# Atomistically-informed cross-slip and slip-system transition modeling in uranium dioxide

Jules-Elémir Suchorski, Adrien Pivano, Laurent Dupuy, Jonathan Amodeo

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Jules-Elémir Suchorski, Adrien Pivano, Laurent Dupuy, Jonathan Amodeo. Atomistically-informed cross-slip and slip-system transition modeling in uranium dioxide. International Journal of Plasticity, 2026, 199, pp.104638. 10.1016/j.ijplas.2026.104638. hal-05525058

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# Atomistically-informed cross-slip and slip-system transition modeling in uranium dioxide

Jules-Elémir Suchorski<sup>a,b</sup>, Adrien Pivano<sup>a ID</sup>, Laurent Dupuy<sup>c</sup>, Jonathan Amodeo<sup>b ID</sup>,*

## ARTICLE INFO

### Keywords:

Cross-slip

Nuclear fuel

Plasticity

Slip systems

Dislocation dynamics

Simulation

## ABSTRACT

Understanding the mechanical behavior of uranium dioxide (UO₂) at high-temperature is essential for assessing the structural integrity of nuclear fuel under reactor operating conditions. The pronounced plastic anisotropy observed in UO₂ single crystals is believed to originate from the complex behavior of 1/2<110> screw dislocations and their ability to cross-slip between the various slip modes of the fluorite structure, as initially suggested by Sawbridge and Sykes more than fifty years ago. However, direct evidence supporting this hypothesis, either from atomic-scale simulations or transmission electron microscopy, has so far been lacking, leaving the microscopic origin of cross-slip and its possible contribution to the plastic anisotropy unresolved. In this work, we combine molecular dynamics, theoretical modeling and discrete dislocation dynamics simulations to investigate the mobility of the 1/2<110> screw dislocation in UO₂ as a function of temperature, stress and orientation. Atomistic simulations reveal that above ~1500 K, the 1/2<110> screw dislocation progressively transitions from planar glide in {001} or {110} slip systems to frequent cross-slip in 1/2<110>{111}, ultimately leading to a full slip-system transition near 1700 K. This behavior correlates with a temperature-induced reorganization of the dislocation core, from a configuration electrostatically-confined within {001} to a {111}-spreading state as the oxygen sublattice disorders. An atomistically-informed stochastic model of cross-slip and slip-system transition is then formulated and implemented in the discrete dislocation dynamics simulation, successfully reproducing atomistic simulations, experimental slip traces and the plastic anisotropy trends observed in UO₂.

## 1. Introduction

Understanding the elementary processes of plastic deformation in uranium dioxide (UO₂) is essential for modeling the mechanical behavior of nuclear fuel under both nominal and transient operating conditions of pressurized water nuclear reactors. UO₂, which adopts the fluorite structure consisting of a face-centered cubic (FCC) uranium sublattice interpenetrated by a simple cubic oxygen network, exhibits a remarkably complex plastic response, characterized by a strong plastic anisotropy and a pronounced temperature dependence at the single-crystal level. Experimental observations have revealed the activation of several slip systems beyond the brittle-to-ductile transition (T > 800 K) including primary 1/2<110>{001} and secondary 1/2<110>{110} and 1/2<110>{111} (sometimes referred as mode I, II and III, respectively), whose relative activity depends on temperature, stoichiometry, and crystallographic orientation (Byron, 1968; Nadeau, 1969; Yust and McHargue, 1969; Sawbridge and and, 1971; Keller et al.,

<footnote><sup>a</sup> CEA, DES, IRESNE, DEC, Saint-Paul-Lez-Durance, F-13108, France</footnote>

<footnote><sup>b</sup> Aix-Marseille Université, Université de Toulon, CNRS, IM2NP, Marseille, F-13013, France</footnote>

<footnote><sup>c</sup> Université Paris-Saclay, CEA, Service de Recherche en Matériaux et procédés Avancés, Gif sur Yvette, F-91191, France</footnote>

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: jonathan.amodeo@cnrs.fr (J. Amodeo).</footnote>
1988a,b). This hypothetic multi-slip behavior is commonly associated with the high lattice friction and thermally-activated glide of dislocations, as well as to the cross-slip process which is enhanced due the single Burgers vector family associated to multiple potential slip planes. In UO₂, a key feature of plasticity within the primary 1/2<110> {001} slip systems is the transition from a lattice-friction dominated regime where dislocation glide proceeds via kink-pair nucleation and migration, reminiscent of the behavior observed in body-centered cubic (BCC) metals, to a viscous glide regime as the temperature rises above approximately 1600–1700 K (Sawbridge and others, 1971; Lunev et al., 2017; Portelette et al., 2018; Soulié et al., 2018; Borde et al., 2023a). In this high-temperature range, dislocation glide becomes easier and is increasingly characterized by wavy or curved trajectories, no longer confined to a single crystallographic plane. Sawbridge and Sykes originally interpreted such behavior as a signature of extensive cross-slip and composite slip (combined activation of multiple slip modes with shared Burgers vectors), in which dislocations continuously transfer between {001}, {110} and {111} glide planes (Sawbridge and others, 1971). These mechanisms are believed to control the orientation-dependent flow stress, strain hardening, and viscoplastic anisotropy observed experimentally. However, their precise origin at the dislocation scale remains uncertain due to the limited knowledge of the dislocation core structure in UO₂ and the lack of direct evidence for cross-slip at the atomic scale.

Recent advances in both experimental characterization and multi-scale modeling have provided new insights into these long-standing questions (Portelette et al., 2018, 2020; Madec et al., 2023; Onofri et al., 2024). Discrete dislocation dynamics (DDD) and crystal plasticity finite element (CPFE) simulations have shown that the anisotropic mechanical response of UO₂ can only be reproduced when cross-slip between slip modes is allowed. Madec and collaborators proposed a composite slip model in which cross-slip is assumed to be a stress-driven and easily-activated mechanism, allowing screw dislocations to transfer spontaneously from a primary slip system to alternative systems sharing the same Burgers vector (Madec et al., 2023). This framework successfully reproduced the anisotropic flow stress measured by Sawbridge and Sykes, thereby supporting the idea that cross-slip plays a central role in UO₂ high-temperature plasticity. Borde and coworkers extended the multi-scale modeling of UO₂ by coordinating molecular dynamics (MD) and DDD simulations (Borde et al., 2023a,b). In particular, they studied the interactions between primary 1/2<110> {001} dislocations and radiation-induced {110} prismatic loops, showing that the 1/2<110> screw dislocations give rise to the most strongly-locked configurations under irradiation. Some of the configurations examined also suggested that the cross-slip of the 1/2<110> screw dislocation can be readily activated in highly-stressed region, influencing the radiation-induced hardening processes. More recently, Cai et al. also used DDD to investigate dislocation density evolution and plastic anisotropy in UO₂ with a cross-slip model adapted from the approach of Chaussidon, applied to {001} and {110} slip systems (Chaussidon et al., 2006; Cai et al., 2025). Despite their different levels of description, these approaches mostly rely on purely phenomenological and stress-controlled treatments of the cross-slip activation, without any quantitative calibration against atomistic or experimental data. In the following, we adopt an alternative framework and show that cross-slip in UO₂ is primarily governed by thermal activation rather than by stress alone. We further demonstrate that temperature plays a decisive role in enabling cross-slip-mediated transfer towards 1/2<110> {111} slip systems, which become increasingly active at high temperature and critically influence the macroscopic plastic response.

At the atomistic scale, Soulié and Lunev carried out the first detailed studies of the 1/2<110> screw dislocation in UO₂ using molecular simulations (Soulié et al., 2018; Lunev et al., 2018). Building on these pioneering works, we subsequently performed new investigations that revealed a pronounced effect of temperature on the screw dislocation core structure (Suchorski et al., 2025). At low temperature, the core primarily spreads in a {001} plane, while it progressively evolves towards a rather tridimensional configuration, including instant spreading in the {111} crystallographic planes above a critical temperature $T_c > 1600$ K. This transition was attributed to a disordering of the oxygen sublattice within the dislocation core upon increasing temperature, corresponding to an anticipated Bredig-like transition localized within the dislocation core. Moreover, a quantitative correlation was established between the evolution of the fraction of oxygen point defects within the core and the ability of the dislocation to switch its most favorable spreading plane during MD fluctuation simulations performed at constant temperature and zero applied stress. While conducted under zeros stress conditions, these simulations offered the very first natural explanation for both the activation of cross-slip and the single-to-multi slip transition observed near 1600 K in the experiments. All together, these studies have established that cross-slip is a key component of UO₂ plasticity, but a quantitative connection between the dislocation core evolution, the cross-slip probability, and the mesoscopic response under stress remains missing.

In this work, we propose a physics-based model for cross-slip and slip-system transition in UO₂ that directly bridges the atomistic and macroscopic scales. We first perform MD simulations of the 1/2<110> screw dislocations under applied shear stress and temperature to analyze dislocation trajectories for various geometries characterized by the maximum resolved shear stress plane (MRSSP). These simulations reveal a strong correlation between the occurrence of cross-slip and the temperature-dependent core spreading observed by Suchorski et al. (2025). Based on these results, we have developed a stochastic model in which the cross-slip probability scales with the local propensity of the screw dislocation core to spread over alternative glide planes increasing the temperature. This model is then implemented in a DDD framework, enabling us to study its effect on dislocation mobility, slip system activation, and overall mechanical response. Our results which are mostly temperature-driven are compared with both atomistic data and experimental flow stress measurements by Sawbridge and Sykes, as well as with the earlier stress-drivel models developed by Madec and Borde (Sawbridge and others, 1971; Madec et al., 2023; Borde et al., 2023b).

