# Some practical methodologies to assess the overall high temperature (one-sided) steam oxidation protectiveness of chromium-based coatings on a zirconium-based substrate, as Enhanced – Accident Tolerant (Nuclear) Fuels (E-ATF) claddings

Jean-Christophe Brachet, Thomas Guilbert, Stephane Urvoy, Elodie Rouesne, Thierry Vandenberghe, Cédric Prou, Thai Le Hong

► To cite this version:

Jean-Christophe Brachet, Thomas Guilbert, Stephane Urvoy, Elodie Rouesne, Thierry Vandenberghe, et al.. Some practical methodologies to assess the overall high temperature (one-sided) steam oxidation protectiveness of chromium-based coatings on a zirconium-based substrate, as Enhanced – Accident Tolerant (Nuclear) Fuels (E-ATF) claddings. Journal of Nuclear Materials, 2025, 606, pp.155620. 10.1016/j.jnucmat.2025.155620. cea-05185406

HAL Id: cea-05185406

https://cea.hal.science/cea-05185406v1

Submitted on 24 Jul 2025

HAL is a multi-disciplinary open access
archive for the deposit and dissemination of sci-
entific research documents, whether they are pub-
lished or not. The documents may come from
teaching and research institutions in France or
abroad, or from public or private research centers.

L'archive ouverte pluridisciplinaire HAL, est
destinée au dépôt et à la diffusion de documents
scientifiques de niveau recherche, publiés ou non,
émanant des établissements d'enseignement et de
recherche français ou étrangers, des laboratoires
publics ou privés.

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

Distributed under a Creative Commons Attribution 4.0 International License
<figure><img src="image_2.png" /></figure>

<figure><img src="image_3.png" /></figure>

# Some practical methodologies to assess the overall high temperature (one-sided) steam oxidation protectiveness of chromium-based coatings on a zirconium-based substrate, as Enhanced – Accident tolerant (Nuclear) fuels (E-ATF) claddings

Jean-Christophe Brachet*, Thomas Guilbert, Stéphane Urvoy†, Elodie Rouesne, Marion Peyret, Thierry Vandenberghe, Cédric Prou, Thai Le Hong©, Didier Hamon‡

Paris-Saclay University, CEA, DES/ISAS/DRMP/SRMA, 91191 Gif-Sur-Yvette Cedex, France

## HIGHLIGHTS

* Critical overview and suggestion of efficient experimental methodologies to assess the overall Cr-based coatings protectiveness on a Zr-based substrate upon High-Temperature one-sided oxidation,

* Use of Weight Gain coupling with oxide scale thickness to deduce the averaged overall oxygen quantity having diffused into the metallic substrate. Use of post-quenching micro-hardness measurements less costly to correctly estimate through-wall clad oxygen diffusion profiles,

* Necessity of taking into account the additional prior-β<sub>Zr</sub> hardening due to Cr diffusion into the Zr-based substrate upon HT transients, which depends on the oxidation temperature and final cooling scenario.

## ARTICLE INFO

**Keywords:**

Chromium coating

Zirconium alloys

High-temperature steam oxidation

Enhanced-accident tolerant nuclear fuel

claddings

Evaluation of coatings protectiveness

ECR

## ABSTRACT

Among various Enhanced-Accident Tolerant nuclear fuel claddings studied worldwide, Cr-coated Zr-based claddings are largely considered as a reference with evolutionary design and resulting ability to improve the high-temperature (HT) oxidation resistance. Most of the numerous studies carried out so far on the HT oxidation behavior of coated claddings have focused on detailed localized *post-mortem* examinations. The present paper aims to suggest complementary and sufficiently simple experimental methodologies to evaluate the overall protectiveness of Cr-based coated materials upon HT steam oxidation. The different proposed approaches have been assessed by comparing their capacity to characterize the Cr-coating loss of protectiveness and the associated oxygen diffusion into the sub-coating metallic wall-clad thickness. The latter is then compared to the results obtained by “more direct” Electron-Probe-Micro-Analysis measurements. The overall oxygen ingress into the tested coated claddings, evaluated by different proposed approaches, can be correlated to the Post-Quenching strength of the claddings, using a novel “Equivalent Cladding Reacted” parameter - called “Equivalent Chromium Reacted” ECrR – recently developed at CEA.

## 1. Introduction

Following the major nuclear accident at the Fukushima-Daiichi nu-
clear power plant in Japan, significant international efforts have been
carried out these last years to develop and evaluate Enhanced-Accident

Tolerant Fuels (E-ATFs) for nuclear Light Water Reactors (LWRs). Among various E-ATF candidates investigated worldwide, chromium (Cr)-coated zirconium (Zr)-based claddings appear an evolutionary design which, among others, take advantage of faster development and assessment based on long-term feedback on the in-service and

<footnote>* Corresponding author.</footnote>

<footnote>E-mail address: jean-christophe.brachet@cea.fr (J.-C. Brachet).</footnote>

<footnote>† present address: Paris-Saclay University, CEA, DES/ISAS/DRMP/SEMI, 91,191 Gif-Sur-Yvette Cedex, France</footnote>

<footnote>‡ Now retired from CEA</footnote>
<figure><img src="image_4.png" /><figcaption>Fig. 1. Refined scheme of Step 1 of the HT oxidation process of Cr-coated Zr and associated typical through-wall clad oxygen profiles (derived from Fig. 14 of [13]).</figcaption></figure>

hypothetical accidental behaviors of reference (uncoated) Zr-based nu-
clear fuels claddings.

Due to its intrinsic higher oxidation resistance as compared to con-
ventional zirconium alloys, one of the main objectives of the (Cr) coating
is to bring additional copying time before the loss of cladding integrity
upon High-Temperature (HT) oxidation typical of hypothetical LOss-of-
Coolant-Accident (LOCA) scenario. Additionally, upon Design Based
Accident (DBA-LOCA) and slightly beyond conditions (up to 1300°C),
the coating is able to delay the gaseous hydrogen production, thus
reducing the explosion risk and may have beneficial effects on the Peak-
Cladding-Temperature (PCT) by reducing/delaying the contribution of
the exothermic Zr substrate oxidation.

A literature survey has shown that, so far, more than one hundred
fifty papers exclusively dealing with studies on (mostly Cr-based) coated
Zr-based materials' behavior upon HT oxidation have been published
these last ten years. A detailed review of these numerous papers is out of
the scope of the present article, but one can refer to several "review
papers" [1-12] that give an overall picture of the R&D developed so far.
From this huge open literature, it is now well established that, among
other potentially influencing parameters, the HT oxidation coating
protectiveness may depend on its initial thickness as well as the coating
deposition process parameters inducing specific coating
microstructures.

However, the HT oxidation process of the Cr-coated Zr-based clad-
dings differs significantly from that of the uncoated reference Zr-based
claddings. Among other numerous papers published more recently,

such a specific behavior was discussed in detail in [13]. At HT and after
the first period for which the coating is fully protective, the Cr-coating
progressively loses its protectiveness, associated with a progressive in-
crease of the inward oxygen flux into the β<sub>Zr</sub> substrate, as schematically
illustrated in Fig. 1 to Fig. 3, which have been derived and refined from
Fig. 14 of [13]. Based on recent complementary examinations, the HT
oxidation process of the Cr-coated cladding can be subdivided now into
5 steps:

- Step 1-(a) corresponds to full protectiveness of the coating with no oxygen ingress into the underlying β<sub>Zr</sub> substrate. However, due to its significant solubility in the HT β<sub>Zr</sub> phase, chromium starts to diffuse into the substrate, thus contributing to the overall Cr-coating consumption upon HT incursion [12,25]. Additionally, according to Zr-Cr inter-diffusion and consistently with the phase diagram prediction, a ZrCr<sub>2</sub> intermetallic layer thickening is observed at the prior Zr-Cr interface, reaching a thickness of a few micrometers before being retransformed into pure chromium and zirconium oxide by the increasing inward flux of oxygen (Step2-(b)).

- Step 1-(b) corresponds to partial coating's protectiveness. In that case, the inward oxygen diffusion into the substrate already occurs
<figure><img src="image_5.png" /><figcaption>**Fig. 2.** Refined scheme of Step 2 of the HT oxidation process of Cr-coated Zr and associated typical through-wall clad oxygen profiles (derived from Fig. 14 of [13]).</figcaption></figure>

but remains limited and the oxygen solubility of the $\beta_{Zr}^{-1}$ is not achieved at the coating/substrate interface.

