# Radiation-induced modifications in ZrN: effect of micro vs. nano crystallites

Connor Beer, Mounib Bahri, Gianguido Baldinozzi, Manabu Ishimaru, Yanwen Zhang, Matthew Sharpe, Callum Mcaleese, Xinyue Chen, Fabio Bohns, Troy Manning, et al.

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## To cite this version:

Connor Beer, Mounib Bahri, Gianguido Baldinozzi, Manabu Ishimaru, Yanwen Zhang, et al. Radiation-induced modifications in ZrN: effect of micro vs. nano crystallites. Acta Materialia, 2025, 298, pp.121391. 10.1016/j.actamat.2025.121391. hal-05201034

HAL Id: hal-05201034

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Submitted on 7 Nov 2025

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UN, has been used as a model system for evaluating the performance of UN in damaging radiation conditions. ZrN exhibits a stable cubic rock-salt structure that can accommodate substantial variations in its stoichiometry due to its ability to tolerate both metal and non-metal interstitials [27].

Previous work has confirmed that ZrN is highly resistant to radiation-induced structural alterations [1–17,28,18–22]. Results have shown that the response to radiation damage can depend on (1) the type and energy of incident radiation and (2) the grain size and stoichiometry of the system. Investigations into bulk microcrystalline ZrN (typically samples which have not been deposited on a substrate via PVD methods) have shown that the lattice parameter of ZrN can readily increase [7,11, 15,20] or decrease [5,9,10,28,18]. Furthermore, conflicting evolutions in crystallinity have been observed following irradiation, some displaying an increase [9,10] while others decrease [15,28]. The hardness evolution in irradiated specimens of ZrN has been reported to increase in bulk micro-crystalline ZrN [3,6,29,19], while others report radiation-induced softening in nanocrystalline material [9,11,18].

Advances in nanotechnology have led to a growth in interest surrounding nanostructured materials. Systems possessing grain sizes far below 100 nm are now being applied in electronics, membranes and coatings, among others [30,31]. Nanocrystalline materials have been the focus of recent studies based on the promise of possessing increased radiation tolerance owing to their high density of large-angle grain boundaries and interfaces acting as effective defect sinks [5,7–9,11,16, 17,28,18,20]. Lu et al. [5] irradiated nanocrystalline ZrN with swift heavy ions (SHIs), depositing energy predominantly in the electronic stopping regime, and observed no structural changes. Following this, an irradiation to study nuclear-stopping interactions was completed, and a unique structural contraction of up to 1.42% was observed. The authors attribute this to the accumulation of nitrogen vacancies and related strain relaxation. Janse van Vuuren et al. [7] also studied the response of nanocrystalline ZrN films to SHI irradiation. No appreciable changes in structural or mechanical properties were reported, labelling the system highly tolerant to the effects of high-energy irradiation. Jiao, Yu, and Chen [9] explored the effects of 900 keV Fe ions incident on nanocrystalline ZrN. The system displayed a clear dependence on grain size, with smaller grains on the order of 10 nm effectively curtailing growth. Craciun et al. [11] investigated the effect of 800 keV Ar ion irradiation on nanocrystalline ZrN. An increase in the unit cell parameter was observed, accompanied by a significant decrease in hardness and Young modulus. Wu et al. [18] reported that irradiation with 600 keV Kr³⁺ resulted in the relief of internal stress accompanied by a unit cell contraction. This behaviour appears to be characteristic of the nanocrystalline ZrN film configuration. Grain size was also observed to increase while both the measured hardness and elastic modulus decreased.

Current literature emphasises the suitability of the ZrN system as a model to investigate the impacts of radiation damage in isostructural UN. It also presents promising applications in various nuclear and engineering fields. Notably, its radiation tolerance is exceptionally high, with no reports of amorphisation under any conditions. However, its safe implementation in nuclear power applications requires a deeper understanding of its radiation response mechanisms. In this study, we investigate the effects of two distinct yet interconnected radiation response mechanisms. Firstly, the impact of grain size on radiation tolerance (structural stability and mechanical properties), and secondly, the mechanism behind nitrogen displacement and the subsequent oxidation of Zr in irradiated environments. Ion beams are employed to generate 600 keV Ar ions and irradiate microcrystalline and nanocrystalline ZrN specimens at room temperature with fluences ranging from 1 × 10¹⁵ to 1 × 10¹⁷ ions/cm². The relation between damage accumulation at room temperature and reactor operating temperatures (300–1000 °C) is different due to the strong temperature dependence of defect mobility and recovery. At room temperature, radiation-induced point defects are largely immobile, resulting in higher defect retention and a greater likelihood of clustering or amorphisation. In contrast, the

elevated temperatures present in reactors promote defect recombination and annihilation at sinks such as grain boundaries, leading to lower net damage and enhanced structural stability. Consequently, materials irradiated at room temperature often exhibit exaggerated damage compared to those exposed under in-reactor conditions. This work seeks to understand the role of interstitial defects rather than accurately simulate operational conditions. Room temperature studies of nc-ZrN therefore provide insights into fundamental defect accumulation mechanisms in the absence of thermally activated recovery. These experiments reveal the material's maximum damage retention capacity. Observations such as defect saturation and amorphisation resistance under extreme conditions offer a conservative baseline for assessing material stability. Practically, these studies inform the design of radiation-tolerant materials by clarifying how nanocrystalline architectures respond to transient low-temperature exposures or early-stage fuel qualification. This paper also outlines the foundation for an experimental campaign systematically studying mc-ZrN, nc-ZrN, and UN with varying ion types, energies, and fluences to determine optimal grain structures for radiation resistance. This campaign explores a diverse range of nuclear to electronic stopping interaction ratios ($S_e/S_n$) to map structural and mechanical changes across different energy deposition regimes.

## 2. Methods

### 2.1. Sample preparation

Microcrystalline ZrN (mc-ZrN) specimens were synthesised using powder (99.0 % purity) procured from Goodfellow Cambridge Ltd. The pellets of ZrN (thickness (~ 3 mm, diameter (~ 10 mm) were consolidated using spark plasma sintering (SPS). The SPS system (DCS10, Thermal Technology LLC, Minden, NV) used graphite foil to isolate the powder and prevent the specimen from binding to the die. Following sintering, the graphite foil was removed from the sample by grinding with abrasive silicon carbide paper. The sintering process applies a load of 60 MPa before increasing the temperature to 1700 °C at a rate of 200 °C/min with a dwell time of 5 minutes. Temperature was monitored using a pyrometer focused on a hole in the centre of the graphite die. The SPS parameters were adapted from the work of Lee et al [32]. ZrN pellets were subject to a fine polish with a Buehler Automet 250 polishing system using MetaDi fluid to a final polishing step comprising 0.02 µm colloidal silica. Pellet densities were measured using the Archimedes method with water as the working fluid. The as-fabricated pellet densities were > 99.0 % compared to the theoretical density of 7.09 g/cm³.

Nanocrystalline ZrN (nc-ZrN) thin films, (~ 1000 nm in thickness, were deposited on (001) silicon substrates with an (~ 5 nm SiO2 buffer layer. An ion beam-assisted deposition (IBAD) system (Mill Lane Engineering, Lowell, MA) at the Nanotechnology Laboratory of the University of Nebraska Medical Centre was used. The deposition process is detailed in a previous publication [31]. Throughout the remainder of this manuscript, microcrystalline and nanocrystalline ZrN will be referred to as mc-ZrN and nc-ZrN, respectively.

### 2.2. Irradiation

Ion irradiations were performed at the UKNIBC: Surrey Ion Beam Centre using a 2 MV Van de Graff accelerator. The mc-ZrN pellets and nc-ZrN films were subjected to 600 keV Ar⁺ ion irradiation at room temperature at fluences of up to 1 × 10¹⁷ cm⁻² with a maximum ion flux of 1 × 10¹² ions/cm²/s. The irradiations were performed in a vacuum with a maximum pressure of 9.7 × 10⁻⁶ Pa with the ion beam incident normal to the sample surface. A summary of the irradiation performed in this work is summarised in Table 1.

Estimated damage in displacements per atom (dpa) and at. % Ar concentration was calculated using the Stopping Range of Ions in Matter (SRIM) code [33]. The calculation used the full damage cascades
<caption>**Table 1**</caption>

A summary of ZrN samples irradiated with 600 keV Ar ions at room temperature to various fluences.