## 2. Methods and theory

### 2.1. Molecular simulations

Molecular simulations of the 1/2<110> screw dislocation in UO₂ are carried out using the LAMMPS code (Thompson et al., 2022), with atomic interactions described by the empirical potential developed by Cooper, Rushton and Grimes (CRG) (Cooper
et al., 2014). The CRG potential combines classical pair interactions, including Coulomb interactions, with the many-body embedded atom method (Daw and Baskes, 1984). It is known as one of the most reliable interatomic potentials for the computation of thermomechanical properties in UO₂, having been successfully applied to the calculation of lattice and elastic constants (Cooper et al., 2014; Balboa et al., 2017), stacking-fault energy (Skelton and Walker, 2017; Amodeo and Bourasseau, 2025), 1/2<110>~{001} edge dislocation core and mobility (Lunev et al., 2017), point defects (Cooper et al., 2014; Amodeo and Bourasseau, 2025) and grain boundaries (Bourasseau et al., 2022). In particular, the potential fitting database includes high-temperature properties such as elastic constants, thermal expansion and specific heat, making it transferable over a wide temperature range, as confirmed by its accurate predictions of the melting temperature ($T_m \sim 3050 \pm 50$ K) and the Bredig transition temperature ($T_B \sim 0.8T_m$), both in good agreements with experimental data (Manara et al., 2005; Ralph and Hyland, 1985; Hutchings, 1987; Hiernaut et al., 1993; Ronchi and Hyland, 1994; Pavlov et al., 2017). Recently, we employed the CRG potential to investigate the 1/2<110> screw dislocation core structure in UO₂ (Suchorski et al., 2025), yielding particularly consistent results when compared to the more sophisticated variable-charge SMTB-Q model (Sattonnay and Tétot, 2013; Mbongo et al., 2020).

Here, the behavior of the 1/2<110> screw dislocation under constant temperature and applied shear stress is studied using the periodic array of dislocations (PAD) setup (Osetsky and Bacon, 2003), as illustrated in Fig. 1a. After generating the initial simulation cell made of a UO₂ fluorite perfect crystal ($a_0=5.45$ Å) with ATOMSK (Hirel, 2015), a single 1/2<110> screw dislocation is introduced at its center using the isotropic elasticity displacement field. The dislocation line (l) and Burgers vector (b) are oriented along the y-axis, the z-axis is normal to the MRSSP and the x-axis aligns with the Peach-Koehler force ($F_{PK}^{app}$) resulting from the applied shear stress (when non-zero). After the dislocation insertion, the y-vector of the cell is tilted along the x-direction by $\delta x = b/2$ to accommodate the plastic shear induced by the dislocation and recover the periodicity along the x-direction. Indeed, periodic boundary conditions (PBCs) are used along the x and y directions during the simulation, whereas free-BCs are used along z. On each side of the z boundary, a layer of atoms of thickness 13 Å is frozen in the z-direction during the simulation to emulate an infinite elastic medium and reduce spurious surface effects. It is worth noticing that the PAD configuration is preferred here over the dipolar or quadrupolar arrangements (see e.g., Ref. Clouet et al. (2009), Clouet (2012)), in order to minimize spurious elastic interactions between dislocations. Such interactions, which can be amplified by cross-slip in an auto-catalytic manner, are particularly significant for screw dislocations at constant temperature. Nevertheless, the drawback of this setup is the presence of surface-induced image forces, which may still affect the dislocation behavior. This aspect will be further discussed in the next sections.

In this study, Coulomb interactions are computed using the Wolf summation method (Wolf et al., 1999), with a damping coefficient $0.3$ Å⁻¹. A uniform cutoff radius of 11 Å cutoff is applied to all the interaction terms of the CRG interatomic potential, in accordance with the original parameterization (Cooper et al., 2014). The CRG potential combined with the Wolf summation method has been validated via transferability tests across a range of bulk UO₂ properties showing good agreement with Ewald-based summation methods and referenced data (see Supplementary Information). Finally, the Wolf method is particularly appropriate here as surfaces relying on free-BCs are bounded by a frozen atom layer over a thickness chosen larger than the interatomic potential cutoff distance.

Three simulation cell orientations are designed to investigate the influence of stress on the dislocation behavior with the (001), (110) or (111) plane as MRSSP (see Fig. 1c-e). For all configurations, the simulation box length along the y direction $L_y$ ($l \parallel b$ direction) is equal to 90 Å (~24 b), while other dimensions being approximately $L_x = 230$ Å and $L_z = 260$ Å for both (001) and (110) MRSSP cases, or $L_x = 240$ Å and $L_z = 270$ Å for the (111) MRSSP case. The cell dimensions are chosen in accordance with convergence tests performed in the studies of Borde et al. (2023a) and Lunev et al. (2018). Additional simulation cell orientation details are provided in Fig. 1.

After the sample design, the potential energy of the system is first minimized using the conjugate gradient method through an iterative process alternating energy minimization and pressure relaxation. During this process, the virial stress tensor is computed excluding atoms in the boundary layers and is adjusted via cell vector length and angle adjustments, with atomic position rescaled accordingly. A final energy minimization run is performed down to a force norm of $10^{-4}$ eV/Å using the FIRE algorithm (Bitzek et al., 2006).

Then, MD simulations at constant temperature are carried out within the NVT ensemble using the Nosé-Hoover thermostat with an integration timestep of 1 fs and a damping parameter of 100 fs. To this aim, the sample is first heated up to a target temperature at a maximum heating rate of 50 K.ps⁻¹. The dimensions of the sample are then rescaled using a second order polynomial adjustment of the lattice parameter $a_0$ evolution with temperature ($\Delta a_0/a_0=1.002 + 4.975 \times 10^{-6} \times T + 3.205 \times 10^{-9} \times T^2$) (Cooper et al., 2014), and a final set of short MD runs, interspersed with cell dimensions adjustments, are performed iteratively to relax residual stresses induced by the lattice expansion. Here, a convergence criterion requiring all components of the stress tensor to be below 20 MPa on average over 5 ps is used. Subsequently, a per-atom force $f_y^{app} = \pm \frac{L_x L_y}{N_{surf}} \tau_{yz}$ is applied to the $N_{surf}$ atoms of the boundary layers. The force is ramped up progressively over 100 ps, after which a constant shear stress and temperature MD simulation is performed for an additional 4 ns.

Dislocations are characterized using OVITO (Stukowski, 2009). The dislocation extraction algorithm (DXA) (Stukowski et al., 2012) is applied to the FCC uranium sublattice in order to compute the dislocation position, averaging results along the dislocation line every 10 ps. Additional python homemade routines are used to compute dislocation trajectories and slip traces. For the cross-slip model, dislocation core spread identification is performed computing disregistry profiles as a function of temperature as detailed in Suchorski et al. (2025).
<figure><img src="image_3.png" /><figcaption>**Fig. 1.** Dislocation simulation setup and cell orientation. (a) Molecular dynamics (MD) setup. For clarity, only atoms near the dislocation core are shown. Two boundary layers are defined near the z-surfaces to mitigate surface effects and apply the shear stress $\tau_{yz}$. (b) Discrete dislocation dynamics (DDD) setup. For both MD and DDD simulation setups, the z-axis is normal to the maximum resolved shear stress plane (MRSSP) while the dislocation line and Burgers vector directions (l and b, respectively) are parallel to the y-axis. Both setups employ semi-periodic boundary conditions with periodic boundary conditions (PBC) along the x and y directions while free boundary conditions (Free-BC) are used along z. (c-e) Simulation cell orientation with MRSSP as (c) (001), (d) (110) and (e) (111).</figcaption></figure>

## 2.2. Discrete dislocation dynamics simulation

DDD simulations are carried out using the *Numodis* code (Gonzalez Joa et al., 2023), which has already been employed to investigate the influence of radiation-induced defects in metals (see e.g., Refs. Shi et al. (2015), Li et al. (2019), Dupuy et al. (2024)) and, more recently, in UO₂ (Borde et al., 2023b), in direct comparison to MD simulations. *Numodis* is a nodal DDD code i.e., dislocations are represented by nodes connected through segments, each associated to a glide plane and a Burgers vector. In the present work, nodal forces are computed using the non-singular theory of dislocation (core width of 1b) (Cai et al., 2006), via the analytical expressions obtained by Arsenlis et al. (2007).

The DDD modeling framework for UO₂ relies on the lattice parameter $a_0 = 5.45 Å$ assuming isotropic elasticity, with temperature-dependent shear modulus (G) and Poisson's ratio (ν) derived from the work of Porto et al. (2024). This choice ensures consistency with the atomistic simulations performed using the CRG interatomic potential (see Supplementary Information for more details). Also, node velocities in the DDD are determined using a variational approach (Weygand et al., 2002; Bulatov and Cai, 2006) and a viscous law is used for the dislocation mobility $v(\tau_i) = \tau_i \cdot b/B_i$, where $B_i$ is the drag coefficient and $\tau_i$ the resolved shear stress in slip plane i. Drag coefficients $B_{(001)} = 4.5 \times 10^{-3}$ Pa.s, $B_{(110)} = 4.0 \times 10^{-1}$ Pa.s and $B_{(111)} = -5.0 \times 10^{-5} \times T + 0.115$ Pa.s were adjusted on MD simulations for temperatures ranging from 1400 to 2200 K. A simulation timestep $\delta t=1$ ps is used in all DDD simulations with minimum and maximum segment lengths $l_{min} = 5$ Å and $l_{max} = 20$ Å, respectively. Dislocation reactions, such as junction formation or annihilation, are handled using an elasticity-based criterion and a collision detection algorithm is implemented to anticipate dislocation interactions following the approach proposed by Bulatov and co-workers (Arsenlis et al., 2007; Bulatov and Cai, 2006).

In the following, we propose an atomistically-informed Arrhenius-based model for cross-slip, adapted to the multiple slip mode conditions of the fluorite crystal structure. All relevant cross-slip systems i.e., {001}, {110} and {111} (all with $b=1/2<110>$) are provided as inputs to the simulation. As we will show, the model explicitly captures the temperature-dependent evolution of the oxygen sublattice within the screw dislocation core depicted in our previous work (Suchorski et al., 2025), making cross-slip primarily governed by thermal activation. In *Numodis*, cross-slip probabilities are evaluated if the stress criterion $\tau_{css} > \tau_s / \text{tol}$ is
satisfied, where $\tau_s$ and $\tau_{css}$ are the shear stresses in the primary and cross-slip slip systems, respectively, and $\text{tol}$ is a user-defined tolerance. In our thermally-driven cross-slip approach, a large tolerance ($\text{tol}>10$) is adopted to ensure that no cross-slip events are missed (see parametric study presented in the Supplementary Information). The cumulative length of the screw segment $l_s$ meeting this condition is then computed, and the dislocation segment is moved accordingly to a Monte-Carlo (MC) scheme. The details of the cross-slip model and its parameterization are provided in the results section, and the full implementation of the DDD algorithm is documented in the Supplementary Information.

