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# Polyepitaxial grain matching to study the oxidation of uranium dioxide

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Jacek Wasik<sup>1, ∗</sup>, Joseph Sutcliffe<sup>1</sup>, Renaud Podor<sup>1, †</sup>, Jarrod Lewis<sup>1, ‡</sup>, James Edward Darnbrough<sup>1,3</sup>, Sophie Rennie<sup>1</sup>, Syed Akbar Hussain<sup>1</sup>, Christopher Bell<sup>1</sup>, Daniel Alexander Chaney<sup>1,4</sup>, Gareth Griffiths<sup>1</sup>, Lottie Mae Harding<sup>1</sup>, Florence Legg<sup>1</sup>, Eleanor Lawrence Bright<sup>1,4</sup>, Rebecca Nicholls<sup>1</sup>, Yadukrishnan Sasikumar<sup>5</sup>, Angus Siberry<sup>1</sup>, Philip Smith<sup>1</sup> & Ross Springell<sup>1, ‡</sup>

Although the principal physical behaviour of a material is inherently connected to its fundamental
crystal structure, the behaviours observed in the real-world are often driven by the microstructure,
which for many polycrystalline materials, equates to the size and shape of the constituent crystal
grains. Here we highlight a cutting edge synthesis route to the controlled engineering of grain
structures in thin films and the simplification of associated 3-dimensional problems to less complex 2D
ones. This has been applied to the actinide ceramic, uranium dioxide, to replicate structures typical in
nuclear fission fuel pellets, in order to investigate the oxidation and subsequent transformation of
cubic UO₂ to orthorhombic U₃O₈. This article shows how this synthesis approach could be utilised to
investigate a range of phenomena, affected by grain morphology, and highlights some unusual results
in the oxidation behaviour of UO₂, regarding the phase transition to U₃O₈.

The nature of grains and grain boundaries is crucial to our understanding of the behaviour of polycrystalline materials<fnref n="1" />⁻⁴. The relationship between these defect-dense regions and the misaligned crystal grains can be affected by geometries, such as tilt and twist angles, and by impurities, which tend to precipitate and segregate⁵,⁶. All this complexity can have dramatic consequences on critical properties, such as mechanical strength, electrical and thermal conductivities, and corrosion resistance⁵⁻¹⁰. In this paper, we demonstrate an approach to controlled grain engineering, which reduces a three dimensional problem to two dimensions, yielding new results in an important applied materials system in the nuclear industry.

Uranium dioxide is the most prevalent fission fuel and is used world-
wide in over 400 nuclear reactors¹¹. Typical fuel pellets are small cylinders of
1 cm height and diameter and are placed inside fuel pins several metres long.
There are tens of thousands of fuel pins in a reactor core, resulting in more
than 10 million fuel pellets in an average, water-moderated reactor¹¹. This
means that there are more than 30,000 tonnes of UO₂ fuel currently in
operation. This fuel is most commonly manufactured by sintering pow-
dered UO₂ at high temperature and pressure to achieve densities
approaching 97%¹².

During operation the fission process causes severe structural trans-
formations, driven by the kinetic energy of daughter nuclei and emitted
neutrons. There is an enormous amount of heat deposited in the fuel pellets,
which is transported away by the coolant in the core, however, the

fundamental thermal conductivity of UO₂ is particularly poor and this
results in large thermal gradients across only a few mm. Centreline tem-
peratures can reach 1500 °C with the outer edge of the fuel pellet just a
300 °C¹³. These effects cause grain growth in the centre of the pellet, with
average grain sizes of about 50 microns, becoming smaller radially outwards,
down to an average size of 10 microns in the so-called rim-structure¹²,¹⁴⁻¹⁷
(see Fig. 1A, C). The grain structure is particularly important in two crucial
stages of the fuel lifetime; (1) affecting the thermal transport and migration
of fission products during operation and (2) affecting oxidation and cor-
rosion post-operation during interim and final storage, especially in aqu-
eous environments¹²,¹⁴⁻¹⁹.

