. The number of helium bubbles in the ion deposition zone decreases with an increasing number of incoherent Cu/V interfaces upon irradiation with He ions [19][20][21]. All these indicate that multilayer coatings can be used as an advanced nuclear and aerospace material.
The largest structural discrepancy of crystal lattices has hexagonal-cubic systems (hcp/bcc and hcp/fcc). Moreover, hcp/bcc systems have excellent potential for creating radiation-resistant composites since a high discrepancy of crystal lattices is present. The large discrepancy allows the incoherent and semi-coherent interfaces of hcp/bcc systems

Sample Preparation
Polished monocrystal Si (111) plates with a thickness of 0.5 mm and a roughness of <0.015 µm were used as a substrate. The Zr/Nb multilayers were deposited by magnetron sputtering using two independent cathodes on a specialized facility developed at the Weinberg Research Center of the School of Nuclear Science and Engineering TPU. The individual layer thickness was 100 ± 20 nm, and the total coating thickness is 1.1 ± 0.1 µm.
The Zr/Nb NMC were irradiated with 900 keV H + -ions using an ESG-2.5 MeV electrostatic generator (Research Institute of Nuclear Physics TPU, Tomsk, Russia) in TPU with a beam current of 2 µA for 30 min (0.07 dpa according to "The Stopping and Range of Ions in Matter" (SRIM) [30]). The stability of the beam current and energy is 0.02%. To obtain a specified ion implantation profile in the Zr/Nb NMC multilayers, according to SRIM calculations, an aluminum degrader with a thickness of 11 µm was used.

Experimental Methods
The structural and phase state was investigated by X-ray diffraction analysis (XRD) on a Shimadzu XRD-7000S diffractometer (Shimadzu, Kyoto, Japan) in Bragg-Brentano geometry with Cu K α radiation (λ = 0.154 nm) [31]. The diffraction analysis was performed at scan speed of 5.0 • /min and sampling pitch of 0.0143 • with an exposure time of 42.97 s. To obtain the structural parameters of the Zr/Nb NMC, XRD patterns were studied by Rietveld analysis, where accuracy factors were~8 for R p and~11 for R wp . For microstrain and coherent domain size calculation, the Williamson-Hall technique was used. Instrumental broadening of 0.14 • was taken into account. All calculations were performed in POWDER CELL program [32]. The samples' cross-sectional structure was studied by transmission electron microscopy on a JEM-2100F (JEOL, Tokyo, Japan) transmission electron microscope with a sample preparation system [33]. The samples were prepared by ion milling using an EM-09100IS Ion Slicer (JEOL, Akishima, Tokyo, Japan). Argon was used as the working gas, the accelerating voltage was 8 kV, and the etching angle was 1.5-4 • .
The layer-by-layer distribution of the Zr, Nb, and H was investigated using glow discharge optical emission spectrometry (GD-OES) on a GD-Profiler 2 spectrometer (Horiba Scientific, France) [34,35]. A two-method approach was used to correct the GD-OES spectra: an initial instrumental selection of the sputtering parameters and a subsequent correction of the spectrum by exponential dependence approximation [36]. NMC Zr/Nb's study was carried out at the following parameters using a radio-frequency source: pressure 650 Pa, power 40 W, frequency 1 kHz, duty cycle 25%.
The Doppler broadening spectroscopy using variable energy positron beam (DBS-VEP) was conducted in the Dzhelepov Laboratory of Nuclear Problems, JINR in Dubna, Russia [37,38]. The samples were measured in initial and irradiation state without any additional preparation at room temperature. The obtained spectras include around 250k counts. Positrons with energies varying from 0.1 to 30 keV were used, which correspond to mean implantation depth from around 10 nm to 1 um, respectively. In the experiment, a broadening of the annihilation line caused by the Doppler effect was monitored and the line shape changes described by three commonly used parameters: S, standing for positron annihilation with low energy, valence electrons; W, responsible for annihilation with core electrons; and the R parameter for the graphical presentation of S = f(W) function [39][40][41][42][43]. The S parameter was calculated as the area below the central part of the annihilation line to the total area below 511 keV line. Similarly, the W parameter was obtained as the area of the wing part of 511 keV line divided by the total number of counts in the annihilation line. The S and W parameters and their errors were obtained using an SP program using the Gauss fitting option after background subtraction. These parameters change when the positron is trapped in defect, and are also sensitive to the different chemical surroundings of annihilation sites. For example, when a positron is trapped in a vacancy where there is lower probability of annihilation with a high momentum electron, the S parameter increases and W decreases. All measurements were performed using HPGe detector model GEM25P4-70 (AMETEK ORTEC, Oak Ridge, TN, USA) with 1.20 keV resolution at 0.511 MeV line.