<table><thead><tr><th>Fluence (ions/cm²)</th><th>Sample</th></tr></thead><tbody><tr><td>1 × 10¹⁵</td><td>mc-ZrN, nc-ZrN</td></tr><tr><td>5 × 10¹⁵</td><td>mc-ZrN</td></tr><tr><td>1 × 10¹⁶</td><td>mc-ZrN, nc-ZrN</td></tr><tr><td>5 × 10¹⁶</td><td>mc-ZrN, nc-ZrN</td></tr><tr><td>1 × 10¹⁷</td><td>mc-ZrN, nc-ZrN</td></tr></tbody></table>

calculation and threshold displacement energies (E<sub>d</sub>) of 40 eV for Zr and N. Both atoms were set with a lattice binding energy of 3 eV for consistency and ease of comparison to other works. While there are works that show E<sub>d</sub> values for ZrN of 33 and 29 eV for Zr and N [21], respectively, 40 eV has been used in this work. This is because it is typical for groups that work on radiation effects in ceramics to use between 40 and 50 eV, accounting for all elements [34–36]. Fig. 1 shows these estimates as a function of depth from the sample surface, with peak dpa occurring at approximately 250 nm, coincident with peak nuclear-stopping power. Peak Ar concentration sits just beyond 350 nm. Electronic energy loss is projected to steadily decrease from approximately 1.4 keV/nm at the surface, while nuclear energy loss rapidly increases from about 1 to approximately 30.8 keV/nm before falling rapidly once most of the ion energy has been lost. Nuclear-stopping interactions dominate throughout the entire ion range beyond

<figure><img src="image_3.png" /><figcaption>**Fig. 1.** (a) SRIM calculated damage dose and ion concentration profile and (b) stopping powers as a function of depth for 600 keV Ar irradiated ZrN to a fluence of 1 × 10¹⁶ ions/cm². Peak damage occurs at approximately 250 nm from the surface, and peak Ar concentration is at a depth of approximately 350 nm.</figcaption></figure>

approximately 20 nm, becoming the dominant interaction mechanism at an implantation depth of approximately 250 nm. The projected range of 600 keV Ar range is estimated to be approximately 550 nm. During irradiation, the ion beam was rastered at 499 Hz in x and 505 Hz in y over the sample surface to ensure uniform irradiation.

## 2.3. Characterisation

Glancing incidence X-ray diffraction (GIXRD) was used to characterise the phase stability of ZrN post-irradiation. ‘Glancing incidence’ refers to the diffraction geometry whereby the diffraction angle is in the plane normal to the sample surface. This term is used in contrast to ‘grazing incidence’, whereby the diffraction angle is in the plane of the sample surface. A Rigaku Smartlab diffractometer operating at 40 kV and 200 mA with a rotating Cu anode X-ray source (Cu<sub>Kα</sub> = 1.540593 Å) was used for diffraction measurements. Parallel beam optics were used at a constant incident angle of ω = 2.8° and scanning from 20° to 120° in 20 with a 0.03° step size. The incident angle was chosen to restrict the penetration of X-rays to a depth of approximately 600 nm from the specimen surface, thus minimising signal contribution from the underlying pristine material [37]. Incident optics comprise 5° Soller slits, a 0.13 mm height limiting slit and a 5 mm length limiting slit. Receiving optics comprise 5° Soller slits followed by open receiving slits, allowing for the use of a Hypix 3000 detector in 1D mode. By using the detector in this scan configuration, enhanced by the X-ray flux produced by the rotating anode source, we can collect profiles with spatial resolution to a limit of ~ 0.03° FWHM (full-width at half-maximum). Profiles were collected for a total of 1 hour per specimen.

Rietveld refinements were carried out using the *xnd* program [38] to treat GIXRD data and extract crystallite size and microstrain as well as accurate peak parameters such as 2θ position and FWHM. This is achieved by accounting for various instrumental parameters and fundamental thermal factors in the refinement fitting, which must be extracted from peaks to determine quantitative changes. ZrN has no refinable positional parameters. The refined parameters comprised: overall scale, phase ratios, 14 linear background parameters, an overall thermal displacement, finite size broadening, and microstrain. The least-square correlation matrix of the parameters does not show significant correlations (e.g., no coupling of parameters with a correlation element larger than 90 %). *xnd* provides an R-expected value that estimates the statistical quality of each measured pattern, representing the minimum achievable R-factor based on counting statistics. Alternatively, this can be expressed by noting that the refinement’s chi-square is not limited by data quality; a chi-square approaching 1 indicates the residuals are dominated by counting statistics rather than systematic errors. Conventional Rietveld refinement programs do not account for the line broadening introduced using the GIXRD configuration.

Most cross-sectional TEM specimens were prepared using a Helios 5CX FIB/SEM instrument (30 keV gallium ion source). To minimise damage induced by the FIB preparation process, a 0.5 µm thick carbon strap was deposited using the electron beam followed by a 1 µm thick carbon strap and then a 1 µm thick platinum strap, both deposited using the ion beam, providing a barrier to surface damage from gallium ions. Specimens were milled using routine TEM sample preparation whereby a trench is created around the region of interest before proceeding with the conventional *in-situ* lift-out procedure with a foil of approximately 1 µm thickness. After the specimen is attached to a copper grid, it is polished on both sides with progressively reducing voltages and currents to remove damage caused by the former. The final cleaning step was performed using a 2 keV gallium beam with a beam current of 23 pA to remove any remaining damage or contamination. Some of the cross-sectional TEM specimens were prepared by mechanical polishing using a tripod polisher and Ar ion milling (PIPS Model 691, Gatan). The ion milling energy was initially set at 4 kV and then gradually reduced to 2 kV. Difficulties associated with effectively polishing a high-hardness material like ZrN have been highlighted in previous work by Gan
et al. [2].

Most TEM imaging was performed using a JEOL 2100+ (200 keV, LaB₆ emitter) to evaluate the microstructure of pristine and irradiated ZrN samples. Images were collected using a Gatan RIO CMOS camera. A JEOL JEM-3000F operated at 300 kV was also used for TEM observations, and the images and diffraction patterns were recorded on imaging plates.

Mechanical properties of the pristine and irradiated mc-ZrN specimen surfaces were probed with a KLA-Tencor G200 nanoindenter (KLA Tencor, Milpitas, California, USA), equipped with an XP head and a Berkovich diamond tip (Synton-MDP, Nidau, Switzerland). Each indent was performed up to a depth of 800 nm with a constant surface approach velocity of 10 nm/s. Indentation positions were chosen manually using an optical microscope attached to the instrument. Positions were chosen primarily to test larger flat surfaces devoid of cracks, precipitate phases, and contamination. The distance between sets of indents was an order of magnitude greater than the depths of the indents to avoid interactions with induced strain fields. Between 6 and 12 indents were made per depth tested to obtain a statistically good average, however, the exact number of indents varied as some locations did not yield any viable data. Elastic modulus and hardness were determined using the Oliver-Pharr method, and a Poisson ratio of 0.25 was used for the elastic modulus calculation [39].

Surface topography mapping was completed using a Bruker Dimension Icon AFM (Billerica, Massachusetts, USA). Images were acquired using Peak Force tapping mode AFM with a ScanAsyst Air probe (tip radius 2 nm). Raw data were processed and analysed using Bruker NanoScope Analysis software.

The elemental composition as a function of depth from the sample surface was measured using time-of-flight elastic recoil detection analysis (ToF-ERDA) at the UKNIBC Surrey Ion Beam Centre. Samples were probed using 16 MeV $^{127}$I$^{8+}$, with incident and exit angles of 69.5° and a scattering angle of 41°, which forward recoil atoms elastically from the sample into two time-of-flight detectors. Time-of-flight and energy of recoil atoms were measured with 2 timing foils and a gas ionisation chamber, respectively [40]. Depth profile analysis of ToF-ERD coincidence spectra used the Potku software [41].

## 3. Results

Fig. 2 shows the GIXRD profiles of pristine and irradiated mc-ZrN. The pristine peaks were indexed to the fcc crystalline structure of

<figure><img src="image_4.png" /><figcaption>Fig. 2. GIXRD profiles of mc-ZrN irradiated with 600 keV Ar ions to fluences ranging from $1 \times 10^{15}$ to $1 \times 10^{17}$ ions/cm$^2$. The measurements were performed at an incident angle of $\omega = 2.8°$, probing the entire damaged subsurface region. The corresponding fluence is labelled beneath each dataset.</figcaption></figure>

cubic ZrN ($Fm\bar{3}m$). Samples were probed with X-rays at a glancing angle of $\omega = 2.8°$, corresponding to a projected penetration depth of ~ 600 nm, generating a signal from the entire damage region. Fig. 2 shows minor peak position and shape changes with increasing fluence. Peak positions remain relatively constant, but a slight asymmetry develops on the left of peaks with increasing damage. mc-ZrN is extremely tolerant to displacive radiation damage, as indicated by the absence of amorphisation. Minor changes in oxide phases are visible, with monoclinic transforming to tetragonal.