To allow the one-to-one comparison with MD simulations, the DDD simulation setup is designed particularly close to the MD PAD one, with same geometry and dimensions, as shown in Fig. 1b. Image forces induced by free surfaces are taken into account using the Hirth and Lothe approach for screw dislocation (Hirth and Lothe, 1982). A preliminary study focusing on the impact of the number of image dislocations and periodic replicas is detailed in the next section.

## 2.3. Image forces

In the following, the influence of free-BCs along $z$ and PBCs along $x$ is investigated to discuss the impact of image forces on the screw dislocation mobility in $UO_2$. Note that the presence of a uniform applied shear stress does not impact the following discussion.

In the MD simulations, atoms in the boundary layer near the $z$ surfaces are fixed perpendicular to the surface thereby enforcing zero normal stress condition ($\sigma_{zz}=0$). However, since they are free to move along $x$ and $y$, the shear components at the $z$ surface, $\tau_{xz}$ and $\tau_{yz}$, are allowed to vary during the whole simulation process. Since the screw dislocation stress field is only characterized by $\tau_{xy}$ and $\tau_{yz}$, the $z$-boundary can be considered as a free surface for the purpose of image stress calculation (Hirth and Lothe, 1982). Note that in the simulation, the strain incompatibility between the screw dislocation core and the boundary layer prevents the dislocation from escaping the simulation cell. This part of the simulation is not accounted in the following analysis and will not be further discussed.

The mirrored stress field of a screw dislocation in the presence of free surfaces along the $z$-direction and assuming an infinite medium along the glide direction $x$ can be computed using the image dislocation approach (Hirth and Lothe, 1982; Cai et al., 2003). This approach involves two series of image dislocations (see, e.g., Hirth and Lothe (1982), Liu et al. (2017)), whose coordinates $z_{nz}^+$ and $z_{nz}^-$ are given by,

$$z_{nz}^{+} = (-1)^{n_z+1}(n_z L_z - z_0) \quad (1)$$

$$z_{nz}^{-} = (-1)^{n_z}(n_z L_z + z_0) \quad (2)$$

where $z_0$ is the $z$-coordinate of the original screw dislocation and $n_z$ is the mirror dislocation index. The Burgers vector of the $n_z^{th}$ image dislocation in each series is given by $b_{nz} = (-1)^{n_z}b$. Due to symmetry, the $yz$ component of the mirror stress, $\tau_{yz}^{mirror}$, vanishes at the position of the original dislocation ($x_0, z_0$). Therefore, only the $xy$ component, $\tau_{xy}^{mirror}$, needs to be considered. It is worth mentioning that, in our case, the $xy$ component of the mirror stress tensor influences the out-of-MRSSP component of the Peach-Koehler force, possibly impacting the cross-slip of the screw dislocation. $\tau_{xy}^{mirror}$ is obtained by summing the contributions from all mirror dislocations assuming an infinite continuum along $x$,

$$\tau_{xy}^{mirror}(x, z) = \frac{Gb}{2\pi} \sum_{n_z \ge 1} (-1)^{n_z} \left( \frac{z - z_{nz}^{+}}{(x - x_0)^2 + (z - z_{nz}^{+})^2} + \frac{z - z_{nz}^{-}}{(x - x_0)^2 + (z - z_{nz}^{-})^2} \right) \quad (3)$$

Due to cancellations between the $z_{nz}^+$ and $z_{nz}^-$ terms, the $n$th term is of order $\frac{1}{n_z^2}$ and the sum is absolutely convergent.

Thus, the resulting image stress $\tau_{xy}^{im}$ that accounts for the use of free-BCs along the $z$-direction and PBCs along $x$ is obtained summing the $\tau_{xy}^{mirror}$ contributions over all $x$-periodic replicas of the mirror dislocations,

$$\tau_{xy}^{im}(x, z) = \sum_{n_x = -\infty}^{\infty} \tau_{xy}^{mirror}(x + n_x L_x, z) \quad (4)$$

Note that the $\tau_{xy}^{mirror}$ contributions of the original dislocation replicas located at $(x_0 + n_x L_x, z_0)$ are omitted from the sum, since it vanishes at the position of the original dislocation. In general, such a resulting series may not be convergent as individual terms may decay slowly. To solve this issue, a linear correction term was proposed by Cai et al. (2003) to restore the periodicity when necessary. However, the $n_x^{th}$ term in our case decays as $\frac{1}{n_x^3}$ which ensures absolute convergence, and such a correction is not required.

Due to the translational invariance along the $x$-direction, the image stress at the dislocation position, $\tau_{xy}^{im}(x_0, z_0)$ does not depend on the $x_0$-coordinate of the original dislocation, but only on $z_0$. The variation of $\tau_{xy}^{im}$ at the dislocation location with respect to $z_0$ is shown in Fig. 2 for different total numbers of periodic replicas $N_x$ and mirror dislocations $N_z$, as compared to the case without PBCs along $x$. The calculation is performed for a cell with dimensions $L_x = 240$ Å and $L_z = 270$ Å, which correspond to the same dimensions as for the MRSSP(=111) case, in MD and DDD simulations. In a nutshell, results show that the dislocation is subjected to image forces only when leaving the MRSSP with $\tau_{xy}^{im}$ reaching about one hundred MPa (the minimum applied stress used later in MD and DDD simulations) for a deviation along $z$ of about 3 to 4 nm from the simulation cell center (depending on $N_x$ and $N_z$). Also, the image stress calculation accuracy significantly depends on the number of mirrored and replica dislocations, respectively $N_z$ and $N_x$, accounted for. In Fig. 2a, a large value of $N_z$ is used to assess convergence with respect to $N_x$ which is achieved rapidly i.e., $N_x = 2$ is sufficient to obtain image stress values accurate to within 1%. Convergence with respect to the number of mirror dislocations $N_z$ is examined further in Fig. 2b. With $N_z = 100$, image stress values are accurate within 2%. Hence considering two periodic replicas
<figure><img src="image_4.png" /><figcaption>**Fig. 2.** Analytical modeling of the image stress $\tau_{xy}^{im}$ induced by a screw dislocation in a simulation cell with free surfaces along $z$ and PBCs along $x$, (a) scheme showing the original screw dislocation located at $z_0$. Mirror dislocations with opposite Burgers vectors are located at $z_n^-$ and $z_n^+$ coordinates (blue and red), on each side of the free surfaces, (b) effect of the number of replica $N_x$ (constant $N_z=10000$), (c) Effect of the number of mirror dislocations $N_z$ (constant $N_x=2$). $\tau_{xy}^{im}$ is evaluated at the original dislocation location ($x_0, z_0$) using Eq. (4) for various $z_0$ values. The origin of the $z$-axis is set at the center of a 2D simulation cell, with dimensions $L_x = 240$ Å and $L_z = 270$ Å.</figcaption></figure>

$(N_x=2)$ and one hundred image dislocations ($N_z=100$) on each sides ensure accurate image forces calculation. These preliminary results confirm that image forces might have a significant impact on the screw dislocation behavior in situations where glide occurs out of the MRSSP. This aspect, which is considered in the present study, will be discussed further in the following.

## 3. Results

### 3.1. Preliminary atomic-scale evidences of cross-slip and slip-system transition in UO$_2$

MD simulations under constant shear stress ($\tau_{yz}$ ranging from 100 to 900 MPa) and temperature ($T=1400-2200$ K) are performed to investigate the mobility of the $1/2<110>$ screw dislocation, for the three cell orientations described above. Analysis of dislocation trajectories are summarized in Fig. 3 and reveal a complex evolution that strongly depends on the temperature and applied shear stress.

At a representative low temperature of $T=1400$ K, the screw dislocation glides in the (001) plane for MRSSP=(001) and (111) configurations (Fig. 3a,c). In contrast, for MRSSP=(110), intermittent glide in the (110) plane is observed, alternating with fluctuations in (001) (Fig. 3e). Note that in this latter case, Schmid factors in $1/2<110>\{001\}$ slip systems are null. Overall, these configurations reflect the low Peierls barrier associated to $1/2<110>\{001\}$ slip systems, that are the most favorable at low temperature in UO$_2$.

Starting from $T=1500$ K, cross-slip in $1/2<110>\{111\}$ slip systems becomes increasingly frequent with increasing temperature. Above a transition temperature $T_c \sim 1700 \pm 100$ K, the screw dislocation predominantly glides in the {111} planes, regardless of the cell orientation, as illustrated at $T=2000$ K in Fig. 3b,d,f. This confirms the progressive transition of the $1/2<110>$ screw dislocation behavior, from {001}-dominated slip activity for $T < 1500$ K, to a composite slip regime, and finally a complete slip-system transition to $1/2<110>\{111\}$ for $T > T_c$. Overall, these results demonstrate that temperature, rather than stress, emerges as the dominant parameter controlling cross-slip and slip-system transitions. Some residual short slip events are still observed in (001) for MRSSP=(001) up to $T=1800$ K, but vanish at higher temperatures. For $T > T_c$, short glide events are recorded successively on the two {111} planes, each extending over a few atomic distances before the dislocation cross-slips to the other plane. Owing to the symmetric orientation of both {111} planes relative to MRSSP=(001) and MRSSP=(110), the dislocation motion averages out along the MRSSP for both configurations. By contrast, for MRSSP=(111), the combination of glide events in the two {111} planes leads to a wavy slip trace, whose orientation depends on the shear stress: being nearly-aligned with (110) a low stress and gradually rotating towards the MRSSP (111) as the stress increases.