The use of thin film deposition techniques to synthesise fission fuel
materials is gaining traction as a viable and, in some cases, preferred route
for test and experimentation<sup>20,21</sup>. The bulk fuel material is complex making it
difficult to deconvolve the many contributing factors. Spent nuclear fuel
presents even more of a problem as the handling of such highly radioactive
material is only possible in a handful of specialist laboratories. A thin film
approach to fuel synthesis creates a macroscopic surface with limited
thickness, typically tens to hundreds of nanometres, where radioactivity
levels are very low; a 100-nm thick, 1 cm<sup>2</sup> sample is 1.5 Bq. It is then possible
to control structural and chemical parameters, such as crystallographic
orientation, strain, and stoichiometry<sup>21</sup>. We present a unique mode of
growth that can engineer the grain size of high temperature ceramics, such

<footnote n="1">¹IAC, School of Physics, University of Bristol, Tyndall Avenue, Bristol, Avon, UK. ²ICSM, Bagnols-sur-Cèze, Site de Marcoule, France. ³Department of Materials, University of Oxford, 21 Banbury Road, Oxford, UK. ⁴European Synchrotron Radiation Facility, 71 Avenue des Martyrs, Grenoble, France. ⁵Oak Ridge National Laboratory, 1 Bethel Valley Road, Oak Ridge, TN, USA. ✉e-mail: jacek.wasik@bristol.ac.uk; phrss@bristol.ac.uk</footnote>
<figure><img src="image_3.png" /><figcaption>**Fig. 1 | Variation of UO₂ grain size within nuclear fuel and replication method.**
A A cross-section of irradiated nuclear fuel¹⁵. B EBSD data collected from UO₂ thin films samples deposited onto un-annealed, 2 h and 8 h-annealed polycrystalline YSZ substrates, with an average grain size of 14.2 ± 0.2 µm, 31.7 ± 0.3 µm, and 45.8 ± 0.4 µm, respectively, demonstrating the ability to engineer samples with</figcaption></figure>

desired grain structure. C Schematic representation of thermal gradient (red 1500 °C → blue 300 °C) across the fuel pellet and resulting increase in the uranium dioxide grain size. D Schematic of the reduction of a three-dimensional granular structure into two dimensions by producing nanometre-thick columns of micon-size diameter grains.

as UO₂, without requiring the extremely high temperatures required in bulk materials. This process relies on a 1:1 substrate match, combined with careful heat treatment and surface preparation of substrate materials, before deposition.

The pre-treatment of the substrate material can be used to selectively engineer desired average grain sizes (see Fig. 1C). The columnar-type growth of these epitaxial grains means that although the lateral grain dimensions mimic those of the underlying polycrystal, the in-plane microstructure is uniform along the growth direction. This approach has distinct advantages for any depth-dependent studies, as a complex 3D problem is effectively reduced to two dimensions (see Fig. 1D). This approach can be adapted to a wide variety of materials, where grain size-dependent effects are of interest. We have applied this method to uranium dioxide, which is most familiar in a small cylindrical sintered pellet form for nuclear fuel. Here we present the first images of UO₂ synthesised in this way in order to investigate the grain orientation-dependent oxidation to U₃O₈. This process follows the oxidation sequence as: UO₂ ⇒ U₄O₉ ⇒ U₃O₇ ⇒ U₃O₈, with nucleation and growth kinetics⁶,¹⁰,¹²,²²⁻²⁹.

This is of direct industrial interest in the initial processing and eventual disposal of nuclear fuel. The precise mechanism of UO₂ oxidation is still not

fully understood and there are outstanding questions regarding the transformation of UO₂ to U₃O₈. Here we have discovered some peculiar results, which could shed light on a specific crystallographic route for this phase transition.

## Results

### Sample synthesis and grain engineering

The UO₂/YSZ system (YSZ is yttria-stabilised zirconia) is an example of a 1:1, or so-called cube on cube matched epitaxial relationship, used to produce single crystal films to study surface and interfacial effects on pristine fuel material³⁰,³¹. The ability to match grains of UO₂ onto YSZ is demonstrated in Fig. 1B, which shows electron-backscatter images from the surface of 100-nm polycrystalline films of stoichiometric UO₂, where the final deposited surface requires no further mechanical treatment (avoiding possible influence of mechanical polishing and defect formation), typically necessary to achieve high-quality electron backscatter diffraction (EBSD) patterns from bulk material. Polycrystalline substrates of YSZ were heat-treated and polished flat before reactive dc magnetron sputtering of UO₂, which required no further treatment to get the observed patterns. The three samples shown here are chosen to represent three positions across a fuel
pellet and have grain orientation distribution akin to that found in bulk material, but with average grain sizes of 14.2 ± 0.2 µm, 31.7 ± 0.3 µm and 45.8 ± 0.4 µm (errors here are represented by the standard error). In a bulk material any measurement of the surface properties that is any deeper than the first few atomic layers will be affected by the grain structure along the surface normal. Our polypoitaxial growth mechanism avoids these issues, creating columnar grains matched along the underlying YSZ grain direc- tions, where the lateral grain structure remains constant as a function of depth (see Fig. 1D). This provides an ideal scenario to study and engineer grain effects for a wide range of materials properties. Here we focus on the effect of grain size, orientation and grain boundary volume on the oxidation of UO2.