Ab Initio Calculations
Ab initio calculations were carried out within density functional theory using the optimized norm-conserving Vanderbilt pseudopotential method [44], as implemented in the ABINIT code [45,46]. The generalized gradient approximation (GGA) in the form of Perdew, Burke, and Ernzerhof [47] was used to describe the exchange and correlation effects. The cutoff energy for the plane wave basis was set to 15 Ha in the structural optimization. The atoms in the system were assumed to be in the equilibrium configuration when the force on each atom was below 10 −3 Ha/bohr. The present calculations were performed for pure Zr and Nb, Zr 36 H and Nb 36 H solid solutions with hydrogen in tetrahedral or octahedral interstitial sites (Figure 1), Zr 63 Nb 40, and Zr 63 Nb 40 H slabs ( Figure 2). The interface in Zr 63 Nb 40 slab was formed by Zr (002) and Nb (111) surfaces. The Zr and Nb slabs consist of 7 and 10 atom monolayers, respectively. In Zr 36 H, Nb 36 H, and Zr 63 Nb 40 H systems H atom is located in the either tetrahedral (T) or octahedral (O) interstitial sites. For a more convenient discussion of results, the T and O sites in Figure 2b,c are enumerated. To carry out the structural optimization and relaxation of the system considered, a cell with 36 Zr or 36 Nb atoms was adopted, and the k meshes were chosen to be 3 × 3 × 3 for the hcp and bcc structures. For multilayer structures Zr 63 Nb 40 , k meshes were chosen 3 × 3 × 1. To analyze the structural stability of the systems under consideration, the binding energy of hydrogen in metals was calculated:  To analyze the structural stability of the systems under consideration, the binding energy of hydrogen in metals was calculated:  It is clearly seen that a structure with alternating layers with a thickness of 100 ± 20 nm was obtained as a result of deposition. According to EDS data, the formed structure is represented by layers of zirconium and niobium with clearly distinguishable boundaries (Figure 3b,c). A more detailed study of the transverse sections of the as-deposited Zr/Nb NMC was carried out using transmission electron microscopy ( Figure 4). It is clearly seen that a structure with alternating layers with a thickness of 100 ± 20 nm was obtained as a result of deposition. According to EDS data, the formed structure is represented by layers of zirconium and niobium with clearly distinguishable boundaries (Figure 3b,c). A more detailed study of the transverse sections of the as-deposited Zr/Nb NMC was carried out using transmission electron microscopy ( Figure 4).
The corresponding selected area electron diffraction patterns (SAED) (Figure 4b) show an appreciable number of reflections distributed over a circle. The SAED patterns contain reflections from the different planes of the α phase of Zr and reflections of the Nb(211)β plane. The microstructure of each Zr and Nb layer is represented by nanoscale columnar grains, the average size of which varies from 20 to 50 nm. Grains in layers grow perpendicular to the substrate (Figure 4a). High-resolution TEM micrographs show an incoherent interface between Zr and Nb ( Figure 5).
