Atomistic Level Study of Ce3Si2 Oxidation as an Accident Tolerant Nuclear Fuel Surrogate

The oxidation mechanisms of Ce 3 Si 2 (as a surrogate for U 3 Si 2 accident tolerant nuclear fuel) have been studied to 750°C in ﬂ owing air. Discrepancies between the terminal and theoretical mass gains for complete oxidation, similar to previous works on the U 3 Si 2 system are attributed to incomplete oxidation of Si, which forms amorphous SiO 2 and regions of nano-crystalline free Si. Nano-grained (<5nm) CeO 2 also formed after oxidation leading to pulverisation, the resulting composition and nano-structure may have highly deleterious consequences for the fuel material under oxidative atompsheres, increasing the rate of oxidation, spallation and ﬁ ssion product release.


Introduction
Since the Fukushima accident in 2011 there has been a large international effort in the development of accident tolerant fuel (ATF) and cladding materials for water cooled reactors [1].Uranium silicide intermetallics are being considered as ATF candidates to replace UO 2 currently used in light water reactors (LWRs) [2].Primarily U 3 Si 2 is being examined, due to its good thermal conductivity (∼15 W/m/K at 500 °C) [3], higher U metal density, (giving economic advantages by requiring lower 235 U enrichments) and reduced radiation swelling compared to the higher U containing silicides such as U 3 Si [4,5].As its thermal conductivity is much higher than UO 2 (∼4 W/m/K at 500 °C [6]) it will reduce the centreline temperature of the fuel and reduce hotspot build up, limiting detrimental thermally active processes, such as; fission product migration, differential swelling and cracking across the fuel pellet, and by reducing the operating temperature of the fuel increase the safety margin in an accident scenario.
However, little is known on the oxidation behaviour of U 3 Si 2 under air and steam environments which could be experienced under a loss of coolant accident (LOCA) scenario with water and/or air ingress into the fuel pin.Sooby Wood et al. [7] performed some of the initial works on the behaviour of U 3 Si 2 up to temperatures of 1100 °C.The authors [7] found an onset temperature of ∼384 °C and that the terminal mass gain of U 3 Si 2 was 21 wt.% after oxidation, compared to a theoretical 25 wt.% gain assuming a complete reaction to U 3 O 8 and SiO 2 , indicating an incomplete reaction.The authors state the oxidation to UO 2 is more energetically favourable than Si and so the formation of a passivating SiO 2 layer is prohibited, leading to complete oxidation of U. The authors report no evidence of SiO 2 formation from x-ray diffraction (XRD) and scanning electron microscopy (SEM), assuming only Si is present.However, the oxidation to U 3 O 8 and Si would have a theoretical weight gain of ∼17 wt.% and so suggest that some Si must oxidise to SiO 2 to achieve the terminal weight gain of 21 wt.%.In further work, Sooby Wood et al. [8] examined the effect of 1.8 at.%Al additions to the oxidation of U 3 Si 2 , noting again a lower than expected terminal weight gain compared to the theoretical calculations by ∼4 or 5 wt.%.Some recent works have also focused on the corrosion of U 3 Si 2 [9,10] under steam where pulverisation of pellets has been observed.Johnson et al. [11] have oxidised arc melted U 3 Si 2 fragments to 800 °C under flowing air, finding an onset of 470 °C, much higher than that reported by Sooby Wood et al. [7].The authors [11] also reported a mass gain of 20 wt.% (similar to that reported by Sooby Wood) which they attribute to the formation of U 3 O 8 and SiO.However, no experimental evidence of SiO is given in Ref. [11] and as it is a metastable solid that rapidly dissociates into Si and SiO 2 [12] this will be highly unlikely to form.Yan et al. [13] studied the room temperature oxidation of U 3 Si 2 using in-situ x-ray photoelectron spectroscopy (XPS) under oxygen atmospheres of between 2.5-5 × 10 -8 mbar.The authors [13] report the formation of a UO 2 followed by uranium silicates, although which type is not identified, based on the shift of the Si 2 s peak being in-between that of free Si https://doi.org/It is clear from the disparity in terminal mass gains compared with theoretical calculations that the mechanism(s) of oxidation in U 3 Si 2 , the resulting microstructure and compositions of the as-oxidised materials are still poorly understood, which will be required before these materials can be used in-core.Oxidation of U 3 Si 2 has also been observed under in-situ Xe ion irradiation at 300 °C with the formation of nanocrystallite UO 2 grains, and a Si rich amorphous phase, however, no details of the fate of Si were given [14].In similar systems, such as UC [15,16], as well as other carbide and nitride ceramics [17,18], oxidation occurs via the formation of nano-crystallite oxide grains, with release of C or N as CO 2 or N 2 , respectively.However, in Si containing material it is important to understand the fate of the Si during oxidation as to predict any effects this may have on the resulting material properties.To examine this, we have performed an atomistic level study of the oxidation of Ce 3 Si 2 as a surrogate for U 3 Si 2 using high resolution scanning and conventional transmission electron microscopy (S/TEM).