- Step 2-(a) corresponds to the early formation of a zirconium phase incursions enriched in oxygen, $\alpha_{Zr}(O)$, beneath the residual Cr coating when the inward oxygen diffusion into the substrate is significant enough to locally achieve the oxygen solubility of the $\beta_{Zr}$ substrate. As illustrated in Fig. 7 of [13], it is worthwhile to notice that $\alpha_{Zr}(O)$ incursions into the $\beta_{Zr}$ substrate occur in the continuity of the Cr coating grain boundaries, suggesting that oxygen mainly diffuses throughout the intergranular thin ZrO$_2$ network previously formed at prior Cr grain boundaries. At that time, except at the locations with the presence of $\alpha_{Zr}(O)$ incursions, continuous ZrCr$_2$ intermetallic interlayer is still observed at the Cr-Zr interface.

- Due to the further increase of oxygen diffusion into the $\beta_{Zr}$ substrate, Step 2-(b) is characterized by the thickening of a continuous oxygen-stabilized $\alpha_{Zr}(O)$ layer beneath the residual Cr coating. During this period, the intermetallic Zr-Cr layer is progressively transformed into zirconium oxide and metallic chromium. This can be explained by

the higher thermodynamic stability of zirconia over chromia oxides. Such a reaction can be clearly observed in Fig. 4, thanks to results recently obtained using Electron BackScatter Diffraction (EBSD) analyses.

However, the inward flux of oxygen is not enough to achieve the $\alpha_{Zr}(O)$ phase oxygen saturation (~29–30 at.%) in the vicinity of the Cr-Zr interface, and to induce ZrO$_2$ formation at this stage.

- Step 3 consists in a subsequent oxygen diffusion into the Zr-based metallic substrate. In this case, the inward flux of oxygen is high enough to attain the $\alpha_{Zr}(O)$ phase oxygen saturation (~29–30 at.%) at the coating/Zr interface, leading to the growth of an additional and continuous sub-coating ZrO$_2$ interlayer. This final step results in the full loss of the residual Cr-coating protectiveness and the acceleration in the oxidation kinetics of the cladding.

As a result of the different steps described above, the complex microstructural and micro-chemical evolutions upon HT oxidation make it difficult to assess the step-by-step loss of the coating protectiveness, i. e., that corresponding to the transition from Step1-(a) to Step1-(b). Thus, this paper attempts to discuss and define some practical and robust methodologies, simple but able to determine the overall Cr-coating

<footnote n="1">¹ ~0.55 to ~0.6 wt.% at 1200°C on as-received materials, depending on the Zr alloy substrate, but higher in the case of significant hydrogen uptake [17–24].</footnote>
<figure><img src="image_6.png" /><figcaption>Fig. 3. Scheme of Step 3 of the HT oxidation process of Cr-coated Zr and associated typical through-wall clad oxygen profiles (derived from Fig. 14 of [13]).</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Fig. 4. SEM-EBSD phase maps of 10 µm-thick Cr-coated Zr-based clad segments after one-sided oxidation at 1200°C for different oxidation times.</figcaption></figure>

protectiveness evolution upon HT oxidation, typical of hypothetical accidental (i.e., LOCA and beyond) situations.

Additionally, recalling that, by setting aside the additional and specific effect of hydrogen uptake,<fnref n="2" /> since the early works of Sawatzki et al. [15] and Chung & Kassner [16] the quenching resistance, as well as the Post-Quenching (PQ) strength and ductility, of the HT oxidized Zr-based claddings strongly rely on the through-wall clad oxygen diffusion profile. Indeed, it has been shown that, following direct water quenching from the HT $\beta_{Zr}$ temperature range, the Ductile-to-Brittle Transition

(DBT) occurs at Room Temperature (RT) for a typical prior-$\beta_{Zr}$ oxygen content of 0.4 wt.% for both Low-Tin Zircaloy-4 and $M_5\text{Framatome}$<fnref n="3" /> alloys [17,24].

In summary, practical methodologies able to assess the overall Cr-coating protectiveness upon oxidation at HT have:

- to be sensitive enough to detect the early oxidation time corresponding to the transition from Step1-(a) to Step1-(b);

- to be able to characterize the (under-coating) through-wall clad thickness oxygen diffusion, which could be further used to establish the relationship with the residual mechanical clad properties.

Finally, it has to be mentioned that the Zirconium-Chromium

<footnote n="2">² In the case of significant hydrogen uptake due to the in-service corrosion, and/or post-breakaway situation, and/or secondary hydriding phenomena (following the cladding burst occurrence), the additional hydrogen embrittlement effect must be taken into account [17–21]</footnote>

<footnote n="3">³ $M_5\text{Framatome}$ is a trademark or a registered trademark of Framatome or its affiliates in the USA or other countries</footnote>
<figure><img src="image_8.png" /><figcaption>**Fig. 5.** Optical micrographs and typical EPMA oxygen (and chromium) through-wall clad concentration profiles of uncoated and Cr-coated Zircaloy-4 claddings after one-sided or two-sided steam oxidation for 600s at 1200°C / i.e., corresponding to Baker-Just ECR<sub>BJ</sub> of 17 and 32% respectively (data partially replotted from figs 11 to 16 of [27]).</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Fig. 6.** Arrhenius plot of the rate constant of the parabolic law describing the growth of Cr₂O₃ for: Cr coated M5<sub>Framatome</sub> under steam environment (CEA/ Framatome data from [13]), and bulk chromium under air environment (Taneichi et al., [29]).</figcaption></figure>

eutectic reaction occurring at slightly above 1300°C causes ultimate failure of the Cr coating protective effect and hence, being beyond the temperature range discussed in this paper.

## 2. Brief critical overview of experimental methodologies to assess the coating's protectiveness upon ht oxidation

In the following, experimental methodologies commonly used to evaluate coated Zr-based alloys' behavior upon HT oxidation are briefly presented and discussed, especially in view of their ability to assess the overall protectiveness of the coatings relying on the oxygen diffusion into the Zr-substrate (if any).

### 2.1. One-sided versus two-sided HT oxidation

So far, numerous institutes have carried out two-sided HT oxidation testing to study the behavior of coated Zr-based claddings. This experimental protocol is fast and easy to apply. Furthermore, it has been proven to be an efficient way to characterize the reference uncoated cladding behavior upon HT oxidation (see for example [26], and many other references....). However, for the evaluation of the behavior of Cr-coated claddings, two-sided oxidation displays some following practical limitations:

- Firstly, at HT (up to 1300°C), Cr coating oxidation kinetics is at least one order of magnitude slower than that of the Zr-based substrate. Therefore, this suggests that, as far as the Cr-coating is protective, in the two-sided oxidation conditions, more than 90% of the overall measured Weight Gain (WG) can be attributed to the inner (uncoated) clad surface oxidation. This makes it difficult to accurately estimate the WG contribution exclusively due to the outer-coated surface oxidation.

- Secondly, along with the oxidation, the through-wall clad oxygen diffusion from the inner uncoated clad surface into the Zr substrate

rapidly takes place, as shown in Fig. 5 (replotted from [27]). This figure depicts optical micrographs as well as typical Electron-Probe-Micro-Analysis (EPMA-WDS) through-wall clad concentration profiles of oxygen (and chromium) on uncoated and Cr-coated claddings that were steam-oxidized in one or two-sided configurations for 600s at 1200°C.<fnref n="4" /> It is obvious that, for two-sided oxidation conditions, the oxygen diffusion coming from the inner clad surface oxidation reaches the outer Cr-Zr interface after only a few hundred seconds. Then, this outward oxygen diffusion from the inner clad surface may interact with the outer Zr-Cr interface, also with the inward flux of oxygen – if any – from the outer Cr-coated clad surface. The latter takes place at longer oxidation times, as previously discussed in Section 1.

In summary: to assess accurately the Cr-coating protectiveness upon HT oxidation, it is recommended to perform preferentially one-sided oxidation conditions. However, two-sided oxidation may be used to reproduce hypothetical LOCA transients, including relevant post-oxidation mechanical properties following the clad ballooning and burst. Indeed, in this case, inner-clad oxidation also occurs due to steam ingress through the burst opening. It is highly important to notice that, due to steam starvation leading to different oxidation kinetics along with secondary hydriding phenomenon, the two-sided oxidation experiments could be only representative of the localized two-sided oxidation conditions that potentially occur in the vicinity of the burst opening.

### 2.2. (Flowing) steam versus air and/or nitrogen-containing environments

For the sake of simplicity, several institutes performed HT oxidation in an air environment instead of flowing steam (more representative of hypothetical DBA-LOCA and slightly beyond transients) to assess the HT oxidation behavior of coated Zr-based claddings. The question then arises of a potential effect of nitrogen-containing atmospheres on the HT oxidation kinetics of Cr-coated claddings, comparing to a LOCA prototypical steam environment.