Fig. 3 shows GIXRD profiles of pristine and irradiated nc-ZrN. The pristine nc-ZrN has a preferred growth orientation in the (200) plane, which is removed with increasing fluence. ZrN peak position initially shifts to higher $2\theta$ upon irradiation to the lowest fluence, indicating a contraction of the unit cell. Following this, peaks shift back to lower $2\theta$ values with the unit cell parameter recovering to its as-prepared value. Significant crystallization of $ZrO_2$ starts at a fluence between $1 \times 10^{15}$ and $1 \times 10^{16}$ ions/cm$^2$ and increases with fluence. Oxygen incorporation likely originated from residual gas species (e.g., $H_2O$, $O_2$) in the deposition chamber, despite the low-pressure environment. The native $ZrO_2$ layer on the Zr target and potential impurities in the sputtering gas may have further contributed. The energetic ion bombardment in IBAD enhances surface reactivity, facilitating oxygen uptake even at low partial pressures. Additionally, long-term storage, over 10 years in ambient conditions in this case, likely caused surface oxidation. Given the high grain boundary density in nanocrystalline films, oxygen diffusion into subsurface regions in the films during the storage may have also occurred, leading to deeper oxidation. Furthermore, ZrN peak intensities decrease as the irradiated fluence increases. The fluence of $5 \times 10^{15}$ ions/cm$^2$ was not investigated due to experimental limitations. The changes in GIXRD profiles in Figs. 2 and 3 represent structural modifications that are quantified using Rietveld refinement via the xnd program [38]. All errors are calculated via the xnd program. Fig. 4 shows the strain in pristine and irradiated ZrN samples. Rietveld refinement gives the strain, $\epsilon$, from the refined measured peak positions in $2\theta$, $a$, and the undamaged ZrN system's lattice parameter, $a_0$, using

$$ \epsilon = (a - a_0)/a_0. \qquad (1) $$

mc-ZrN has a compressive strain that is relieved by irradiation, increasing with fluence. nc-ZrN has a tensile strain that is relieved by irradiation to the lowest fluence, before accumulating additional tensile strain with increased fluence.

<figure><img src="image_5.png" /><figcaption>Fig. 3. GIXRD profiles of nc-ZrN irradiated with 600 keV Ar ions to fluences ranging from $1 \times 10^{15}$ to $1 \times 10^{17}$ ions/cm$^2$. The measurements were performed at an incident angle of $\omega = 2.8°$. The corresponding fluence is labelled beneath each dataset. The rate of oxide crystallisation accelerates between fluences of $1 \times 10^{15}$ and $1 \times 10^{16}$ ions/cm$^2$.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>**Fig. 4.** Strain variation as a function of fluence for ZrN samples irradiated with 600 keV Ar ions, extracted from GIXRD data in Figs. 2 and 3 using Rietveld refinement. Data is shown for mc-ZrN (black dots) and nc-ZrN (red dots) samples.</figcaption></figure>

Fig. 5 shows the variation in crystallite size with fluence of mc-ZrN and nc-ZrN. At the lowest fluence, mc-ZrN reduces by 83 %, decreasing by a further 33 % at the highest fluence. The crystallite size of nc-ZrN gradually decreases before tapering off at approximately 6 nm at $5 \times 10^{16}$ ions/cm$^2$. At the highest fluence, the crystallite size of nc-ZrN dramatically decreases by 78 %.

Fig. 6 shows the evolution of microstrain in irradiated mc-ZrN and nc-ZrN specimens. The microstrain is the variance of strain (the second moment of the strain distribution), producing a specific angular dependence of the peak width broadening. It describes the heterogeneity of the strain distribution over the volume probed by GIXRD. In other words, more localised damage leads to higher microstrain. The micro-strain in mc-ZrN shows minor fluctuations despite increasing fluence. The nc-ZrN system exhibits significantly larger microstrain variations than mc-ZrN, indicating the pronounced impact of radiation-induced defects on its microstructure. The nano-scale crystallite size of nc-ZrN makes accurate GIXRD profile fitting challenging, as many variables influence peak shape. Consequently, errors in crystallite size and microstrain are substantial.

AFM topography maps were measured to reveal the swelling be-

<figure><img src="image_7.png" /><figcaption>**Fig. 6.** Microstrain variation as a function of fluence for ZrN samples irradiated with 600 keV Ar ions, extracted from GIXRD data in Figs. 2 and 3 using Rietveld refinement. Data is shown for mc-ZrN (black dots) and nc-ZrN (red dots) samples.</figcaption></figure>

tween the pristine/irradiated interface to confirm the validity of the lattice parameter increase produced by the Rietveld refinement models. Fig. 7 shows the AFM topography map of the interface region of mc-ZrN irradiated to a fluence of $1 \times 10^{17}$ ions/cm$^2$. The contrast between dark gold and light gold between pristine and irradiated regions corresponds to a measurable swelling. Five line scan measurements across different locations on the interface reveal an average volumetric swelling of $6.4 \pm 2.4$ % in the damaged depth. Smaller grains with clearer boundaries were always observed in the irradiated areas, but not in the unirradiated areas. Significantly higher local surface roughness is observed in nano-scale-sized bumps (~10s nm) in the irradiated regions.

GIXRD provides insight into the average long-range structure and its changes due to radiation damage. TEM helps understand how these radiation-induced effects manifest at the micro- and nanoscale, from micron-scale ion tracks to nano-scale defect clusters and impact the average structure and bulk materials' properties. GIXRD results show that the rate of change of sample crystallinity is effectively constant. Hence, we investigate the samples with the highest irradiated fluence in TEM, enabling visualisation of the most contrasting microstructural changes.

The pristine mc-ZrN surface is shown in S.F. 1. The mc-ZrN samples are well-crystallised, with an ~20 nm thick surface layer contrasting the underlying bulk. This layer likely contains contaminants and defects from mechanical polishing.

TEM imaging and SAED analysis of as-prepared nc-ZrN (S.F. 2) confirm nano-scale crystallites, leading to diffuse polycrystalline rings. Indexing these rings shows a cubic lattice with a lattice parameter of ~4.8 Å, midway between ZrN (4.586(2) Å) and ZrO$_2$ (5.135(9) Å). Over a long range, the lattice parameter could average to the midpoint due to high atomic concentrations of O and N shown in ToF-ERDA data (Fig. 11).

S.F. 3 shows the microstructure of mc-ZrN irradiated with 600 keV Ar at a fluence of $1 \times 10^{17}$ ions/cm$^2$. The darker contrast in the BF image is due to lattice distortion. Large, sparsely distributed defect loops are present near the surface (≤ 200 nm), while beyond this depth, defect density increases and loop size decreases. This is attributed to dense displacive collision cascades caused by Ar ions and an increase in point defects. SAED patterns display degrees of periodic and incommensurate ordering at intermediate sites in the irradiated specimens. S.F. 3(a) shows the irradiated surface, including additional cubic spots offset toward the main beam from the ZrN spots, which correspond to crystalline ZrO$_2$ at the surface. There also exist spots condensing at intermediate

<figure><img src="image_8.png" /><figcaption>**Fig. 5.** Crystallite size variation as a function of fluence for ZrN samples irradiated with 600 keV Ar ions, extracted from GIXRD data in Figs. 2 and 3 using Rietveld refinement. Data is shown for mc-ZrN (black dots) and nc-ZrN (red dots) samples.</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 7.** (a) Optical and (b) 50 µm × 50 µm AFM image of mc-ZrN irradiated with 600 keV Ar ions to a fluence of $1 \times 10^{17}$ ions/cm$^2$. The irradiated and pristine regions are labelled, with a distinct interface visible in both images. The AFM image includes an averaged height profile (dashed white square) across the interface, with the corresponding height profile plotted below in (c). The AFM scan reveals a height difference of approximately 38.4 nm between the irradiated and pristine regions, indicating surface swelling or restructuring. A colour scale bar on the right represents the surface height variation in nm.</figcaption></figure>

positions. S.F. 3(b) and 3(c) show the variation in these SAED patterns with increasing depth from the surface. Cubic ZrO$_2$ peaks are no longer present, while the spots that are present at intermediate positions become diffuse and reduce in intensity with increasing depth from the surface. No Ar bubbles or amorphisation was detected in mc-ZrN at the highest irradiation fluence. Unlike other materials, this sample lacks a defect-free denuded zone, suggesting interstitial defect migration to the surface is inhibited by the ZrO$_2$ layer.