Experimental
Ce 3 Si 2 samples were fabricated from Ce ingot (Alfa Aesar, 99 wt.%, main impurity 0.15 wt.% La) and Si pieces (Alfa Aesar, 99.9999 %) which were arc melted (Centorr 5SA arc furnace) in an Ar atmosphere glovebox (Kiyon) with < 5 atomic parts per million (appm) O 2 .Phase analysis was performed on an x-ray diffractometer (Malvern Panalytical Empyrean) operated at 40 kV and 40 mA, using CuK α radiation and a zero-background single crystal Si sample holder.In-situ phase analysis during oxidation was measured using high temperature (HT)-XRD (Bruker D8), performed at 40 kV and 30 mA, using CuK α radiation where samples were heated to 750 °C at a ramp rate of 0.1 °C/min under flowing synthetic air (approximately 100 mL/min) using a Al 2 O 3 (corundum) sample holder.In-situ scans were acquired over 75 min each, resulting in a 7.5 °C temperature increase throughout the duration of the scan.XRD patterns were indexed using powder diffraction files (PDF) from the International Centre for Diffraction Data (ICDD).Crystallite size of the as-oxidised materials was determined from the full width half maximum (FWHM) and HighScore Plus (version 4.8) based on the Scherrer equation [19] assuming all peak broadening arises from crystallite size and there are negligible strain effects.Crystallite size is reported as the average determined from three CeO 2 reflections and the error is the standard deviation.Intrinsic instrument line broadening was accounted for using a 650c Si standard from the National Institute of Standards and Technology (NIST).
A fragment of as-melted Ce 3 Si 2 was oxidised under flowing synthetic air (50 mL/min) using thermogravimetric analysis coupled with differential scanning calorimetry (TGA/DSC) (STA-449-F1; Netzsch) up to 750 °C at a ramp rate of 5 °C/min.A correction run (identical temperature profiles) was performed prior to the sample oxidation run, with empty sample and reference crucibles to account for intrinsic errors in the instrument (such as buoyancy effects from gas flows).In addition, the crucibles were baked before the correction and sample runs in the TGA/DSC furnace under flowing synthetic air at 800 °C for 1 h, as to remove any adsorbed moisture or gases.Off-gassing from the sample was not observed during the TGA/DSC runs, likely due to the preparation of the material by solid state synthesis in an Ar atmosphere glovebox prior to oxidation.The error in the balance stability has been measured to be ± 0.01 mg over a two-hour period, the initial starting mass of the Ce 3 Si 2 fragment was 53.8 mg and so this corresponds to an error of ± 0.02 % in the reports of relative mass gains during oxidation.
As-melted and as-oxidised samples were prepared for S/TEM characterisation by standard crushing in a pestle and mortar followed by dispersion in distilled water before drop casting onto carbon film grids.S/TEM and energy dispersive x-ray spectrometry (EDS) were performed on a FEI Talos X-FEG (operated at 200 kV) equipped with Super X EDS with 4 windowless silicon drift detectors (SDDs).Energy filtered (EF)TEM was performed on a Hitachi H-9500 (operated at 300 kV) equipped with electron energy loss spectroscopy (EELS), and a Gatan Imaging Filter (Quantum SE model 963).Selected area diffraction patterns (SADP) and fast Fourier transforms (FFTs) of high resolution (HR)TEM images were indexed by matching d hkl values with reference patterns from the Inorganic Crystal Structure Database (ICSD) and calculated SADPs using SingleCrystal software (version 2.2.9 CrystalMaker Software Limited).