On the one hand, a limited number of oxidation tests carried out so far at CEA in an air environment at 1100–1200°C have shown that the overall HT oxidation kinetics of the Cr-coating were not drastically different from the ones observed under a pure (flowing) steam environment, provided that the coating is nearly fully protective. This is consistent with the observation made by previous studies on bulk Cr oxidized in an air environment [29,30]. It should be noted that these last-mentioned studies were conducted for short enough oxidation times. This can be observed in Fig. 6 in which the rate constant of the parabolic law describing the growth of Cr₂O₃ is plotted as a function of the inverse of the temperature for Cr-coated M5<sub>Framatome</sub> HT oxidation data under a steam environment [13], compared to those of bulk chromium under an air environment from [29]. In this last work, one of the author's conclusions was that: "...The changes of oxide layer thickness with heating time in pure chromium (and two other chromium alloys) were controlled by a parabolic rate law, and the rate controlling process for (air) oxidation was Cr cation diffusion through the outer Cr₂O₃ oxide layer...". This proposed mechanism tends to explain the relative insensitivity of pure chromium to HT air oxidation for short oxidation times as compared to pure oxygen and/or steam environments. However, after a certain oxidation duration, significant nitridation of the bulk Cr (sub-oxide) due to the occurrence of outer Cr₂O₃ cracking and/or spalling was observed [29] [30]. It must be mentioned that, even if the oxidation time applied by the authors was sometimes much longer than that generally applied on Cr-coated Zr alloys, it was observed that, in the case

<footnote n="4">⁴ For 0.58 mm thick claddings, this oxidation time at 1200°C results in a calculated Equivalent Cladding Reaction (ECR<sub>Bj</sub>) of 17 and 32% for one- and two-sided oxidation, respectively, using reference “Baker-Just” [28] correlation.</footnote>
<figure><img src="image_10.png" /><figcaption>Fig. 7. Weight Gain evolution as a function of the square root of oxidation time at 1200 and 1300°C of reference uncoated and first-generation coated $M_5Framatome$ (or Zircaloy-4) claddings. Comparison between continuous WG measurements (dotted line) and discontinuous PQ WG measurements for different applied oxidation times and direct water quenching down to Room Temperature (triangles).</figcaption></figure>

of bulk Cr, the relative contribution of Cr volatilization on the WG evolution was limited, as also discussed in [13] for Cr-coated Zr materials in a steam environment. In summary, for typical DBA-LOCA and slightly beyond conditions, the overall HT oxidation kinetics of the Cr-coated Zr alloys should be comparable between steam and air environments, but only for short-enough oxidation times, i.e., as long as the Cr-coating is still nearly protective. For longer oxidation times, outer $Cr_2O_3$ scale cracking/spalling and associated Cr nitridation may occur and could accelerate the overall HT oxidation by modifying the underlying oxidation/nitridation mechanisms.

On the other hand, it has been well documented that zirconium-based cladding is tremendously sensitive to nitrogen-containing environments, which induce early “breakaway” oxidation kinetics, and thus

significant oxidation acceleration, as compared to a pure steam atmosphere (among others, see numerous papers from KIT [31-35] and from IRSN [36-41]). Moreover, under air or nitrogen-containing atmospheres, strong effects of the tested sample geometric singularities should be underlined: edge effects or scratches on the metal surface inducing locally early “breakaway” and further localized acceleration of the oxidation. This means that, in the case of preexisting coating defects (scratches...) and/or sample ends/edges where the coating is thinner and potentially poorly protective, contribution of the geometric singularities on the overall oxidation kinetics could be more visible for air oxidation tests than for pure steam oxidation conditions. Indeed, unlike nitrogen-containing environments, it has been observed by numerous authors that, in a pure steam atmosphere, the localized oxidation of the
<figure><img src="image_11.png" /><figcaption>**Fig. 8.** Evolution of post-oxidation hydrogen content of Cr-coated $M_5\text{Framatome}$ claddings, at several initial Cr-coatings thicknesses (Cr₀) and having gone one-sided steam oxidation at 1100–1300°C, as a function of: (a) measured WG, (b) calculated one-sided Baker-Just ECR, and (c) new “ECrR” parameter.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>Fig. 9. Through-wall clad thickness oxygen diffusion profiles measured by EPMA within the residual metallic part (i.e., corresponding to the α<sub>Zr</sub>(O) and prior-β<sub>Zr</sub> layers) of uncoated M5<sub>Framatome</sub> claddings having gone one-sided steam oxidation at 1100, 1200, and 1250°C.</figcaption></figure>

Zr-based substrate at coating crack positions neither induces accelerated (localized) oxidation, nor early Cr-coating delamination (as long as the as-received Zr-Cr interface displays good metallurgical/crystallinity continuity and thus a good bonding [42] [43],).

Finally, a recent paper compared the HT oxidation behaviors of Cr-coated Zircaloy-4 claddings tested in air and steam environments [44]. The authors observed that the overall weight gain kinetics of the Cr-coated materials were higher under an air environment, with a more defective outer chromia scale, as compared to the one formed under a

steam environment.

In summary, to reproduce DBA-LOCA and slightly beyond conditions on Cr-coated zirconium-based nuclear fuel claddings, it is recommended to use a pure (flowing) steam environment. However, some specific hypothetical accidental scenarios may induce air and/or nitrogen ingress (i.e., Severe Accident (SA), Spent Fuel Pool Accidents, SB-LOCA, or Station Black Out events with accumulators (containing N<sub>2</sub>) discharge...). For these particular cases, air or steam-air mixture environments should be used to better take into account the specific effects
<figure><img src="image_13.png" /><figcaption>**Fig. 10.** Evolutions of (a) Weight Gain and (b) ZrO₂ thickness of one-sided steam oxidized uncoated M5<sub>Framatome</sub> cladding as a function of the square root of oxidation time at 1100, 1200, and 1250°C.</figcaption></figure>

of nitrogen on the HT oxidation of both outer chromium coating and Zr-based substrate. In these last scenarios, considering the sensitivity of the Zr-based substrate to nitrogen, even at a low fraction in the atmosphere, it is expected that, in the case of defective or cracked Cr coatings, the localized oxidation of the substrate could be more penalizing than that eventually observed under a (flowing) steam environment.

## 2.3. Sample geometry effects

Numerous institutes have used coated sheet samples as a surrogate for tubular clad geometry, to assess the Cr-coated Zr behavior upon HT oxidation. Indeed, when coated on their two surfaces, such samples are easier to test. Preliminary CEA experiments on first-generation coated materials [12] extensively used this kind of sample geometries; and this approach was very useful to carry out accelerated corrosion tests for both in-service conditions (e.g., autoclave test under pressurized water or steam at 360 or 415°C) and upon HT steam environment (LOCA and

beyond conditions). Moreover, such geometry is more adapted to Thermo-Gravimetric Analysis (TGA) experiments and facilitates the machining of conventional specimens for mechanical testing, such as tensile tests. However, for coated materials, using flat (sheet) samples instead of the ones with more representative tubular clad geometry may lead to some drawbacks as follows:

- Due to the coating deposition process, the use of non-representative samples may induce different grain morphology and crystallographic textures, as well as different internal stress states within the coatings in as-received conditions, as compared to those likely obtained for samples with tubular clad geometry. Such specific properties in the coating may have an influence on their subsequent behaviors upon HT.

- As already mentioned, compared to tubular clad geometries, the coating deposition on a flat substrate could result in some potential detrimental “sample-end effects”. Indeed, it is generally observed
<figure><img src="image_14.png" /><figcaption>Fig. 11. (a) Evaluation of WG estimated from the measured ZrO2 thickness and “extra WG” due to the oxygen diffusion into the metallic Zr substrate as a function of measured WG, (b) Overall quantity of oxygen that has diffused into the Zr-based metallic substrate of uncoated $M_5_{\text{Framatome}}$. Comparison of calculated values, estimated from WG and $ZrO_2$ thickness, with averaged values deduced from through-wall clad thickness integration of experimental EPMA oxygen diffusion profiles for oxidation temperature range of 1100–1250°C.</figcaption></figure>

<figure><img src="image_15.png" /><figcaption>Fig. 12. Calculated averaged quantity of oxygen that has diffused into the Zr-based metallic substrate of uncoated reference $M_5_{\text{Framatome}}$ as a function of square root of oxidation time at 1100, 1200, and 1250°C.</figcaption></figure>

that HT oxidation occurs much earlier at geometric singularity locations, thus inducing accelerated and heterogeneous oxidation.

In summary, even if flat coated samples have been proven useful for (qualitative) comprehensive studies and specific testing purposes, representative nuclear fuel cladding tubular geometries should be preferred to assess their behavior upon HT oxidation.