Fig. 8 shows the microstructure of nc-ZrN irradiated with 600 keV Ar ions to fluences of $1 \times 10^{15}$ ions/cm$^2$ and $1 \times 10^{17}$ ions/cm$^2$. In nc-ZrN irradiated to a fluence of $1 \times 10^{15}$ ions/cm$^2$, the contrast pattern through the depth of the film has changed compared to the pristine specimen (Fig. 2), indicating a radiation-induced response from the nanocrystallites. No defect loops or other characteristic signs of irradiation damage are present. Figs. 8(a)-(d) reveal minor crystallite coarsening, which becomes most apparent closer to the implanted surface. With both cubic phases of ZrN and ZrO$_2$ present, it can be reasonably said that the dominating contributor to the diffraction patterns is ZrO$_2$,

considering it is the majority phase identified in GIXRD (Fig. 3). Curiously, the end of the ion range is not observable in the microstructure, unlike the specimen irradiated to a fluence of $1 \times 10^{17}$ ions/cm$^2$. However, the effect of the ion range is observable through SAED patterns, with a clear drop off in crystallite coarsening in Figs. 8(c) and 8(d), which are taken at locations beyond the maximum ion depth. Figs. 8(e)-(h) reveal radiation-induced ZrO$_2$ crystallite growth. ZrO$_2$ dominates diffraction rings at the surface, with ZrN being the minority phase. Beyond the ion range, the system remains nanocrystalline, with diffuse rings persisting. Damage accumulation resulting from increased fluence has a pronounced effect on crystallite size via oxidative mechanisms and gas accumulation. These form large blisters, subsurface voids, and cracking, resulting in the ultimate failure of the specimen integrity.

TEM investigation of the surface of the irradiated mc-ZrN is presented in Fig. 9. The sample is tilted slightly from the [112] zone axis of cubic ZrN to obtain a clear contrast. The BF image of Fig. 9(a) shows sparse extended dislocation lines and loops. A continuous layer with a thickness of ~ 20 nm is formed on the topmost layer, which is also observed in the BF image of S.F. 3. The nanobeam electron diffraction patterns, taken with a diameter of ~ 80 nm, from the surface and interior of the sample are shown in Figs. 9(d) and 9(e), respectively. The pattern in Fig. 9(e) is consistent with the (112) reciprocal lattice plane of ZrN. In Fig. 9(d), extra reflections appear inside the ZrN reflections. These reflections can be indexed as the cubic ZrO$_2$, consistent with the GIXRD results. To examine the spatial distribution of ZrO$_2$, dark-field (DF) TEM images of Figs. 9(b) and 9(c) were taken using the 311-type spots due to ZrO$_2$ and ZrN, respectively. These DF images were taken from the identical regions. Only the bulk ZrN shows the bright contrast in Fig. 9 (c), while a continuous bright layer is observed on the topmost layer of ZrN in Fig. 9(b): the surface layer is identified as cubic ZrO$_2$. The ZrO$_2$ exhibits a cubic fluorite structure with an epitaxial relationship with the underlying bulk ZrN of the form $(1\bar{1}1)_{ZrN} // (1\bar{1}1)_{ZrO_2}$, [112]$_{ZrN} // [1\bar{1}2]_{ZrO_2}$.

S.F. 4 shows SAED patterns as a function of depth below the irradiated mc-ZrN surface. Basic cubic reflections of ZrN and ZrO$_2$ are present near the surface. Additional high-intensity peaks exist at commensurate regions (e.g., 1/2 1/2 1/2) below the surface until a depth of ~ 300 nm. Surrounding these spots, we observe diffuse scattering that condenses near incommensurate positions. This diffuse scattering is stronger at depths between 100 nm and 300 nm.

Optical colour variations were observed in specimens, which can indicate changes in the sys-tem's stoichiometry. ToF-ERDA was used to analyse the elemental composition as a function of depth into the sample and identify suspected stoichiometry changes. Fig. 10 presents the ToF-ERDA coincidence spectra and corresponding 1D elemental depth profile for nc-ZrN irradiated with 600 keV Ar to a fluence of $1 \times 10^{17}$ ions/cm$^2$. This technique does not probe the full depth of the film. If it is assumed that these specimens are the theoretical density of ZrN, then the maximum analysis depth would be ~ 200 nm for oxygen and nitrogen to be completely separated in the coincidence spectra. However, due to the high concentration of oxygen inclusions, an accurate estimate of density is difficult. Likely, specimen density changes as a function of irradiated fluence; therefore, the maximum analysis depth will vary between specimens. The regions corresponding to each element show the relationship between the measured energy and time-of-flight. The relationship between kinetic energy, mass, and velocity is responsible for the parabolic nature of the curves. Significant oxidation has occurred at the highest irradiation fluence at the expense of displacing nitrogen in a possible redox reaction. Oxygen, and thus ZrO$_2$ phases, are dominant at the surface of the film, reducing in concentration rapidly from ~ 65 at. % to ~ 45 at. %. Throughout the remainder of the film, zirconium and nitrogen remain at ~ 25 at. %. The impurities of hydrogen and carbon never exceed ~ 5 at. % within the film.

Fig. 11 shows the elemental ratios in nc-ZrN samples extracted from ToF-ERDA data in Figure 10. When calculating the elemental ratios, the
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Cross-sectional TEM micrograph of nc-ZrN irradiated with 600 keV Ar ions to a fluence of (a) $1 \times 10^{15}$ ions/cm$^2$ and (b) $1 \times 10^{17}$ ions/cm$^2$. Corresponding SAED patterns are shown on the right. (c)-(f) and (g)-(j) present the SAED patterns with increasing depth from the surface of samples irradiated to fluences of $1 \times 10^{15}$ ions/cm$^2$ and $1 \times 10^{17}$ ions/cm$^2$, respectively. Slight ZrO$_2$ crystallite coarsening is observed in the damaged region of the sample irradiated to $1 \times 10^{15}$ ions/cm$^2$, and modest coarsening is observed in the sample irradiated to $1 \times 10^{17}$ ions/cm$^2$, with the most pronounced effect at the surface.</figcaption></figure>

surface contamination region was ignored to not be distorted by impu-
rities, so a depth range of 500–1000 thin film units (TFU) was used,
allowing for a good representation of the bulk. All errors are calculated
using the standard quoted errors, being 10 % for hydrogen and 5 % for
all other elements. The trend with increasing implant fluence samples
from the pristine nc-ZrN shows the hydrogen and nitrogen ratios to
zirconium decrease while the oxygen to zirconium ratio increases. The
carbon to zirconium ratio remains unchanged. In all nc-ZrN samples, an
oxide-rich layer is measured at the surface. The trends of O/Zr increase
and N/Zr decrease with fluence in Fig. 11 corroborate the findings of
GIXRD in Fig. 3. The pristine and highest irradiated fluence mc-ZrN

samples were measured using ToF-ERDA. Minimal compositional
changes are confirmed with nitrogen and oxygen ratios to zirconium
remaining similar, with a slight reduction in surface contaminants in the
irradiated specimen. An oxide layer at the surface was also previously
identified in Fig. 9. The oxide layer thickness in mc-ZrN is 1/4 of that in
nc-ZrN. The thickness of these oxide layers differs due to the variations
in their crystallinity, with mc-ZrN forming a distinct crystalline layer
and nc-ZrN forming crystallites dispersed amongst ZrN crystallites.

To explore the correlations between the radiation-induced micro-
structural alterations observed in the TEM images and the long-range
structural changes identified using GIXRD, we investigate the bulk
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** TEM micrographs and nanobeam diffraction analysis of mc-ZrN irradiated to a fluence of $1 \times 10^{17}$ ions/cm$^2$. (a) BF image showing contrasted extended defects. (b) DF image selecting the (331) diffraction spot of ZrO$_2$ showing a structure on the sample surface. (c) DF image of the damaged ZrN matrix. (d) Nanobeam electron diffraction pattern of the surface region in the [112] zone axis containing both the ZrO$_2$ layer and bulk matrix. (e) Nanobeam electron diffraction pattern in the [112] zone axis of only the damaged bulk matrix. Both SAED patterns show evidence of diffuse scattering at intermediate regions between ZrN spots, indicating the formation of short-range ordered structures.</figcaption></figure>

material property changes using nanoindentation. Fig. 12 presents nanoindentation results for mc-ZrN samples post-irradiation. The indent depths ranged from 200 nm to 800 nm.

Notably, the hardness and elastic modulus decreased by approximately 35 % with increasing indent depth in all specimens. All errors were calculated via the KLA Nanosuite software. The most significant change is observed in the sub-surface region of 200 nm, where the hardening increases substantially. The size effect likely explains the general decrease in measured values with depth. However, the relative changes in absolute values and gradient can be attributed to radiation damage. The specimen irradiated to a fluence of $1 \times 10^{17}$ ions/cm$^2$ experiences the most pronounced hardening. It exhibits a maximum hardness increase of 40.5 % compared to the pristine specimen at the same depth. This corresponds to an 8.3 % increase in the modulus. Peak hardening occurs at 200 nm, just before the peak damage region depicted in Fig. 1. Interestingly, the modulus experiences the most significant increase after irradiation to a fluence of $1 \times 10^{15}$ ions/cm$^2$. Beyond this point, the irradiated specimens exhibit higher variability and lower modulus values.

## 4. Discussion

As-deposited nc-ZrN films exhibit tensile strain, which relaxes with irradiation, increasing the lattice parameter with fluence, consistent with prior studies [7,11,15,20]. This may stem from substrate interface

effects or enhanced strain fields around nanoscale crystallites. Initial strain relaxation suggests irradiation-induced densification, annihilating pre-existing defects. At higher fluences, point defect formation and cluster accumulation, coupled with oxygen incorporation replacing nitrogen vacancies (Fig. 11), drive lattice expansion. Once existing defects have been annihilated, the accumulation of point defects and clusters can begin again, increasing the lattice parameter.

Conversely, mc-ZrN shows compressive strain in pristine samples, relaxing to an unstrained state post-irradiation (Fig. 4), with minimal further changes. This swelling, confirmed by AFM (Fig. 7), likely results from point defect accumulation. High defect mobility or self-healing mechanisms in ZrN may enable defect recombination at room temperature, minimising strain [5,9]. However, some studies report increased compressive stress post-irradiation, possibly due to differing defect dynamics or ion energies [4,10].

Literature comparisons reveal varied lattice responses: Lu et al. (300 keV $Ne^+$) observed unit cell contraction with nitrogen loss (N/Zr from 1.02 to 0.88, 1.42 % contraction) [5], while Craciun (800 keV $Ar^+$) reported lattice expansion with hardness and modulus reduction [20]. Wu (600 keV $Kr^+$) noted contraction, grain growth, and softening [18]. The nc-ZrN shows lattice expansion despite significant nitrogen loss (N/Zr = 0.68, Fig. 11), suggesting oxygen-driven oxide formation (corroborated by Fig. 3), not nitrogen vacancies, dominates lattice changes.

Unlike reported radiation-induced grain growth in nc-ZrN and FeZr [9,10,42], GIXRD indicates crystallite size reduction with increasing
<figure><img src="image_12.png" /><figcaption>Fig. 10. (a) ToF-ERDA coincidence spectra and (b) corresponding elemental depth profile for nc-ZrN irradiated with 600 keV Ar to a fluence of 1 × 10¹⁷ ions/cm². A logarithmic colour scale represents the density of detection counts in the coincidence spectra, which can be related to the population of the element. The 1D elemental depth profile was created from the ToF-ERDA coincidence spectra via Potku software [41], showing the relative elemental concentration as a function of depth, shown in thin film units (TFU).</figcaption></figure>

fluence (Fig. 5). However, TEM reveals coarsening (Fig. 8), with ZrO₂ rings intensifying near the surface. GIXRD (Fig. 3) reveals the dominance of the ZrO₂ phase with increasing fluence. Therefore, it is the larger ZrO₂ crystallites that are dominating observations in Fig. 8. These crystallites should not be confused with ZrN phases, which, as Rietveld refinement reveals, are reducing in size. Grain boundaries, rather than acting as defect sinks, accumulate defects, facilitating a redox reaction:

$$2\text{ZrN} + 3\text{O}_2 = 2\text{ZrO}_2 + 2\text{NO} \quad (2)$$

forming ZrO₂ crystallites at the expense of ZrN. Two concurrent processes are occurring: the first is radiation damage that reduces the correlation length of the pristine phase, which makes the nc-ZrN rings more and more diffuse. The second is the chemical reaction with oxygen occurring close to the surface that transforms nitrides into oxides. ZrO₂ grows inward, and its amount continues to increase as long as the surface region is continuously damaged, so the chemical reaction progresses with the fluence. ZrO₂ never passivates the surface in the nanocrystalline system as it has a different volume from ZrN. Therefore, the ZrO₂ film breaks as soon as it grows beyond a given thickness of the order of 20–30 nm. We observe a broadening of the ZrN phase (Fig. 8, more diffuse ring) and a spot sharpening from crystallisation and growth of ZrO₂ (diffuse rings of the early oxidation become bright spots when

<figure><img src="image_13.png" /><figcaption>Fig. 11. Elemental ratios extracted from ToF-ERDA analysis of nc-ZrN irradiated to a maximum fluence of 1 × 10¹⁷ ions/cm² with 600 keV Ar. Oxygen displaces nitrogen and hydrogen in the zirconium matrix to form oxides with increasing fluence. These ratios concern only the bulk material as the surface is excluded to ignore surface contamination.</figcaption></figure>

<figure><img src="image_14.png" /><figcaption>Fig. 12. Room temperature hardness and modulus evolution with depth probed from the surface of nc-ZrN irradiated with 600 keV Ar ions. Linear fits coloured to their corresponding data set are shown to guide the eye.</figcaption></figure>
the film grows inward). The nanoscale grains in nc-ZrN enhance oxygen diffusion, promoting oxidation and nitrogen accumulation, leading to blisters (S.F. 5) or effusion. The mc-ZrN, with larger grains, shows minimal redox, limiting ZrO₂ formation. The SAED pattern in Fig. 8(h) shows a cubic structure with a lattice parameter of ≈ 4.74 Å, 3.9 % larger than pristine mc-ZrN and 2.2 % larger than literature [5]. The diffuse rings, resulting from the majority of the sample, contribute little to the GIXRD signal (Fig. 3). They form the diffuse background at low 2θ angles, while the sharp peaks result from better-crystallised ZrN and ZrO₂ phases near the sample's surface. This explains why we extract a unit cell parameter of 4.59 Å from Rietveld refinement (predominantly from larger crystallites near the surface) instead of 4.74 Å from localised nanocrystallites in the cross-section. Microstrain in nc-ZrN decreases with damage homogenisation (Fig. 6), independent of stoichiometry (0.01 < x < 0.45) [43]. There is difficulty surrounding accurate Rietveld refinement of the pristine mc-ZrN specimen. The GIXRD technique probes a shallow area beneath the surface. An accurate estimate of the diffracted intensities requires a statistical representation of all crystallite orientations. This criterion is well satisfied when the crystallites are significantly smaller than the projected penetration depth of the X-ray beam and the sample is not textured, which is the case of the diffraction patterns obtained in the irradiated samples. On the other hand, these requirements are not satisfied in the pristine samples where the crystallite size typically exceeds the penetration depth of the X-ray beam. Therefore, it is not possible to perform an accurate extraction of the structural parameters, suggesting that the initial decrease in crystallite size in mc-ZrN may be due to a limitation of this technique.

In mc-ZrN, a cubic ZrO₂ epitaxial layer forms (Fig. 9), stabilised by nanoscale size and low interfacial energy [1]. GIXRD (Fig. 2) and TEM (Fig. 9) confirm the cubic phase, with SAED patterns showing short-range ordering and diffuse scattering from nitrogen-vacancy ordering [1,44]. The epitaxial match between ZrN (111) and ZrO₂ (111) (Fig. 13, 12 % mismatch) facilitates crystallisation. Egeland et al. [1] reported similar SAED patterns along the ZrN

[110] zone axis (Fig. 9(d)), attributing split peaks to double-diffraction or superstructure reflections. Our analysis attributes these to epitaxial cubic ZrO₂ on bulk ZrN. Robertson [19] observed nanocrystalline cubic ZrO₂ in mc-ZrN under 4 MeV Au ions, supporting surface-layer formation. Energy deposition in the top ~ 20 nm enhances O and N anion diffusion, promoting recrystallisation. Cubic ZrO₂, stable above 2370 °C, is stabilised by nanoscale size, epitaxial match, and nitrogen incorporation [1,44]. The uniform nanoscale layer (TEM, Fig. 9) and low interfacial energy reduce stabilisation energy. Mechanisms include localised amorphisation and healing or lattice recovery above a damage threshold [1]. As an oxygen conductor, ZrO₂ offers limited oxidation protection, but the denser cubic polymorph may impede oxygen diffusion, limiting growth [1]. No denuded zones appear in mc-ZrN (S.F. 3, Fig. 9), unlike typical damaged materials. The ZrO₂

layer likely blocks larger dislocation loops, increasing diffusion barriers, while point defects and smaller loops annihilate at the phase boundary [1].

