The Bright-Field TEM micrographs shown in Fig. 1 display the nanostructure of the as-deposited Al 2 O 3 thin films. The dark contrast spots correspond to randomly-oriented ultra-fine nanocrystalline γ-Al 2 O 3 domains (6 ± 4 nm), whereas the bright contrast results from the presence of the amorphous phase of Al 2 O 3 . The appearance of a relatively sharp ring, together with a diffused intensity halo in the diffraction pattern (DP) confirms that the the material consists of a dual phase structure, with the amorphous phase dominating the overall structure. The volume fraction of γ-Al 2 O 3 nanocrystalline domains is very low at approximately 1% (see Supplementary Information).

Figure 1 BF-TEM micrograph, and high-resolution (HR) close-up (inset) of the nanostructure of the as-deposited Al 2 O 3 thin films showing a homogeneous dispersion of a low volume fraction of randomly-oriented nanocrystalline Al 2 O 3 domains in an amorphous Al 2 O 3 matrix. Full size image

The main advantage of this type of dual structure over a fully nanocrystalline structure is that it confers an unusual ensemble of metal-like mechanical properties (Young’s modulus E = 195 ± 9 GPa, ν = 0,29 ± 0,02) and moderate hardness (H = 10 ± 1 GPa) to the Al 2 O 3 thin films. In particular, the amorphous matrix precludes grain sliding, enables plastic deformation and inhibits crack nucleation29. Despite the lower hardness compared to single crystal sapphire (H sapphire = 27.6 ± 2 GPa)33, the Al 2 O 3 thin films are still significantly harder than most metallic materials. Moreover, the resulting H/E ratio (i.e. 0.051) is comparable with the H/E ratios of superhard nanocomposite coatings for tribological applications34. This may be beneficial for example during fuel rod insertion or grid-to-rod fretting during operation.

The ADF-STEM micrographs in Fig. 2 show the structural features of the as-deposited and the irradiated thin films. These images indicate that a fully nanocrystalline structure is realized upon irradiation, and that extended irradiations induce grain growth as the dpa levels are increased. The average grain size increases from 6 ± 4 nm to 101 ± 56 nm at 20 dpa, 153 ± 62 nm at 40 dpa and 293 ± 85 nm at 150 dpa (Fig. 2b–d). The crystallization and grain growth observed manifest as an evolution of the DPs from a diffused intensity halo to rings and isolated spots. The crystalline phases present in the irradiated nanoceramic are γ-Al 2 O 3 up to 40 dpa, and both γ-Al 2 O 3 and α-Al 2 O 3 at 150 dpa (see Supplementary Information). It is worth highlighting that the irradiation did not induce any loss of adhesion or delamination effects at the thin film-substrate interface. The combined effect of irradiation and strain imposed by the substrate (for instance, due to thermal and irradiation creep or swelling) is beyond the scope of this study, and warrants further investigation.

Figure 2 ADF-STEM micrographs and DPs showing as-deposited (a) and irradiated Al 2 O 3 thin films after 20 dpa (b), 40 dpa (c) and 150 dpa (d) at 600 °C. The coarsening induced by irradiation releases excess free energy due to the interaction between point defects and GBs26. Full size image

It is likely that temperature plays an important role in determining the kinetics of the structural evolution. However, we attribute crystallization and grain growth to the sole effect of irradiation (see Supplementary Information). The initial crystallization is expected to occur readily upon irradiation, and may be homogeneous35, epitaxial36,37, or both. The subsequent coarsening effect can be explained in terms of a fast disorder-driven mechanism, which is available even below room temperature25,26, and which is governed by the capture of interstitials by GBs22,23,24. The incident ions introduce a large amount of local disorder through atomic displacement cascades. The disordered regions interact with GBs, releasing excess free energy and leading to an overall growth. It is also interesting to notice that the extent of grain growth is strongly influenced by the total amount of energy injected by the ions into the material27. In the energy range investigated, the energy of the ions is transferred to the material both by electronic excitations and displacive damage (i.e. nuclear collisions). The effect of these different kinds of energy loss may be additive, synergistic or even competing. In the case of oxide nanoceramics, the effect is generally additive27. The plot in Fig. 3a shows the dependence of grain growth both on the total amount of energy injected into the material (keV per target atom), and on displacive radiation damage (displacements per atom). The graph indicates that radiation-induced grain growth is a self-limiting process, which follows a sublinear dependence on damage exposure, in good agreement with previous results concerning other nanocrystalline oxides25,26,27,28.

Figure 3 Grain growth in the Al 2 O 3 thin films as a function of total energy injection and displacive radiation damage (a). The grain coarsening is accompanied by the formation of twin boundaries (b), which release accumulated mechanical energy. The presence of a mirror plane in both the HR-TEM micrograph (c) (indicated by arrows), and in the DP inset confirms the twin relationship of the adjacent grains. Full size image

In the irradiated material, grain growth is accompanied by the formation of planar defects with two parallel flat boundaries. These defects are found occasionally, and their presence is independent of damage exposure. The defects are identified as twins, and an example is shown in Fig. 3b. The presence of a mirror-plane both in the high-resolution TEM (HR-TEM) micrograph and in the DP inset in Fig. 3c confirms that the defects observed are indeed twins. The formation of twins in nanocrystalline solids can be understood in terms of mechanisms such as nanoscale multiplane shear38 or stacking fault formation led by Shockley partial dislocations39. From an energy balance perspective, the formation of twins may be explained by the need to release the excess mechanical energy accumulated during incoherent grain coarsening.

The structural rearrangements induced by the irradiations (i.e., crystallization and grain growth) bring about changes in the mechanical properties of the material. These changes are plotted in Fig. 4 as a function of the average grain size. The reduced Young’s modulus E r increases monotonically with grain size (i.e. E r,20dpa = 205 ± 7 GPa, E r,40dpa = 222 ± 10 GPa, and E r,150dpa = 245 ± 19 GPa). Accordingly, the Young’s modulus E increases from E 20dpa = 235 ± 10 GPa, to E 40dpa = 262 ± 15 GPa and E 150dpa = 301 ± 31 GPa (Fig. 4a). The hardness H (Fig. 4b) peaks at moderate damage exposures, varying from H 20dpa = 17.8 ± 0.9 GPa, to H 40dpa = 17.2 ± 1.2 GPa, and H 150dpa = 15.9 ± 1.6 GPa. Notably, the trend is well described by the Hall-Petch effect, whereby a material’s strength and hardness decrease as the average grain size increases. The Hall-Petch relationship describes the measured hardness H v according to the formula H v = H 0 + kD−1/2, where H 0 is the intrinsic hardness dependent on frictional lattice resistance to dislocation motion, k is the material-specific strengthening coefficient, and D is the average grain size. In this work, the best linear fit of H v as a function of D−1/2 yields H 0 = 13.255 GPa and k = 46.638 GPa.nm1/2, with a coefficient of determination equal to R2 = 0.9756. Below the so-called strongest grain size (typically in the range 10–20 nm40), the strengthening effect is balanced by GB shear, which yields a reduction of hardness for decreasing grain size. This effect is usually referred to as the inverse Hall-Petch effect. A detailed overview of the mechanisms that yield an enhancement of hardness in nanoceramics has been recently reported by Veprek40.

Figure 4 Effect of radiation-induced grain growth on the mechanical properties of Al 2 O 3 nanoceramic thin films, namely the Young’s modulus E (a), the hardness H (b) and the hardness to Young’s modulus ratio H/E (c). The trend of hardness is well-described by the Hall-Petch effect, due to the increase of grain size with increasing damage exposures. Full size image

A direct comparison between the mechanical properties of the irradiated thin films and polycrystalline α-Al 2 O 3 is difficult. The mechanical properties of the latter vary depending on the grain size, the presence of impurities and on the processing route41. However, comparisons can be made with bulk nanocrystalline α-Al 2 O 3 (bnc-alumina). The reported hardness and stiffness for bnc-alumina with a grain size of 150 nm are H bnc-alumina = 25.5 ± 0.3 GPa and E bnc-alumina = 403 GPa42,43. Here, the maximum hardness is reached when the average grain size is 101 nm (H 20dpa = 17.8 ± 0.9 GPa). The corresponding Young’s modulus is E 20dpa = 235 ± 10 GPa. The differences observed are probably due to: (i) the presence of different phases (γ-Al 2 O 3 versus α-Al 2 O 3 ), (ii) the measurement method (Berkovich nanoindentation versus Vickers microindentation), (iii) the presence of radiation-induced point defects, or (iv) combinations thereof.

Another important implication of the observed irradiation-induced crystallization is that the H/E ratio of the thin films is enhanced in response to irradiation. The H/E ratio peaks at moderate damage exposures (when the volume fraction of GBs is the highest), varying from 0.051 for the as-deposited condition, to 0.076, 0.066 and 0.053 for 20, 40 and 150 dpa, respectively (Fig. 4c). These results suggest an improvement in service of the robustness of the thin films against wear. This is of particular interest concerning liquid metal erosion44 or rod-to-grid fretting45.

Fracture toughness of ceramic materials is typically determined from nanoindentation tests by measuring the length of surface radial cracks emanating from the corner of imprints. This type of measurement is not possible here because cracks are not observed in any case, due to the low load and the low film thickness. However, an indirect estimation of fracture toughness is given by the H/E ratio34. The as-deposited material lacks long-range order and nanostructural defects (such as dislocations) that may shield stress and suppress crack openings. Thus, the attainable plasticity in the wake of a crack tip is limited, and any opening would be likely accommodated by unstable crack propagation. Accordingly, the H/E ratio of the pristine material is comparatively low. In contrast, the mechanical response of the irradiated material is mainly driven by GBs. The large volume fraction of GBs makes new energy dissipation mechanisms available (e.g., twinning). The resulting H/E ratio is higher, which suggests an enhancement of fracture toughness. Although thin film coatings are not structural components, an enhancement in service of fracture toughness is desirable because future design rules might rely entirely on the presence of a coating for the correct operation of a reactor. In this perspective, a certain extent of cracking may be acceptable during the extended exposure to neutron radiation fields, while unstable crack propagation would certainly not be an option.

Additional qualitative evidence in support of the enhancement of fracture toughness is provided experimentally by nanoimpact tests. In these tests, a cube-corner diamond tip is periodically blasted against the surface of the thin films. The impact depth is the highest for the as-deposited thin films (see Supplementary Information), which suggests that the impact energy is dissipated more efficiently in the irradiated samples. As a matter of fact, the impact response of the as-deposited and the irradiated thin films is radically different. Figure 5 displays the cross-sectional images of representative nanoimpact imprints for as-deposited and irradiated samples. Two selected area (SA) DPs are acquired for each cross-section, both distant from (white box) and within (yellow box) the impact zone (below and in the vicinity of the impact imprint). In the unirradiated samples (Fig. 5a), impact energy is dissipated through shear banding, and no major structural rearrangements are induced by the impact loading. This observation is confirmed by the fact that the SADPs gathered distant from and within the impact imprint appear identical, as shown in Fig. 5d. Figure 5b,c show the cross-section of nanoimpact imprints in samples exposed to 20 dpa (corresponding to the peaks of H and H/E ratio in Fig. 4b,c) and 150 dpa (end-of-life exposure), respectively.

Figure 5 Cross-sectional TEM micrographs of representative nanoimpact imprints on the Al 2 O 3 nanoceramic thin films before (a) and after irradiation up to 20 dpa (b) and 150 dpa (c). No major structural rearrangements are induced by impact loading in the unirradiated samples, as confirmed by the identical SADPs gathered distant from and below the impact imprint (d). The appearance of arcs and rings in the SADPs beneath the impact zones in the irradiated samples is due to energy dissipation through bending of the lattice planes. Another energy dissipation mode present is localized amorphization, which is indicated by arrows in (b,c), and shown in high-resolution in (e,f). The FFT insets in (e,f) confirm that the bright contrast corresponds to the amorphous phase, and that the dark contrast corresponds to the crystalline phase. Full size image

The appearance of arcs and rings in the SADPs beneath the impact zones is due to the bending of lattice planes, which denotes plastic strain as one of the main energy dissipation mechanisms. Another energy dissipation mechanism present is localized amorphization. Notably, crystalline-to-amorphous phase transitions are often described as toughening mechanisms9. The HR-TEM micrographs in Fig. 5e,f show the localized amorphization indicated by arrows in Fig. 5b,c, respectively. The FFT insets show a diffused halo where the contrast is bright, and diffraction spots where the contrast is dark, confirming the amorphous nature of the bright contrast band, and the long-range order of the contiguous zone. The close alignment of the crystal lattice on either side of the amorphous band rules out the formation and subsequent rebonding of two cracked surfaces. Localized amorphization has been observed in sapphire46 and in other unirradiated ceramics exposed to shock loading, such as B 4 C47,48, SiO 2 49, Y 2 Si 2 O 7 50 or B 6 O51. The onset of the phenomenon has been explained by shock-induced plastic waves46 and shear instability48,49. These phenomena are relevant at the extremely high stresses and strain rates induced by shock loading. Further causes include the coalescence of dislocation loops under high shear stresses, as occurs upon quasi-static mechanical load50,51, and adiabatic shear, which is governed by elastic strain energy in brittle solids52. Arguably, both the high strain rate induced by impact loading, and the coalescence of defects and defect clusters formed during irradiation may play an important role in the amorphization process observed here. However, the impact speed in this study (≈500 μm/s) is several orders of magnitude lower than the impact speed in shock loading experiments (≈18 km/s)46,47,48,49. This fact suggests that shock-induced plastic waves and shear instability are unlikely as the main driving forces.

It is worth noting that the utility of oxide nanoceramics as radiation tolerant materials is often thought to be limited by grain growth. Indeed, GBs are usually considered as the actual source of radiation tolerance due to their efficient behavior as defect sinks. The problem is that the density of GBs decreases inexorably as grain growth proceeds. Thus, the radiation tolerance of oxide nanoceramics is expected to fade away for increasing radiation damage exposures. However, radiation tolerance can be defined in many ways. For example, radiation tolerance can also be conceived in terms of the expected lifecycle of a given component. From this point of view, nanoceramic thin films can be utilized as radiation tolerant coatings indeed. As a matter of fact, the thin films in this work are able to withstand radiation damage up to 150 dpa without suffering catastrophic failure nor delamination. It is also worth noting here that coarse-grained polycrystalline α-Al 2 O 3 suffers void swelling and releases the resulting stresses through cracking at much lower damage exposures53,54.