## 2.4. Measurements of weight gain and/or gaseous hydrogen production

Weight gain measurement (WG) is an efficient and popular experimental protocol to assess the overall HT oxidation of many materials, including nuclear fuel claddings. As already discussed in Section 2.1, for such measurements, two-sided oxidation is not recommended to avoid a major contribution of the inner uncoated surface on the total weight gain, except when the samples (e.g., sheet geometry) are efficiently coated on their two sides. WG measurement methodologies can be classified into two types, which are complementary to one another:

<figure><img src="image_16.png" /><figcaption>Fig. 13. Calculated (*) averaged quantity of oxygen that has diffused into the Zr-based metallic substrate of first-generation Cr-coated $M_5_{\text{Framatome}}$ as a function of square root of oxidation time (*) using the same approach as that used for the uncoated materials illustrated in previous figures, but considering the relative outer $Cr_2O_3$ thickening contribution to the measured WG (instead of $ZrO_2$).</figcaption></figure>

### 2.4.1. On-line WG measurements

Thermo-Gravimetric Analysis (TGA) is frequently used to follow continuously the WG evolution upon HT transients in an oxidizing environment. For a given material, such experiments allow the rapid establishment of the oxidation kinetics at different temperatures and upon heating/cooling (at moderate rates). However, there exists some limitations of these experiments, especially in the case of coated Zr-based materials:

- anisothermal transients and associated measurements are limited to relatively slow heating/cooling (generally less than 1°C/s). Then, it is not possible to reproduce prototypical LOCA (or other dynamic accidental) transients;

- for most commercially available TGA facilities, the maximum steam flow is quite limited. This experimental configuration may sometimes result in steam starvation and/or may not be able to extract efficiently the heat induced by a brief temperature exothermic/
<figure><img src="image_17.png" /><figcaption>Fig. 14. Evolution of averaged quantity of oxygen that has diffused into the Zr-based metallic substrate of uncoated and Cr-coated $M_5\text{Framatome}$ as a function of Baker-Just ECR (calculated for uncoated materials), for different oxidation temperatures and initial Cr-coating thicknesses.</figcaption></figure>

endothermic reaction event. Moreover, in TGA experiments, the thermocouples are generally not welded directly on the samples. This makes it impossible to correctly and precisely follow the temperature evolution on the sample surface in particular conditions. Indeed, such experimental boundary conditions could sometimes lead to uncontrolled (not measured) brief temperature overshoot during the heating phase or at the early beginning of the HT isothermal oxidation step, due to the exothermic oxidation reaction of zirconium at high oxidation temperatures.

- due to the limited size of the tested samples, and as already discussed, (poorly coated) potential “sample-end” effects may have a significant contribution to the overall WG measured by TGA. This is illustrated in Fig. 7, in which the weight gain obtained by TGA experiments carried out at CEA on first-generation Cr-coated Zr sheet samples is compared with that from post-quenching (PQ) weighing measurements (see also the discussion in the following Section 2.4.2). Indeed, it is observed on Cr-coated materials that, at both 1200°C and 1300°C, the weight increase kinetics derived from TGA is slightly faster, as compared to PQ weighing data derived from testing several centimeters long clogged clad segments. This is believed to be mainly attributed to the small (sheet) sample’s geometry/surface used for TGA, enhancing the relative contribution of “sample-end” effects.

Alternatively or complementarily, the quantification of the overall quantity of released gaseous hydrogen could be an interesting tool, which is able to follow continuously the oxidation progress upon HT transients. Indeed, except for some specific conditions, such as “post-breakaway” or “steam starvation”,<fnref n="5" /> in a flowing steam environment, the hydrogen uptake of advanced Zr alloy-based claddings can be negligible [47]. This could also be valid for Cr-coated materials as no significant hydrogen uptake has been observed so far on these materials in a flowing steam environment, except for some limited transient hydrogen uptake in which the coating definitely loses its protectiveness (see Section 2.5). This means that almost all the gaseous hydrogen produced by the water molecule dissociation due to the oxidation reaction is released in the steam environment (not absorbed by the materials). Thus, on-line measurements of the $H_2$ production kinetics can be easily and directly related to the overall oxidation progression. Furthermore, contrary to TGA (limited to relatively slow heating/cooling conditions), on-line

hydrogen production measurements make it possible to follow dynamic transients (as done extensively by KIT for more than two decades to support their “Quench” bundle tests [48]) and appears to be an efficient tools when applied to the study of HT oxidation kinetics of both uncoated and coated Zr.

### 2.4.2. Post-Quenching (PQ) WG measurements

The PQ WG measurements are generally performed on clad segments of a few centimeters long, allowing afterward PQ mechanical (and metallurgical) evaluations. After the HT oxidation step, the final cooling scenario applied can be either a direct water quenching down to Room Temperature (RT) or a “step cooling” corresponding to a first cooling at an intermediate cooling rate (1–10°C/s) followed by a final water quenching down to RT from a temperature range of 700 and 800°C. This last cooling scenario has been believed to be more representative of the LOCA transient, as compared to a direct quenching from the HT (i.e., from the $\beta_{Zr}$ phase temperature range) that could be a more penalizing scenario in terms of PQ mechanical behaviors, not only for the uncoated [17–19], but also for Cr-coated Zr-based claddings [52]. However, the direct quenching scenario allows the “freeze” of the cladding’s metallurgical state at the end of the HT oxidation process. Thus, such a cooling scenario is useful to better understand the HT oxidation process, without an additional effect of the on-cooling metallurgical evolution at moderate cooling rates, inducing quite complex micro-chemical partitioning during the allotropic $\beta_{Zr} \rightarrow \alpha_{Zr}$ phase transformation and due to the associated Secondary Phases Precipitation (SPP).

To summarize, compared to the HT continuous weighing measurement by TGA, PQ weighing and/or continuous $H_2$ release quantification are likely more LOCA prototypical and allows the evaluation of the PQ mechanical (and metallurgical) properties of HT oxidized nuclear fuel claddings, which may depend on the final cooling scenario. However, such an approach is more costly in terms of “time and materials consuming”, and must be preferentially conducted under a uniform temperature as well as a sufficiently high steam flow environment to avoid heterogeneous oxidation conditions along the tested clad segments.

### 2.4.3. Partial conclusions on WG (and/or $H_2$) measurements

WG (and/or gaseous $H_2$ release) measurements are useful to quantify the overall HT oxidation kinetics of coated zirconium claddings. However, with the final objective to make the link with the residual mechanical clad properties, such experiments failed to give a full picture of the through wall clad oxygen diffusion/partitioning of the Cr-coated claddings, which strongly differs from the reference uncoated Zr-based claddings [14].

## 2.5. Microstructural and micro-chemical characterizations

The major part of published works mostly dealt with multiscale examinations of post-oxidation and post-quenching Cr-coated Zr materials. As the main objective of the present paper is to discuss the overall evaluation of the Cr coating protectiveness, a deep and thorough review of all results and experimental methodologies applied so far is out of our scope; but, as already mentioned, overviews can be found in refs [1–12]. Beyond the overall HT oxidation kinetic evaluation discussed in previous sections, advanced PQ microstructural and micro-chemical characterizations developed so far have different objectives as follows:

- to better understand the relationship between the coating deposition parameters and the resulting Cr-coating microstructure as well as its impacts on the subsequent Cr-coated Zr HT oxidation behavior;

- to better understand the underlying metallurgical and diffusion mechanisms, from which physical-based modeling could be derived for the specific HT oxidation processes of Cr-coated zirconium-based alloys;

However, even if post-oxidation microstructural and micro-chemical characterizations gave some valuable and fundamental insights into the specific HT oxidation and underlying diffusion mechanisms, such

<footnote n="5">⁵ For such specific conditions, it has been shown that up to more than 50% of the $H_2$ quantity produced by the oxidation process can be entrapped into the Zr-based substrate (i.e., the Fraction of Absorbed Hydrogen (FHA) is higher than 50% [20],[45],[46])</footnote>
<figure><img src="image_18.png" /><figcaption>Fig. 15. (a) PQ "residual ductility", and (b) "maximum load" derived from PQ RCT of Cr-coated $M_5Framatome$, after being exposed to one-sided steam oxidation at 1200°C, as a function of the new "ECrR" parameter [51]; (c) Comparison with the evolution of the averaged sub-coating metallic substrate oxygen content deduced from measured overall WG and $Cr_2O_3$ thickness for various initial Cr-coating thicknesses, oxidation times, and oxidation temperatures.</figcaption></figure>

localized examinations are not able to provide an overall picture of the Cr-coating protectiveness evolution upon HT oxidation, especially in the case of:

- non-uniformity of the initial coating thickness;

- pre-existence of freshly formed (through-coating) defects/cracks due to the clad ballooning;

- non-uniformity of oxidation conditions along the tested sample and "sample-end" effects...

Moreover, such approaches are "time-consuming" and more difficult to apply systematically in hot-cell environments for Post-Irradiation Examinations (PIEs).