It is clearly seen that a structure with alternating layers with a thickness of 100 ± nm was obtained as a result of deposition. According to EDS data, the formed structure represented by layers of zirconium and niobium with clearly distinguishable boundar (Figure 3b,c). A more detailed study of the transverse sections of the as-deposited Zr/N NMC was carried out using transmission electron microscopy ( Figure 4).    Figure 6 presents the diffraction patterns for as-deposited and irradiated Zr/N NMC obtained for samples' surface.  (b) (c) Figure 5. Cross-sectional HRTEM images of the Zr/Nb multilayer (a) bright-field image of the Zr/Nb multilayers (b) corresponding fast Fourier transform (FFT) for the HRTEM images of the Zr/Nb (c) and the Nb/Zr interfaces. The blue lines show the interface between Zr and Nb. Figure 6 presents the diffraction patterns for as-deposited and irradiated Zr/Nb NMC obtained for samples' surface.  Diffraction analysis showed that Zr layers are more susceptible to the formation of macro-and microstrains during deposition and irradiation. It can be seen from the pattern that the diffraction peaks are considerably broad, and values of full width on half maxima (FWHM) for Zr(100) and Zr(002) peaks are larger than for Nb (110), 0.9 • and 0.67 • against 0.58 • . Peak broadening is usually attributed to large microstrain and small coherent diffraction domain size. For as-deposited Zr layers, the calculated microstrain and coherent diffraction domain sizes were 0.0073 and 9.9 ± 2.5 nm. Due to the high (110) texture of Nb layers, it is impossible to separate the contribution of microstrain and coherent diffraction domain size to peak broadening. Structural analysis from diffraction data revealed the presence of distortions in the crystal lattice, indicating macrostrains in as-deposited and irradiated Zr/Nb NMC. Macrostrains are especially pronounced for Zr layers; the hexagonal lattice parameters for as-deposited NMC were the following: a = 3.162 Å, c = 5.129 Å, which is 2.2% and 0.4% lesser than ideal lattice parameters for bulk Zr (a = 3.232 Å, c = 5.147 Å) [48]. For Nb layers, cubic a parameter was 3.278 Å (0.9% lesser than ideal a = 3.306 Å [49]). After 30 min of irradiation, minor changes in the diffraction pattern were noticed. A slight decrease in the lattice parameters was observed for the Zr and Nb layers, the values of which are as follows: a = 3.158 Å and c = 5.126 Å for Zr, and a = 3.276 Å for Nb. Additionally, an increase in FWHM value was noticed for Zr(100) to 1.07 • , for Zr(002) to 0.7 • , and for Nb(110) to 0.6 • . Thus, for irradiated Zr layers, microstrain increased to 0.008, and coherent diffraction domain size decreased to 9.1 ± 2.4 nm.

First-Principles Calculations of Zr/Nb Multilayer Coatings
The calculated lattice parameters of pure Zr and Nb, Zr 36 H, and Nb 36 H solid solutions are presented in Table 1. These theoretical results are in good agreement with the results of experiments and previous theoretical works [22,23]. To form the supercell of Zr 63 Nb 40 slab, we decreased the theoretical lattice parameters of pure Nb and increased the lattice parameters of pure Zr until the total supercell energy became minimal. As a result, the supercell Zr slab have parameters a = 3.165 Å and c = 5.160 Å, and the Nb slab have parameter a = 3.341 Å. The distance between the Zr (002) and Nb (111) atomic layers in the interface was optimized. The equilibrium distance is equal to 2.215 Å, which is slightly lower than the result obtained in the previous theoretical work [24]. Finally, the atom positions in the three Zr and five Nb atomic layers nearest to the interface were relaxed. It was found that the Zr atom shifts near the interface are significantly higher than the Nb atom shifts. For example, the maximum atom shifts in the first Zr and Nb atomic layers are 1.090 Å and 0.348Å, respectively. In the second Zr and Nb atomic layers, these shifts reach 0.869 Å and 0.229Å, respectively. The Zr and Nb atom shift values in other relaxed atomic layers are close to each otherand do not exceed 0.2 Å. This fact partially explains microstrain studies' results: the Zr layer's deformation is noticeably higher than that of the Nb. The results of the hydrogen binding energy calculation in solid solutions and slab are presented in Table 2. The calculated hydrogen binding energy in zirconium is in good agreement with the results of previous theoretical work [22]. The maximum binding energy corresponds to the tetrahedral coordination of a hydrogen atom in both metals. It