X-ray diffraction investigation of oxidation process

The lattice parameter of the as-grown samples was 5.476(2) Å, similar to
values reported for the bulk material<sup>12,32-34</sup>. The polycrystalline character of
the sample was confirmed by comparing measured integrated intensities
with predicted structure factor intensities. Based on the (111) being the
strongest intensity 100%, the measured values for (220) and (002) were
48 ± 3% and 34 ± 3%, with expected theoretical values 47% and 36%,
respectively, exhibiting minimal preferential texture. The lattice parameter
as a function of x in the UO<sub>2+x</sub> system was observed to have non-Vegard
character as reported by Elorrieta et al.<sup>35</sup>. The oxidation of polypoitaxial UO<sub>2</sub>
thin films was performed in two stages. In the first stage, the sample was
exposed to 200 mbar of O<sub>2</sub> at 150 °C for 16 h (Supplementary Fig. 1). The
lattice constant first shrinks in the region below x = 0.13, and then slightly
expands for 0.13 < x < 0.17 before shrinking again for x above 0.17. The final
product at this stage was identified as tetragonal U<sub>3</sub>O<sub>7</sub>, as c is 1.9% longer
than a<sup>36-38</sup>. No further oxidation was observed, as expected in this tem-
perature range<sup>10,22</sup>.

In the second stage of the experiment, the pressure of oxygen was kept constant at 200 mbar, and the temperature was increased to 300 °C. The Longitudinal scans were collected every 0.5 h for 48 h, and a selection of these are shown in Fig. 2A. A general decrease in the intensity is observed for (111), (002) and (220) reflection assigned to the UO<sub>2+x</sub> structure as the cubic/tetragonal structure undergoes a phase transition to the orthorhombic structure of U<sub>3</sub>O<sub>8</sub>. This coincides with an increase in the observed intensities of the U<sub>3</sub>O<sub>8</sub> reflections within the XRD pattern. To better illustrate this process, an area under each peak was extracted, normalised to the highest value and plotted against time as shown Fig. 2B.

The kinetics of oxidation from UO₂+x to U₃O₈ presented in Fig. 2B. are
in good agreement with the sigmoidal nucleation-and-growth mechanism
described in literature for this phase transition¹²,²²,²⁵,³⁹⁻⁴². The induction
period, where the initial rate of oxidation is very low, takes place during first
the 10 h of the oxidation process. After this time, the oxidation rate gradually
increases to a maximum, and this linear region¹² is observed from around 13
to 21 h. When the reaction approaches completion, it tails off, finishing after
44 h with the sample completely oxidised to U₃O₈.

Data presented here is in contradiction to what was reported by Allen et al.<sup>23,24</sup>, where the oxidation rate of UO<sub>2</sub> was suggested to be highest for the [111] orientation due to epitaxial relations with U<sub>3</sub>O<sub>8</sub>, followed by [110] and slowest for [001]. The [001] oriented grains within the polypoitaxial sample oxidised quickest, and no UO<sub>2+x</sub> {002} reflection was observed after 23 h. There is no significant difference between kinetics observed for the [111] and [110] oriented grains. This might suggest that there is potentially crystallographic relationship between the [001] orientation of UO<sub>2+x</sub> and orthorhombic U<sub>3</sub>O<sub>8</sub>.

In situ electron observation of oxidation

To further explore the potential of a polypoitaxial system, a UO₂ sample was oxidised in situ in a High Temperature Environmental Scanning Electron Microscope (HT-ESEM). Prior to oxidation an area of sample was marked and mapped using EBSD. Selected images taken during this in situ oxidation experiment are shown in Fig. 3A–C. Clear cracking of the surface is observed as a result of the 36% volume expansion<sup>12,22,29,35</sup> associated with transition to U₃O₈. The kinetics extracted from this process as the percentage of the damaged area to the entire area are provided in Fig. 3E. The characteristic nucleation and grow character is also independently observed in this experiment.