### 3. Suggestions for practical methodologies to assess the overall cr coating's protectiveness upon ht (steam) oxidation

As already discussed, the loss of Cr-coating protectiveness upon HT steam oxidation is a multi-step process, and preferred practical methodologies to assess the overall Cr-coating protectiveness upon (steam) oxidation at HT should be:

- sufficiently sensitive to detect the early oxidation time corresponding to transition from Step1-(a) to Step1-(b) (Fig. 1);

- able to characterize the (under-coating) through-wall clad thickness oxygen diffusion, to further derive the relationship with the residual clad mechanical properties;

- simple enough to allow fast evaluation and, in view of Post-Irradiation Examinations (PIEs), able to be easily transposed in hot-cell environments.
<figure><img src="image_19.png" /><figcaption>Fig. 16. Typical (linear) hardening relationships of prior-β<sub>Zr</sub> of M5<sub>Framatome</sub> (and for some Zr1Nb(O)+Cr bulk model alloys [52]) after direct water-quenching from the β<sub>Zr</sub> temperature range down to RT, as a function of: (a) oxygen content [17], and (b) chromium content in wt.%.</figcaption></figure>

<figure><img src="image_20.png" /><figcaption>Fig. 17. Weight Gain evolution upon one-sided steam oxidation at 1100°C of uncoated and first generation (14 µm thick) Cr-coated M5<sub>Framatome</sub>.</figcaption></figure>

3.1. Post-oxidation hydrogen uptake measurements

It has been recalled previously that, upon one-sided HT steam oxidation, a transient hydrogen uptake occurs when the coating loses definitely its protectiveness [13], which corresponds to the transition between Step2-(b) (Fig. 2) and Step3 (Fig. 3).

Fig. 8 depicts some updated results from [13]. Fig. 8(a) shows that there is no tendency between the PQ hydrogen contents and the measured WG. Nevertheless, as already mentioned in [13], when plotting the data as a function of the calculated Baker-Just ECR (ECR<sub>BJ</sub>), see Fig. 8(b), hydrogen uptake starts to become significant for a calculated one-sided Baker-Just ECR of nearly 50% for 10–20 µm thick Cr-coated materials. It should be mentioned that this particular ECR value corresponds to the transition between Step-2(b) and Step3. However, in the case of lower and higher Cr-coating thicknesses, i.e., smaller than 10 µm or greater than 20 µm, it is shown that the Baker-Just ECR thresholds for significant hydrogen uptake strongly depend on the initial Cr-coating thickness, as presented in Fig. 8(b). As already discussed in [14], this confirms that the use of the conventional regulatory Baker-Just ECR

parameter (used for several decades on reference uncoated claddings) seems no longer to be an appropriate metric to anticipate the PQ mechanical behavior of Cr-coated claddings having undergone HT steam oxidation. Then, to overcome this problem, a new “Equivalent Chromium Reacted” (ECrR) parameter has been recently proposed to take into account the effect of the initial coating thickness [51]:

$$ECrR (\%) = 6.73 \cdot 10^5 \cdot \exp \left( - \frac{15784}{T} \right) \frac{t^{1/2}}{X_0} \quad (1)$$

where T is the oxidation temperature in Kelvin, t is the oxidation time in seconds, and X₀ is the initial Cr-coating thickness in µm.

PQ hydrogen contents for Cr-coated claddings are then plotted as a function of this new “ECrR” parameter in Fig. 8(c). Interestingly, it can be shown that regardless of the initial coating thickness and the oxidation temperature (ranging between 1100 and 1300°C), the transition of hydrogen uptake is found along with an “ECrR” value of ~70%. It should be recalled that this last “ECrR” value has been coincidently observed as a “critical” threshold for the PQ Ductile-to-Brittle Transition (DBT) of
<figure><img src="image_21.png" /><figcaption>(c) Partially protective coating / Step 1 – (b) – just before reaching Step 2 – (a) (i.e., oxygen content reaching oxygen solubility of the sub-coating β<sub>Zr</sub> layer may further induce α<sub>Zr</sub>(O) incursions - Figure 2)</figcaption></figure>

<caption>**Fig. 18.** Through-wall clad O and Cr diffusion profiles (from EPMA) of first generation 14 µm-thick Cr-coated M5<sub>Framatome</sub> after one-sided steam oxidation at 1100°C for different oxidation times and after direct water-quenching.</caption>

the failure mode at RT, as detailed in [51]. Thus, PQ hydrogen measurement could be an appropriate experimental indicator, sufficiently sensitive to determine the definitive loss of the Cr-coating protectiveness upon HT steam oxidation. It is worth noticing that a similar criterion (i. e., based on a PQ hydrogen uptake of 200 wt.ppm as a threshold value for two-sided oxidation) has been proposed as an experimental indicator of time for “breakaway” occurrence upon steam oxidation around 800°C and 1000°C on reference uncoated nuclear fuel claddings materials [49].

In summary, measurements of Cr-coated clad hydrogen uptake following HT (one-sided) steam oxidation are shown as an “overall experimental indicator” for the oxidation time corresponding to (1) the

definitive loss of the Cr-coating protectiveness and (2) the associated PQ DBT occurrence. However, such criterion is not able to detect the early transition from Step1-(a) (i.e., fully protective coating) to Step1-(b) (i.e., partially protective coating inducing early oxygen diffusion into the underlying β<sub>Zr</sub> substrate, without sub-coating formation of neither α<sub>Zr</sub>(O) nor ZrO<sub>2</sub>). This is why, in the following section, two additional overall approaches are proposed.
<figure><img src="image_22.png" /><figcaption>**Fig. 19.** Through-wall clad micro-hardness ($HV_{0.025}$) profiles of first generation 14 µm-thick Cr-coated $M_5\text{Framatome}$ after one-sided steam oxidation at 1100°C followed by two types of final cooling scenarios: Direct water-quenching or two-step cooling.</figcaption></figure>
<figure><img src="image_23.png" /><figcaption>Fig. 20. Evaluation of the oxygen concentration profiles in the residual prior-β<sub>Zr</sub> inner layer (for Cr content lower than 1.5 wt.%) deduced from the measured through-wall clad micro-hardness profiles, by subtracting the Cr-induced hardening contribution, of first-generation 12–15 µm-thick Cr coated M5<sub>Framatome</sub> after one-sided steam oxidation at 1100°C and direct water-quenching.</figcaption></figure>

## 3.2. Overall Cr-coating protectiveness evaluation, upon HT (one-sided) oxidation, based on weight gain and outer oxide scale (ZrO<sub>2</sub> or Cr<sub>2</sub>O<sub>3</sub>) thickness measurements

### 3.2.1. Assessment of the methodology on uncoated (reference) Zr-based claddings: estimation of averaged oxygen content in the metallic substrate

It is first recalled that, by setting aside the additional and specific effect of hydrogen (coming for example from the in-service corrosion [50]), the residual mechanical strength and ductility of Zr-based claddings following HT oxidation and quenching strongly depend on oxygen diffusion into the Zr-based metallic substrate. Indeed, through-wall clad oxygen diffusion induces (1) the thickening of the outer brittle layers, constituted of zirconium oxide ZrO<sub>2</sub> and oxygen-enriched metallic phase α<sub>Zr</sub>(O), as well as (2) the hardening and embrittlement of the residual inner prior-β<sub>Zr</sub> layer. Consequently, preferred “practical” methodologies to evaluate the overall Cr-coating protectiveness upon HT (one-sided) oxidation should be able to quantify the oxygen diffusion profiles (if any) throughout the wall-clad thickness following HT oxidation. For this purpose, Electron Probe Micro-Analysis (EPMA) has been extensively used these last decades at CEA, as illustrated in Fig. 9, and in other research institutes. However, experimental procedures require accurate sample preparation and, in some cases, tricky calibration, especially due to oxygen surface contamination of the analyzed samples....

The following part of this section aims to use EPMA as an

experimental methodology reference to assess other (simplified) alternative ones.

The first proposed approach consists of deriving the overall quantity of oxygen, which has diffused into the residual metallic part of the HT oxidized cladding, from (1) the measured Weight Gain (WG) and (2) the outer oxide scale thicknesses. These two parameters are plotted in Fig. 10 for uncoated M5<sub>Framatome</sub>, as a function of the square root of (one-sided) steam oxidation time. As expected, for the oxidation temperature range investigated in this study, i.e., 1100–1250°C, the HT steam oxidation kinetic of the uncoated reference industrial material (M5<sub>Framatome</sub>) is parabolic as the relationship between WG or ZrO<sub>2</sub> thickness and the square root of oxidation time is nearly linear (see Fig. 10).

On the basis of the relative change in volume density between the formed oxide scale ZrO<sub>2</sub> (or Cr<sub>2</sub>O<sub>3</sub>) and the former Zr (or Cr) metallic substrate, the following equation is proposed to describe the relationship between the measured oxide thickness (X<sub>Ox</sub>) and relevant WG:

$$X_{Ox} = K \cdot WG \quad (2)$$

with K~0.15 and 0.16 µm/(mg/cm²) for ZrO<sub>2</sub> and Cr<sub>2</sub>O<sub>3</sub> respectively.

Fig. 11(a) presents the comparison of the calculated WG derived from the oxide thickness using Eq. (2) and the overall one measured directly after oxidation tests. It can be seen that Eq. (2) tends to underestimate the measured WG. This is logical as, apart from the formation of the oxide scale, a certain amount of oxygen, which also
<figure><img src="image_24.png" /><figcaption>**Fig. 21.** Comparison between the evolutions, as a function of new “ECrR” parameter, of prior-β<sub>Zr</sub> oxygen content (derived from the different approaches already described) and PQ cladding strength (derived from RCT at 135°C) of first-generation (14 µm-thick) Cr-coated M5<sub>Framatome</sub> after one-sided steam oxidation at 1100°C followed by direct water-quenching.</figcaption></figure>

contributes to the overall WG, diffuses into the underlying metallic layer. This part of WG, named as “extra” WG in the following, is then evaluated by subtracting the calculated WG due to the oxide scale formation from the overall measured WG, as depicted in Fig. 11(a). The next step consists of transforming the deduced “extra WG” to an oxygen content having diffused into the metallic part of the HT oxidized cladding. It is important to notice that the last value of the oxygen content is an averaged one that is normalized to the residual wall-clad metallic thickness. It can be seen in Fig. 11(b) that the overall quantity of oxygen diffusing into the metallic substrate, estimated using the methodology above-described, is consistent with the “reference” value deduced from the experimental EPMA oxygen diffusion profiles illustrated in Fig. 9.

The overall (averaged) oxygen quantity, which has diffused into the Zr-based metallic substrate of uncoated M5<sub>Framatome</sub>, is plotted in Fig. 12 as a function of the square root of oxidation time at 1100, 1200, and 1250°C. At this point, it should be kept in mind that this parameter remains an overall oxygen content averaged on the whole residual metallic substrate, and cannot provide information about the diffusion and partitioning of oxygen between the formed α<sub>Zr</sub>(O) and residual prior-β<sub>Zr</sub> layers. Thus, this parameter should only be seen as an overall indicator of the oxygen ingress into the residual metallic layer. Therefore, it is insufficient to deduce accurately the relationship with the residual PQ strength and ductility of HT oxidized Zr-based claddings using only this “overall (averaged oxygen concentration) indicator value”.

### 3.2.2. Application of the methodology to Cr-coated materials for one-sided<sup>6</sup> oxidation conditions

A similar approach can be used on the Cr-coated materials, but this

time by considering the relative outer Cr<sub>2</sub>O<sub>3</sub> thickening contribution to the measured WG (instead of ZrO<sub>2</sub>). Fig. 13 illustrates averaged quantity of oxygen that has diffused into the Zr-based metallic substrate of first-generation Cr-coated M5<sub>Framatome</sub> (calculated using the same method as described in Section 3.2.1) as a function of the square root of one-sided oxidation time for oxidation temperature range of 1100–1300°C.

Fig. 14 compares averaged quantity of oxygen that has diffused into the Zr-based metallic substrate of uncoated with that of Cr-coated M5<sub>Framatome</sub> as a function of Baker-Just ECR<sub>BJ</sub> (calculated for uncoated materials). It is confirmed that 12–17 µm Cr coating significantly mitigates the oxygen ingress into the sub-coating metallic Zr-based cladding, as compared to the historical regulatory ECR<sub>BJ</sub> = 17% threshold, which has been derived for reference uncoated Zr-based claddings. In other words, it is confirmed that the Cr coating with typical thicknesses ranging from 10 to 20 µm is beneficial in delaying the critical oxidation time for reaching PQ clad embrittlement after one-sided steam oxidation up to 1300°C. Nevertheless, it is shown in Fig. 14 that too thin Cr coatings (with initial thickness < 10 µm) bring only limited protection upon HT oxidation. This is consistent with the previous results and observations discussed in Section 3.1. Besides, the ECR<sub>BJ</sub> parameter (calculated for the uncoated materials) is no longer able to capture this initial Cr-coating thickness effect.

Fig. 15(a) depicts the PQ mechanical “ductility” at RT of Cr-coated M5<sub>Framatome</sub> (derived from Ring Compression Testing, RCT) after being exposed to one-sided steam oxidation at 1200°C as a function of new “ECrR” parameter, taking into account several initial Cr-coating thicknesses. As discussed in [51], the Ductile-to-Brittle transition at RT corresponds to a “critical” ECrR value of ~70%.

Using the same ECrR parameter, averaged sub-coating metallic substrate oxygen content deduced from the measured overall WG and Cr<sub>2</sub>O<sub>3</sub> thicknesses is plotted in Fig. 15(c) for different initial Cr-coating thicknesses, different oxidation temperatures and times. Interestingly, it can be observed that the overall (averaged) metallic substrate oxygen content to achieve conventional PQ “clad embrittlement” threshold (i.e., less than 2% of “residual ductility” derived from RCT) is close to 0.4–0.5 wt.%. Such a value is consistent with the critical oxygen content in prior-

<footnote n="6">⁶ In case of using uncoated end plugs to carry out the one-sided oxidation experiments, the extra WG measured at the end plugs due to the uncoated Zr-based surface area has to be corrected by subtracting their relative contribution to the overall WG measured. Such a correction may induce some additional uncertainties and then, the use of Cr-coated end plugs should be preferred.</footnote>
<figure><img src="image_25.png" /><figcaption>**Fig. 22.** Evaluation of the oxygen concentration profiles in the residual prior-β<sub>Zr</sub> inner layer (corresponding to Cr content lower than 1.5 wt.%) from the measured through-wall clad micro-hardness profiles, by subtracting the additional hardening due to Cr diffusion, of first-generation 12–17 µm-thick Cr coated M5<sub>Framatome</sub> after one-sided steam oxidation at 1200°C followed by direct water-quenching.</figcaption></figure>

β<sub>Zr</sub> inducing Ductile-to-Brittle transition of this layer at RT in the case of uncoated zirconium alloys [17] [24],. However, this value appears to be lower than that averaged on the whole metallic part, including α<sub>Zr</sub>(O) and prior-β<sub>Zr</sub>, for uncoated claddings, ranging between ~0.7 and 1.0 wt. % when reaching ECR<sub>BJ</sub>=17% (Fig. 14). As discussed in [14], this is

fundamentally due to the difference in relative partitioning of oxygen within the HT oxidized residual metallic substrates of uncoated and Cr-coated claddings:
<figure><img src="image_26.png" /><figcaption>**Fig. 23.** Chromium and oxygen diffusion profiles (left) and EPMA chemical maps (right) of 6–7 µm-thick Cr coated $M_5\text{Framatome}$ after one-sided steam oxidation at 1200°C followed by direct water-quenching. Remark: for a higher oxidation time, the scattering of the EPMA data observed within the outer part of the wall-clad thickness (left) is due to the partitioning of chromium and oxygen between the $\alpha_{Zr}(O)$ incursions and the residual prior-$\beta_{Zr}$, as depicted in the EPMA chemical maps (right).</figcaption></figure>

- for uncoated claddings, as shown in Fig. 9, a significant part of the overall quantity oxygen having diffused into the metallic substrate is “entrapped” in the $\alpha_{Zr}(O)$ layer, only a limited amount of oxygen diffuses through the $\alpha_{Zr}(O)/\beta_{Zr}$ interface towards the residual (inner) prior-$\beta_{Zr}$ layer. In other words, the “critical” value of the overall (averaged) oxygen content in the metallic part must be higher than the prior-$\beta_{Zr}$ layer “critical” one (~0.4 wt.% at RT) inducing prior-$\beta_{Zr}$ Ductile-to-Brittle transition.

- in the case of the coated materials, the Cr-coating strongly delays the oxygen ingress and $\alpha_{Zr}(O)$ formation within the sub-coating metallic substrate. Then, except for longer oxidation times potentially leading to significant $\alpha_{Zr}(O)$ incursion occurrence, in the early steps of the HT oxidation process (i.e., before reaching Step 2 – (b)), most of oxygen diffuses into the sub-coating prior-$\beta_{Zr}$ layer, inducing Brittle-to-Ductile transition at RT once the oxygen concentration reaches about 0.4–0.5 wt.% within the through-wall (prior-$\beta_{Zr}$) clad thickness.

Finally, as detailed in [14], the discussion above explains again why the overall WG, and, thus, the conventional Baker-Just ECR<sub>BJ</sub> parameter derived for the uncoated reference claddings are not fully appropriate to describe the specific behavior of coated Zr-based claddings upon HT oxidation as well as its consequences on PQ mechanical cladding properties.

### 3.3. Use of Post-Quenching micro-hardness measurements as a surrogate for EPMA

The practical methodology described above (based on measured overall WG and outer oxide scale thicknesses, i.e., $ZrO_2$ or $Cr_2O_3$) provides the first overall picture of Cr-coating protectiveness evolution upon HT oxidation. However, it fails to give a full picture of oxygen diffusion and partitioning within HT oxidized coated claddings to refine the link with the residual PQ mechanical properties. As already mentioned, EPMA is an efficient tool to quantify accurately the through-wall clad oxygen diffusion profiles following HT oxidation. However, this technique is strongly impacted by the presence of native oxides or hydroxides potentially formed during sample polishing and preparation, which requires careful calibration during experiment manipulation. Then, in the following, a simplified experimental methodology based on PQ micro-hardness measurements is proposed by taking advantage of the relationships between the water-quenched prior-$\beta_{Zr}$ hardening and the oxygen and/or chromium concentration(s) locally achieved within the residual metallic part of the HT oxidized claddings. Indeed, as already discussed in [17] and in [52], for dilute alloys, both oxygen and chromium additions induce a more or less linear water-quenched prior-$\beta_{Zr}$ hardening. Fig. 16 presents typical (linear) hardening relationships observed on PQ Prior-$\beta_{Zr}$ of $M_5\text{Framatome}$ as a function of the respective oxygen and chromium contents. Except for data obtained on
“bulk” Cr-doped model alloys [163], the Cr-induced hardening relationship lines in Fig. 16(b) have been derived by associating the PQ micro-hardness profiles with the Cr concentration profiles measured by EPMA. This method is similar to that used to establish the oxygen-induced hardening relationship replatned from [17] in Fig. 16 (a). However, to estimate the chromium contribution itself, only short oxidation times were considered, corresponding to Step 1 – (a) of Fig. 1 (i.e., fully protective Cr-coating), to avoid any additional influences of oxygen ingress. It is worth noticing that the PQ prior-β<sub>Zr</sub> hardening due to chromium depends on the temperature from which the water quenching is performed.

It is also important to notice that, for chromium, the linear PQ "solid-solution"-induced hardening is only valid for concentrations lower than ~1.5 – 2.0 wt.%, likely corresponding to a "critical" concentration inducing PQ DBT of the failure mode. For greater Cr concentrations, a sharp hardness increase is observed and relevant prior-β<sub>Zr</sub> microstructure is much more complex (its detailed description is out of the scope of the present paper and will be published elsewhere). In the case of oxygen, nearly linear prior-β<sub>Zr</sub> hardening could be observed up to ~1.0 wt. % [20].

Finally, according to the early studies of Rumball [53], it is assumed,
in the following, that oxygen and chromium have an additive linear
contribution to the water-quenched prior-β<sub>Zr</sub> hardening. This strong
assumption means that for Zr-(Cr-O) dilute alloys, inserted oxygen and
chromium atoms within the water-quenched (martensitic?) prior-β<sub>Zr</sub>
structure contribute "independently" to the observed macroscopic "solid
solution-induced" hardening. This hypothesis probably needs more
fundamental studies to be physically supported. However, as shown
further, it appears that this "linear prior-β<sub>Zr</sub> hardening additivity rule"
between the respective oxygen and chromium contributions seems to be
operating for the above-mentioned oxygen and chromium concentration
ranges.

3.3.1. Case study: one-sided steam oxidation at 1100°C of 14 µm-thick Cr-coated M5<sub>Framatome</sub>

As depicted in Fig. 17, after steam oxidation at 1100°C, the WG evolution of a first generation 14 µm-thick Cr-coated M5<sub>Framatome</sub> as a function of the square root of oxidation time is limited and very smooth. Furthermore, PQ microstructural examinations did not reveal occurrence of sub-coating α<sub>Zr</sub>(O) incursions even for the coated samples oxidized for up to 18000s (5h), indicating that Step 2 (Figure 2) of the oxidation process was not achieved. Indeed, after being oxidized for 18000s at 1100°C, nearly half of the initial metallic Cr thickness remains, which is consistent with the calculated “ECrR” of 67%.

To go further, Fig. 18 displays the through-wall clad O and Cr diffusion profiles obtained by EPMA, for three typical oxidation times at 1100°C (i.e., 1900, 6000, and 18000s respectively). From this figure, it appears that the Cr-coating remains fully protective at least up to 1900s, while after 6000s, a certain amount of oxygen starts to diffuse into the prior-β<sub>Zr</sub> metallic substrate (Step 1 – (b)). Interestingly, after oxidation for 18000s, the full (prior-β<sub>Zr</sub>) wall-clad thickness reaches the oxygen solubility (saturation) of the β<sub>Zr</sub> phase at 1100°C. In other words, this particular oxidation time may correspond to the end of Step 1 – (b) of the HT oxidation process, just before reaching Step 2 – (a) (i.e., occurrence of sub-coating α<sub>Zr</sub>(O) incursion - Fig. 2 – (a)).

Associated to Cr and O diffusion profiles presented in Fig. 18, Fig. 19
depicts micro-hardness (HV<sub>0.025</sub>) profiles for two types of final cooling
scenarios, including (1) a direct water-quenching from 1100°C, and (2) a
two-step cooling consisting of a slow cooling from 1100°C down to
700°C followed by a final water-quenching from this temperature. Note
that, even if the first part of the two-step cooling is quite slow (cooling
rate < 1°C/s), this cooling scenario can be considered more LOCA pro-
totypical than direct water-quenching from the β<sub>Zr</sub> temperature range.
As discussed in [52], Fig. 19 confirms that during the first step of slow
cooling down to 700°C, most of the available chromium within the
parent β<sub>Zr</sub> has enough time to precipitate as coarse ZrCr<sub>2</sub> intermetallic

particles, thus inducing significant prior-β<sub>Zr</sub> “softening” effect. In the
following, only the micro-hardness data obtained after direct
water-quenching from 1100°C are used.

On the basis of the additive contribution of oxygen and chromium assumption, residual prior-β<sub>Zr</sub> hardening due exclusively to oxygen is plotted in Fig. 20(a) after subtracting the Cr-induced hardening contribution, using the β<sub>Zr</sub> hardening relationship identified after direct water-quenching from 1100°C plotted in Fig. 16(b). Then, using the oxygen hardening relationship of Fig. 16(a), the corresponding prior-β<sub>Zr</sub> oxygen concentration profiles can be deduced, as plotted in Fig. 20 (b). It is worth noticing that the comparison between the oxygen profiles measured by EPMA and those deduced from the micro-hardness measurements (subtracting the Cr contribution) are consistent, even if EPMA seems to slightly overestimate the oxygen content in this particular case (likely due to the contribution of the native hydroxide formed on the surface during the samples preparation (polishing) process...).

Finally, Fig. 21 allows a comparison between the evolutions of the averaged prior-β<sub>Zr</sub> oxygen content, derived from the different approaches above-described, as a function of the new “ECrR” parameter (Eq. 1), for first-generation Cr-coated M5<sub>Framatome</sub> having gone one-sided steam oxidation at 1100°C for up to 18000s (5 hours). This figure also illustrates PQ cladding strength deduced from RCT at 135°C in addition to that obtained at RT. It is worth noticing that:

- The three practical “methodologies” applied to evaluate the aver-
  aged oxygen content in the prior-β<sub>Zr</sub> phase are in good agreement.

- The PQ cladding strength evolution is consistent with the evolution of the evaluated averaged prior-β<sub>Zr</sub> oxygen content.

3.3.2. Extension of the methodology to higher oxidation temperatures
(1200–1300°C)

As illustrated in Fig. 22 for oxidation at 1200°C, the extension of the methodology described above to higher oxidation temperatures seems not to be straightforward. Indeed, due to the increase of the chromium solubility and higher thermal diffusion coefficient within the β<sub>Zr</sub> substrate, a significant part of the cladding wall is affected by the Cr diffusion above the critical Cr content (i.e., 1.5 wt.%), inducing a sharp and non-monotone increase of the prior-β<sub>Zr</sub> hardening (and DBT at RT after direct water-quenching). Consequently, the methodology described above may be only applied to the residual part of the wall-clad thickness in which the chromium content is lower than 1.5 wt.%.

However, as shown in Fig. 22(c), the proposed methodology based
on PQ micro-hardness measurements seems to remain operating, as far
as only the inner wall-clad thickness is considered. Note that for a higher
oxidation time, some additional oxygen diffusion occurs from the un-
coated inner clad surface due to some residual partial pressure of oxygen
inside the tested clogged clad segments. However, This particular situ-
ation should be more or less representative of inner oxygen diffusion
coming from reduction of the thin zirconium oxide formed at the inner
clad surface after irradiation at Medium- up to High-Burnup.

Finally, Fig. 23 presents the behavior of a thin (6–7 µm) Cr-coated M5<sub>Framatome</sub> upon one-sided steam oxidation at 1200°C. In this particular case, consistently with its specific WG evolution plotted in Fig. 14, the thin 6–7 µm Cr coating loses its protectiveness only after a few minutes at 1200°C. Thus, this last case study underlines the sensitivity of the proposed methodology (using PQ micro-hardness measurements) to the influence of the initial Cr-coating thickness and confirms that Cr-coatings of less than 10 µm provides limited protectiveness at 1200°C and above.

**4. Conclusions and further work**

Most of the numerous studies carried out so far on the HT oxidation
behavior of Cr-based coated zirconium have focused on detailed local-
ized *post-mortem* examinations. Even if these (*post-oxidation*) micro-
structural and micro-chemical characterizations gave valuable and
fundamental insights into the specific HT oxidation and underlying diffusion mechanisms, such localized examinations do not give an overall picture of the Cr-coating protectiveness evolution upon HT oxidation. After a brief critical overview and discussion on existing approaches in the literature, some simple but efficient practical methodologies to assess the overall protectiveness of Cr-based coatings on zirconium-based substrate upon HT (one-sided) steam oxidation are suggested. Some of the proposed approaches are assessed by comparing the results with direct PQ chemical measurements using EPMA measurements.

Firstly, PQ measurements of the Cr-coated clad hydrogen uptake can be viewed as an "overall experimental indicator" of the oxidation time corresponding to definitive loss of Cr-coating protectiveness and the associated PQ DBT occurrence. However, such criteria is not able to detect the early transition from Step1-(a) (i.e., fully protective coating) to Step1-(b) (i.e., partially protective coating inducing early oxygen diffusion into the underlying β<sub>Zr</sub> substrate, without sub-coating formation of neither α<sub>Zr</sub>(O), nor ZrO<sub>2</sub>). It is why, two additional overall approaches have been proposed.

The second proposed approach is simply based on a comparison between the measured overall Weight Gain and the actual chromium scale thickness, which gives an overall picture of the averaged oxygen quantity that has diffused into the (Prior-β<sub>Zr</sub>) sub-coating metallic substrate (if any). Note that such a simple approach may be only used for experiments with good-enough HT one-sided oxidation uniformity and long-enough clad segment to mitigate "geometric singularity/sample-end" effects on the measured overall WG.

As a surrogate for EPMA, the third proposed approach takes advantages of through-wall clad micro-hardness measurements, following direct water-quenching from the β<sub>Zr</sub> temperature range. The methodology needs to take into account for the Cr diffusion from the coating within the sub-coating β<sub>Zr</sub> layer and its induced PQ hardening.

Overall, it has been shown that the two last approaches are efficient to characterize the quantity of oxygen that has diffused into the metallic Zr-based substrate following HT steam oxidation. They are sensitive enough to detect the early oxygen ingress into the Zr-substrate when the Cr-coating just begins to lose its protectiveness, before the occurrence of sub-coating α<sub>Zr</sub>(O) incursions. Thus, these approaches can be relied to the PQ mechanical properties of Cr-coated claddings following HT steam one-sided oxidation and direct water-quenching.

However, further work should extend and then adapt these approaches to other prototypical hypothetical accidental conditions, such as:

- Two-sided oxidation conditions, which may occur in the vicinity of burst opening (if it happens) or due to oxygen diffusion from the HT reduction of the zirconium oxide formed at the inner clad surface at Medium and High Burnup;

- Final "cooling scenario" effect, i.e., for more LOCA prototypical "two-step" cooling conditions;

- Potential additional hydrogen embrittlement effects coming from in-service corrosion or in the case of HT "post-breakaway" situations and/or secondary hydriding following cladding burst.

### CRediT authorship contribution statement

**Jean-Christophe Brachet:** Writing – original draft, Visualization, Validation, Supervision, Methodology, Formal analysis, Data curation, Conceptualization.
**Thomas Guilbert:** Methodology, Investigation.
**Stéphane Urvoy:** Methodology, Investigation.
**Elodie Rouesne:** Methodology, Investigation.
**Marion Peyret:** Investigation.
**Thierry Vandenberghe:** Investigation.
**Cédric Prou:** Investigation.
**Thai Le Hong:** Writing – review & editing.
**Didier Hamon:** Methodology, Investigation.

### Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

### Acknowledgments

Numerous people at CEA took part other past decade in the R&D on Cr-coated Zr-based nuclear fuel claddings and the authors really want to warmly thank them all.

The presented works have been supported by the CEA-EDF-FRAMATOME "Innovation COMBustible" project and the authors want to thank the industrial partners for their funding contribution and fruitful discussions.

### Data availability

Data will be made available on request.

### References

[1] M.Stueber Tang, H.J. Seifert, M. Steinbrueck, Protective coatings on zirconium-based alloys as accident tolerant fuel (ATF) claddings, Corros. Rev. 35 (2017) 141-165.

[2] G. Geelhood, WG. Luscher, PNNL Report n°28437, 2019. June.

[3] H Chen, X Wang, R. Zhang, Application and Development progress of cr-based surface coating in nuclear fuel elements: ii. current status and shortcomings of performance studies, Coatings 10 (9) (2020) 835, 28 août.

[4] K Krejčí, J Kabátová, F Manoch, J Kočí, L Cvrček, J Málek, et al., Development and testing of multicomponent fuel cladding with enhanced accidental performance, Nucl. Eng. Technol. 52 (3) (2020) 597-609. March.

[5] H Yeom, K. Sridharan, Cold spray technology in nuclear energy applications: a review of recent advances, Annal. Nucl. Ener 150 (2021) 107835. January.

[6] E Kasharov, B Afornu, D Sidelev, M Krinitcyn, V Gouws, A. Lider, Recent advances in protective coatings for accident tolerant Zr-based fuel claddings, Coatings. 11 (5) (May 2021) 557.

[7] J Yang, M Steinbrück, C Tang, M Große, J Liu, J Zhang, et al., Review on chromium coated zirconium alloy accident tolerant fuel cladding, J Alloy. Compoun 895 (2022) 162450. February.

[8] RH Khlifa, NN. Nikitenkov, Protective coatings for accident tolerant fuel claddings - a review, J. Nucl Fuel Cycle. waste Technol 21 (1) (2023) 115-47. March.

[9] Z Li, X Wang, H Chen, R Zhang, T Wei, L Yao, et al., Review on performance of chromium-coated zirconium and its failure mechanisms, Front Nucl Eng 2 (2023) 1212351. June.

[10] Ling Sun, Yuchen Xiao, Weijiu Huang, Baifeng Luan, Baoan Wu, Tang Huiyi, Review on the Preparation and High-Temperature Oxidation Resistance of Metal Coating for Fuel Cladding Zirconium Alloys, Nucl. Sci. Eng. (2023), https://doi.org/10.1080/00295639.2023.2219802.

[11] Martin Steinbrueck, Mirco Grosse, Chongchong Tang, Juri Stuckert, Hans Juergen Seifert, An overview of mechanisms of the degradation of promising atf cladding materials during oxidation at high temperatures, High Temp. Corr. Mater. (2024), https://doi.org/10.1007/s11085-024-10229-y.

[12] J.C. Brachet, I. Idarraga-Trujillo, M. Le Flem, M.Le Saux, V. Vandenberghe, S. Urvoy, E. Rouesne, T. Guilbert, C. Toffolon-Masclet, M. Tupin, C. Phalippou, F. Lomello, F. Schuster, A. Billard, G. Velisa, C. Ducros, F. Sanchette, Early studies on Cr-Coated Zircaloy-4 as enhanced accident tolerant nuclear fuel claddings for light water reactors, J. Nucl. Mater. 517 (2019) 268-285.

[13] J.C. Brachet, E. Rouesne, T. Guilbert, J. Ribis, M.Le Saux, G. Nony, H. Palancher, A. David, J. Bischoff, J. Augereau, E. Pouillier, High temperature steam oxidation of chromium-coated zirconium-based alloys: kinetics and process, Corros. Sci. 167 (2020) 108537, https://doi.org/10.1016/j.corsci.2020.108537.

[14] J.C. Brachet, M.Le Saux, J. Bischoff, H. Palancher, R. Chosson, E. Pouillier, T. Guilbert, S. Urvoy, G. Nony, T. Vandenberghe, A. Lequien, C. Miton, P. Bossis, Evaluation of equivalent cladding reacted parameters of cr-coated claddings oxidized in steam at 1200°c in relation with oxygen diffusion/partitioning and post-quench ductility, JNM 533 (2020) 152106, https://doi.org/10.1016/j.jnucmat.2020.152106.

[15] A. Sawatzki, G.A. Ledoux, S. Jones, Oxidation of zirconium during a high temperature transient, in: A.L. Lowe Jr., G.W. Parry (Eds.), Oxidation of zirconium during a high temperature transient, Zirconium in the Nuclear Industry, ASTM STP 633 (1977) 134-149.

[16] H.M. Chung and T.F. Kassner, "Embrittlement criteria for zircaloy fuel cladding applicable to accident situations in light-water reactors", Summary Report, NUREG/CR-1344, ANL-7948, (1980).

[17] J.C. Brachet, V. Maillot, L. Portier, D. Gilbon, A. Lesbros, N. Waeckel, J.P. