Microstructural evolution of neutron irradiated
ultrafine-grained austenitic stainless steel

Frederic Habiyaremye, Solène Rouland, Bertrand Radiguet, Fabien Cuvilly, Benjamin Klaes, Benoit Tanguy, Joël Malaplate, Christophe Domain, Diogo Goncalves, Marina Abramova, et al.

<figure><img src="image_1.png" /></figure>

To cite this version:

Frederic Habiyaremye, Solène Rouland, Bertrand Radiguet, Fabien Cuvilly, Benjamin Klaes, et al. Microstructural evolution of neutron irradiated ultrafine-grained austenitic stainless steel. Journal of Nuclear Materials, 2025, 607, pp.155710. 10.1016/j.jnucmat.2025.155710. hal-04970952

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Submitted on 10 Mar 2025

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<figure><img src="image_2.png" /></figure>
Microstructural evolution of neutron irradiated
ultrafine-grained austenitic stainless steel.

Frederic Habiyaremye¹, Solène Rouland¹,⁷, Bertrand Radiguet¹,⁷, Fabien Cuvilly¹,
Benjamin Klaes¹, Benoit Tanguy², Joël Malaplate³, Christophe Domain⁴,⁷, Diogo
Goncalves⁴, Marina M. Abramova⁵, Nariman A. Enikeev⁵,⁶, Xavier Sauvage¹, Auriane
Etienne¹,⁷

¹ Univ Rouen Normandie, CNRS, INSA Rouen Normandie, Groupe de Physique des Matériaux UMR 6634, F-76000 Rouen, France

² Université Paris-Saclay, CEA, Service d'Étude des Matériaux Irradiés, 91191, Gif-sur-Yvette, France

³ Université Paris-Saclay, CEA, Service de Recherche en Matériaux et procédés Avancés, 91191, Gif-sur-Yvette, France

⁴ EDF Lab, Département MMC, Les Renardières, F-77250 Moret sur Loing, France

⁵ Laboratory of Metals and Alloys under Extreme Impacts, Ufa University of Science and Technology, 32 Zaki Validi Str., Ufa 450076, Russia

⁶ Saint Petersburg State Marine Technical University, Lotsmanskaya Str. 10, St. Petersburg 190121, Russia

⁷ Laboratoire Commun EDF-CNRS Etude et Modélisation des Microstructures pour le Vieillissement des Matériaux (EM2VM), France

**Corresponding author:** auriane.etienne@univ-rouen.fr

**Keywords**

Ultrafine-grained materials, austenitic stainless steel, neutron irradiation, radiation-induced segregation, microstructure, electron microscopy, atom probe tomography
**Abstract**

Austenitic stainless steels utilized in in-core components of pressurized water reactors are prone to radiation-induced segregation, which leads to the degradation of microstructure and mechanical properties. To improve irradiation resistance, one possible solution is to increase the number density of point defect sinks, such as grain boundaries. For this purpose, ultrafine-grained or nanostructured microstructures are recommended due to their high density of grain boundaries. This paper investigates the microstructural changes in ultrafine-grained 316 austenitic stainless steel exposed to neutron radiation up to 3.9 dpa in irradiation conditions representative of light water reactors. The microstructure at different length scales was analyzed using electron backscattered diffraction, transmission electron microscopy, and atom probe tomography before and after neutron irradiation. The study compares its findings with those of existing literature on coarse-grained austenitic stainless steels to evaluate the benefit of ultrafine-grained 316 austenitic stainless steels regarding irradiation ageing in representative conditions of light water reactors.
# 1. Introduction

Austenitic stainless steels exhibit an unique combination of strength and ductility at both low and high temperatures, as well as a good resistance to creep, fatigue, corrosion, and irradiation [1]. These properties make them ideal for use as structural materials for in-core components of light water reactors (LWRs) such as baffle bolts, baffle plates, etc. [2] [3]. Additionally, they have potential applications in Generation IV and fusion reactors [2] [4] [5]. In LWRs, the aforementioned in-core components serve to support the core weight, facilitate primary coolant circulation, position the core and fuel assembly, and protect the reactor pressure vessel steel from irradiation embrittlement [6]. Due to their proximity to the core, these components are subjected to neutron irradiation and to thermal and mechanical loads, in a corrosive and radioactive environment. Consequently, they are susceptible to deterioration of mechanical properties, which is attributed to the nanoscale evolution of the microstructure [7] [8]. The microstructural alterations are initiated by the interactions of highly energetic neutrons from fission reactions with lattice atoms, resulting in the formation of collision cascades that generate defects (vacancies, self-interstitials and point defect clusters) [9]. The elimination of these excess point defects is achieved through the processes of mutual recombination, clustering, and annihilation at point defect sinks such as grain boundaries (GBs), dislocations, free surfaces, etc. [10] [11] [12] [13]. It is a well-established phenomenon that the diffusion of interstitials or vacancies towards point-defect sinks often results in concomitant preferential fluxes of atoms. This phenomenon, known as radiation induced segregation (RIS), redistributes solute elements at both short and long length scales [13]. RIS has significant implications for the use of stainless steels in core components of nuclear reactors [14]. In a corrosive reactor environment and under an applied stress, depletion of Cr at GB due to RIS can contribute to irradiated assisted stress corrosion cracking (IASCC) [15], which may cause the failure of bolts. Various studies have investigated the IASCC of pressurized water reactor (PWR) baffle/former bolts made from standard (coarse-grained) austenitic stainless steels [14] [15] [16].

In order to reduce the radiation damage, one solution is to increase the number density of internal sinks for radiation-induced point defects. One method involves the introduction of ultra-high-density of dispersion-strengthening nanofeatures [17]. The large fraction of interfaces within these materials act as sinks for point defects. The
use of this method has been successful in Oxide Dispersion Strengthened (ODS) Eurofer and 14YWT steels, as demonstrated by observations in [17] [18]. Another approach is to increase the density of GBs. The latter reduce radiation damage by acting also as point defect sinks [19]. To achieve a sufficient density of GBs, it is necessary to decrease the grain size and produce ultrafine-grained (UFG) materials with a grain size lower than 1 µm [20]. Preliminary researches indicate that, compared to coarse-grained (CG) materials, UFG Cu-0.5Al₂O₃, Nickel, and Tungsten demonstrate superior resistance to radiation damage when subjected to ion or proton irradiation [21] [22]. These results opened the door to exploring more advanced nanostructured and UFG high strength steels. Studies on conventional CG and UFG 316 steel, ion irradiated up to 10 dpa (displacements per atom) at 350°C, demonstrate that UFG 316 steel exhibits reduced segregation within grains and at GBs [23] [24]. Sun et al. [25] examined the effects of 100 keV He ion irradiation at a fluence of 6 × 10²⁰ ions/m² on the irradiation hardening in UFG austenitic steels. Their findings indicated that the density of He bubbles and dislocation loops was diminished in comparison to the CG steels as a result of the introduction of high-angle grain boundaries (HAGBs) and a high density of dislocations, acting as sinks [25].

Literature also reports an excellent radiation tolerance of ultrafine-grained materials in the case of ferritic/martensitic steels [26] [27] [28], however controversial data have been reported on the low activation steels with refined grain structure which were ion irradiated to huge damaging doses approaching several hundreds of dpa. UFG EK-181 or T91 steels exhibited trends for faster swelling development as compared to their conventional structural states [29] [30]. This can be attributed to the fact that imitational ion irradiation studies cannot fully reproduce the effects of a material degradation under neutron irradiations occurring in fission reactors [31] [32].

Actually, only few studies are available in literature on neutron irradiated, in representative reactor conditions, UFG austenitic steels. Shamardin et al. comparatively analyzed the changes in microstructure and mechanical performance in a conventional and an UFG metastable 08C-18Cr-10Ni-Ti reactor steel, which was neutron irradiated in the sodium-cooled fast reactor BOR-60 up to 12 dpa at 350°C and up to 15 dpa at 450°C [33] [34]. It was demonstrated that grain refinement provides reduced accumulation of Frank loops, about 3 times lower than it was reported for the conventional steel. The changes in microstructure correlated with the difference in mechanical properties in the UFG steel proved a less pronounced radiation-induced
hardening and higher both strength and plasticity [33]. In addition, higher resistance for crack initiation has also been demonstrated [34]. However, this metastable steel also showed a tendency for undesired irradiation-induced phase transformation [33]. At the same time, the data on microstructure evolution of neutron irradiated stable UFG austenitic steels, such as UFG 316 steels used in PWR is totally missing. It is essential to know what are the microstructural evolutions caused by neutron irradiation in UFG stainless steels, moreover in LWRs representative irradiation conditions, in order to evaluate the potential use of UFG 316 steels as a material for in-core components. To address this gap, an UFG 316 steel has been elaborated by Equal Channel Angular Pressing (ECAP), in order to have massive samples, and has been neutron irradiated for several years, up to 3.9 displacements per atom (dpa), in an experimental nuclear reactor. Therefore, the irradiated microstructure of the UFG 316 steel was characterized using several techniques, Electron BackScattered Diffraction (EBSD), Transmission Electron Microscopy (TEM) and Atom Probe Tomography (APT). The main objective of this paper is to present the microstructural evolutions due to neutron irradiation and to compare results with existing literature on conventional CG 316 steels.

## 2. Material and methods

### 2.1. Materials, irradiation conditions and preparation

A cylindrical block of 316 steel with the nominal composition given in Table 1 was annealed at 1050°C for 1 hour and then quenched in water. The ultrafine-grained microstructure was processed by ECAP with 4 passes at 400 °C and 1 pass at 500 °C using the Bc route. The Bc route consists of rotating the cylindrical block by 90° counter-clockwise before each pass [35] [36] [37]. In the following, this material will be named as 316-ECAP.

<table><caption>Table 1: The nominal composition of 316 stainless steel (wt.% and at. %) used in this study. Fe is the balance.</caption><thead><tr><th></th><th>C</th><th>Cr</th><th>Ni</th><th>Si</th><th>Mn</th><th>P</th><th>Mo</th><th>Cu</th><th>Co</th><th>S</th></tr></thead><tbody><tr><th>wt. %</th><td>0.026</td><td>16.75</td><td>12.06</td><td>0.64</td><td>1.80</td><td>0.017</td><td>2.55</td><td>0.10</td><td>0.05</td><td>0.016</td></tr><tr><th>at. %</th><td>0.12</td><td>17.92</td><td>11.44</td><td>1.27</td><td>1.82</td><td>0.03</td><td>1.48</td><td>0.09</td><td>0.05</td><td>0.028</td></tr></tbody></table>
After ECAP, the microstructure is fully austenitic as confirmed by X-ray diffraction experiment (See Supplementary materials).

Tensile test specimens were machined from 316-ECAP samples and irradiated in OSIRIS material testing reactor (CEA Saclay, France). The average irradiation temperature was ~ 365°C at the location of the specimens. The fast flux at the irradiation device location was 2.07×10¹⁴ n.cm².s⁻¹ (E > 1MeV). The irradiation of the studied tensile specimen took place during 9 reactor cycles cumulating a total duration of 183 Equivalent Full Power Days corresponding to a total fluence of 3.28×10²¹ n.cm² (E>1MeV). The total dose evaluated at the specimen location is equal to 3.9 dpa and the average dose rate was equal to 3.10⁻⁷ dpa. s⁻¹. After mechanical tests, a small rod has been cut in an irradiated tensile test specimen head and prepared for microstructural analysis. This specimen was finely ground with diamond suspensions down to 1 µm and then subjected to chemico-mechanical polishing using colloidal silica suspension for 20 minutes. In addition, only for unirradiated samples, a final step with Struers A2 solution was used for electro-polishing, lasting 20 seconds at a voltage of 26 V to achieve a fine surface.

## 2.2. Methods

### 2.2.1. Electron backscattered diffraction

An EBSD map of an unirradiated specimen was acquired using an EDAX Hikari camera and APEX™ EBSD software (EDAX, AMETEK Inc., USA) within a JEOL 7900F Field Emission Gun Scanning Electron Microscope (FEGSEM) (Tokyo, Japan). The acquisition parameters comprised an acceleration voltage of 20 kV, a probe current of 3 nA, a step size of 10 nm, and a working distance of 17 mm.

EBSD maps of the irradiated samples were obtained using an EDAX Octane Super camera and APEX™ EBSD software (EDAX, AMETEK Inc., USA) on a Zeiss XB540 Focused Ion Beam Scanning Electron Microscope (FIBSEM) (Oberkochen, Germany). The acquisition parameters consist of an acceleration voltage of 20 kV, a probe current ranging of 5 nA, a step size of 10 nm and a working distance of 12.6 mm.

For both unirradiated and irradiated samples, the post processing and data treatment were performed using the OIM Analysis™ 8.5.1002 software (EDAX, AMETEK Inc., USA).
Grain size has been determined using the linear intercept method implemented in OIM Analysis™, using a tolerance angle of 5° and a minimum number of pixels of 5. The standard deviation is used to calculate the uncertainties.

High angle grains boundaries (HAGBs) and low angle grain boundaries (LAGBs) are
defined as interfaces between grains with misorientation angle greater than 15° and
5°–14°, respectively [38].

2.2.2. Atom probe tomography

Tips for APT were prepared by lift-out followed by annular milling using the Zeiss
XB540 FIBSEM. The lift-out and annular milling procedures are detailed in Ref. [39].
The annular milling's last cleaning step was performed at 2 kV and 90 pA to remove
the Ga-ion damaged layer. The final radii of the tips were less than 50 nm.

The selection of the region of interest from which the tips for APT were obtained was
based on the grain structure observed in SEM. This approach was undertaken to
ascertain that the extraction of the tips occurred in regions characterized by ultrafine
grains.

APT tips were analyzed employing a CAMECA LEAP 5000 XR (AMETEK Inc., USA) utilizing voltage or laser mode. In the voltage mode, datasets were obtained at a temperature of 60 K using a pulse fraction, α, of 20%, pulse rate of 200 kHz, and detection rate of 0.2%. In laser mode, the data sets were collected at a temperature range of 50 to 55 K. The pulse energy was adjusted in order to have an equivalent pulse fraction α of 20% of the standing voltage. The pulse energy was kept constant during analysis. It was set between 20 and 30 pJ. The pulse rate was 200 kHz, and the detection rate varied between 0.1% and 0.2%. Reconstructions of the acquired volumes were carried out using the Cameca's software Integrated Visualization and Analysis Software (IVAS) v3.8.16 (AMETEK Inc., USA) [40] [41]. The 3D reconstruction was performed using the voltage method. In voltage mode, the reconstruction parameters (k<sub>f</sub> the geometric field factor and ξ the image compression factor) were determined using crystallographic poles and interplanar spacing. For laser experiments, k<sub>f</sub> was corrected by dividing the field factor determined in voltage mode by (1+α). Further data treatment was carried out using Gpm_3dSAT software (IDDN.FR.001.430017.000. S.P.2020.000.10000). Overlapped peaks on mass spectrum were identified and decomposed using the natural abundance of the elements, after background noise subtraction. The decomposed data were used to
calculate the global composition of the samples. The matrix composition corresponds to the composition of the volume without all identified objects (segregated grain boundaries, segregated dislocations, clusters....). The isoposition method (IPM) was used to identify the clusters in the unirradiated and irradiated samples. Details about IPM can be found in Ref. [42] [43].

Segregation at a grain boundary was characterized by plotting concentration profiles across a box perpendicularly to it as illustrated in Figure 1a. The box was first divided into small bins, and the composition value quantified in each bin. The experimental profiles were then fitted with a Gaussian function (Figure 1b), which enables the composition of the peak/trough, full width at half maximum (FWHM) and the relative position of the GB to be determined. The peak concentration $Y_c$ is determined using the relation:

$$Y_c = Y_o + \frac{A}{FWHM \sqrt{\frac{\pi}{4\ln(2)}}}$$

where $Y_o$ is the base concentration and A is the area under the curve (Figure 1b). It is worth noting that the results depend on the size of the sampling box, the orientation and curvature of the grain boundary.

The Gibbsian interfacial excess ($\Gamma_i$), which represents the excess number of solute atoms associated with the GB ($N_{excess}$), was obtained by plotting cumulative composition profiles of all atoms against cumulative composition profile of each solute atom for a box normal to the GB as in Figure 1a [44]. Subsequently, the value of $\Gamma_i$ is evaluated using the relation:

$$\Gamma_i = \frac{N_{excess}}{\eta S}$$

where $\eta$ is the detection efficiency of LEAP 5000 XR (equal to 0.52) and S is the cross section of the box using to plot the profile.

### 2.2.3. Transmission electron microscopy

Scanning transmission electron microscopy (STEM) was performed on JEOL ARM 200 CF (Tokyo, Japan) at 200 kV. The TEM lamella of the unirradiated sample was extracted by lift-out technique using a Thermo Scientific Helios 5 PFIB (Plasma
Focused Ion Beam) Dual Beam. The TEM lamella for the sample irradiated at 3.9 dpa was extracted on the Zeiss XB540 FIBSEM.

Lamellae were mounted onto copper TEM grids and thinned to 160 nm via a PFIB at a low voltage of 5 kV and via XB540 at 2 kV. The samples were subsequently reduced to a thickness of 80 nm by back-polishing under a low-voltage Argon ion beam (0.5 KeV) for 35 minutes using the Precise Ion Polishing System II (PIPS II) from Gatan [45]. The thickness of samples during the thinning process was gauged using Energy- Filtered TEM method on a Gatan Imaging Filter (GIF) Quantum ER 965.

STEM bright field (STEM-BF) micrograph were acquired with a semi collection angle of 30 mrad. STEM annular bright field (STEM-ABF) micrograph acquired with an inner semi collection angle of 1.3 mrad and an outer semi collection angle of 5.3 mrad.

Grain size (d) was determined from STEM-BF micrographs using line intercept method. Pattern of horizontal and vertical lines were drawn on the micrographs. The total number of intersections (N) of the pattern of lines and GBs was counted (intersection at the edge and triple points were scored 0.5 and 1.5 respectively whereas other normal intersections were scored as 1). The mean intercept length (d) was calculated using the relation:

$$d = \frac{L}{N}$$

where L is the total length of the lines in the pattern.

Similarly, the dislocation density (ρ) was determined using the line intercept method, as described in Ref. [46], from STEM-BF micrographs. A grid was superimposed on the micrograph using ImageJ [47] and the number of intersections between the grid and the dislocations (n) was counted. The ρ value was calculated via the intercept method, using the following formula:

$$\rho = \frac{2n}{Lt}$$

where L is the total length of the grid, and t is the thickness of the foil.

A correction in order to take into account the projection of dislocations from volume to 2D has been applied to this formula [48]:
$$ \rho_{corr} = \frac{4}{\pi} \rho $$

Frank dislocation loops have been imaged using Dark-field TEM rel-rod method [49] [50] at a zone axis of [011]. Number density and loop size were determined from TEM micrographs. At least 100 Frank loops have been measured.

Cavities have been imaged in TEM using the under/over focus technique. The defocus used in this study is ± 1 µm. The diameter of the cavities has been measured by plotting intensity profiles on TEM images.

For the densities and size, TEM uncertainties are determined from the calculation of the standard deviations of the values sets, meaning the dispersion of the mean values. For Frank loops size, the standard deviation is large, as the loop size distribution is large. So, the statistical uncertainty of type A has also been measured.

# 3. Results

## 3.1. Grain boundaries and dislocations

Microstructure of the samples observed using EBSD before and after irradiation are depicted in Figure 2. Inverse pole figure Z overlaid with HAGB maps show elongated grains with preferential orientation. In Figure 2a, the grain orientation is preferentially around 45° to the transversal direction (Y-direction of the map). Figure 2b shows grains and non-indexed pixels, which appear in white. Non-indexed pixels are prevalent in high deformed sample because small changes in crystal lattice orientation due to high density of dislocations lead to degraded patterns. They were excluded in the analysis. The mean grain size significantly increased from 245 ± 305 nm in the unirradiated sample to 358 ± 457 nm in the irradiated sample as reported in Table 2. The microstructure exhibits heterogeneities, wherein some regions contain ultrafine grains (grain size ranging from 0.1 µm to 1 µm) and other regions contain larger grains (grain size ranging from 1 µm to 10 µm). This represents a typical bimodal grain structure, with elongated deformed grains surrounded by fine equiaxed grains that been reported in others studies [51]. Ultrafine grains have average sizes of 208 ± 169 nm and 257 ± 228 nm, whereas larger grains have average sizes of 1.5 ± 0.6 µm and 1.6 ± 0.5 µm, for unirradiated sample and irradiated sample at 3.9 dpa, respectively. Fraction of HAGBs decreased while the fraction of LAGB increased from the unirradiated state to irradiated state. As expected, no secondary phase was observed from EBSD map.
<table><caption>Table 2: Grain size and percentages of high angle grain boundaries and low angle grain boundaries before and after irradiation at 3.9 dpa. Errors for grain size correspond to the standard deviation.</caption><thead><tr><td></td><td>unirradiated</td><td>3.9 dpa</td></tr></thead><tbody><tr><td>Average grain size (nm)</td><td>245 ± 305</td><td>358 ± 457</td></tr><tr><td>HAGB (%)</td><td>86</td><td>69</td></tr><tr><td>LAGB (%)</td><td>14</td><td>31</td></tr></tbody></table>

Figure 3 displays a STEM-BF micrograph of the unirradiated sample. As observed by EBSD, the microstructure exhibits regions with a significant number of elongated nanostructured grains and sub-grains, accompanied by few larger equiaxed grains. Some of the grains are separated by distinct high-angle grain boundaries marked in dotted red rectangles, while others are separated by non-equilibrium grain boundaries with a high density of dislocations marked in dotted yellow rectangles. The average grain size measured in the STEM micrographs was 280 nm ± 50 nm, which is in good agreement with the EBSD results.

Figure 4a depicts a STEM annular bright-field (STEM-ABF) micrograph illustrating the distribution of dislocations within the grains acquired by orienting the lamellae on [114] zone axis as shown by the SAED pattern in Figure 4b. It highlights the heterogeneous distribution of dislocations within the grain. It also reveals interwoven dislocation cells, as indicated by the red arrows. To determine the dislocation density ($\rho$), a section of the large grain shown by red rectangle in Figure 4c and portrayed in Figure 4d was used, avoiding the dislocation cells. The dislocation density within the grain was calculated as $5 \times 10^{14} \text{ m}^{-2}$. This value is comparable to the dislocation density of $6 \times 10^{14} \text{ m}^{-2}$ determined by the x-ray diffraction method for a sample processed via high pressure torsion at pressure of 6 GPa and temperature of 400°C with 10 die set rotations [52].

After irradiation, the microstructure still exhibits regions with a significant number of elongated nanostructured grains and sub-grains, accompanied by few larger equiaxed grains. As depicted on the STEM-BF image on Figure 5, distinct high-angle grain boundaries and non-equilibrium grain boundaries with a high density of dislocations are still visible. The average grain size measured in the STEM micrographs was (240
± 50) nm. This value is lower than the one measured by EBSD (358 ± 457) nm. The relatively low number of measured grains, either on the EBSD map and on the STEM image, can explain the difference. In any case, the sample after irradiation still exhibits an ultrafine-grained structure.

The dislocation density within the grains, measured from TEM images, is 4.6×10¹⁴ m⁻².
This value is similar to the one before irradiation, indicating that no dislocation recovery
has occurred during irradiation.

As it can be seen on Figure 5b, Frank loops have been formed during irradiation and
are visible on TEM images. Frank loop density after irradiation at 3.9 dpa is equal to
(1.8 ± 0.2)×10²² m⁻³. The average size of the Frank loops is estimated to be 15 nm.
The standard deviation from the mean value is equal to 12 nm. The smallest Frank
loop measures 2.6 nm whereas the largest one measures 60 nm. The statistical
uncertainty is equal to 2.4 nm.

Some small cavities have also been observed by TEM (Figure 5c). The number density of cavities is (8.0 ± 1.4)×10²¹ m⁻³. The average size of the cavities is (1.6 ± 0.3) nm. The smallest cavity measures 1.2 nm whereas the largest one measures 2.1 nm.

## 3.2. Radiation-induced segregation

### 3.2.1. Unirradiated 316-ECAP sample

Several APT analyses were performed on the unirradiated sample; a total volume of 3.4×10⁻²² m³ for a length of 896 nm was analyzed. All elements were randomly distributed within the APT volumes as depicted on the atom maps of one analyzed volume in Figure 6. The global and matrix compositions obtained for the unirradiated sample are reported in Table 3.

Table 3: APT global and matrix composition of the unirradiated 316-ECAP sample.
Global average, minimum and maximum concentrations are given for each element.
Fe is the balance. Fe is the balance.
<table><thead><tr><th></th><th>C</th><th>Cr</th><th>Ni</th><th>Si</th><th>N</th><th>Mn</th><th>P</th><th>Mo</th><th>Cu</th><th>Co</th><th>S</th><th>V</th><th>Al</th></tr></thead><tbody><tr><td>Average (at.%)</td><td>0.76</td><td>18.1</td><td>12.0</td><td>0.90</td><td>0.09</td><td>1.7</td><td>0.02</td><td>1.4</td><td>0.014</td><td>0.04</td><td>-</td><td>0.03</td><td>0.003</td></tr><tr><td>Min (at.%)</td><td>0.09</td><td>16.5</td><td>11.1</td><td>0.74</td><td>0</td><td>1.5</td><td>0.01</td><td>1.2</td><td>0.004</td><td>0.005</td><td>-</td><td>0.02</td><td>0</td></tr><tr><td>Max (at.%)</td><td>1.42</td><td>19.8</td><td>13.9</td><td>1.21</td><td>0.18</td><td>1.8</td><td>0.05</td><td>1.8</td><td>0.02</td><td>0.06</td><td>-</td><td>0.03</td><td>0.005</td></tr><tr><td>Matrix (at.%)</td><td>0.77</td><td>18.1</td><td>12.0</td><td>0.90</td><td>0.09</td><td>1.7</td><td>0.02</td><td>1.4</td><td>0.014</td><td>0.04</td><td>-</td><td>0.03</td><td>0.003</td></tr></tbody></table>

The global and matrix compositions are similar as no cluster or segregated dislocation have been observed in the unirradiated sample. Only one segregated grain boundary within the unirradiated sample has been intercepted in APT volume.

By comparing the nominal and global compositions, except for C, Si and Cu, concentration values between both are similar. An excess of C is measured by APT; C concentration varies largely between the different analyzed volumes as it can be seen in Table 3. Measured Si concentration is lower than expected. Traces of Al and V are measured by APT.

One segregated grain boundary has been intercepted in unirradiated sample. A concentration profile in Figure 7 shows a depletion of Fe atoms. This depletion is compensated by the enrichments of Mo, Cr, C and V, as shown in Figure 7b – d. This kind of segregation has been already observed in a 316 steel deformed by high pressure torsion [52] and is discussed on the discussion section. The width of the GB, as quantified by FWHM, exhibits considerable variability across different elements as shown in Table 4. The mean value is 2.5 nm. The Fe atoms depleted, with a Gibbsian interfacial excess of -5.0 atoms/ nm², are substituted with atoms of Mn, Cr, C, and V atoms as shown in Table 4.

<table><caption>Table 4: Characteristics of the intercepted grain boundary in unirradiated 316-ECAP sample. n.m. means not measurable.</caption><thead><tr><th></th><th>Fe</th><th>Cr</th><th>Mo</th><th>C</th><th>V</th></tr></thead><tbody><tr><td>Peak (at. %)</td><td>60</td><td>22.2</td><td>3.2</td><td>0.4</td><td>0.3</td></tr><tr><td>Γ<sub>i</sub> (atoms/ nm<sup>2</sup>)</td><td>-5.0</td><td>2.8</td><td>1.8</td><td>0.4</td><td>0.1</td></tr><tr><td>FWHM (nm)</td><td>1</td><td>0.6</td><td>1.5</td><td>1.2</td><td>n.m.</td></tr></tbody></table>
### 3.2.2 316-ECAP sample irradiated at 3.9 dpa

Six APT volumes of the sample irradiated at 3.9 dpa were acquired. The global composition and the matrix composition obtained are given in Table 5.

<table><caption>Table 5: APT global and matrix composition of 316-ECAP irradiated at 3.9 dpa (at. %). Global average, minimum and maximum concentrations are given for each element. Fe is the balance.</caption><thead><tr><th></th><th>C</th><th>Cr</th><th>Ni</th><th>Si</th><th>N</th><th>Mn</th><th>P</th><th>Mo</th><th>Cu</th><th>Co</th><th>V</th><th>Al</th></tr></thead><tbody><tr><th>Average (at.%)</th><td>3.0</td><td>17.6</td><td>12.1</td><td>0.73</td><td>0.1</td><td>1.5</td><td>0.007</td><td>1.25</td><td>0.03</td><td>0.1</td><td>0.05</td><td>0.003</td></tr><tr><th>Min (at.%)</th><td>0.15</td><td>17.1</td><td>11.8</td><td>0.52</td><td>0.01</td><td>1.4</td><td>0.002</td><td>0.89</td><td>0.02</td><td>0.08</td><td>0.04</td><td>0</td></tr><tr><th>Max (at.%)</th><td>4.3</td><td>19.3</td><td>12.5</td><td>0.86</td><td>0.06</td><td>1.8</td><td>0.04</td><td>1.4</td><td>0.06</td><td>0.23</td><td>0.06</td><td>0.01</td></tr><tr><th>Matrix (at.%)</th><td>2.9</td><td>17.5</td><td>12.0</td><td>0.66</td><td>0.11</td><td>1.5</td><td>0.006</td><td>1.23</td><td>0.03</td><td>0.10</td><td>0.05</td><td>0.003</td></tr></tbody></table>

By comparing the average global composition of the irradiated sample at 3.9 dpa with the one of the unirradiated sample, a clear excess of C is observed. One of the volumes of the sample irradiated at 3.9 dpa shows a large C concentration (4.3 at. %) compared to others. If this volume is not taken into account, the average C concentration is 0.43 at.%. This value is comparable to the one obtained for the unirradiated sample. The measured global Si concentration is lower in the sample irradiated at 3.9 dpa than in the unirradiated sample.

Looking at the differences between the global and the matrix compositions of the sample irradiated at 3.9 dpa, excepted a small difference for Si, concentrations look slightly similar. The lower concentration of Si in the matrix can be explained by the presence of Si enriched grain boundaries.

Different features are observed in APT volumes i.e., GBs, segregated dislocations, dislocation loops and carbides enriched with Cr and Mo (in only one volume). Contrary to what is expected [53], no Ni-Si enriched clusters are observed in APT volumes. A typical volume is depicted in Figure 8.
In one of the volumes, carbides are observed in the matrix (see supplementary materials). They are enriched in Cr and Mo. Compositions of these carbides as well as enrichment factors F are reported in Table 6. The enrichment factor is given by the ratio of the concentration in the feature over the matrix concentration. Their average diameter is equal to 6.3 nm and their number density is relatively low and equal to (1.0 ± 0.4) ×10²² m⁻³. It should be noted that these carbides have been observed in the volume with the higher C concentration. In the same volume, dislocations with C, Mo and Cr are observed. Their characteristics are also given in Table 6. Carbides and Mo-Cr dislocations are enriched with the same elements and enrichment factors are slightly similar.

In other volumes, dislocations segregated with Ni and Si are observed. The density of segregated dislocations is equal to 4.5×10¹⁴ m⁻².

Some small dislocation loops decorated with Ni and Si and with a diameter of about 3.4 nm are also present in the sample. There number density is relatively low and equal to (4 ± 2)×10²¹ m⁻³. As well as other objects observed in the sample irradiated at 3.9 dpa, the composition of dislocation loops is given in Table 6. Dislocation lines and dislocation loops show similar segregations in terms of elements and enrichment factors are closed for both lines and loops. Dislocation lines are a bit more enriched in P and in Si.

<table><caption>Table 6: Average density, size and composition obtained by APT of different objects observed in the sample irradiated at 3.9 dpa (at. %). The enrichment factor F is also reported. Fe is the balance.</caption><thead><tr><th></th><th>Density</th><th>Size (nm)</th><th></th><th>C</th><th>Cr</th><th>Ni</th><th>Si</th><th>N</th><th>Mn</th><th>P</th><th>Mo</th><th>Cu</th><th>Co</th><th>V</th><th>Al</th></tr></thead><tbody><tr><td rowspan="2">Mo-Cr carbides</td><td rowspan="2">10<sup>22</sup> m<sup>-3</sup></td><td rowspan="2">6.3</td><td>C (% at.)</td><td>15.9</td><td>41.2</td><td>2.9</td><td>0.25</td><td>0.1</td><td>0.6</td><td>0.03</td><td>4.7</td><td>0.06</td><td>0.04</td><td>0.1</td><td>-</td></tr><tr><td>F</td><td>5.4</td><td>2.4</td><td>0.25</td><td>0.4</td><td>0.9</td><td>0.4</td><td>4.9</td><td>3.8</td><td>2.3</td><td>0.4</td><td>2</td><td>-</td></tr><tr><td rowspan="2">Cr-Mo dislocations</td><td rowspan="2">2.2×10<sup>15</sup>m<sup>-2</sup></td><td rowspan="2">-</td><td>C (% at.)</td><td>10.6</td><td>32.0</td><td>6.8</td><td>0.5</td><td>0.05</td><td>1.2</td><td>0.02</td><td>3.6</td><td>0.02</td><td>0.03</td><td>0.2</td><td>-</td></tr><tr><td>F</td><td>3.6</td><td>1.8</td><td>0.6</td><td>0.8</td><td>0.5</td><td>0.8</td><td>3.3</td><td>2.9</td><td>0.8</td><td>0.3</td><td>4</td><td>-</td></tr><tr><td rowspan="2">Ni-Si dislocations</td><td rowspan="2">4.5×10<sup>14</sup>m<sup>-2</sup></td><td rowspan="2">-</td><td>C (% at.)</td><td>0.3</td><td>13.2</td><td>24.4</td><td>2.8</td><td>0.3</td><td>1.1</td><td>0.14</td><td>0.9</td><td>0.1</td><td>0.3</td><td>0.02</td><td>0.02</td></tr><tr><td>F</td><td>0.1</td><td>0.75</td><td>2.0</td><td>4.3</td><td>2.8</td><td>0.7</td><td>23</td><td>0.7</td><td>3.9</td><td>3</td><td>0.5</td><td>-</td></tr><tr><td rowspan="2">Dislocation loops</td><td rowspan="2">4.1×10<sup>21</sup>m<sup>-3</sup></td><td rowspan="2">3.4</td><td>C (% at.)</td><td>0.2</td><td>14.2</td><td>24.9</td><td>1.3</td><td>0.15</td><td>1.8</td><td>0.05</td><td>0.9</td><td>0.05</td><td>0.35</td><td>-</td><td>0.05</td></tr><tr><td>F</td><td>0.1</td><td>0.8</td><td>2.0</td><td>2.0</td><td>1.4</td><td>1.2</td><td>8</td><td>0.7</td><td>2</td><td>3.5</td><td>-</td><td>-</td></tr></tbody></table>
Six grain boundaries with solute segregation have been intercepted in atom probe volumes in the sample irradiated at 3.9 dpa. Grain boundaries are either enriched with Ni and Si, or enriched with Cr and Mo. Four kinds of grain boundary segregation are observed: (i) Three of these boundaries have a homogeneous distribution of segregated solutes, on the area of the grain boundary. One example is depicted in Figure 9(c and d). They are enriched in Ni and Si, and depleted in Cr and Fe (ii) One GB has a heterogenous distribution with enrichment in Cr, Mo and C elements, with no specific structure. This grain boundary is visible on Figure 9 (a and b). (iii) One GB exhibits a heterogenous segregation of Ni and Si and with a depletion in Cr and Fe, with no specific structure. (iv) Finally, as depicted on Figure 9 (e and f) one shows a dislocation structure, indicating that it is a LAGB. The dislocation structure looks like dislocation cells with parallel dislocations in one direction. The mean distance between dislocations is 4.2 nm.

An example of a concentration profile plotted perpendicularly to a homogeneous Ni-Si segregation of a grain boundary is given on Figure 10. Ni and Si enrichments are clearly visible. The pic concentration of Ni reaches 27 at. % and the one of Si 8.5 at.%. In the same time, Fe concentration decreases to 46 at. % and Cr concentration to 12 at. % at the grain boundary. Enrichments in P and C are also observed. Mo profile looks like a W-shape profile, with an enrichment at the grain boundary and depletion on each side. Whereas Ni, Si, Cr, and Fe profiles are extended over maximum 2-3 nm, Mn profile is larger than 5 nm, as well as the Mo W-shape profile.

The average characteristics of grain boundaries in the sample irradiated at 3.9 dpa are reported in Table 7 for the 4 types of observed intergranular segregation, i.e., heterogeneous segregation with enrichment in Cr, Mo and C, heterogeneous segregation with enrichment in Ni and Si, homogeneous segregation with enrichment in Ni and Si, and low angle grain boundary. For each case, the elemental concentrations at peak, the full width at half maximum and the Gibbs excess are given.
Note that some grain boundaries are also enriched in P (Figure 10) but the fitting of concentration plot was difficult and so no fitting values can be given.
<table><caption>Table 7: Characteristics of different grain boundaries intercepted by APT in the sample irradiated at 3.9 dpa. Elemental concentrations at GB peak, Gibbs excess and GB full width at half maximum are given. Empty cases mean that there is no enrichment or depletion.</caption><thead><tr><th></th><th>Fe</th><th>Cr</th><th>Ni</th><th>Si</th><th>Mo</th><th>Mn</th><th>C</th></tr></thead><tbody><tr><td colspan="8">Heterogenous GB with Cr and Mo enrichment</td></tr><tr><td>Peak (at. %)</td><td>55.4</td><td>20.9</td><td></td><td></td><td>3.0</td><td></td><td>3.2</td></tr><tr><td>Γ<sub>i</sub> (atoms/ nm<sup>2</sup>)</td><td>-13.9</td><td>7.1</td><td></td><td></td><td>2.2</td><td></td><td>5.0</td></tr><tr><td>FWHM (nm)</td><td>3.3</td><td>2.9</td><td></td><td></td><td>2.1</td><td></td><td>4.5</td></tr><tr><td colspan="8">Heterogeneous GB with Ni and Si enrichment</td></tr><tr><td>Peak (at. %)</td><td></td><td>12.0</td><td>19.7</td><td>3.6</td><td>1.0</td><td>0.6</td><td></td></tr><tr><td>Γ<sub>i</sub> (atoms/ nm<sup>2</sup>)</td><td></td><td>-26.9</td><td>23.7</td><td>13.9</td><td>-11.4</td><td>-7.8</td><td></td></tr><tr><td>FWHM (nm)</td><td></td><td>1.6</td><td>1.6</td><td>1.9</td><td>5.5</td><td>2.8</td><td></td></tr><tr><td colspan="8">Homogeneous GB with Ni and Si enrichment</td></tr><tr><td>Peak (at. %)</td><td>48.8</td><td>11.5</td><td>27.6</td><td>7.7</td><td>1.5</td><td>0.7</td><td></td></tr><tr><td>Γ<sub>i</sub> (atoms/ nm<sup>2</sup>)</td><td>-15.3</td><td>-16.0</td><td>26.8</td><td>9.7</td><td>-2.2</td><td>-4.6</td><td></td></tr><tr><td>FWHM (nm)</td><td>1.7</td><td>4.7</td><td>2.6</td><td>2.5</td><td>4.4</td><td>6.4</td><td></td></tr><tr><td colspan="8">Low Angle Grain Boundary</td></tr><tr><td>Peak (at. %)</td><td>55.8</td><td>11.5</td><td>25.0</td><td>6.6</td><td>0.5</td><td>0.5</td><td></td></tr><tr><td>Γ<sub>i</sub> (atoms/ nm<sup>2</sup>)</td><td>-14.0</td><td>-22.3</td><td>32.4</td><td>15.1</td><td>-6.8</td><td>-5.7</td><td></td></tr><tr><td>FWHM (nm)</td><td>2.1</td><td>5</td><td>2.7</td><td>2.6</td><td>13.3</td><td>5.9</td><td></td></tr></tbody></table>

The 3 kinds of Ni-Si intergranular segregation exhibit the same trends in behavior: enrichment for Ni and Si and depletion for Fe, Cr and Mn. Mo is enriched in one GB and depleted in two GBs.

Gibbs interfacial excess shows that Ni and Si enrichments are balanced by Cr, Mo, Mn and Fe depletion. Fe and Cr exhibit the largest negative Gibbs excess, Ni and Si the largest positive Gibbs excess. Even if the peak concentration for Mo and Mn are not so low in comparison to the matrix composition, as their concentration profile is spread on several nanometers (up to 13.3 nm for Mo), their Gibbs excess is not negligible.

### 3. Discussions

A heterogeneous UFG fully austenitic microstructure consisting of elongated grains and sub-grains was observed after ECAP with 4 passes at 400°C and 1 pass at 500°C with a high density of dislocations and an average grain size of (245 ± 305) nm from
EBSD map. The STEM micrograph showed highly deformed microstructure with an average grain size of (280 ± 50) nm. The difference between both values is explained by a lower number of grains measured in STEM. Similar microstructures have been reported in previous studies [54] [55] [56]. Mal'tseva et al. [54] observed domains with sizes ranging from 100 – 250 nm using STEM after ECAP with 4 passes at 400°C in metastable 03Kh14N11K5M2YuT and 03Kh14N11M2YuT austenitic stainless steels. Hajizadeh et al. [55] also found a wide grain size distribution ranging from less than 20 nm to 200 nm after 8 passes at 350°C in a 316 steel. Similarly, an average grain size of 350 nm was observed after ECAP with 8 passes at 400°C in austenitic stainless steel [56]. The variation in the mean grain size among various studies has resulted from various factors, such as the site chosen for sample extraction, the techniques employed for data acquisition (e.g., EBSD or (S)TEM), the methods used for determining grain size. Based on our results and previous studies, it is evident that the sample produced through ECAP exhibits an ultrafine-grained microstructure with an average grain size smaller than 300 nm.

Our findings also confirm that there is no secondary phase formed after 4 passes of ECAP at 400°C and 1 pass at 500°C. These results are consistent with previous studies after ECAP at intermediate temperature [54] [55] [56]. Conversely, strain-induced martensite was detected as a secondary phase after severe plastic deformation by ECAP at room temperature in ultra-low carbon austenitic steel [57] and (0.08%C–18.3%Cr–9.8%Ni–0.6% in wt.%) austenitic stainless steel [58]. It is possible that the absence of strain-induced secondary phase after ECAP at intermediate temperatures is linked to the decrease in strain hardening caused by high temperatures. The latter also minimizes the nucleation sites such as stacking faults, intersections of slip bands, and twins on which secondary phase can nucleate [59].

After neutron irradiation at 3.9 dpa, grain size measured by EBSD is (358 ± 457) nm. The grain size measured by STEM is smaller; (240 ± 50 nm). The statistics relatively low in STEM (small number of grains) can explain the difference between both values. Looking only at grain sizes obtained by EBSD, a grain growth was observed after neutron irradiation. Grain growth after ion, proton and neutron irradiation was observed in nanostructured and ultrafine-grained materials in numerous studies. Radiguet et al. [23] observed grain growth in 316 steel that underwent high pressure torsion (HPT) at room temperature, and subsequently irradiated with 160 keV Fe ions at 350°C. The
grain size grew from 40 nm prior to irradiation to 60 nm post-irradiation. No grain growth was observed after annealing at 350°C. Grain growth was also observed in Cu-0.5Al₂O₃ processed by ECAP and HPT [60]. The average grain size increased from 178 nm before irradiation to 493 nm subsequent to proton irradiation up to 0.91 dpa at room temperature. For Cu-0.5Al₂O₃, annealing begins at 175°C. Given that the irradiation occurred at room temperature, grain growth can be attributed solely to proton irradiation. Grain growth induced by He ion irradiation up to 24.6 dpa was also reported in nanostructured ferrite/martensite steel produced by surface mechanical attrition treatment [61]. The irradiation temperature of 365°C in our study is below the annealing temperature required to induce grain growth. Thus, the observed grain growth can be attributed to neutron irradiation rather than thermal effects. Grain growth caused by irradiation is linked to thermal peaks that occur during the formation of displacement cascades and sub-cascades during irradiation. The migration of GBs is triggered by these thermal peaks, ultimately leading to grain growth under irradiation [62].

The absence of intragranular segregation in the unirradiated sample in this study is consistent with the findings of previous studies [54].

Regarding APT results and the total length and volume analyzed, grain boundaries in the unirradiated state are either not segregated or enriched with Cr and Mo. In the literature, following severe plastic deformation, minor segregation of elements on grain boundaries was reported in nanostructured steel. As observed in this study, in a 316 steel, Abramova et al. [52] observed enrichment of Mo, Si and Cr at grain boundaries following high-pressure torsion at 400°C. However, no such segregation was observed at room temperature [52]. Similarly, after high pressure torsion at 430°C and annealing at 450°C for 5 hours, enrichment in Cr, Si, P, C and Mo was reported in the resulting nanostructured 316 steel [63]. Given that no evidence of segregation at the grain boundaries has been observed in the literature at room temperature, the segregation observed in this work (4 passes at 400°C and 1 pass at 500°C) may be related to the thermal diffusion of solute elements during ECAP.

After neutron irradiation at 3.9 dpa, Frank loops and cavities are observed in TEM. In addition, Ni and Si segregation on dislocation loops is observed thanks to atom probe tomography. The number density of dislocation loops and their size measured in APT are smaller (4×10²¹ m⁻³ and 3.4 nm, respectively) to what observed on TEM micrographs (1.8×10²² m⁻³ and 15 nm, respectively). The difference in number density
can be explained by the following two reasons: First, the number of Frank loops observed in APT is very low, this leads to a lack of statistics in APT. For the difference in size, due to the small acquired volumes in APT, the largest loops are certainly partially intercepted in volumes and are not considered as loops. Secondly, in TEM, the number density and size of dislocation loops have been measured on largest grains with a size larger than 1 µm. Depending on the size of the grain in which the defects are counted, a difference of size and number density can be observed, the smaller the grain size, the lower the number density of defects [64] [10]. As shown by Matsuoka et al. [20], effects of the nanostructuration on damage limitation (hardness measurements) are mostly visible for grain size smaller than 300 nm. In APT, looking at the number of grain boundaries intercepted, the size of grains analyzed is lower than 1 µm, and/or the analyzed volumes are very close to grain boundaries and so can be located in the denuded zone, a defect free zone near the grain boundaries [65] [66].

While comparisons of samples that have undergone different irradiation conditions can be challenging, the consistency between our findings and those reported in [33] for UFG 321 type steel, which was irradiated to 12 dpa at 350 °C in sodium-cooled fast reactor BOR-60, is evident. Given the higher dose, the number density of frank loops was (3.0 ± 0.7)×10²² m⁻³, which is slightly higher than (1.8 ± 0.2)×10²² m⁻³ for our study as anticipated. When similar conditions were applied to coarse grain 321-type steel, the number density of frank dislocation loops was found to be 2-3 higher, thereby confirming the effectiveness of UFG steel in reducing intragranular segregation. This is depicted on Figure 11 (a and b). Although our study did not examine coarse steel subjected to comparable conditions to those of our ECAP, it is anticipated that the number density of Frank loops would be greater in such sample.

At this irradiation fluence, it is common in austenitic stainless steels to observed these irradiation defects and it is well reported in the literature in conventional steels [67] [68] [50] [69] [70] [71] [33] [72] [73]. Frank Loops in conventional 316 steel are usually present with a high number density of 10²²-10²³ m⁻³. In comparison to literature results, the number density of Frank loops measured by TEM in our 316-ECAP steel is in the lower range of values found in the literature for conventional 316 steel neutron irradiated in experimental reactors or PWR (Figure 11a). The size of the Frank loops (Figure 11b) observed in TEM in the 316-ECAP is in the upper range of values found for conventional 316 steel.
Looking at cavities, results found in this study are compared to those found in neutron irradiated conventional steel on (Figure 11c and d). Cavity number density and size are in the lower range of what is observed in neutron irradiated conventional 316 steel.

Compared with the literature results, this work shows that the irradiation defects (dislocation loops and cavities) are relatively few in number.

After neutron irradiation, different phases are known to precipitate or form, like carbides, phosphides, Ni and Si rich phases or intermetallic in conventional neutron irradiated 316 steel [53] [50] [74] [75] [76]. In this study, formation of phases in the 316-ECAP material after neutron irradiation has been characterized thanks to APT. Mo-Cr carbides heterogeneously distributed, have been observed. Their composition, measured in APT, is between $M_2C$ and $M_{23}C_6$ composition. The presence of carbide is consistent with the findings of Hoshimoto et al. [67] who observed a low number density of $M_{23}C_6$ of ~$10^{19} m^{-3}$ after neutron irradiation to 3 dpa in CG 316 steel at 90°C and 250°C. The discrepancy in values between the two studies can be attributed to two factors: firstly, the slightly higher dose in the current study, and secondly, the fact that the current study employed APT for data collection, whereas the other study utilized TEM imaging.

No Ni-Si clusters have been observed by APT after neutron irradiation. The absence of Ni-Si enriched phases in grains indicates that the formation of such phases is limited by the high density of grain boundaries and the high density of dislocations.

To summarize, after neutron irradiation, Frank loops and cavities are present in the 316-ECAP materials in a less extend than in conventional 316 steel. Some carbides are present in grains of the 316-ECAP. No Ni-Si enriched clusters are observed. All these observations indicate a limited irradiation damage in the 316-ECAP, thanks to the nano-sized grains and to the high density of dislocations.

Regarding radiation induced segregation at GBs for the 316-ECAP irradiated at 3.9 dpa, deux kinds of grain boundaries segregation are observed: (i) in the majority of GBs, a Si and Ni enrichment with Cr and Fe depletion and (ii) in one case, Cr, Mo and C enrichment at a GB due to the presence of carbides. Ni and Si radiation induced segregation at GBs has been relatively well documented in the literature in coarse-grained 316 steels following neutron irradiation. The values of $R_{Cr}$, $R_{Ni}$ and $R_{Si}$ ($R_i$ is the ratio between the average peak concentration of the element $i$ at the GB and the
nominal concentration of i in the alloy) as a function of irradiation dose for conventional
316 steels [73] [77] [70] [78] are reported on Figure 12. As the composition of stainless
steel can vary from one study to another, the composition of a given element in a GB
is dependent of its nominal composition. Therefore, the Rᵢ-value is introduced to
remove such bias. Results for the six GB analyzed in APT in this work are also reported
on Figure 12. The depletion in Cr observed in the 316-ECAP is a bit higher than what
is reported in the literature. In the same extent, the enrichment in Ni and Si are higher
in the 316-ECAP.

Our results on the GBs in the ultrafine-grained 316 steel illustrates that the depletion
of Cr and enrichment of Ni and Si are slightly higher than in coarse-grained 316 steel.
Some discrepancies between literature results can be induced by the fact that most of
the literature results on RIS at GBs have been obtained using TEM-EDS (Energy
Dispersive X-Rays Spectroscopy). On Figure 12, literature data obtained by TEM-EDS
and using APT are represented with circle marks and triangle marks respectively. Its
seems that APT data show higher radiation-induced segregation than TEM.

Higher RIS in 316-ECAP than in coarse-grained 316 steel contrasts with the
anticipated outcome based on the ion irradiated nanostructured 316 steel [23] [24].
Unlike our 316-ECAP sample with a heterogeneous microstructure and average grain
size of ~250 nm, ion irradiation was performed on samples processed by high pressure
torsion with a homogeneous microstructure and average grain size of 40 nm. As shown
by [79] [80] [10], the smaller grain size, the higher the limitation of radiation damage,
and so the higher the limitation of radiation-induced segregation at GBs.
## 4. Conclusions

This study examines the microstructural evolutions in ultrafine-grained 316 austenitic stainless steel after exposure to neutron irradiation at 3.9 dpa in real conditions of light water reactor. Microstructure was examined using electron backscatter diffraction, transmission electron microscopy and atom probe tomography. The results can be summarized in three key points:

* Comparative analysis conducted pre- and post-irradiation indicated that neutron irradiation at 350°C increases grain size.

* Frank loops, cavities and Mo-Cr carbides are observed after neutron irradiation. Number density and size of Frank loops and cavities are in the lower range of magnitude to what can be found in literature in conventional 316 steels.

* No Ni-Si clusters have been observed in grains after neutron irradiation.

* Radiation-induced segregation measured by APT of the major elements Ni, Si and Cr in the ultrafine grained was slightly higher than in the coarse-grained 316 stainless steels.

Our findings indicate that the contribution of grain refinement and the high density of dislocations induced by ECAP after irradiation is slightly significant. It limits the formation of point defect clusters (cavities and Frank loops) and solute clusters (Ni-Si) in the grains. Nevertheless, there is no evidence of limitation of radiation-induced segregation at GBs in comparison to coarse-grained 316 stainless steels.

## 5. Acknowledgements

The authors acknowledge the financial support from CARNOT-ESP in the framework of RESIRAN project.

Experiments have been performed on GENESIS and GPM-microscopy platform in GPM laboratory in Rouen, France. GENESIS is supported by the French National Research Agency under the "Investissements d'avenir" as part of the program with the reference ANR-11-EQPX-0020, the Normandy Region, the Métropole Rouen Normandie, the CNRS via LABEX EMC3. TEM acquisitions on RIO camera have been performed thanks to CATHY project. The European Union and Région Normandie have contributed to the realisation of CATHY project. CATHY is co-funded via the European Regional Development Fund. LEAP experiments have been performed
thanks to GENESIS+ project. European Union and Région Normandie have
contributed to the realisation of GENESIS+ project. GENESIS+ is cofund via the
European Regional Develompent Fund.

Irradiation was performed in the frame of a CEA, EDF funded project.

This work was partially supported by the CNRS Federation IRMA-FR 3095.

N. Enikeev acknowledges partial support by the Ministry of Science and Higher
Education of the Russian Federation as part of World-class Research Center program:
Advanced Digital Technologies [contract No. 075-15-2022-312 dated 20 April 2022].

6. CRediT authorship contribution statement

Frederic Habiyaremye – Methodology, Formal analysis, Investigation, Writing -
Original Draft, Visualization

Solène Rouland – Methodology, Formal analysis, Investigation, Writing - Review & Editing

Bertrand Radiguet - Conceptualization, Methodology, Writing - Review & Editing,
Funding acquisition, Visualization

Fabien Cuvilly - Formal analysis, Investigation, Writing - Review & Editing

Benjamin Klaes - Software, Investigation, Writing - Review & Editing

Benoit Tanguy - Conceptualization, Methodology, Formal analysis, Investigation, Writing - Review & Editing

Joël Malaplate - Conceptualization, Methodology, Formal analysis, Investigation, Writing - Review & Editing

Christophe Domain – Conceptualization, Writing - Review & Editing

Diogo Goncalves – Conceptualization, Writing - Review & Editing

Marina M. Abramova – Resources, Writing - Review & Editing

Nariman A. Enikeev – Resources, Writing - Review & Editing
Auriane Etienne – Conceptualization, Methodology, Formal analysis, Investigation, Writing - Review & Editing, Funding acquisition
7. References

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8. Supplementary materials
<figure><img src="image_3.png" /><figcaption>Figure 1</figcaption></figure>
<figure><img src="image_4.png" /><figcaption>Figure 2</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>Figure 3</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>Figure 4</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>Figure 5</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>Figure 6</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>Figure 7</figcaption></figure>

<figure><img src="image_10.png" /></figure>

<figure><img src="image_11.png" /></figure>

<figure><img src="image_12.png" /></figure>
<figure><img src="image_13.png" /><figcaption>Figure 8</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>Figure 9</figcaption></figure>
<figure><img src="image_15.png" /><figcaption>Figure 10</figcaption></figure>

<figure><img src="image_16.png" /></figure>

<figure><img src="image_17.png" /></figure>

<figure><img src="image_18.png" /></figure>
<figure><img src="image_19.png" /><figcaption>Figure 11</figcaption></figure>
<figure><img src="image_20.png" /><figcaption>Figure 12</figcaption></figure>
Microstructural evolution of neutron irradiated
ultrafine-grained austenitic stainless steel.

**<u>Supplementary materials</u>**

Frederic Habiyaremye¹, Solène Rouland¹,⁷, Bertrand Radiguet¹,⁷, Fabien Cuvilly¹,
Benjamin Klaes¹, Benoit Tanguy², Joël Malaplate³, Christophe Domain⁴,⁷, Diogo
Goncalves⁴, Marina M. Abramova⁵, Nariman A. Enikeev⁵,⁶, Xavier Sauvage¹, Auriane
Etienne¹,⁷

¹ Univ Rouen Normandie, CNRS, INSA Rouen Normandie, Groupe de Physique des Matériaux UMR 6634, F-76000 Rouen, France

² Université Paris-Saclay, CEA, Service d'Étude des Matériaux Irradiés, 91191, Gif-sur-Yvette, France

³ Université Paris-Saclay, CEA, Service de Recherche en Matériaux et procédés Avancés, 91191, Gif-sur-Yvette, France

⁴ EDF Lab, Département MMC, Les Renardières, F-77250 Moret sur Loing, France

⁵ Laboratory of Metals and Alloys under Extreme Impacts, Ufa University of Science and Technology, 32 Zaki Validi Str., Ufa 450076, Russia

⁶ Saint Petersburg State Marine Technical University, Lotsmanskaya Str. 10, St. Petersburg 190121, Russia

⁷ Laboratoire Commun EDF-CNRS Etude et Modélisation des Microstructures pour le Vieillissement des Matériaux (EM2VM), France

**Corresponding author:** auriane.etienne@univ-rouen.fr

**Keywords**

Ultrafine-grained materials, austenitic stainless steel, neutron irradiation, radiation-
induced segregation, microstructure, electron microscopy, atom probe tomography
<figure><img src="image_21.png" /><figcaption>Figure supplementary material 1: XRD pattern of the 316 ECAP in the unirradiated state.</figcaption></figure>
<figure><img src="image_22.png" /><figcaption>Figure supplementary material 2: APT volume of 316-ECAP irradiated at 3.9 dpa. C and Mo atoms are represented. Highlighted regions are regions with Mo and Cr concentrations higher than 25.5 at.%.</figcaption></figure>
<figure><img src="image_23.png" /></figure>