Irradiation induces hardening in mc-ZrN (40.5 % at 200 nm, Fig. 12), driven by defect clusters resisting dislocation propagation [2,3]. The nc-ZrN softens due to nitrogen vacancy production [9]. Hardness tapers without saturation, unlike [22] (50 dpa plateau), due to dynamic annealing reducing defect density. The cubic ZrO₂ layer in mc-ZrN could enhance near-surface hardness, acting as a deformation-resistant coating [1]. Modulus increases most at low fluences (25 GPa, Fig. 12), linked to stoichiometry changes (Fig. 11). Numerous conventional mesoscale material characteristics, such as grain size, crystallographic texture, residual stress, and surface roughness, alter a material's strength [45–47]. Mechanical property changes are attributed to radiation-induced effects, including induced defect structure, Ar ion accumulation, and surface oxide phase formation. The implanted Ar modifies the near-surface properties of mc-ZrN, but the mechanical property changes seem to reduce proportionally to ion energy deposition, suggesting they result from induced defect structures rather than Ar ion incorporation.

ZrN's rocksalt structure tolerates high vacancy concentrations, absorbing damage without amorphisation [1]. Oxygen diffusion and nitrogen-vacancy ordering in mc-ZrN prevent strain accumulation, supporting radiation tolerance. The presence of a vacancy defect structure alters atomic-scale factors, including energy band states and electronic structures, which directly affect bonding type, strength, and physical properties like electronic and thermal conductivity [48]. Nitrogen interstitials are mobile at room temperature, while vacancies are immobile. Radiolysis due to electronic stopping and N₂ gas effusion softens the surface, while the surface hardens when the cation sublattice dissipates incoming ionising energy. The low level of ionising damage in this work minimises radiolysis and N₂ effusion, supporting the induced defect structure as the primary factor in material hardening. ZrN is an intermetallic compound which, with high thermal conductivities and ionic-covalent bonding types, is able to effectively dissipate the thermal energy induced by excitations from ionising interactions with 600 keV Ar. That is not to say that higher energy incident ions will not produce an effect. Previous works have investigated ZrN irradiated with SHI's of 200, 167, and 100 MeV Xe and observed distinct effects on defect creation and annealing as a result [49,15,17]. These studies investigate the effects of electronic stopping interactions, whereas this work has attempted to minimise this by imparting an electronic stopping power an order of magnitude lower (1 keV/nm, as opposed to 30 keV/nm). If material synthesis and processing techniques can control the nitrogen vacancy concentration, tuning these systems to exhibit desirable properties for specific applications may be possible.

Diffuse intensity maxima in mc-ZrN's [121] and [111] zone axes (Fig. 9, S.F. 4) match short-range ordering observed in ZrN [6,44] and

<figure><img src="image_15.png" /><figcaption>**Fig. 13.** Proposed epitaxial relationship between ZrN (111) and cubic ZrO₂ (111). Green circles represent Zr atoms of ZrN, while red circles represent Zr atoms of ZrO₂. A 12 % mismatch exists between the primitive unit cells of ZrN and ZrO₂ in this relationship.</figcaption></figure>
transition metal carbides [50,51], as theoretically analysed [52]. The [112] zone axis lacks long-range order, indicating a disordered phase where vacancy probability is influenced by site arrangement, reflecting short-range order [53]. Nitrogen interstitials cause non-stoichiometry in pristine ZrN due to synthesis, but irradiation induces nitrogen-vacancy ordering in the nitrogen sublattice, absent in unirradiated mc-ZrN [1, 6,44,52]. Ion irradiation generates nitrogen vacancies via displacement damage, exacerbated by oxygen displacing two nitrogen atoms, enhancing non-stoichiometry and short-range order. This order occurs between the crystalline ZrO₂ surface layer and unirradiated ZrN beyond the ion range. Oxygen diffusion, decreasing with depth, explains the varying SAED intensity. TEM could quantify the transition region, assessing oxygen migration [54]. A bixbyite-type Zr₂ON₂ (Ia₃) may account for new SAED maxima [55]. No strain accumulation in mc-ZrN suggests compositional adjustments mitigate radiation-induced strain. Nanobeam diffraction is needed to observe the transition range, but higher-energy irradiations may clarify changes by inducing a larger transition area. This effect will be explored in subsequent work. Oxygen introduction typically alters dielectric properties, chemical bonding, and bulk modulus. While bulk modulus increases, lattice parameter changes are minimal, as oxygen substitution counteracts nitrogen vacancy effects [6].

Cubic nitrides, with their rocksalt structure and high vacancy tolerance, readily accommodate nitrogen defects across a wide stoichiometry range [1]. The robust cation lattice absorbs significant radiation damage, preventing defect cluster growth and promoting correlated disorder as the energetically favourable state, underscoring their suitability for long-term nuclear applications.

## 5. Conclusion

GIXRD, TEM, nanoindentation, AFM, and ToF-ERDA revealed radiation-induced microstruc-tural and mechanical changes in ZrN under 600 keV Ar ion irradiation up to $1 \times 10^{17}$ ions/cm². Microcrystalline ZrN (mc-ZrN) maintains its crystalline framework and properties, showing exceptional tolerance to displacive nuclear interactions, while nanocrystalline ZrN (nc-ZrN) is more susceptible due to crystallite size effects. A radiation-induced redox reaction, detected by GIXRD and ToF-ERDA, favoured ZrO₂ formation in nc-ZrN over mc-ZrN. Rietveld refinement of GIXRD data indicated increased unit cell parameters and reduced crystallite sizes in both systems, with AFM confirming swelling in mc-ZrN.

Nanoindentation showed uniform hardness and modulus increases, most pronounced near the surface. Cross-sectional TEM and SAED revealed a cubic ZrO₂ epitaxial layer on mc-ZrN, aligned as (111)ZrN//(111)ZrO₂, [112]ZrN//[112]ZrO₂. SAED patterns confirmed correlated disorder in the damaged matrix, with diffuse scattering matching literature predictions.

This work forms the foundation for assessing ZrN's suitability as a model system of UN, while highlighting the susceptibility of nc-ZrN in displaceive damage environments. Further investigation of oxide crystallisation and defect ordering is needed to explore their physical effects on ZrN in extreme radiation environments. This deeper understanding provides insights into ZrN's suitability as a nuclear structural material or inert fuel matrix.

### CRediT authorship contribution statement

**Connor Beer:** Writing – review & editing, Writing – original draft, Visualization, Validation, Project administration, Methodology, Investigation, Funding acquisition, Formal analysis, Data curation, Conceptualization.
**Mounib Bahri:** Writing – review & editing, Validation, Methodology.
**Gianguido Baldinozzi:** Writing – review & editing, Supervision, Software, Methodology, Formal analysis.
**Manabu Ishimaru:** Writing – review & editing, Resources, Methodology, Investigation,

Formal analysis, Data curation. **Yanwen Zhang:** Writing – review & editing, Resources. **Matthew K Sharpe:** Writing – review & editing, Software, Methodology, Investigation, Data curation. **Callum D McAleese:** Formal analysis, Data curation. **Xinyue Chen:** Writing – review & editing, Validation, Methodology, Investigation, Formal analysis, Data curation. **Fabio Bohns:** Methodology, Investigation. **Troy Manning:** Resources, Methodology. **Nigel D Browning:** Writing – review & editing. **Karl Whittle:** Writing – review & editing, Funding acquisition. **Ross Springell:** Writing – review & editing. **Riaz Akhtar:** Resources. **Maulik Patel:** Writing – review & editing, Visualization, Validation, Supervision, Resources, Project administration, Funding acquisition, Conceptualization.

### Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgements

The authors thank Surrey Ion Beam Centre staff Luke Antwis and Adrian Cansell for running the irradiations and Vlad Palitsin for help in the ion beam analysis laboratory.

The authors wish to acknowledge the Henry Royce Institute for Advanced Materials, funded through EPSRC grants EP/R00661X/1, EP/S019367/1, EP/P02470X/1 and EP/P025285/1, for "Bruker Dimension Icon AFM" access at The University of Sheffield.

The authors wish to acknowledge the University of Liverpool for funding through the W. M.

Horsmann endowment.

The authors wish to acknowledge the EPSRC ATLANTIC consortium (EP/S011935/1).

YZ was supported by the Canada Excellence Research Chair (CERC) program.

The authors thank Jonathan Alaria in the Stephenson Institute for Renewable Energy for use and support with the XRD instrument.

### Supplementary materials

Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.actamat.2025.121391.

### References

[1] G.W. Egeland, Radiation Damage and Fission Product Release in Zirconium Nitride, Ph.D. thesis, Los Alamos National Laboratory (08 2005). doi:10.21 72/884966. https://www.osti.gov/biblio/884966.

[2] J. Gan, Y. Yang, C. Dickson, T. Allen, Proton irradiation study of GFR candidate ceramics, Journal of Nuclear Materials 389 (2) (2009) 317–325, https://doi.org/ 10.1016/j.jnucmat.2009.02.021.