Oxidative mass gain
Fig. 1 shows the mass gain of Ce 3 Si 2 under flowing air up to 750 °C, showing the onset temperature of oxidation is around 400 °C and breakaway reaction occurs at around 500 °C.This onset is higher than that reported for U 3 Si 2 in Ref. [7].This is likely due to the lower Gibbs free energy of formation for UO 2 as compared to CeO 2, shown in the excerpt of the Ellingham diagram reproduced from Refs.[8,20,21] in Fig. 2. The mass gain reaction follows a single rate step and so likely proceeds by the direct formation of CeO 2 with this being the highest oxidation state available for Ce, unlike in U 3 Si 2 where multiple reaction  steps may be observed due to the formation of higher oxides from UO 2 such as, U 3 O 7 , U 4 O 9 and U 3 O 8 .Fig. 1 also shows the terminal mass gain of the oxidation in this work is 30.0 wt.% which is inconsistent with the theoretical mass gains for the possible reactions, assuming reaction products of CeO 2 and Si, SiO or SiO 2 .This discrepancy with the terminal mass being lower than the theoretical mass for the full oxidation to CeO 2 and SiO 2 is similar to previous works on the uranic system in [7,8] indicating similar mechanistic effects.

XRD analysis
From the XRD pattern in Fig. 3a it can be seen the starting powder before oxidation was mainly tetragonal Ce 3 Si 2 with space group P mbm 4/ (PDF 04-004-6138 [22]), analogous with the U 3 Si 2 crystal structure) with a minor phase of tetragonal Ce 5 Si 4 (space group P41212 PDF 01-078-5346 [23]).Although Ce 5 Si 4 is present as a minor phase and would alter the terminal mass gain from oxidation as compared to the theoretical calculations assuming pure Ce 3 Si 2 , its presence would give a higher than expected mass gain and so cannot account for the lower than expected terminal mass gain observed in TGA data.After oxidation, the XRD pattern (Fig. 3b) indexed as face centred cubic (FCC) CeO 2 with space group Fm m 3 ¯(PDF 96-900-9009 [24]), with no evidence of crystalline Si or SiO 2 .However, FCC Si has the diamond structure (space group Fd m 3 ¯) and the lattice spacings of FCC Si and CeO 2 are very similar (a 0 for CeO 2 = 0.5411 nm [24]and 0.5431 nm for Si [25]) and so it is difficult to conclude the presence of crystalline Si from the XRD pattern.
Fig. 4 shows XRD patterns taken during in-situ heating experiments and shows the starting material is a major Ce 3 Si 2 phase with minor Ce 5 Si 4 impurities and additional peaks from the Al 2 O 3 sample holder.It can be seen that the intensity of the Ce 3 Si 2 peaks begin to decrease from 350 °C and these are not visible from 450 °C.Between 450−550 °C, there are no discernible peaks except from a very diffuse peak at around 26°2θ corresponding to (111) CeO 2 reflection, showing there is no formation of crystalline Ce 2 O 3 during the oxidation process and the Ce proceeds directly to CeO 2 .Beyond a temperature of 550 °C the CeO 2 reflections become more pronounced and sharper, indicating an increase in their crystallinity and crystallite size.Due to the similarity in the lattice parameters and crystal structure of CeO 2 and Si it is again difficult to conclude the presence of crystalline Si from the XRD patterns and no evidence for crystalline SiO 2 formation was observed.Significant peak broadening was observed after the reaction completed at 750 °C (similar to material oxidised in TGA/DSC experiments shown in Fig. 3b).Fig. 5a shows an excerpt of the in-situ XRD patterns of the (220) reflection for CeO 2 , showing the grains crystallise around 600 °C and increase in crystallinity and size as the temperature is increased.Fig. 5b plots the crystallite size as a function of temperature calculated from the FWHMs using the Scherrer equation [19] and it can be seen these are around 1.8 nm when initially observed at 600 °C and grow to around 4.5 nm when the oxidation reaction completes at 750 °C.