The ignition centres, observed as cracks related to oxidation, were extracted and are indicated by red markers on the overlay of SEM and EBSD maps in Fig. 3D. In 75% of cases the oxidation-related cracking was observed at the grain boundaries, which is expected and has been previously demonstrated that oxidation preferentially occurs along the grain boundaries, and then advances from boundary to the centre of the grains<sup>6,12</sup>. This shows that faster oxidation rate correlate to regions of UO₂ with higher grain boundary density. Post-oxidation examination of the sample revealed U₃O₈ grains within the polypoitaxial sample that did not disintegrate, Fig. 3F, G. This would suggest a unique crystallographic relationship between UO₂ and U₃O₈ along the UO₂ [001] direction.

<figure><img src="image_4.png" /><figcaption>**Fig. 2** | **In situ XRD during oxidation of polypoitaxial UO<sub>2</sub> thin film.** A Example longitudinal XRD patterns collected during oxidation polypoitaxial UO<sub>2</sub> thin film at 200 mbar of oxygen at 300 °C over 48 h. The right peak on the UO<sub>2</sub> (220) column is associated with the XRD sample stage. Data is represented by open circles and fit by</figcaption></figure>

line. B Normalized Bragg's peaks area for 3 main orientations of cubic UO₂ system and orthorombic U₃O₈ plotted over time. The data shows faster oxidation along (002) orientation, and characteristic nucleation and growth kinetics.
<figure><img src="image_5.png" /><figcaption>**Fig. 3** | In situ HT-ESEM oxidation of polypitaxial UO₂ thin films. A SEM image of marked area of polypitaxial UO₂ thin film before oxidation, B example of an SEM image taken during partial oxidation to U₃O₈, and C SEM image taken after full oxidation to U₃O₈. The cracking of the thin films is attributed to the volume expansion associated with this phase transition. D an overlap of EBSD map and SEM image. Red markers indicates locations where cracking related to formation of U₃O₈ started. From a total of 28 points where the cracking started, 21 points were located on the grain boundaries (75%). This indicates that not only the crystallographic</figcaption></figure>

orientation of the grain has to be taken into consideration, but also the density of the grain boundaries. E Kinetics extracted from the SEM images based on the ratio of the fractured area to the entire area. With reference to (A–C), data are represented by open circles, and the fit is represented by a solid red line. Typical nucleation and growth mechanism is observed. F SEM image taken after sample oxidation showing some grains that did not lose integrity. G Higher magnification of a U₃O₈ grain mirroring the initial UO₂ grain.

## Discussion

The combined use of XRD, SEM and EBSD on polypitaxial UO₂ deposited onto specially engineered YSZ substrates has demonstrated that it is possible to mimic the grain structure of the nuclear fuel. This was achieved through different heat annealing treatments of the substrate producing similar variations in the grain size morphology of the thin film to that observed in pristine and irradiated sintered pellets of UO₂. The polycrystalline character of the system, was confirmed by XRD. EBSD results showed that the grain size could be controlled within 10–50 µm without needing to expose UO₂ to high temperatures. Although, other have discussed grain-to-grain epitaxy<sup>43,44</sup> or domain-matching epitaxy<sup>45</sup>, none has realised the power of this method to controllably engineer grain structure to study material properties in the way presented here. This new high-quality system with constant lateral grain structure, and ability to acquire desired grain boundary density by controlling the average grain size, could open new areas of further investigation.

Here we have used it to look at the UO₂ ⇒ U₃O₈ oxidation transformation. Data collected using in situ XRD and SEM approaches has shown good agreement with the characteristic kinetics of the transition¹², and the importance of the grain boundary density¹,⁴⁶⁻⁴⁹. The SEM data confirms recent simulation studies, were it has been shown that oxidation predominantly takes place along the grain boundaries before advancing from the boundaries towards the centre of grains⁶. Our data shows that the [001] orientation is most readily oxidised, in contradiction to previous studies which report a [111] > [110] > [001] order of oxidation rate²³,²⁴. The sample system demonstrated here can be used for further studies to confirm faster oxidation in fuel with finer grain due to the faster oxygen diffusion in the grain boundary³,⁶,¹².

Notably, although most grains transforming from UO₂ to U₃O₈ experienced a high degree of disintegration due to the extreme volume expansion (36%), a small number of grains remained laterally unaffected,

and appeared to expand along the surface normal. This suggests a possible crystallographic relationship between UO₂ and U₃O₈ that needs further study.