[3] Y. Yang, C.A. Dickerson, T.R. Allen, Radiation stability of ZrN under 2.6 MeV proton irradiation, Journal of Nuclear Materials 392 (2) (2009) 200–205, https:// doi.org/10.1016/j.jnucmat.2009.03.040.

[4] V. Skuratov, V. Uglov, A. Sokhatsky, S. Zlotski, A. van Vuuren, J. Neethling, J. O’Connell, in: Radiation stability of nanocrystalline zrn coatings irradiated with high energy xe and bi ions, 2011, pp. 20–22, 9th International Conference “Interaction of Radiation with Solids”, Vol. 46, Minsk, Belarus, https://inis.iaea. org/search/search.aspx?orig q=RN:46045111.

[5] F. Lu, M. Huang, F. Yaqoob, M. Lang, F. Namavar, C. Trautmann, H. Sun, R. C. Ewing, J. Lian, Displacive radiation-induced structural contraction in nanocrystalline ZrN, Appl Phys Lett 101 (4) (2012), https://doi.org/10.1063/ 1.4738772.

[6] G.W. Egeland, J.A. Valdez, S.A. Maloy, K.J. McClellan, K.E. Sickafus, G.M. Bond, Heavy-ion irradiation defect accumulation in ZrN characterized by TEM, GIXRD, nanoindentation, and helium desorption, Journal of Nuclear Materials 435 (1–3) (2013) 77–87, https://doi.org/10.1016/j.jnucmat.2012.12.025.

[7] A. Janse Van Vuuren, V.A. Skuratov, V.V. Uglov, J.H. Neethling, S.V. Zlotski, Radiation tolerance of nanos-structured ZrN coatings against swift heavy ion irradiation, Journal of Nuclear Materials 442 (1–3) (2013) 507–511, https://doi. org/10.1016/j.jnucmat.2013.02.047.
[8] A. Janse Van Vuuren, J.H. Neethling, V.A. Skuratov, V.V. Uglov, S. Petrovich, The effect of He and swift heavy ions on nanocrystalline zirconium nitride, Nuclear Instruments and Methods in Physics Research, Section B: Beam Interactions with Materials and Atoms 326 (2014) 19–22, https://doi.org/10.1016/j.nimb.2013.10.063.

[9] L. Jiao, K.Y. Yu, D. Chen, C. Jacob, L. Shao, X. Zhang, H. Wang, Radiation tolerant nanocrystalline ZrN films under heavy dose heavy-ion irradiations, J Appl Phys 117 (14) (2015), https://doi.org/10.1063/1.4917381.

[10] S. Khan, I. Ahmed, N. Khalid, M. Mehmood, A. Waheed, M. Malik, Carbon ions irradiation induced modifications in structural and electrical resistivity characteristics of ZrN thin films, Mater Sci Semicond Process 39 (2015) 530–535, https://doi.org/10.1016/j.mssp.2015.05.062.

[11] D. Craciun, G. Socol, G. Dorcioman, D. Simeone, D. Gosset, S. Behdad, B. Boesl, V. Craciun, Ar ions irradiation effects in ZrN thin films grown by pulsed laser deposition, Appl Surf Sci 336 (2015) 129–132, https://doi.org/10.1016/j.apsusc.2014.10.032.

[12] K.B. Kuznetsov, I.A. Kovalev, A.N. Nechaev, A.I. Ogarkov, S.V. Shevtsov, A. S. Chernyavskii, K.A. Solntsev, Stability of the structure of compact zirconium nitride ceramics to irradiation with high-energy xenon ions, Inorganic Materials 52 (12) (2016) 1235–1239, https://doi.org/10.1134/S0020168516120062.

[13] A.I. Ogarkov, S.V. Shevtsov, K.B. Kuznetsov, I.A. Kovalev, A.S. Chernyavskii, K. A. Solntsev, Irradiation of titanium, zirconium, and hafnium nitrides with high-energy ions, Inorganic Materials 52 (6) (2016) 561–565, https://doi.org/10.1134/S002016851606011X.

[14] W. Bao, S. Robertson, J.X. Liu, G.J. Zhang, F. Xu, H. Wu, Structural integrity and characteristics at lattice and nanometre levels of ZrN polycrystalline irradiated by 4 MeV Au ions, J Eur Ceram Soc 38 (13) (2018) 4373–4383, https://doi.org/10.1016/j.jeurceramsoc.2018.05.013.

[15] S. Takaki, M. Takano, N. Ishikawa, Lattice expansion in ZrN and DyO.3Zr0.7N irradiated by 200 MeV Xe ions, Progress in Nuclear Science and Technology 5 (0) (2018) 204–207, https://doi.org/10.15669/pnst.5.204.

[16] V.V. Uglov, G. Abadias, S.V. Zlotski, N.T. Kvasov, I.A. Saladukhin, A. A. Malashevii, Blister formation in ZrN/SiN multilayers after He irradiation, Surface and Coatings Technology 344 (2018) (October 2017) 170–176, https://doi.org/10.1016/j.surfcoat.2018.02.095.

[17] T.F. Mokgadi, M.J. Madito, M. Mlambo, V.A. Skuratov, S.V. Motloung, T. T. Hlathswayo, Slow and swift heavy ions irradiation of zirconium nitride (ZrN) and the migration behaviour of implanted Eu, Nuclear Instruments and Methods in Physics Research, Section B: Beam Interactions with Materials and Atoms 461 (2019) 63–69, https://doi.org/10.1016/j.nimb.2019.09.031 (June).

[18] Z. Wu, Y. Wu, Q. Wang, A comparative investigation on structure evolution of ZrN and CrN coatings against, Heliyon 5 (3) (2019) e01370. doi:10.1016/j.heliyon.2019.01370.

[19] S. Robertson, Impact of Ion Irradiation Damage On SiC and ZrN Mechanical Properties, Ph.D. thesis, Loughborough University (2019). file://ccfepc/Home/F-J/gkarimi/Downloads/s.robertsonthesis.pdf.

[20] D. Craciun, B.S. Vasilie, E. Lambers, H. Makino, V. Craciun, Microstructural investigations of 800 keV Ar ions irradiated nanocrystalline ZrN thin films, Surface Engineering 36 (3) (2020) 326–333, https://doi.org/10.1080/02670844.2019.1668677.

[21] M.M. Rahman, Study On the Production and Aggregation of Defects in Zirconium Nitride, Ph.D. thesis, Graduate School of Engineering Kyushu University, Fukuoka, Japan (2021).

[22] A.J. Terricabras, L. Wang, A.M. Raftery, A.T. Nelson, S.J. Zinkle, Properties and microstructure evolution of silicon nitride and zirconium nitride following Ni ion irradiation, Journal of Nuclear Materials 563 (2022) 153643, https://doi.org/10.1016/j.jnucmat.2022.153643.

[23] Y. Arai, K. Nakajima, Preparation and characterization of PuN pellets containing ZrN and TiN, Journal of Nuclear Materials 281 (2–3) (2000) 244–247, https://doi.org/10.1016/S0022-3115(00)00393-7.

[24] M. Streit, F. Ingold, M. Pouchon, L.J. Gauckler, J.P. Ottaviani, Zirconium nitride as inert matrix for fast systems, Journal of Nuclear Materials 319 (2003) 51–58, https://doi.org/10.1016/S0022-3115(03)00133-8.

[25] J.H. Kim, G.Y. Jeong, S. Kim, Y.J. Jeong, D.S. Sohn, Effect of coating thickness and annealing temperature on ZrN coating failure of U-Mo particles under heat treatment, Journal of Nuclear Materials 507 (2018) 347–359, https://doi.org/10.1016/j.jnucmat.2018.05.019.

[26] T.R. Allen, D. Kaoumi, J.P. Wharry, A.T. Motta, Comprehensive Nuclear Materials 2 (2012) includes the theoretical density of uranium nitride (UN) as 14.32 g/cm³, https://www.elsevier.com/books/comprehensive-nuclear-materials/konings/978-0-08-056033-5.

[27] S. Yu, Q. Zeng, A.R. Oganov, G. Frapper, B. Huang, H. Niu, L. Zhang, First-principles study of Zr-N crystalline phases: phase stability, electronic and mechanical properties, RSC Adv 7 (8) (2017) 4697–4703, https://doi.org/10.1039/c6ra27233a.

[28] K. Mahmood, S. Bashir, F. ul Haq, M. Akram, A. Hayat, M.S. Rafique, A. Mahmood, Surface, structural, electrical and mechanical modifications of pulsed laser deposited ZrN thin films by implantation of MeV carbon ions, Nuclear Instruments and Methods in Physics Research, Section B 448 (October 2018) 61–69, https://doi.org/10.1016/j.nimb.2019.04.013 (2019).