should be noted that in the case when the hydrogen atom is located in the O site of niobium, even its small displacements from the center of this site lead to the appearance of force that shifts the H atom to the nearest T site. This indicates the instability of this hydrogen position in niobium. As a result, in the Zr 63 Nb 40 H slab, we considered the location of an H atom in T and O sites of the Zr layer and in T sites of the Nb layer. The maximal value of hydrogen binding energy is observed in the T6 site in the first Zr atomic layers (see Figure 2c). It should be noted that the binding energy values are varied in a wide range for each metal atomic layer. In the first Zr atomic layers, this range is from 0.653 to 1.035 eV for tetrahedral coordination and from 0.758 to 0.970 eV for octahedral coordination of hydrogen atoms. However, the average value of hydrogen binding energy per metal atomic layer is decreased with an increase in distance between the atomic layer and interface (Figure 7). In the Zr layer, this decrease is slower than in the Nb layers.       Figure 8 shows a GD-OES profile of the Zr/Nb multilayer system consisting of five Zr-layers approximately 80 ± 10 nm thick and five Nb-layers with 100 ± 10 nm thickness, which agrees well with the STEM results ( Figure 3). The transient areas between layers present in the GD-OES profiles are related to the crater effect [36,50,51]. The Zr/Nb interface near the Si substrate was not considered. Trace amounts of hydrogen are present in the unirradiated material, and hydrogen is mainly localized in the Nb/Zr interface area. The main changes in the relative GD-OES curves are associated with the accumulation of hydrogen ions. The hydrogen concentration maximum occurs at a depth of ~800 nm, which corresponds to the Bragg peak for H + + ions, according to SRIM calculations [15].  Figure 8 shows a GD-OES profile of the Zr/Nb multilayer system consisting of five Zr-layers approximately 80 ± 10 nm thick and five Nb-layers with 100 ± 10 nm thickness, which agrees well with the STEM results ( Figure 3). The transient areas between layers present in the GD-OES profiles are related to the crater effect [36,50,51]. The Zr/Nb interface near the Si substrate was not considered. Trace amounts of hydrogen are present in the unirradiated material, and hydrogen is mainly localized in the Nb/Zr interface area. The main changes in the relative GD-OES curves are associated with the accumulation of hydrogen ions. The hydrogen concentration maximum occurs at a depth of~800 nm, which corresponds to the Bragg peak for H + ions, according to SRIM calculations [15]. The hydrogen distribution has a multimodal character, with local maxima of hydrogen concentration observed at the Nb/Zr interfaces, while the accumulation on Zr/Nb is much lower. The intensity ratio of the H line for the Nb/Zr and Zr/Nb interfaces is approximately 1:2, and the hydrogen concentration at the Zr/Nb interface is almost equal to the hydrogen concentration in the bulk of niobium. Moreover, hydrogen predominantly localizes in the zirconium layers, which agrees well with the first-principles calculations (Figure 7), which showed that near the Zr/Nb interface, the hydrogen binding energy in the zirconium atomic layers is significantly higher than in the niobium atomic layers.

Depth Profiling and Defect Characterization of Zr/Nb Nanoscale Multilayers after H + Irradiation
The DBS-VEP analysis of the annihilation line in Zr/Nb NMCs is shown in Figure 9.
Metals 2021, 11, x FOR PEER REVIEW 11 of 18 concentration observed at the Nb/Zr interfaces, while the accumulation on Zr/Nb is much lower. The intensity ratio of the H line for the Nb/Zr and Zr/Nb interfaces is approximately 1:2, and the hydrogen concentration at the Zr/Nb interface is almost equal to the hydrogen concentration in the bulk of niobium. Moreover, hydrogen predominantly localizes in the zirconium layers, which agrees well with the first-principles calculations (Figure 7), which showed that near the Zr/Nb interface, the hydrogen binding energy in the zirconium atomic layers is significantly higher than in the niobium atomic layers. The DBS-VEP analysis of the annihilation line in Zr/Nb NMCs is shown in Figure 9.
(a) (b) Figure 9. DBS-VEP depth profile of the S (a) and W (b) parameters for Zr/Nb 100. The S parameter stands for positron annihilation with low momentum electrons (it increases with defects concentration and is higher for Zr than Nb), and W is responsible for positron annihilation with high momentum electrons.
Because the depth of positron implantation nonlinearly changes with increasing energy, only the first layers (from the top surface) of niobium and zirconium have been entirely examined. The S parameter for niobium is about 0.507 ± 0.004 and increases when positrons annihilate in Zr. In unirradiated samples at positron energies from 4 to 26 keV, the S parameter value is about 0.525 ± 0.005. The W parameter also changes weakly in this energy range, which shows that positrons beginning from energy 4 keV start annihilate in the zirconium area. The dependence S = f(W), shown in Figure 10, also confirms this assumption. Because the depth of positron implantation nonlinearly changes with increasing energy, only the first layers (from the top surface) of niobium and zirconium have been entirely examined. The S parameter for niobium is about 0.507 ± 0.004 and increases when positrons annihilate in Zr. In unirradiated samples at positron energies from 4 to 26 keV, the S parameter value is about 0.525 ± 0.005. The W parameter also changes weakly in this energy range, which shows that positrons beginning from energy 4 keV start annihilate in the zirconium area. The dependence S = f(W), shown in Figure 10, also confirms this assumption.
The number of positron annihilation in Zr may be higher due to the selective positron trapping [52] in the multilayer system caused by the positron affinity of zirconium which is 26% bigger than for niobium [53]. However, suppose the thickness of individual layers is about 100 ± 20 nm, and the average positron diffusion length is comparable with nanocrystallites' size-32 ± 15 nm. In that case, it is unlikely that all positrons annihilate only in Zr. Therefore, the obtained value of 0.525 ± 0.005 should be assigned to the mixed state of the S parameter value of pure Zr and Nb, which is closer to pure Zr value due to lower Nb positron affinity. when positrons annihilate in Zr. In unirradiated samples at positron energies from 4 to 26 keV, the S parameter value is about 0.525 ± 0.005. The W parameter also changes weakly in this energy range, which shows that positrons beginning from energy 4 keV start annihilate in the zirconium area. The dependence S = f(W), shown in Figure 10, also confirms this assumption. To reveal the influence of the Zr/Nb interface on peculiarities of the distribution of positron annihilation, a comparative analysis of the electron density distribution in the interface vicinity of Zr 63 Nb 40 slabs was performed. Figure 11 shows that the electron density distribution near the interface is not homogeneous on both the Zr and Nb layers.
There are areas of a reduced electron density (*) in the Zr atomic layers near the interface. With increasing distance from the interface, the number of these areas increases, but their size decreases. The niobium atomic layers are almost entirely devoid of such areas. The appearance of such areas in the Zr atomic layers is due to high Zr atom shifts to the interface. In the fourth Zr atomic layer, where atoms are located at the ideal hcp lattice site, such areas are not observed. Thus, positron annihilation near the interfaces will mainly occur in the zirconium area. A detailed study of reduced electron density areas requires a positron lifetime spectroscopy analysis.
After irradiation with H + ions for 30 min in Zr/Nb NMC, a slight decrease in the S parameter and W parameter growth in the Zr and Nb first layers are observed. That indicates a reduction of free volume in these layers. Moreover, in the Zr layer, these changes are more pronounced. A further increase in the positron implantation depth does not lead to significant variations in the DBS parameters. The indicated relative changes in DBS parameters (S↑W↓) are observed during hydrogen-vacancy complexes formation in hcp materials [54,55]. However, in this case, S 0 < S H < S VAC , where S 0 corresponds to the initial material, S H to the hydrogenated material, and S VAC to the material with vacancies. It is also shown in [55] that the S parameter decreases and W increases compared to the initial value is possible with H accumulation in the material.
In order to reveal the influence of hydrogen on the interaction between zirconium and niobium atoms in the interface, a comparative analysis of the electron density distribution in the interface vicinity of Zr 63 Nb 40 H slabs with H atom in the most energetically favorable octahedral (O1) and tetrahedral (T6) positions was performed (Figures 12 and 13). only in Zr. Therefore, the obtained value of 0.525 ± 0.005 should be assigned to the mixed state of the S parameter value of pure Zr and Nb, which is closer to pure Zr value due to lower Nb positron affinity.