One can notice that a hybrid composite slip regime is observed in a narrow temperature range around $T = T_c$, where the glide behavior of the $1/2<110>$ screw dislocation depends on the cell orientation. In this temperature regime, the screw dislocation glides exclusively in (001) for MRSSP=(001), but additional cross-slip events into {111} are noticed for MRSSP=(111).
<figure><img src="image_5.png" /><figcaption>Fig. 3. Slip traces and DXA trajectories of the 1/2<110> screw dislocation under constant stress and temperature computed using MD and the CRG interatomic potential, (a,b) MRSSP=(001), $\tau_{yz}=300$ MPa, respectively at $T=1400$ and 2000 K, (c,d) MRSSP=(111), $\tau_{yz}=300$ MPa, respectively at $T=1400$ and 2000 K, (e,f) MRSSP=(110), $\tau_{yz}=500$ MPa, respectively at $T=1400$ and 2000 K. Atoms are colored according to their displacement field $u_j$ (normal to the plane of the figure) computed with respect to the simulation initial conditions.</figcaption></figure>

## 3.2. Cross-slip and slip-system transition modeling

Following approaches commonly used for metals (see e.g., Refs. Zbib and Rubia (2002), Malka-Markovitz et al. (2021)), we employ an Arrhenius-type formulation to describe the probabilistic nature of cross-slip and slip-system transition of the 1/2<110> screw dislocation in UO$_2$,

$$P_i(T, \bar{\bar{\sigma}}) = \frac{l}{l_0} \frac{\delta t}{\delta t_0} \exp \left( - \frac{\Delta G_i(\tau_i, T)}{k_B T} \right), \quad (5)$$

where $\Delta G_i(\tau_i, T)$ is the Gibbs free-energy activation barrier to cross-slip that depends on the temperature and shear stress in slip plane $i$, $\bar{\bar{\sigma}}$ is the stress tensor, $l$ is the screw dislocation segment length under consideration, $\delta t$ is the simulation timestep, $l_0$ and $\delta t_0$ are reference counterparts and $k_B$ is the Boltzmann constant. As we believe that cross-slip in UO$_2$ is primarily a temperature-driven process, we decompose $\Delta G_i(\tau_i, T)$ into a dominant thermal contribution and a mechanical work term,

$$\Delta G_i(\tau_i, T) = \Delta G_{0,i}(T) - \tau_i V, \quad (6)$$

With $\Delta G_{0,i}(T)$ the stress-free contribution to the Gibbs free energy that contains most of the physics (internal energy variations, entropic effects) associated with the rearrangement of the oxygen sublattice within the dislocation core under constant temperature, while the mechanical work contribution $-\tau_i V$ (with $V$ the cross-slip activation volume) only enhances the probability of a mechanism that is already thermally accessible.

Thus, the probability $P_i(T, \bar{\bar{\sigma}})$ of cross-slip towards plane $i$ from any other plane can be expressed as,

$$P_i(T, \bar{\bar{\sigma}}) = \frac{l}{l_0} \frac{\delta t}{\delta t_0} P_i(T) \exp \left( \frac{\tau_i V}{k_B T} \right), \quad (7)$$

With $P_i(T) = P_0 \exp[-\Delta G_{0,i}(T)/k_B T]$ the normalized probability of cross-slip under zero applied stress, reflecting the sole effect of temperature on cross-slip.

As shown in our previous work (Suchorski et al., 2025), the temperature-driven evolution of the 1/2<110> screw dislocation core in UO$_2$ has a direct impact on the preferred slip plane in which the dislocation will fluctuate, in constant temperature and zero applied stress MD simulations. Precisely, we have demonstrated a quantitative correlation between the disordering of the oxygen sublattice within the core increasing the temperature and the dislocation most favorable fluctuation plane i.e., the number of oxygen point defects has shown to increase sharply near $T \sim 1600$ K, which coincides precisely with the temperature at which (i) the dislocation core change its favorable core spreading plane from exclusive {001} to more hectic configurations including core extension within {111} and (ii) the dislocation originally fluctuating in {001} transits in {111} slip planes, driven solely by thermal activation (zero applied stress simulations). To better understand the interplay between dislocation core disordering, core spreading and favorable slip plane, core configurations were sampled during MD fluctuation simulations, and the most favorable core spreading plane was characterized on-the-fly using a filtering criterion based on the core spreading disregistry parameters $\zeta_i$ among the various crystallographic plane $i$ (see technical details in Ref. Suchorski et al. (2025)). As this temperature-dependent behavior is further corroborated by MD simulations under constant stress applied stress presented in Fig. 3, we propose a cross-slip model whose zero-stress component $P_i(T)$ is directly informed by core-spreading statistics obtained from MD fluctuation simulations computed
<figure><img src="image_6.png" /><figcaption>Fig. 4. Core spreading and cross-slip probabilities of the 1/2<110> screw dislocation computed under zero stress, (a) dislocation core spread probability $P_i^{core}$ as a function of temperature. Symbols correspond to MD-derived raw data, while the solid lines correspond to fits calculated using Eq. (8), (b) temperature-dependent cross-slip and slip-system transition probability $P_i(T)$, calculated using Eqs. (9) and (10). Note that for both $P_i^{core}(T)$ and $P_i(T)$, probabilities in {111} planes have to be counted twice due to symmetry.</figcaption></figure>

<caption>Table 1</caption>

Fitting parameters for the core spreading probability $P_i^{core}(T)$ (Eq. (8)). Note that probabilities in {111} planes have to be counted twice due to symmetry.

<table><thead><tr><th></th><th colspan="3">Slip plane <i>i</i></th></tr><tr><th></th><th>{001}</th><th>{110}</th><th>{111}</th></tr></thead><tbody><tr><th><i>P</i><sub><i>i</i></sub><sup>core</sup>(0)</th><td>~1</td><td>~0</td><td>~0</td></tr><tr><th><i>P</i><sub><i>i</i></sub><sup>core</sup>(∞)</th><td>0.18</td><td>0.16</td><td>0.33</td></tr><tr><th><i>T</i><sub><i>c</i></sub><sup><i>i</i></sup></th><td>1732</td><td>1717</td><td>1736</td></tr><tr><th><i>T</i><sub><i>w</i></sub><sup><i>i</i></sup></th><td>57</td><td>58</td><td>54</td></tr></tbody></table>

without applied stress. To this end, the temperature-dependent core spreading probability, $P_i^{core}(T)$, is defined as the ratio between the number of configurations in which the dislocation core extends in plane $i$ and the total amount of sampled configurations, obtained from zero-stress MD simulations. This statistical analysis is performed at each temperature in the 1400–2200 K temperature range, with a 200 K step. The evolution of $P_i^{core}$ as a function of temperature is shown in Fig. 4a where data are interpolated using a logistic sigmoid function (Eq. (8)) that allows to describe the progressive and continuous transition between the two distinct regimes of the screw dislocation core i.e., from planar core spreading in {001} to more tridimensional configurations, as the temperature increases,

$$P_i^{core}(T) = P_i^{core}(0) + \frac{P_i^{core}(\infty) - P_i^{core}(0)}{1 + e^{-(T-T_c^i)/T_w^i}}, \quad (8)$$

Bounding probabilities $P_i^{core}(0)$ and $P_i^{core}(\infty)$ as well as adjusted transition temperature $T_c^i$ and sharpness $T_w^i$ are provided in Table 1. In detail, the value $P_{\{001\}}^{core}(0) \sim 1$, combined with near-zero contributions in {110} and {111} at the lower bound, indicates that the screw dislocation core primarily extends within {001}. This confinement, in turn, constrains the glide of the screw dislocation predominantly in this slip system at low temperature. On the other hand, the decrease of $P_{\{001\}}^{core}$ accompanied by the raise of $P_{\{110\}}^{core}$ and $P_{\{111\}}^{core}$ at the higher bound, reflects the progressive 3D spreading of the dislocation core as the temperature increases (see Ref. Suchorski et al. (2025) for more details). As expected, fitting values of $T_c^i=1726.5 \pm 9.5K$ are very close for all slip systems and fairly align with the dislocation core transition temperature c.a., in the 1600–1800 K range, estimated in our previous work (Suchorski et al., 2025). The fitted width parameters $T_w^i$, on the order of 55 K, confirm a transition spread over a few hundreds of Kelvin.

Although a cross-slip hypothesis based on the screw core spread probability $P_i^{core}$ may hold below $T_c$, the results presented in Fig. 4a become inconsistent with the slip activity observed in MD dislocation simulations at higher temperature, both with and without an applied stress. Above $T_c$, MD simulations show that (i) the screw core intermittently extends onto all available slip planes due thermal fluctuations (Suchorski et al., 2025) whereas (ii) the dislocation glides predominantly in the 1/2<110>{111} slip systems (see e.g., Fig. 3) under stress. This confirms that, for $T > T_c$, there is no longer straightforward correlation between the screw dislocation multi-planar core spreading and the active slip systems.

Importantly, this correlation breakdown is not unexpected, nor does it indicate a limitation of the approach, but rather follows naturally from the same crystallographic arguments that govern the core-spreading transition itself (Suchorski et al., 2025). Specifically, the disordering of the oxygen sublattice occurring in the dislocation core for $T > T_c$ resembles a local phase transformation (anticipated-Bredig transition) (Dworkin and Bredig, 1968; Fossati et al., 2024; Suchorski et al., 2025). During this process, the local crystallographic environment evolves from the fluorite structure, characterized by core spreading into the softer {001} slip plane and strong electrostatic friction in the {111} and {110} planes, towards a more “metallic-like” configuration,
<caption>**Table 2**</caption>

Parameters of the cross-slip probability $P_i(T, \bar{\sigma})$. Note that probabilities in {111} planes have to be counted twice due to symmetry.

<table><thead><tr><th>l<sub>0</sub> [Å]</th><th>δt<sub>0</sub> [ps]</th><th>T<sub>c</sub> [K]</th><th>T<sub>w</sub> [K]</th><th>V [Å<sup>3</sup>]</th></tr></thead><tbody><tr><td>100</td><td>50</td><td>1700</td><td>50</td><td>50</td></tr></tbody></table>

consisting of an FCC uranium sublattice, with dense and energetically more favorable {111} slip planes, embedded in a sea of deorganized oxygen atoms. Thus, in the same way that the fluorite crystal structure controls the spreading of the screw dislocation core for $T < T_c$, the emergence of a local FCC-like arrangement governs the dislocation behavior for $T > T_c$, justifying the preferential activation of 1/2<110>{111} slip systems while suppressing the activity of 1/2<110>{001} and 1/2<110>{110}. As a consequence, for $T > T_c$, $P_i^{core}$ is no longer an appropriate descriptor of cross-slip. In this high temperature regime, cross-slip is more appropriately described by a model that primarily accounts for the {111} slip planes ($P_i^{fcc}$), in close analogy with the behavior observed in FCC metals. Thus, the probability simplifies to $P_i(T > T_c) = P_i^{fcc}$, with values of 0.5 on each of the two {111} planes, while vanishing for the {001} and {110} planes. In the following, the two regimes ($T < T_c$ and $T > T_c$) are branched using a mixing law,

$$P_i(T) = (1 - f(T))P_i^{core}(T) + f(T)P_i^{fcc} \quad (9)$$

with $f(T)$ a logistic sigmoid function with similar features than the one used in Eq. (8),

$$f(T) = \frac{1}{1 + e^{-\frac{T - T_c}{T_w}}} \quad (10)$$

$f(T)$ varies from 0 ($T \ll T_c$) to 1 ($T \gg T_c$) with a transition at $f(T_c)=0.5$ taking place over a width $T_w$ similar to the one used in Eq. (8). Here we use the same values of $T_c=1700$ K and $T_w=50$ K for the three slip systems, in accordance with the transition described above. The resulting probability of cross-slip and slip-system transition under zero stress $P_i$ is shown as a function of temperature in Fig. 4b. It exhibits a gradual shift from {001} to {111} slip around the transition temperature $T_c=1700$ K, evolving from a composite slip regime with cross-slip among the {001}, {110} and {111} slip planes to a state dominated exclusively by glide in {111}, where the cross-slip is confined to the two {111} slip planes. Note that occasional cross-slip events within {110} are predicted within the 1600–1800 K temperature range, as indicated by slightly above-zero values of $P_{i(110)}(T)$. At higher temperatures ($T > 1800$ K), cross-slip is restricted to {111} planes, reflecting the dislocation behavior observed in MD simulations more accurately than predictions based solely on $P_i^{core}$.

Back to the main definition of the probability $P_i(T, \bar{\sigma})$ described in Eq. (7), the characteristic length and time, $l_0$ and $\delta t_0$, as well as the activation volume $V$ that corresponds to the effective volume of material involved during the cross-slip process, still have to be quantified. Note that, although $l_0$ and $\delta t_0$ have a smaller impact in our cross-slip model when compared to temperature and stress terms, they must still be set properly, as large variations may still affect the dislocation trajectory. Firstly, $l_0=100$ Å is imposed as providing a good approximate of the dislocation length measured in MD simulations. Additionally, a characteristic value of $\delta t_0=50$ ps is derived from the statistical analysis of cross-slip events observed in MD simulations performed under constant stress and temperature (see Supplementary Information for more details). The most critical parameter is the activation volume $V$, which enters through an exponential term, contributing to the stress dependence of $P_i(T, \bar{\sigma})$, thereby influencing the resulting dislocation trajectory. To first order, $V$ is expected to be relatively small (of the order of ~10-100 Å³), reflecting the predominantly thermally driven nature of the cross-slip process in UO₂, and hence its weak sensitivity to stress. To quantify $V$, a simplified MC scheme combining probabilistic cross-slip in the various slip planes and glide events is used to generate synthetic dislocation trajectories (MRSSP=(111) case, high temperature $T=2200$ K), for various guess-values of $V$. More details on this approach are provided in the next section. The resulting angle, $\theta = \arctan(z/x)$, which characterizes the dislocation trajectory, was then compared to that obtained from MD simulations in order to calibrate the value of $V$. It is worth noting that the MRSSP selected here promote wavy yet stress-asymmetric glide conditions, allowing to characterize the trajectory sensitivity to the activation volume $V$ since the effect of temperature on cross-slip is the same on both 1/2<110>{111} slip systems. This geometry is thus particularly suited to evaluate the sole impact of the shear stress on the cross-slip process. As depicted in Fig. 5, a value of $V=50$ Å³ satisfactorily captures the stress sensitivity of the dislocation trajectory, when averaged along 1000 simulations. Such a value of $V$ can rescale $P_i$ by a factor ranging from 1 to 5 for $\tau_i$ increasing from 0 to 1 GPa. Given that the stress range of interest is limited to a few tens to hundreds of MPa, whereas $P_i(T)$ spans several orders of magnitude, we confirm that the influence of stress on the cross-slip process of the 1/2<110> screw dislocation in UO₂ is comparatively minor relative to that of temperature. All cross-slip parameters are summarized in Table 2.

### 3.3. Preliminary tests: theory vs. simulations

To assess the cross-slip and slip-system transition model over relevant temperature and stress ranges, the total cross-slip rate $\Gamma_{cs}$, defined here as the total rate of successful events, is computed using different approaches. First, an analytical formulation of the total cross-slip rate $\Gamma_{cs}^{th}(T, \bar{\sigma})$ is proposed as derived from the probabilistic formulation of $P_i(T, \bar{\sigma})$. Assuming primary slip planes $i$, $\Gamma_{cs}^{th}(T, \bar{\sigma})$ can be derived summing the contributions of all cross-slip planes $j$, $\sum_{j \neq i} \frac{P_i(T, \bar{\sigma})}{\delta t}$ for each plane $i$, multiplied by the fraction
<figure><img src="image_7.png" /><figcaption>Fig. 5. Calibration of the activation volume V used in the cross-slip and slip-system transition model. The angle θ=arctan(z/x) characterizes the dislocation trajectory. It is calculated (i) averaging over a 1000 trajectories computed using a simplified Monte-Carlo (MC) approach (MRSSP=(111), T=2200 K) for different trial values of V (curves with filled circle symbols) and (ii) using a single dislocation trajectory inferred from MD simulations under constant stress and temperature (red cross).</figcaption></figure>

of elapsed time in plane i, $\frac{P_i(T, \bar{\sigma})}{\sum_k P_k(T, \bar{\sigma})}$ obtained by the detailed balance principle, where index k runs over all slip systems (primary and cross-slip). To account for the cross-slip tolerance criterion $\tau_{ess} > \frac{\tau_s}{tol}$, two additional weighting functions are introduced. On the one hand, $A_{ij}=1$ only if $tol \times \tau_{j \neq i} \ge \tau_i$, else $A_{ij}=0$. This weighting function is at the roots of the tolerance criterion that enables whatever slip plane to become a possible cross-slip plane. On the other hand, the criterion is also applied to plane i using the weight function $G_i$. Indeed, $G_i=1$ only if $tol \times \tau_i \ge \tau_{j \neq i}$ for at least one value of j, else $G_i=0$.

Finally, $\Gamma_{cs}^{th}(T, \bar{\sigma})$ is defined as,

$$ \Gamma_{cs}^{th}(T, \bar{\sigma}) = \frac{\sum_i G_i P_i(T, \bar{\sigma}) \sum_{j \neq i} A_{ij} P_j(T, \bar{\sigma})}{\delta t \sum_k P_k(T, \bar{\sigma})} \quad (11) $$

The analytical formulation of $\Gamma_{cs}$ proposed in Eq. (11) is confronted in the following to MD simulations as well as to a Markov-chain model resolved using MC simulation as follow. Assuming a dislocation originally lying in slip plane i, its trajectory under constant applied shear stress and temperature is computed using a MC scheme i.e., by evaluating the cross-slip probabilities $P_{j \neq i}(T, \bar{\sigma})$ at each discrete time step ($\delta t = 1$ ps). At each step, the dislocation position is thus updated either validating a cross-slip event into a new plane j, or remaining on the original plane i, depending on the computed probabilities. Finally, the displacement of the dislocation is integrated over time using the atomistically-informed mobility laws described in the methods section. Note that this latter approach was used to calibrate the cross-slip activation volume V in Section 3.2. Dislocation trajectories simulated using MC are analyzed using two independent counters of cross-slip events. First, the number cross-slip events $N_{cs}$ is counted during the simulation time $t_{MC}$ providing the total cross-slip rate, $\Gamma_{cs}^{MC1} = \frac{N_{cs}}{t_{MC}}$, at the highest time resolution of 1 ps. Second, a coarse-grained number of cross-slip events is derived from the variations of the slope of the simulated dislocation trajectory ($x(t), z(t)$). This last method is particularly relevant as it allows to count cross-slip events occurring during MD simulations under constant temperature and applied shear stress leading to $\Gamma_{cs}^{MD}$ (more details on the slope variation method are provided in Supplementary Information).

Fig. 6 compares the normalized total cross-slip rates $\Gamma_{cs}$ for each MRSSP configuration, obtained by solving Eq. (11) analytically ($\Gamma_{cs}^{th}$), by solving the Markov-chain problem using MC simulation ($\Gamma_{cs}^{MC1}$), by extracting rates from dislocation trajectories computed using MC ($\Gamma_{cs}^{MD}$) and from MD ($\Gamma_{cs}^{MC2}$) simulations. Cross-slip being temperature-driven, $\Gamma_{cs}$ naturally increases with temperature in the low-temperature regime. Near $T_c \sim 1700$ K, $\Gamma_{cs}^{th}$ obtained from Eq. (11) exhibits a pronounced peak for all MRSSP configurations. This peak corresponds to a regime in which the number of accessible cross-slip planes is maximum, as also shown in Fig. 4b. Beyond this temperature, the cross-slip rate progressively decreases as the number of available cross-slip planes is reduced ({{001}} and {{110}} becoming increasingly unfavorable), while the probability functions $P_i(T)$ approach saturation. This behavior is illustrated in Fig. 4b, where $dP_i(T)/dT \sim 0$ at high temperature. At this stage, the temperature dependence of the cross-slip probability $P_i(T, \bar{\sigma})$ (Eq. (7)) is slightly influenced by the activation volume term, leading to a smooth diminution of the total cross-slip rate with increasing temperature. This characteristic trend is also captured solving the Markov-chain problem with MC simulation ($\Gamma_{cs}^{MC1}$), which benefits
<figure><img src="image_8.png" /><figcaption>**Fig. 6.** Normalized total cross-slip rate $\Gamma_{cs}^{th}$ computed as function of temperature and applied shear stress for MRSSP as (a) (001), (b) (1\bar{1}0), and (c) (\bar{1}11). The first column shows the analytical resolution of $\Gamma_{cs}^{th}$ as inferred from Eq. (11) (tol=10), second and third columns emphasize the outcomes of the Monte-Carlo (MC) simulations $\Gamma_{cs}^{MC_1}$ and $\Gamma_{cs}^{MC_2}$ (highly-resolved and coarse-grained, respectively), the fourth columns shows $\Gamma_{cs}^{MD}$ as derived from MD simulations. For this latter, the contour lines are guide for the eyes reproduced from $\Gamma_{cs}^{MC_2}$.</figcaption></figure>

from the same elevated time resolution. $\Gamma_{cs}^{MC_2}$ obtained from MD dislocation trajectories post-treated using the slope variation method exhibit a qualitative match with these results (see technical details in Supplementary Information). While the reduced accuracy of the method does not allow to describe the variations of the total cross-slip rate near $T_c$, it still captures the progressive increase of the rate and the transition, assigning a certain level of confidence to the trajectory slope variation approach. Moreover, the total cross-slip rate estimated using the slope variation method applied to MD simulations ($\Gamma_{cs}^{MD}$) shows a fair agreement with $\Gamma_{cs}^{MC_2}$. In the two simulations, a slight decrease in the transition temperature is observed with increasing the applied shear stress for the both (001) and (1\bar{1}0) MRSSPs (see Figs. 6a and 6b). Additionally, a higher transition temperature is observed for MRSSP=(001) when compared to MRSSP=(1\bar{1}0).

Such a comparison demonstrates that, although the detection resolution is a critical factor when modeling cross-slip using theoretical or simulation approaches, the cross-slip and slip-system transition model developed in Section 3.2 provides an accurate description of the screw dislocation behavior as captured by MD simulations under constant applied shear stress and temperature. The issue of multi-scale time-resolution is addressed using MC simulations: the high-resolution scheme ($MC_1$ model) converges closely to the theoretical predictions, while the lower-resolution scheme ($MC_2$ model) successfully reproduces MD simulations that rely on coarser post-processing. The successful validation of the cross-slip model across a broad range of temperatures and shear stresses paves the way for upscaling, enabling its integration into the DDD simulation.

### 3.4. Discrete dislocation dynamics simulation

This section presents the upscaling of the atomistically-informed cross-slip and slip-system transition model introduced in Section 3.2, as implemented in the DDD code Numodis (Drouet et al., 2014; Gonzalez Joa et al., 2023), along with its first
<figure><img src="image_9.png" /><figcaption>**Fig. 7.** Screw dislocation behavior as computed using DDD under a constant applied stress (MRSSP=(001), $\tau_{yz}=300$ MPa) and constant temperature, (a) 1600 K, the dislocation mostly glides in the MRSSP, few cross-slip events in (111) and (\bar{1}11) slip planes and (b) 2000 K, characterized by consecutive cross-slip and glide events exclusively in the two {111} slip planes. The MRSSP is emphasized using blue dashed-lines. Slip traces are colored according to corresponding slip planes (MRSSP=blue, (111)=green and (\bar{1}11)=red).</figcaption></figure>

applications. The main objective is to enable a more rigorous and physically-grounded modeling of high-temperature plastic deformation of $UO_2$ single crystal. Numodis cross-slip generalities are recalled in Section 2.2, and cross-slip parameters are provided in Tables 1 and 2. The full implementation of the DDD algorithm is documented in the Supplementary Information. The DDD simulation cell and setup were designed to match the MD configuration (Fig. 1b).

Fig. 7 illustrates the behavior of an infinite screw dislocation simulated using DDD at $T=1600$ and 2000 K (below and beyond $T_c$ = 1700 K, respectively) under a constant applied shear stress $\tau_{yz}=300$ MPa, and relying on the MRSSP=(001) geometry. At $T=1600$ K, the dislocation mostly glides in (001) with only brief and occasional excursions in the (111) plane. This behavior reflects the elevated value of $P_{(001)}(T=1600$ K) (see Fig. 4b): once the dislocation cross-slips out of the MRSSP, it rapidly jumps back to it due to the strong preference for (001) core spread and slip. Note that such a behavior is always reinforced by the applied shear stress which always promotes dislocation glide or cross-slip within the MRSSP, this latter via the stress term in Eq. (7). Also, due to the dislocation geometry and the homogeneous applied stress and temperature fields, the dislocation glides as a straight line, remaining within the screw orientation ($\alpha < 2^\circ$) over its entire length, and is thus fully subject to cross-slip. For the second case computed at $T=2000$ K (Fig. 7b), the behavior differs markedly due to the influence of $P_{(111)}^{\text{fcc}}$ at more elevated temperature. In this simulation, the dislocation deviates from the MRSSP by repeatedly cross-slipping between the (111) and (\bar{1}11) slip planes, following a trajectory characterized by the mean angle $\theta$. These first DDD results confirm that the cross-slip model developed and implemented in the Numodis DDD code reproduces, at least qualitatively at this stage, the temperature-dependent evolution of the screw dislocation trajectories observed in MD simulations.

In a more comprehensive manner, trajectories computed using the DDD simulation for the three simulation cell geometries and covering the whole temperature range are shown in Fig. 8, as compared to MD simulations. For the MRSSP=(001) case, the screw dislocation glides exclusively in (001) up to 1600 K in both the MD and DDD simulations. Above this temperature, cross-slip and the gradual transition towards {111} slip originally observed in MD simulation are accurately reproduced by the DDD simulation. These results are in good agreement with those presented in Fig. 7. Similarly, for MRSSP=(110), the transition to {111} slip also occurs for $T \ge 1600$ K, while (110) remains the preferred slip plane at lower temperatures. These findings are also in fair agreement with the MD trajectories computed for this geometry, and originally shown in Fig. 3e,f. For MRSSP=(\bar{1}11), the overall dislocation trajectory is well reproduced by the DDD simulation including exclusive {001} slip at $T = 1400$ K, the onset of the {001}/{111} composite slip at $T = 1600$ K, and the full {001}-to-{111} slip-system transition at $T \ge 1800$ K. This is particularly true when the DDD simulation relies on image forces as described in Section 2.3, here applied using $N_x=2$ and $N_z=100$ (Fig. 8g). For this orientation, DDD simulations performed without image forces (Fig. 8f) already capture, during the early stages, the slip system transition observed in MD simulations, which is primarily temperature-driven rather than stress-driven. This indicates that image forces have a limited influence on the occurrence of the transition itself. However, they become increasingly important once the dislocation significantly moves out from the MRSSP, i.e., for $x \sim 50$ Å, where they noticeably affect the subsequent dislocation trajectory. This result is further supported by the quantification of the mean trajectory $\theta = \arctan(z/x)$ shown in Fig. 9, which confirms that including image forces in the DDD reconciles the mismatch observed in the two initial simulations. A slight discrepancy
<figure><img src="image_10.png" /><figcaption>**Fig. 8.** Trajectory of the 1/2<110> screw dislocation computed under constant applied shear stress and temperature (1400–2200 K range): MD vs. DDD, (a-b) MRSSP = (001) and $\tau_{yz}=100$ MPa, (c-d) MRSSP = (110) and $\tau_{yz}=300$ MPa, e-g) MRSSP = (111) and $\tau_{yz}=100$ MPa (DDD simulations conducted with (w/) and without (w/o) image forces implementation).</figcaption></figure>

still remains concerning the simulation conducted at $T = 1600$ K, where a lower cross-slip rate is noticed in the DDD simulation conducted with image forces. This deviation may stem from the isotropic elasticity assumption used in the DDD simulation, which might overestimates image forces (when compared to MD), or from uncertainties in the dislocation mobility parameters that had to be defined in a simplified manner due to the limited experimental data and the occurrence of cross-slip in the MD analysis. Overall, these results confirm that the atomistically-informed cross-slip and slip-system transition model presented in this study reliably captures the screw dislocation behavior across a wide range of stresses and temperatures, consistent with the operational conditions of $UO_2$ fuel in nuclear reactors.

# 4. Discussion

## 4.1. Cross-slip in $UO_2$: similarities and differences with metallic systems

Cross-slip is a fundamental dislocation mechanism that plays a central role in the plasticity and recovery of crystalline solids. In FCC metals, the cross-slip of 1/2<110> screw dislocations between {111} planes is a well-known thermally activated process that facilitates dynamic recovery via dislocation annihilation at elevated strain, in the so-called stage III that follows the strain hardening regime (Basinski and Basinski, 1979; Jackson, 1985; Kocks and Mecking, 2003; Vivekanandan et al., 2023). In this context, cross-slip in FCC metals is particularly sensitive to stress and temperature, as closely related to the Friedel-Escaig dislocation dissociation/constriction process and the stacking-fault energy (Caillard and Martin, 1989; Duesbery, 1998; Puschl, 2002; Rao et al., 2009; Oren et al., 2017). It is worth noticing that, even if less common, crystallographic cross-slip out of {111} slip planes has also
<figure><img src="image_11.png" /><figcaption>**Fig. 9.** Dislocation trajectory angle $\theta$ between the average glide plane and the (111) MRSSP computed after a simulation time of 2 ns, (a) without image forces in DDD simulations, (b) with image forces (Eq. (4)). Error bars emphasize standard deviation.</figcaption></figure>

been observed in FCC, in relevant temperature and orientation conditions, leading to composite slip (Hazif et al., 1973; Caillard and Martin, 2009).