Fate of Ce during oxidation
This transformation into nano-crystalline materials is further confirmed by TEM analysis in Fig. 6.Fig. 6a and b show bright field (BF) and HRTEM images of the starting Ce 3 Si 2 which was confirmed to be larger (> 100 nm) single crystalline grains as confirmed by the SADP in Fig. 6c.However, after oxidation, it can be seen in Fig. 6d and e the formation of ∼5 nm size crystallites has occurred, which indexed as CeO 2 from the FFT in Fig. 6f, agreeing well with the XRD results and crystallite size measurements.
The formation of nano-crystalline oxide grains has been observed previously in the oxidation of U 3 Si 2 under ion irradiation at 300 °C [14] and air oxidation of UC [15,16], thus, the observation of < 5 nm size CeO 2 crystallites in this work accurately emulates the behaviour in the uranic systems.Grain growth in nano-crystalline CeO 2 begins to occur at temperatures > 500 °C [26].Typically, grain coarsening occurs by an Ostwald ripening mechanism where atomic diffusion is driven by the reduction of crystallite surface area which are regions of high energy [27].However, Ivanov et al. [26] studied nano-particle coarsening in CeO 2 by XRD, TEM and small-angle neutron scattering and report the grain growth occurs by rotation of neighbouring crystals before intergrowth rather than by an Ostwald ripening type mechanism.Due to the short annealing time in the present work above 500 °C, this will restrict grain growth via the alignment and intergrowth of individual crystallites.
Nano-crystalline UO 2 is also observed to form in irradiated nuclear fuel and is referred to as the high burn up structure (HBS).After prolonged irradiation, micron sized grains in the outer part of the fuel pellet (which is cooler than the centreline temperature) undergo grain subdivision resulting in a nano-grained structure alongside the formation of 1 μm size pores [28].The mechanistic understanding of grain subdivision in UO 2 is still fairly poor but is likely based on polygonization (rather than recrystallization) whereby dislocations rearrange into rows, creating low angle grain boundaries and reducing the stored energy of the microstructure [29,30].Whilst the formation of a nano-structure in UO 2 fuels under normal operating and high burn up conditions of LWRs was predicted to have adverse effects on fuel stability and performance, such as a decreased thermal conductivity and increased brittleness.However, recent results have shown that the increased concentration of μm size closed porosity improves fission gas retention [31] and the higher grain boundary surface area may also provide an increased sink concentration for annihilation of radiation induced defects [31].
However, the formation of this nano-structure in oxidised Ce 3 Si 2 accompanied by the volumetric expansion during oxidation inevitably leads to the pulverisation of the Ce 3 Si 2 fragments observed in the current and previous works on densified U 3 Si 2 compacts under air and steam oxidation experiments [9,10].Furthermore, due to the higher oxidation states available in U, the final oxidation product will be U 3 O 8 , leading to higher degrees of volumetric expansion and spallation.This will likely have highly deleterious effects on the integrity of the fuel material with the enormous increase in grain boundary area giving rise to; increased oxygen diffusion beyond the oxide layer increasing oxidation rates, rapid diffusion of fission products (FPs) leading to potential release of gaseous species and higher rates of spallation and exfoliation of fuel material.

Fate of Si during oxidation
Fig. 7a and b show BF-STEM and high angle annular dark field (HAADF) images of a region from the as oxidised powder showing a region of low Z material with nano-particles with higher atomic number.The STEM-EDS maps in Fig. 7 confirm the nano-particles are Ce and O contained within a Si and O rich matrix, likely to be SiO 2 .A HR-TEM image in Fig. 7c shows a higher magnification of one of the nano-particles showing it is crystalline, within an amorphous phase.The FFT shown in Fig. 7d of the Si-O rich region confirms it is amorphous and is likely SiO 2 and the FFT of the crystalline nano-particle in Fig. 7e indexed with CeO 2 , which confirmed the EDS results.Fig. 8a also shows a HAADF-STEM image of another SiO 2 with nanocrystalline CeO 2 grains and this is confirmed to be amorphous in the SADP shown in Fig. 8b, where the diffraction spots arise from the {220} reflections from the nano-crystalline CeO 2 within the Si-O matrix particle.Thus, this work has also shown the presence of amorphous SiO 2 confirming Si does undergo oxidation, an observation that was not observed in previous works using XRD and SEM analysis [7].
Although oxidation of the U-Si system is poorly understood, much work has focused on the oxidation behaviour of accident tolerant cladding materials such as SiC for LWR [32], advanced fuel particle coating applications (such as TRISO [33]) and MAX phase materials (M = transition metal, A = A group element and X = C or N) containing Si [34].Initial studies by Jorgensen et al. [35] examined the oxidation behaviour of SiC using dry air at temperatures between  [27]) formed after an incubation time associated with an increase in the oxidation kinetics.This agreed well with work by Berton et al. [35] who observed amorphous SiO 2 to form in SiC oxidised at 1200 °C and cristobalite at temperatures above 1300 °C noting again an increase in the oxidation kinetics due to stresses between the crystalline and amorphous SiO 2 leading to crack formation and oxygen ingress.Meschter [36] has also studied the oxidation of MoSi 2 under flowing air at temperatures of 400−600 °C, finding the Mo and Si oxidise simultaneously forming monoclinic Mo 9 O 26 and amorphous SiO 2 .In other works on the oxidation behaviour of the MoSi 2 system transformation from an amorphous SiO 2 phase to cristobalite is reported to occur between 1000−1400 °C [37].Thus, the observation of amorphous SiO 2 in  this work is likely due to the lower temperatures (750 °C) of oxidation being below that required for crystallisation of SiO 2 into cristobalite.
Fig. 9a and b show BF and HAADF-STEM images, respectively showing islands of low Z number material.Coupled with STEM-EDS maps of Ce, Si and O in Fig. 9, it can be seen these are ∼10 nm Si rich regions that are deficient in oxygen and likely due to the presence of unreacted Si.Fig. 9c also shows a HRTEM image of these Si rich regions with the FFT of the Si rich region appearing amorphous (Fig. 9d), however, this may also be due to the nano-crystal being off the crystallographic zone axis.It can be seen the surrounding crystalline region indexed as CeO 2 (Fig. 9e), agreeing well with EDS mapping.The disparity in terminal and theoretical mass gain to CeO 2 and SiO 2 can thus  be attributed to an incomplete reaction with the oxidation of Si.Sooby Wood et al. [7,8] attributed the disparity in mass gain to either the incomplete oxidation of Si and free Si remaining in the material or incomplete oxidation of U to U 4 O 9 from XRD analysis where the authors [7,8] could not discern between the two.However, from this work it is evident that un-oxidised Si remains in the material after reaching its apparent plateau in mass gain during oxidation.
Fig. 10a shows a BF-TEM image of the as-oxidised powder showing a darker region of nano-crystallites, likely to be CeO 2 and lighter regions likely to be Si containing.Fig. 10b shows a HR-TEM image of the lighter region showing it is crystalline and Fig. 10b and c show a higher magnification and filtered image of the lattice and FFT, respectively, showing it indexed with a FCC [110] zone axis which could possibly be Si or CeO 2 .Fig. 10e-g show EFTEM maps of Ce, O and Si, respectively, confirming the formation of much larger (> 100 nm) Si regions which again were deficient in O and indicate the crystalline region in Fig. 10b is unreacted Si.It should be noted that FCC Si, and CeO 2 are indistinguishable by XRD or SADP.However, coupled with EFTEM analysis showing Si rich and Ce and O deficient regions, and the high metastability reported for SiO, this strongly indicates the presence of crystalline Si and the HR-TEM image and FFT indexed as close to the FCC Si [110] zone axis, with forbidden {002} reflections observed.These reflections are kinematically forbidden due to the structure factor of diamond FCC crystal, however, these are often observed in TEM due to dynamical scattering events [38], whereby scattered beams from, for example, the (1 ¯11) plane may be re-diffracted by the (11 ¯1) plane giving rise to an apparent excitation in the (002) reflection ( + = 1 ¯11 11 ¯1 002).The crystallisation temperature of Si has been reported to be in the region of 550−700 °C [39] and so it would be expected that the Si observed in this work would be crystalline due to the oxidation temperature reaching a maximum of 750 °C.From Fig. 2, it can be seen that the Gibbs free energy of formation of CeO 2 is much lower than that of SiO 2 meaning that the Ce will begin to oxidise initially and preferentially over the Si.The regions of unreacted Si in Fig. 9 and Fig. 10 indicate that this leads to a separation of Ce and Si into discrete phases and in some cases (Fig. 9) that the CeO 2 may then provide a passivation layer inhibiting full oxidation of the free Si.From the pseudo-binary phase diagram [40], Si and Ce oxides are insoluble with each other (similar to the U-Si-O system [41]) with the exception of a narrow region in the Ce-Si-O system between 18-25 at.%Ce where several CeSiO silicates form.Steam corrosion of Ce 3 Si 2 as a surrogate for U 3 Si 2 has been performed under 9 MPa at 300 °C for 48 h by Urso et al. [42] to replicate LWR operating conditions.The authors [42] noted formation of a ∼10 μm thick Ce 4.67 (SiO 4 ) 3 O layer on the sample surface as characterised by XRD.However, no observation of Ce-Si-O phases were observed in the current work, likely due to the higher O 2 partial pressures and higher oxidation temperature leading to a more complete reaction and to the segregation of both the Si and Ce oxide phases.In the case of this work, Si will likely undergo oxidation to SiO 2 once separated from the Ce, thus providing no protection through formation of a passivation layer, typically desired for accident tolerant fuel and cladding materials [11].This may be more representative of the UO 2 -SiO 2 system [41] where there is no uranium silicate formation anticipated across the entire compositional range from the binary UO 2 -SiO 2 phase diagram.However, the phase diagram [41] is only determined for temperatures above 1200 °C and at lower temperatures a U-Si-O phase may form as reported by Yan et al. [13] during room temperature oxidation of U 3 Si 2 .