## Methods

### Reactive DC magnetron sputtering

Polyepitaxial thin film samples were fabricated using reactive DC magnetron sputtering system at the FaRMS (Facility for Radioactive Materials Surfaces) at the University of Bristol²¹. The system operates at a base pressure of 1 × 10⁻¹⁰ mbar, with 5.5N argon used as the main sputtering gas at a pressure of 7.3 × 10⁻³ mbar. To grow polypitaxial UO₂ a depleted uranium target was used, and a partial pressure of 2 × 10⁻⁵ mbar of O₂. Substrates were held at elevated temperature of 650 °C to assist crystalline growth of UO₂. With calibrated deposition rate, the growth time was controlled to achieve roughly a 100-nm-thick layer of UO₂²¹.

### Substrate preparation

To investigate the effect of grain structure on UO₂ oxidation, polypitaxial UO₂ thin films were grown on polycrystalline ceramic of yttrium stabilized zirconia (Y₂O₃)₀.₀₈(ZrO₃)₀.₉₂ substrates obtained commercially from MTI Corporation, with dimensions 10 × 10 × 0.5 mm, polished to <100 Å root mean square (rms) roughness.

The structure of uranium dioxide is cubic fluorite, with space group of UO₂ is Fm$\bar{3}$m, No. 225, with a lattice constant at room temperature 5.47 Å¹²,³²⁻³⁴. A good match between the thin film and the substrate helps to reduce lattice strain and structural defects.

In the work of Bao et al.⁵⁰ a match between [001] oriented thin film of UO₂ and the [001]-oriented substrates of calcium fluoride (CaF₂) and lanthanum aluminate (LAO) was reported. The yttria-stabilized zirconia (YSZ) substrate to grow a [001] single crystal of UO₂ was reported by Strehle et al.³⁰. Later all three orientations [111], [110], and [001] of YSZ were used
by Rennie et al.³¹ to produce single crystals of UO₂ for corrosion studies. In this project, YSZ was selected due to its suitability for surface preparation, given its ease of mechanical polishing and relatively straightforward induction of grain growth.

Lateral grain sizes of the YSZ substrates was increased by annealing them in a tube furnace. This process was carried out in air atmosphere, at 1600 °C. The heating ramp rate was set to 15 °C/min, and the same rate was used for cooling down. During heat treatment the roughness of the material increased due to changes in the grain size and shape, to obtain a smooth and clean surface, polishing and cleaning processes were applied.

To assist with the surface preparation process, annealed substrates were encapsulated in ClaroCit resin (Struers). The first stage of polishing process for one side polished YSZ was P1200, followed by P2500 and P4000 silicon carbide (SiC) abrasive papers (Buehler). Time needed for each step was between 3 and 20 min, based on the surface quality assessment performed using an optical microscope. The polishers used during this stage were Buehler Metaserv, operated at speed of 600 rpm, with water flow. The ECOMET polishers were used in next stage. Diamond pastes with 3, 1, 0.25, and 0.1 µm sizes were used, together with MetaDi fluid lubricant (Buehler). The three polishers with higher grades, use MasterTex polishing cloths fitted to a wheel, while the last one was fitted with MicroCloth. The last step of the polishing process involves use of a Vibro-Polisher (with 50 nm of Colloidal Silica), at medium vibration speed for 24 h.

The initial cleaning stage involved ultrasonicating the substrate in a 500 ml beaker filled with 250 ml acetone for 1 h to remove larger resin contaminants from the sample. Once this step was completed, the sample was rinsed with acetone and underwent an additional 45 min of ultrasonication in clean acetone. Subsequently, to further eliminate organic contaminants, the substrates were immersed in ethanol and subjected to ultrasonic treatment for 1 h, then rinsed with ethanol. The same procedures were then repeated using isopropanol. Following the cleaning process with these three solvents, the surface was dried using a flow of warm, clean air to prevent residue buildup.

## In situ X-ray analysis

X-ray diffraction (XRD) measurements were performed using a Philips X'Pert PRO MPD diffractometer with a Cu-Kα source, configured for Bragg-Brentano diffraction geometry, with specialized mirror designed to reflect Cu-Kβ away from the detector. The footprint width of the X-ray beam on the sample surface was 10 mm, and the incident programmable divergence slit, located 240 mm away from the sample, was set to 0.5'. The step size was 0.05' and counting time per step was 2s, with sample being held stationary. The in situ oxidation experiment was performed utilizing the Anton Paar HTK 1200 high temperature environment chamber, each time at a constant oxygen pressure of 200 mbar. Lattice parameters were extracted using a least-squares fitting routine of Pseudo-Voigt peak profiles.