In BCC metals, dislocations are generally non-dissociated and cross-slip is thus even more frequent than in FCC, underpinning the complex temperature and stress dependence of plastic flow associated to the screw dislocation core structure in BCC compounds. Dislocation can glide in primary {110} slip planes, but also in {112} or {123} possibly leading to non-crystallographic slip usually referred as pencil glide in BCC (Taylor and Elam, 1926a; Christian, 1983; Naamane et al., 2010; Dezerald et al., 2016; Cereceda et al., 2016; Ters and Shehadeh, 2019). This differentiated propensity for cross-slip between BCC and FCC, associated with the screw dislocation core structure, is also responsible for different pinch-off mechanisms of prismatic dislocation loops, which appear during nanoindentation tests (Vítek and Kroupa, 1966; Hsiung, 2010).

In hexagonal close-packed (HCP) compounds, the number of slip systems with shared $\langle a \rangle$ Burgers vectors allows for the activation of multiple slip systems (Akhtar, 1973; Yoo et al., 2001; Wu and Curtin, 2016; Caillard et al., 2018; Liu et al., 2025). This specificity of the HCP structure concerns not only metals but also oxides, such as e.g., sapphire $(\alpha$-Al$_2$O$_3$), which exhibits somehow a similar plastic behavior to that of HCP metals, when deformed at elevated temperature (Farber et al., 1993; Lagerlof et al., 1994). Nevertheless, the plastic deformation of oxides is generally characterized by atypical dislocation properties, such as unusual high stacking-fault energy or the activity of uncommon, non-packed, slip system, all stemming from the nature of the atomic bonding (ionic, covalent or iono-covalent) (Mitchell et al., 1985; Messerschmidt, 2010; Amodeo et al., 2016). In the case of UO$_2$, which crystallizes within the fluorite structure made of a simple cubic oxygen sublattice surrounded by FCC uranium atoms, the alternation of pure anionic and cationic planes along the <001> directions screens the high repulsive forces between cations that occur when shearing, minimizing the electrostatic energy. This promotes 1/2<110>{001} slip, over secondary 1/2<110>{110} and 1/2<110>{111}, at least up to ~1300 K in UO$_2$ (Sawbridge and Sykes, 1970; Sawbridge and and, 1971). A similar electrostatic argument prevails for other crystal structures typical of oxides, such as the rock-salt structure (Takeuchi et al., 2009; Amodeo et al., 2011), in contrast with the classical elastic argument (close-packed planes and dense directions) that generally rules the slip system activation in most metals.

In our previous work (Suchorski et al., 2025), the importance of charge screening upon the stability of the screw dislocation core in UO$_2$ was emphasized using MD simulations. Indeed, the strong disordering of the oxygen sublattice for $T > T_c$, and the resulting breakdown of the local charge arrangement near the dislocation, has shown to be concomitant with a modification of the dislocation core spreading: from a well-established and thermally stable configuration confined in the {001} plane to a more tridimensional configurations involving the {111} planes. In the current study, we demonstrate that such a temperature-induced changes of the dislocation core structure is responsible for (i) a {001}/{111} composite slip regime in the 1500–1700 K temperature range where cross-slip between the two slip modes is particularly effective and (ii) a full {001}-to-{111} slip-system transition for temperatures beyond. Note that close to MRSSP={110}, a {110}/{111} composite slip regime is also observed with a transition to pure {111} slip increasing the temperature. While the cross-slip and slip system transition towards {111} in UO$_2$ shows some similarities with the well-known {111} slip of FCC metals, it is important to note that no dislocation dissociation is observed in UO$_2$ for any slip system. Consequently, the elementary cross-slip mechanism observed in ours simulations differs from the classical Friedel-Escaig process typical of FCC metals. Instead, it resembles the behavior of BCC metals, exhibiting pencil glide of perfect dislocations and non-crystallographic slip regimes that depend on the crystallographic orientation, temperature, and applied stress. But here again, the analogy should not be overextended. In particular, non-Schmid effects, commonly interpreted as the influence of out-of-plane stress components on the dislocation core and mobility, and typically associated with slip-system changes and pencil glide of BCC metals (Taylor and Elam, 1926b; Ito and Vitek, 2001), were found to be negligible in UO$_2$ (see Supplementary Information). This
confirms that temperature controls the intrinsic accessibility of cross-slip in UO₂, as the first parameter acting on the activation barrier itself. In contrast, the applied shear stress only provides an energetic bias that lowers the barrier once cross-slip is already thermally accessible, without significantly altering the underlying core structure or the stability of the oxygen sublattice. This interpretation is further supported by the small activation volume extracted from MD simulations, indicating a limited sensitivity of the activation barrier to stress. More broadly, this identifies a distinct class of cross-slip mechanism, potentially relevant to other oxides or complex alloys, in which temperature-induced instabilities within a specific atomic sublattice fundamentally reshape the plastic deformation mechanisms.

## 4.2. Can the temperature-driven cross-slip model resolve plastic anisotropy in UO₂ single crystal?

In their pioneer work, Sawbridge and Sykes performed compression tests on UO₂ single crystals for various sets of orientations, at a temperature T=1600 K (Sawbridge and al, 1971). These orientations were rationalized using stereographic projections and the standard triangle, where φ orientations relies on compression axis (CA) orientations along the edge of the unit triangle from the [001] pole (φ=0°) to [1̄11] (φ=54.7°), and Ω orientations ranges from [011] (Ω=0°) to φ=16°, near the [1̄15] orientation. The authors reported a strongly anisotropic viscoplastic response, characterized by orientation-dependent flow stress with significant deviations from the Schmid's law (especially near low values of φ and Ω), as well as by pronounced differences in terms of strengthening rate and lattice rotations. From slip trace analysis, they concluded that such behavior arose from the complex activation of slip systems, possibly involving cross-slip and thereby enabling composite slip processes. This hypothesis of a composite slip process active in UO₂ was recently re-examined by Madec and collaborators using the DDD simulation, in the same conditions of deformation than in the experiment (Madec et al., 2023).

For CA||[1̄11] i.e., φ=54.7°, very thin {001} slip bands are characterized in the experiments with only few deviations from the main slip traces. This orientation leads to a maximum Schmid factor *m* ~ 0.47 in a {001} primary slip plane, followed by *m* ~ 0.27 for two {1̄11} slip planes and *m* = 0 in {1̄10}. This configuration is geometrically equivalent to our MRSSP={001} case. For the same orientation, our simulations show the screw dislocation gliding in {001} up to *T* ~1500 K i.e., close to the 1600 K temperature of Sawbridge and Sykes experiments, above which it starts to cross-slip in {1̄11} (see Fig. 3a,b). These results are in good agreements with Sawbridge & Sykes observations and the few deviations mentioned by the authors align with the initiation of cross-slip observed in our simulations, near *T*=1600 K. In this case, our results also corroborate Madec et al. observations conducted using the DDD simulation combined to a particularly soft cross-slip criterion between slip modes. Indeed, Madec's cross-slip model also builds on a simplified Friedel-Esaig formulation, in which cross-slip is primarily governed by the applied stress, with a deliberately low critical stress set to promote cross-slip activity. For the same CA||[1̄11], Madec's DDD simulations show a particularly weak influence of cross-slip on the plastic deformation and stress response computed at *T* = 1600 K i.e., 1/2<1̄10>{001} slip systems are mostly active with and without cross-slip, in good agreement with our results.

For φ = 2° i.e., near CA||[001], Sawbridge & Sykes characterize the activation of multiple {1̄10} slip systems associated with a maximum Schmid factor *m* = 0.5, while being null in {001}. This orientation is the exact same as our MRSSP={1̄10} geometry, with additional *m*=0.408 within two neighbor {1̄11} slip planes (see Fig. 1d). In this case, plastic deformation in the experiments is pretty heterogeneous with extensive marks of cross-slip that lead to wavy slip traces and non-crystallographic glide. This behavior is particularly well described by our simulations in which the screw dislocation glide in the (1̄10) plane with occasional temperature-induced fluctuations in (001) for *T* < *T*<sub>c</sub>, while cross-slip and exclusive glide in the two symmetric (1̄11) and (1̄11) planes arise when increasing the temperature beyond *T*<sub>c</sub> (see Figs. 3e,f and 8c,d). In this latter regime, the screw dislocation repeatedly cross-slips on both sides of the (1̄10) plane, resulting in a harsh {1̄10} trajectory. In other words, increasing the temperature does not alter the primary trajectory direction but only affects the fine structure of the slip trace, changing it from a well-defined {1̄10} crystallographic trace to a wavy, zig-zagging, {1̄10}-oriented trajectory. Note that at the sharp temperature of *T*=1600 K, the screw dislocation trajectory in the MD is characterized by an hybrid regime made of glide and cross-slip events in the (1̄10) and both {1̄11} slip planes ({1̄10}/{1̄11} composite slip). For this orientation, the DDD simulations of Madec et al. likewise reveal a strong influence of cross-slip on the dislocation microstructure, with the primary dislocation activity transitioning from mode III to mode II slip systems, ultimately resulting in a {1̄10}/{1̄11} composite slip regime consistent with that observed both by Sawbridge & Sykes and our study. It is worth noticing that Madec's DDD model *a priori* assumes a dislocation population of various characters distributed on {1̄11} slip planes, in contrast to our simulations where a single screw dislocation is considered without any predefined slip system. The close agreement between the two is thus particularly striking in view of these distinct simulation protocols.

Finally, Sawbridge & Sykes discussed the case with CA close to [011] (Ω=1°) with elevated Schmid factors on {1̄11} and, in a lesser extent, in {001} slip planes. For this orientation, the authors observed very wavy slip traces on the sample surface that is relevant for the screw character. They identified a macroscopically non-crystallographic behavior that converges towards the MRSSP, mentioned to be probably due to cross-slip between {001} and {1̄11} slip planes. For CA|| [1̄10] and *T*=1600 K, Madec et al. show that {001} slip is favored in the DDD simulation regardless of cross-slip activation, although a {001}/{1̄11} composite slip process becomes more likely when cross-slip is enabled. Without being strictly the same orientation, our MRSSP={1̄11} geometry is sufficiently close to offer relevant information for this latter case. Indeed, our MRSSP={1̄11} geometry is similarly characterized by a maximum Schmid factor in (1̄11), followed by *m*=0.40, 0.28 and 0.16 respectively in the (1̄10), (001) and (1̄11) slip planes. Such a configuration is typical of a <012> compression direction e.g., [012] located on the standard triangle close to the [011] pole, in the direction of the [001] one. For this case, our MD simulations reveal pure screw dislocation glide in {001} at low temperature, with occurrence of cross-slip events in {1̄11} emerging near *T*=1600 K, ultimately leading to a {001}/{1̄11} composite slip and possible non-crystallographic slip traces at higher temperature.
<figure><img src="image_12.png" /><figcaption>**Fig. 10.** Simulation of the 1/2<110> screw dislocation interacting with a {110} irradiation loop at *T* = 2000 K, *τ*<sub>s</sub>=900 MPa, (a) *Numodis DDD* simulation relying on the atomistically-informed cross-slip model, (b) MD (CRG interatomic potential).</figcaption></figure>

Overall, our results are fully consistent with the observations of Sawbridge & Sykes, which find in our approach a comprehensive explanation of the slip traces observed experimentally. In addition, the present atomistic study offers a solid physical basis for the cross-slip mechanisms hypothesized in recent DDD simulations by Madec et al. which relied on a phenomenological treatment of cross-slip in UO₂.

### **4.3. Application to irradiation-induced strain hardening in UO₂: cross-slip mediated interactions between primary dislocations and irradiation loops**

As further demonstrated in this study, cross-slip plays an important role in the plastic deformation of UO₂ at high temperature, facilitating dislocation mobility and allowing complex slip interactions. Interestingly, it is shown that it can also become a key mechanism during irradiation where it contributes to the evolution of point defects, dislocation microstructures (recovery, annihilation, mixed-climb) and the accommodation of radiation-induced stresses (Yust and McHargue, 1969; Onofri et al., 2017; Borde et al., 2023b; Onofri et al., 2024, 2025).

In the work of Borde et al. (2023b), the interaction between 1/2<110>{001} primary dislocation and {110} prismatic loops in UO₂ is investigated by mean of MD and DDD simulations. These unfaulted loops emerges during irradiation through the agglomeration of interstitial point defects (e.g., Frenkel pairs) and are characterized by 1/2<110> Burgers vector normal to their habit plane. {110} Prismatic loops in UO₂ can reach diameter up to about 10 nm before coalescing into larger defect and evolve into dislocation lines (Ye et al., 2011; He et al., 2014). Borde and collaborators investigated the interactions between primary 1/2<110>{001} dislocations (edge and screw characters) and prismatic loops, to identify the elementary mechanisms responsible for irradiation-hardening at *T* = 2000 K. In particular, they showed that the screw character played a dominant role, giving rise to the hardest strengthening configurations. In the case where the Burgers vector of the primary screw dislocation was orthogonal to that of the loop, MD simulations using the MOX-07 interatomic potential (Potashnikov et al., 2011) revealed that the reaction was controlled by a cross-slip event occurring just prior to contact between the two line defects. More specifically, the trajectory of the 1/2<110> screw dislocation, initially gliding in (001) (the Borde configuration being equivalent to our MRSSP={001} case), is characterized by a single cross-slip event into the (11̄1) plane containing one segment of the prismatic loop. Such change in slip system, induced by cross-slip, enables a local contact reaction between parallel segments of the primary dislocation and prismatic loop, ultimately leading to the zipping of a <001> Hirth junction. By adjusting a cross-slip DDD model on their MD simulation results, Borde and collaborators showed that the junction formation was only possible when enabling cross-slip; otherwise, crossed-states between orthogonal dislocation segments, without junction formation, was observed. Although not informed by quantitative activation parameters, the cross-slip model developed by Borde et al. demonstrates that cross-slip can affect irradiation hardening in UO₂.

To test our cross-slip model, the same configuration i.e., a screw dislocation vs. prismatic loop with 90° Burgers vectors, is modeled using the DDD simulation with the exact same atomistically-informed cross-slip parameterization than presented in the previous section (no additional adjustments conducted). The simulation is conducted at *T* = 2000 K and *τ*<sub>s</sub> = 900 MPa, following the original setup of Borde et al. including the same simulation cell geometry and dimensions. Our results (Fig. 10a-c) confirm Borde's observation i.e., the gliding screw dislocation interacts with one of the parallel segment of the prismatic loop via a cross-slip process within the (11̄1) plane. However, a major difference arises between the two simulations. While Borde's DDD simulations showed the mobile screw dislocation following a linear trajectory until a single cross-slip event occurred near the prismatic loop, our results reveal multiple thermally-driven, erratic, cross-slip events in the two symmetric (11̄1) and (1̄11) slip planes prior to the contact reactions. Indeed, in this latter case, the screw dislocation is observed to cross-slip and glide on both sides of the
(001) plane, as also previously described in the MRSSP=(001) case for $T > T_c$ (see Fig. 8a), in good agreement with experimental observations (Sawbridge and and, 1971). This result highlights the influence of temperature on the cross-slip process in $UO_2$, an effect that was not captured in previous models. Only when the screw dislocation approaches closer the prismatic loop does the glide process become more predictable, as it is increasingly governed by the stress field of the loop. When sufficiently close, a cross-slip event in the (111) plane, followed by an extended period of glide within the same plane (whereas the dislocation had only zigzagged between planes up to this point), enabled the contact reaction between the primary dislocation and the loop, leading to the formation of the Hirth junction and its subsequent unzipping (Fig. 10a-c). This sequence illustrates the primary temperature dependence of the cross-slip process in $UO_2$, as well as the secondary role of stress, which becomes relevant only in highly stressed regions. Additional MD simulations of the 1/2<110> screw dislocation vs. prismatic loop with 90° Burgers vectors, conducted using the CRG interatomic potential under the same protocol and deformation conditions exhibit the exact same behavior, further confirming the viability of the approach (Fig. 10d-f).

## 5. Conclusion

In this work, we investigated the influence of temperature on the stress-driven trajectory of 1/2<110> screw dislocation in uranium dioxide by combining atomistic simulations, theoretical modeling, and DDD simulations. MD simulations revealed that the behavior of the screw dislocation is strongly affected by temperature and stress geometry. Above $T \sim 1500$ K, cross-slip becomes gradually activated from 1/2<110>{001} and 1/2<110>{110} slip systems towards 1/2<110>{111}, defining the onset of a composite slip regime. At higher temperature (from $T \sim 1700$ K and above), this evolution culminates in a fully-developed slip-system transition where the screw dislocation predominantly glides and cross-slip within the 1/2<110>{111} slip systems. The cross-slip and slip-system transition are found to correlate directly with the temperature-induced modification of the dislocation core structure in $UO_2$ originally characterized in Ref. Suchorski et al. (2025). Before the transition, the predominance of 1/2<110>{001} slip reflects the strong ionic character of bonding in $UO_2$ and screening effects between polar crystallographic planes allow to confine the dislocation core spread. As temperature increases, the progressive disordering of the oxygen sublattice within the dislocation core reduces these electrostatic constraints and weakens the directional nature of ionic bonding. Locally, $UO_2$ thus becomes increasingly akin to a metallic-like FCC structure, whose plasticity is governed by 1/2<110>{111} slip systems. Consequently, the screw core becomes increasingly prone to spread over {111} planes, ultimately leading to a complete transition of the active slip system to 1/2<110>{111} at high temperature. Building on these atomistic observations, we proposed a physically-based cross-slip model in which the probability of transition between slip systems scales with the core-spreading probability. This approach bridges atomistic and mesoscopic descriptions by introducing a temperature-driven stochastic law of cross-slip that captures the underlying changes in the dislocation core structure. By incorporating the previously missing thermally-activated character, our cross-slip and slip-system transition model not only reproduces the key features of screw dislocation motion observed in MD, such as the gradual loss of planar glide and the emergence of wavy dislocation trajectories at elevated temperatures, but also provides a consistent explanation for experimental observations, both in terms of slip system activation and tortuous traces of dislocation motion, in $UO_2$ at high temperature. By elucidating the elementary mechanisms of temperature-driven cross-slip, this work provides a cornerstone for multiscale modeling approaches of $UO_2$ deformation at high-temperature, bridging dislocation-scale physics with higher-scale descriptions in which dislocation glide and climb, diffusion-assisted processes, and grain-boundary-mediated plasticity collectively govern the high-temperature mechanical response. Beyond $UO_2$, the methodology developed here may serve as a framework for linking atomic-scale core physics to dislocation behavior in other high-friction ionic or covalent materials, where temperature-related sublattice instabilities develop and composite slip conditions are observed.

## CRediT authorship contribution statement

**Jules-Elémir Suchorski:** Writing – review & editing, Writing – original draft, Methodology, Investigation, Data curation.
**Adrien Pivano:** Writing – review & editing, Supervision, Methodology, Investigation, Funding acquisition, Data curation, Conceptualization.

**Laurent Dupuy:** Writing – review & editing, Validation, Software, Data curation.
**Jonathan Amodeo:** Writing – review & editing, Writing – original draft, Visualization, Validation, Methodology, Investigation, Data curation, 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.

## Acknowledgments

This research is part of the EURATOM OperaHPC Project, which was co-funded by the European Union. This work was performed using HPC resources from GENCI TGCC Irène under the allocation 2024-R0151010339 and 2024-A0170906922. L.D. acknowledges the continuous support of the MATIX project for the development of the *Numodis* code. J.A. acknowledges Pr. Rabier (Univ. Poitiers, France) for insightful discussions.
# Appendix A. Supplementary data

Supplementary material related to this article can be found online at https://doi.org/10.1016/j.ijplas.2026.104638.

## Data availability

Data will be made available on request.

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