[29] G.W. EGeland, K. Wheeler, P. Peralta, K.J. McClellan, S.A. Maloy, G.M. Bond, Plastic deformation in zirconium nitride observed by nanoindentation and TEM, Journal of Nuclear Materials 416 (3) (2011) 253–261, https://doi.org/10.1016/j.jnucmat.2010.12.002.

[30] D.J. Norris, A.L. Efros, S.C. Erwin, Doped nanocrystals, Science 319 (5871) (2008) 1776–1779, https://doi.org/10.1126/science.1143802.

[31] F. Namavar, C.L. Cheung, R.F. Sabirianov, W.N. Mei, X.C. Zeng, G. Wang, H. Haider, K.L. Garvin, Lotus effect in engineered zirconia, Nano Lett. 8 (4) (2008) 988–996. doi:10.1021/nl072147v.

[32] G. Lee, M.S. Yurlova, D. Giuntini, E.G. Grigoryev, O.L. Khasanov, J. McKittrick, E. A. Olevsky, Densification of zirconium nitride by spark plasma sintering and high voltage electric discharge consolidation: a comparative analysis, Ceram Int 41 (10) (2015) 14973–14987, https://doi.org/10.1016/j.ceramint.2015.08.042. https://doi.org/10.1016/j.ceramint.2015.08.042.

[33] J.F. Ziegler, J.P. Biersack, The Stopping and Range of Ions in Matter, Springer US, Boston, MA, 1985, pp. 93–129, https://doi.org/10.1007/978-1-4615-8103-1_3.

[34] K.L. Smith, M. Colella, R. Cooper, E.R. Vance, Measured displacement energies of oxygen ions in titanates and zirconates, J. Nucl. Mater. 321 (1) (2003) 19–28, https://doi.org/10.1016/S0022-3115(03)00197-1.

[35] Y.H. Li, B.P. Uberuaga, C. Jiang, S. Choudhury, J.A. Valdez, M.K. Patel, J. Won, Y.-Q. Wang, M. Tang, D.J. Safarik, D.D. Byler, K.J. McClellan, I.O. Usov, T. Hartmann, G. Baldinozzi, K.E. Sickafus, Role of antisite disorder on preamphorization swelling in titanate pyrochlores, Phys. Rev. Lett. 108 (2012) 195504, https://doi.org/10.1103/PhysRevLett.108.195504.

[36] A. Taylor, M.K. Patel, J.A. Aguiar, Y. Zhang, M.L. Crespillo, J. Wen, H. Xue, Y. Wang, W.J. Weber, Bubble formation and lattice parameter changes resulting from the irradiation of defect-fluorite gd2zr2o7, Acta Mater 115 (2016) 115–122, https://doi.org/10.1016/j.actamat.2016.05.045. https://www.sciencedirect.com/science/article/pii/S1359645416303950.

[37] F. Jackson, D. Hawkes, X-ray attenuation coefficients of elements and mixtures, Phys Rep 70 (3) (1981) 169–233, https://doi.org/10.1016/0370-1573(81)90014-4.

[38] J.F. Berar, G. Baldinozzi, Xnd code: from x-ray laboratory data to incommensurately modulated phases. Rietveld modelling of complex materials, CPD Newsletter 20 (1998) 3–5.

[39] N. Frantsevich, E.A. Zhurakovskii, A.B. Lyashchenko, Elastic constants and characteristics of the electronic structure of certain types of refractory compounds obtained by sintering., Izv. Akad. Nauk SSSR, Neorg. Mater, 3: 8–16(Jan. 1967). (01 1967). https://www.osti.gov/biblio/4384248.

[40] H.R. Cox, M.K. Sharpe, C. McAleese, M. Laitinen, J. Dulai, R. Smith, J. England, W. H. Ng, M. Buckwell, L. Zhao, S. Fearn, A. Mehonic, A.J. Kenyon, The role of hydrogen in ReRAM, Advanced Materials 2408437 (2024) 1–10. doi:10.1002/adma.202408437.

[41] J.Julin Arstila, M. Laitinen, J. Aalto, T. Konu, S. Kärkkäinen, S. Rahkonen, M. Raunio, J. Itkonen, J.-P. Santanen, T. Tuovinen, T. Sajavaara, Potku – new analysis software for heavy ion elastic recoil detection analysis, in: 331, 2014, pp. 34–41, https://doi.org/10.1016/j.nimb.2014.02.016, 11th European Conference on Accelerators in Applied Research and Technology, https://www.sciencedirect.com/science/article/pii/S0168583X14002717.

[42] A.T.Motta Kaoumi, R.C. Birtcher, Influence of alloying elements on grain-growth in Zr(Fe) and Cu(Fe) thin-films under in situ ion-irradiation, Journal of Nuclear Materials 382 (2–3) (2008) 184–189, https://doi.org/10.1016/j.jnucmat.2008.04.011.

[43] A.J.M. Alexandr, I. Gusev, Andrej A. Rempel, Disorder and Order in Strongly Nonstoichiometric Compounds, Springer, Berlin, Heidelberg, 2001. https://link.springer.com/book/10.1007/978-3-662-04582-4.

[44] P. Li, J.M. Howe, Short-range ordering and 1,1/2,0 diffuse intensity maxima in ZrNx, Acta Mater 51 (5) (2003) 1261–1270, https://doi.org/10.1016/S1359-6454(02)00506-2.

[45] R.W. Harrison, W.E. Lee, Mechanism and kinetics of oxidation of ZrN ceramics, Journal of the American Ceramic Society 98 (7) (2015) 2205–2213, https://doi.org/10.1111/jace.13575.

[46] W.J. Chou, G.P. Yu, J.H. Huang, Deposition of TiN thin films on Si(100) by HCD ion plating, Surface and Coatings Technology 140 (3) (2001) 206–214, https://doi.org/10.1016/S0257-8972(01)01120-3.

[47] H. Jiménez, E. Restrepo, A. Devia, Effect of the substrate temperature in ZrN coatings grown by the pulsed arc technique studied by XRD, Surface and Coatings Technology 201 (3–4) (2006) 1594–1601, https://doi.org/10.1016/j.surfcoat.2006.02.030.

[48] Z.B. Qi, P. Sun, F.P. Zhu, Z.C. Wang, D.L. Peng, C.H. Wu, The inverse Hall-Petch effect in nanocrystalline ZrN coatings, Surface and Coatings Technology 205 (12) (2011) 3692–3697, https://doi.org/10.1016/j.surfcoat.2011.01.021.

[49] C. Han, G.Q. Lin, C. Dong, K.P. Tai, X. Jiang, Influence of nitrogen vacancy concentration on mechanical and electrical properties of rocksalt zirconium nitride films, Acta Metallurgica Sinica (English Letters) 30 (11) (2017) 1100–1108, https://doi.org/10.1007/s40195-017-0636-x.

[50] J.-M. Costantini, G. Gutierrez, G. Lelong, M. Guillaumet, M.M. Rahman, K. Yasuda, Raman spectroscopy study of damage in swift heavy ion-irradiated ceramics, Journal of Raman Spectroscopy 53 (9) (2022) 1614–1624. arXiv:https://analyticalsciencejournals.onlinelibrary.wiley.com/doi/pdf/10.1002/jrs.6414. 6414.

[51] J. Billingham, P.S. Bell, M.H. Lewis, Vacancy short-range order in substoichiometric transition metal carbides and nitrides with the NaCl structure. I. Electron diffraction studies of short-range ordered compounds, Acta Crystallographica Section A 28 (6) (1972) 602–606, https://doi.org/10.1107/S0567739472001524.

[52] J. Buršík, G.C. Weatherly, A study of short-range order in A F.C.C.-based stoichiometric a:0.5B:0.5 solid solution, Physica Status Solidi (A) Applied Research 176 (2) (1999) 835–842, https://doi.org/10.1002/(SICI) 1521-396X(199912)176:2<835:AID-PSSA835>3.0.CO;2.

[53] E.Parthé Sauvage, Vacancy short-range order in substoichiometric transition metal carbides and nitrides with the NaCl structure. II. Numerical calculation of vacancy
arrangement, Acta Crystallographica Section A 28 (6) (1972) 607–616, https://doi.org/10.1107/S0567739472001536.

[54] G. Kostenko, A.A. Rempel, S.V. Sharf, A.V. Lukoyanov, Short-range order in disordered transition metal oxides, carbides, and nitrides with the B1 structure,

Physics of the Solid State 57 (4) (2015) 637–651, https://doi.org/10.1134/S1063783415040137.

[55] S.J. Clarke, C.W. Michie, M.J. Rosseinsky, Structure of Zr2ON2 by neutron powder diffraction: the absence of nitride-oxide ordering, J Solid State Chem 146 (2) (1999) 399–405, https://doi.org/10.1006/jssc.1999.8372.