Proposed oxidation mechanism
This work has shown during oxidation of Ce 3 Si 2 phase separation of the CeO 2 and Si occurs due to their low mutual solubility.Ce will preferentially oxidise over Si due to the more negative Gibbs free energy of formation of CeO 2 (−950 kJ mol -1 O 2 ) as compared with SiO 2 (−750 kJ mol -1 O 2 ) at the breakaway temperature of 500 °C.The reaction products are CeO 2 , as confirmed by XRD and TEM, and SiO 2 as well as unreacted free Si determined by STEM-EDS and EFTEM analysis and a proposed mechanism for the oxidation is given in Fig. 11.
The associated mass gain during oxidation confirms a partial reaction of the Si to SiO 2, indicated by Si rich regions remaining in the material after the oxidation reaction plateau is observed in the TGA data.The Si rich regions are generally accompanied by a surrounding CeO 2 grain, which may form during the initial segregation of Si from CeO 2 and provides a passivating layer and inhibits the oxidation of Si.However, the Si provides no evidence of forming a passivating layer of SiO 2 to protect the Ce 3 Si 2 surface from further oxidation.CeO 2 grains that form are small (< 5 nm) nano-crystallites that have also been observed to be dispersed in the amorphous SiO 2 particles which likely occurs during the phase segregation as the Ce and Si oxidise.Accompanied with the volumetric expansion of the material during oxidation this leads to pulverisation of the fragments and the production of powder, revealing virgin material which undergoes further oxidation.
Table 1 shows the calculated wt.% fractions from a theoretical full oxidation to CeO 2 and SiO 2 versus the mass gain from experimental TGA data in this work.In the theoretical mass gain calculations in Table 1, all Ce (0.6 mol fraction of starting material) is oxidised to CeO 2 and all Si (0.4 mol fraction of starting material) is oxidised to SiO 2 .To calculate the discrepancy in this work, we assume all Ce (0.6 mol fraction) is oxidised to CeO 2 and the remaining mass gain is due to SiO 2 formation and the mole fraction of SiO 2 in the resultant powder can be calculated (0.343 mol fraction).The discrepancy between the mole fraction of SiO 2 in the theoretical full oxidation (0.4 mol fraction) and in the observed TGA experiment (0.343 mol fraction) can then be used to calculate the amount of unreacted Si (0.057 mol fraction).No observation of SiO has been observed and this will be unlikely due to the metastable nature of SiO which will readily decompose into Si and SiO 2 .From Table 1 it can be seen that the amount of free Si is ∼1.3 wt.% in the material and so approaching the detectability limits of XRD, which may account for its lack of observation in previous works in the U 3 Si 2 system [8,9] if its oxidation processes are similar to Ce 3 Si 2 .