## SEM characterization

The initial sample characterisation, including EBSD, measurement were conducted at an acceleration voltage of 30 keV and aperture of 120 µm using Zeiss Sigma High Definition Variable Pressure Field Emission SEM with EDAX Electron Backscatter Diffraction detector. To analyse the data OIM Analysis™ (EDAX) software was used. The same microscope was used to characterise sample surface post-oxidation.

The in situ SEM oxidation experiment was performed using a HT-ESEM at the Institut de Chimie Separative de Marcoule (ICSM) in Marcoule, France, model: FEI Quanta 200 FEG ESEM, Thermo Fisher Scientific, Massachusetts, USA. The insertion was performed in air. Samples were heated under an atmosphere of 3.5 mbar of oxygen. Once the signs of the oxidation were visible, temperature was held constant. The sample regions underwent continuous monitoring throughout the entire duration of the experiment, with images being recorded at intervals of 3–5 s. All pictures were recorded with the same magnification of ×250, corresponding to an area of 512 × 470 µm².

Fiji ImageJ³¹ software, with in-built functions, was used to analyse the surface changes during the sample treatment. The sample surface cracking

due to oxidation and further expansion of the crack propagation was monitored by measuring the percentage of the damaged surface to the entire ROI. The background of each photo was generated using a Gaussian-blurred version of the image and then subtracting it from the original. To enhance the contrast between the oxidised and the pristine areas of the sample, a series of find edges and smooth functions were used. In the next step Gaussian-blurred or Medina filters were applied.

## Data availability

The data will be available in the University's Data Repository in a form suitable for long-term retention and wider publication. Available from the corresponding author upon request.

Received: 7 December 2023; Accepted: 23 May 2024;

Published online: 11 July 2024

## References

1. Clarke, D. R. Grain boundaries in polycrystalline ceramics. *Annu. Rev. Mater. Sci.* **17**, 57–74 (1987).

2. Watanabe, T. Grain boundary engineering: historical perspective and future prospects. *J. Mater. Sci.* **46**, 4095–4115 (2011).

3. Cantwell, P. R. et al. Grain boundary complexion transitions. *Annu. Rev. Mater. Res.* **50**, 465–492 (2020).

4. Krause, A. R. et al. Review of grain boundary complexion engineering: know your boundaries. *J. Am. Ceram. Soc.* **102**, 778–800 (2019).

5. Yajima, S., Furuya, H. & Hirai, T. Lattice and grain-boundary diffusion of uranium in UO₂. *J. Nucl. Mater.* **20**, 162–170 (1966).

6. Liu, M., Kang, Q. & Xu, H. Grain-scale study of the grain boundary effect on UO₂ fuel oxidation and fission gas release under reactor conditions. *Chem. Eng. Sci.* **229**, 116026 (2021).

7. Sabioni, A. C. S., Ferraz, W. B. & Millot, F. Effect of grain-boundaries on uranium and oxygen diffusion in polycrystalline UO₂. *J. Nucl. Mater.* **278**, 364–369 (2000).

8. Millett, P. C., Tonks, M. R. & Biner, S. Grain boundary percolation modeling of fission gas release in oxide fuels. *J. Nucl. Mater.* **424**, 176–182 (2012).

9. Andersson, D. A. et al. Multiscale simulation of xenon diffusion and grain boundary segregation in UO₂. *J. Nucl. Mater.* **462**, 15–25 (2015).

10. Lim, H. C. et al. Effects of microstructural constraints on the transport of fission products in uranium dioxide at low burnups. *J. Nucl. Mater.* **477**, 24–36 (2016).

11. Bayraktar, M. & Pamik, M. Nuclear power utilization as a future alternative energy on icebreakers. *Nucl. Eng. Technol.* **55**, 580–586 (2023).

12. McEachern, R. J. & Taylor, P. A review of the oxidation of uranium dioxide at temperatures below 400 °C. *J. Nucl. Mater.* **254**, 87–121 (1998).

13. Wilson, P. *The Nuclear Fuel Cycle: From Ore to Wastes* (Oxford University Press, 1996).

14. Devanathan, R. et al. Modeling and simulation of nuclear fuel materials. *Energy Environ. Sci.* **3**, 1406–1426 (2010).