To reveal the influence of the Zr/Nb interface on peculiarities of the distribution of positron annihilation, a comparative analysis of the electron density distribution in the interface vicinity of Zr63Nb40 slabs was performed. Figure 11 shows that the electron density distribution near the interface is not homogeneous on both the Zr and Nb layers. compared to the initial value is possible with H accumulation in the material. In order to reveal the influence of hydrogen on the interaction between zirconium and niobium atoms in the interface, a comparative analysis of the electron density distribution in the interface vicinity of Zr63Nb40H slabs with H atom in the most energetically favorable octahedral (O1) and tetrahedral (T6) positions was performed (Figures 12 and  13).  Hydrogen placed at the favorable T6 ( Figure 12) or O1 (Figure 13) positions has a minor effect on the electron density distribution near the Zr/Nb interface. The reduced electron density area configuration is changed. The presence of hydrogen atoms in O1 or T6 position increases the electron density by~5-10% of this area in the Zr and Nb atomic layers nearest to the interface and slightly reduces the size of these areas in the second and third Zr atomic layers. Such variations will not significantly affect the positron annihilation parameters shown in Figures 9 and 10. All experimental points on the S = f(W) dependence before and after irradiation are on the same straight line, indicating one prevailing type of positron trapping-the free volume near the interface in the vicinity of zirconium.
(b) (c) Figure 12. Electron density distribution in the interface vicinity of Zr63Nb40H slab with H atom in tetrahedral (T6) site (a) and in the first Zr (b) and Nb (c) atomic layers. The blue and gray isosurfaces in panel (a) correspond to a charge density of 0.02 and 0.05 electrons/Bohr 3 , respectively. The areas of a reduced electron density in the Zr (b) and Nb (c) atomic layers are marked with (*). Color gradation scale in the (b,c) panels is given in electrons/Bohr 3 units. Hydrogen placed at the favorable T6 ( Figure 12) or O1 (Figure 13) positions has a minor effect on the electron density distribution near the Zr/Nb interface. The reduced electron density area configuration is changed. The presence of hydrogen atoms in O1 or T6 position increases the electron density by ~5-10% of this area in the Zr and Nb atomic layers nearest to the interface and slightly reduces the size of these areas in the second and third Zr atomic layers. Such variations will not significantly affect the positron annihilation parameters shown in Figures 9 and 10. All experimental points on the S = f(W) dependence before and after irradiation are on the same straight line, indicating one prevailing type of positron trapping-the free volume near the interface in the vicinity of zirconium.

Conclusions
In this work, a Zr/Nb nanoscale multilayer system before and after H+ irradiation was investigated by combining different experimental techniques with first-principles calculations. A multilayer Zr/Nb system with individual layer thicknesses of 80 ± 10 nm and 100 ± 10 nm for Zr and Nb, respectively, was formed using magnetron sputtering.

Conclusions
In this work, a Zr/Nb nanoscale multilayer system before and after H+ irradiation was investigated by combining different experimental techniques with first-principles calculations. A multilayer Zr/Nb system with individual layer thicknesses of 80 ± 10 nm and 100 ± 10 nm for Zr and Nb, respectively, was formed using magnetron sputtering. Microstructures are characterized by nanoscale columnar grains with an average size from 20 to 50 nm. XRD analyses revealed the presence of macro-and microstrains in the Zr/Nb NMC. These strains are noticeably higher for Zr layers because, according to the first-principles calculations, the Zr atom shifts from the ideal lattice site near the interface are significantly higher than Nb atom shifts. Irradiation with 900 keV H+ ions for 30 min leads to an intense accumulation of H in the metal atomic layer near the Nb/Zr interfaces due to the highest hydrogen binding energy. The decrease in the average hydrogen binding energy per metal atomic layer with an increase in distance between the atomic layer and interface is slower in the Zr layer than in the Nb layer. That clearly explains the predominant hydrogen localization in Zr layers and increasing microstrains observed by GD-OES and XRD methods. The DBS-VEP analysis results do not indicate an increase in defect concentration after H+ irradiation across the coating thickness. Before and after irradiation, positrons annihilate near the interfaces, mainly in the Zr vicinity, due to high positron affinity and the presence of reduced electron density area due to high Zr atom shifts to the interface.