Conclusions
The oxidation mechanism of Ce 3 Si 2 as a surrogate for U 3 Si 2 under flowing synthetic air to 750 °C has been examined.Oxidation mass gains and products were characterised using TGA, XRD and S/TEM with EDS and EFTEM chemical mapping.Nano-structured CeO 2 grains form (< 5 nm) which may be potentially deleterious for fuel material due to the high surface area of grain boundaries resulting in increased oxidation and corrosion rates, fission product release and exfoliation of fuel material.The formation of nano-grained material coupled with the large volumetric expansion of the oxidised material leads to pulverisation.Oxidation of Ce occurs preferentially over Si, with unreacted Si present in the material accounting for discrepancies in theoretical against as observed mass gains in this and previous works on U 3 Si 2 oxidation.The majority of Si is converted to SiO 2 , however, this provides no passivation layer to the Ce and due to the mutual insolubility of SiO 2 and CeO 2 these form nano-particles of CeO 2 dispersed within amorphous SiO 2 particles.S/TEM analysis has confirmed the presence of unreacted, crystalline Si accounting for mass gain discrepancies

Fig. 1 .
Fig. 1.Plot of mass gain as function of temperature for Ce 3 Si 2 oxidised under flowing air to 750 °C.

Fig. 3 .
Fig. 3. XRD patterns of a) starting material showing major Ce 3 Si 2 phase with minor Ce 5 Si 4 phase and b) as-oxidised material, showing full conversion to CeO 2 .

Fig. 4 .
Fig. 4. In-situ XRD patterns of Ce 3 Si 2 samples heated under flowing air up to 750 °C.

Fig. 5 .
Fig. 5. a) excerpt of in-situ XRD pattern of the CeO 2 (220) reflection as a function of temperature and b) plot of crystallite size as a function of temperature calculated from the peak broadening.

Fig. 6 .
Fig. 6. a) BF-TEM image of starting Ce 3 Si 2 grain, b) high resolution (HR)TEM of area in a, c) SADP showing single crystalline starting Ce 3 Si 2 material, d) BF-TEM image of as-oxidised powder showing nano-grains, e) HRTEM of oxidised powder showing nanocrystals ∼5 nm in size and f) FFT of image in e, indexed as CeO 2.

Fig. 7 .
Fig. 7. a) BF-STEM and b) HAADF-STEM image of as-oxidised Ce 3 Si 2 material, with Ce, O and Si elemental wt.% STEM-EDS maps c) HR-TEM image of CeO 2 particle in SiO 2 rich region, with FFT showing d) amorphous SiO 2 region e) crystalline CeO 2 .

Fig. 8 .
Fig. 8. a) HAADF-STEM image of SiO 2 particle with CeO 2 nano-crystallite dispersed and b) SADP showing particle is amorphous, with exception of crystalline regions of CeO 2 particle embedded and corresponding Ce, O and Si elemental wt.% STEM-EDS maps.

Fig. 9 .
Fig. 9. a) BF-STEM and b) HAADF-STEM image of as-oxidised Ce 3 Si 2 material with Ce, O and Si elemental wt.% STEM-EDS maps c) HR-TEM image of Si rich region, with FFT showing d) Si rich region e) crystalline CeO 2 .

Fig. 10 .
Fig. 10.a) BF-TEM image of as oxidised Ce 3 Si 2 material, b) HR-TEM image of Si rich region, c) FFT filtered image of HRTEM image showing lattice fringes, d) FFT of image in b, indexed as Si, and EFTEM maps of Ce, O, and Si, respectively.

Fig. 11 .
Fig. 11.Proposed mechanism of oxidation of Ce 3 Si 2 under synthetic air up to 750 °C.
10.1016/j.corsci.2019.108332Received 23 August 2019; Received in revised form 1 November 2019; Accepted 3 November 2019 and SiO 2 .However, no reaction mechanism of UO 2 to uranium silicate is proposed and the low oxidation temperature and partial pressures of O 2 are not representative of typical operating temperatures of ATF.

Table 1
wt.% fraction of CeO 2 , SiO 2 and Si in oxidised Ce 3 Si 2 material based on reaction products observed and oxidative mass gains.