15. Suzuki, M., Sugiyama, T. & Fuketa, T. Thermal stress analysis of high-burnup LWR fuel pellet pulse-irradiated in reactivity-initiated accident conditions. *J. Nucl. Sci. Technol.* **45**, 1155–1164 (2008).

16. Sonoda, T. et al. Transmission electron microscopy observation on irradiation-induced microstructural evolution in high burn-up UO₂ disk fuel. *Nucl. Instrum. Methods Phys. Res. Sect. B: Beam Interact. Mater. At.* **191**, 622–628 (2002).

17. Spino, J., Vennix, K. & Coquerelle, M. Detailed characterisation of the rim microstructure in PWR fuels in the burn-up range 40–67 GWd/tM. *J. Nucl. Mater.* **231**, 179–190 (1996).

18. Shoesmith, D. W. Fuel corrosion processes under waste disposal conditions. *J. Nucl. Mater.* **282**, 1–31 (2000).

19. Springell, R. et al. Water corrosion of spent nuclear fuel: radiolysis driven dissolution at the UO₂/water interface. *Faraday Discuss.* **180**, 301–311 (2015).
20. Vallejo, K. D. et al. Advances in actinide thin films: synthesis, properties, and future directions. *Rep. Prog. Phys.* **85**, 123101 (2022).

21. Springell, R. et al. A review of uranium-based thin films. *Adv. Phys.* **71**, 87–165 (2023).

22. Rousseau, G. et al. A detailed study of UO₂ to U₃O₈ oxidation phases and the associated rate-limiting steps. *J. Nucl. Mater.* **355**, 10–20 (2006).

23. Allen, G., Tempest, P. & Tyler, J. The formation of U₃O₈ on crystalline UO₂. *Philos. Mag. B* **54**, L67–L71 (1986).

24. Allen, G. C., Tempest, P. A. & Tyler, J. W. Oxidation of crystalline UO₂ studied using X-ray photoelectron spectroscopy and X-ray diffraction. *J. Chem. Soc., Faraday Trans. 1: Phys. Chem. Condens. Phases* **83**, 925–935 (1987).

25. Allen, G. & Holmes, N. A mechanism for the UO₂ to α-U₃O₈ phase transformation. *J. Nucl. Mater.* **223**, 231–237 (1995).

26. Geønvold, F. & Haraldsen, H. Oxidation of uranium dioxide (UO₂). *Nature* **162**, 69–70 (1948).

27. Leinders, G. et al. Low-temperature oxidation of fine UO₂ powders: a process of nanosized domain development. *Inorg. Chem.* **55**, 3915–3927 (2016).

28. Idriss, H. Surface reactions of uranium oxide powder, thin films and single crystals. *Surf. Sci. Rep.* **65**, 67–109 (2010).

29. Quémard, L. et al. On the origin of the sigmoid shape in the UO₂ oxidation weight gain curves. *J. Eur. Ceram. Soc.* **29**, 2791–2798 (2009).

30. Strehle, M. M. et al. Characterization of single crystal uranium-oxide thin films grown via reactive-gas magnetron sputtering on yttria-stabilized zirconia and sapphire. *Thin Solid Films* **520**, 5616–5626 (2012).

31. Rennie, S. et al. The role of crystal orientation in the dissolution of UO₂ thin films. *Corros. Sci.* **145**, 162–169 (2018).

32. Barrett, S., Jacobson, A., Tofield, B. & Fender, B. The preparation and structure of barium uranium oxide BAUO₃⁺. *Acta Crystallogr. Sect. B: Struct. Crystallogr. Cryst. Chem.* **38**, 2775–2781 (1982).

33. Swanson, H. E. *Standard X-ray Diffraction Powder Patterns* Vol. 25 (US Department of Commerce, National Bureau of Standards, 1955).

34. Desgranges, L., Baldinozzi, G., Rousseau, G., Niepce, J.-C. & Calvarin, G. Neutron diffraction study of the in situ oxidation of UO₂. *Inorg. Chem.* **48**, 7585–7592 (2009).

35. Elorrieta, J., Bonales, L., Rodríguez-Villagra, N., Baonza, V. & Cobos, J. A detailed raman and X-ray study of UO₂ₓ⁻ oxides and related structure transitions. *Phys. Chem. Chem. Phys.* **18**, 28209–28216 (2016).

36. Smith, D., Scheetz, B., Anderson, C. A. & Smith, K. Phase relations in the uranium-oxygen-water system and its significance on the stability of nuclear waste forms. *Uranium* **1**, 79–111 (1982).

37. Leinders, G. et al. Assessment of the U₃O₇ crystal structure by x-ray and electron diffraction. *Inorg. Chem.* **55**, 9923–9936 (2016).

38. Leinders, G. et al. Charge localization and magnetic correlations in the refined structure of U₃O₇. *Inorg. Chem.* **60**, 10550–10564 (2021).

39. Aronson, S., Roof Jr, R. & Belle, J. Kinetic study of the oxidation of uranium dioxide. *J. Chem. Phys.* **27**, 137–144 (1957).

40. Hoekstra, H., Santoro, A. & Siegel, S. The low temperature oxidation of UO₂ and U₄O₉. *J. Inorg. Nucl. Chem.* **18**, 166–178 (1961).

41. Desgranges, L., Baldinozzi, G., Simeone, D. & Fischer, H. Refinement of the α-U₄O₉ crystalline structure: new insight into the U₄O₉ → U₃O₈ transformation. *Inorg. Chem.* **50**, 6146–6151 (2011).

42. Walker, D. The oxidation of uranium dioxides. *J. Appl. Chem.* **15**, 128–135 (1965).

43. Hwang, C., Geiss, R. H. & Howard, J. K. Imaging of the grain-to-grain epitaxy in NiFe/FeMn thin-film couples. *J. Appl. Phys.* **64**, 6115–6117 (1988).

44. Bhat, A. P. & Ramadurai, R. et al. Grain to grain epitaxy-like nano structures of (Ba, Ca)(ZrTi)O₃/CoFe₂O₄ for magneto-electric based devices. *ACS Appl. Nano Mater.* **3**, 11098–11106 (2020).

45. Estandía, S., Dix, N., Chisholm, M. F., Fina, I. & Sánchez, F. Domain-matching epitaxy of ferroelectric Hf₀.₅Zr₀.₅O₂(111) on La₂/₃Sr₁/₃MnO₃(001). *Cryst. Growth Des.* **20**, 3801–3806 (2020).

46. Campbell, T. K., Gilbert, E. R., White, G. D., Piepel, G. F. & Wrona, B. J. Oxidation behavior of nonirradiated UO₂. *Nucl. Technol.* **85**, 160–171 (1989).

47. Taylor, P., Burgess, E. A. & Owen, D. G. An X-ray diffraction study of the formation of β-UO₂.₃₃ on UO₂ pellet surfaces in air at 229 to 275 °C. *J. Nucl. Mater.* **88**, 153–160 (1980).

48. Guenther, R. et al. Dry storage experience with spent fuel and high-level and transuranic wastes. *Trans. Am. Nucl. Soc.* **46** (1984).

49. Campbell, T. K., Gilbert, E. R., Thornhill, C. K. & Wrona, B. J. Oxidation behavior of spent UO₂ fuel. *Nucl. Technol.* **84**, 182–195 (1989).

50. Bao, Z. et al. Antiferromagnetism in UO₂ thin epitaxial films. *Phys. Rev. B* **88**, 134426 (2013).

51. Schneider, C. A., Rasband, W. S. & Eliceiri, K. W. NIH Image to ImageJ: 25 years of image analysis. *Nat. methods* **9**, 671–675 (2012).

## Acknowledgements

This research was supported by the Bristol Centre for Functional Nanomaterials, Centre for Doctoral Training, the University of Bristol Facility for Radioactive Materials Surfaces (FaRMS), funded by the UK Engineering and Physical Sciences Research Council [EP/V035495/1], and the TRANSCEND consortium on nuclear waste and decommissioning, also funded by the UK Engineering and Physical Sciences Research Council [EP/S01019X/1].

## Author contributions

J.W.: Conceptualisation, substrate preparation, sample fabrication, EBSD data collection and analysis, in situ XRD experiment and data analysis, SEM data collection and analysis, HT-ESEM data analysis, figures, original draft preparation. J.S.: Substrate preparation, sample preparation, EBSD data collection and analysis. R.P.: In-situ HT-ESEM data collection and analysis. J.L.: Sample preparation, XRD data collection. R.S.: Conceptualisation, data analysis, supervision, manuscript review and editing. All authors discussed and contributed to the writing of the paper.

## Competing interests

The authors declare no competing interests.

## Additional information

**Supplementary information** The online version contains supplementary material available at https://doi.org/10.1038/s41529-024-00479-1.

**Correspondence** and requests for materials should be addressed to Jacek Wasik or Ross Springell.

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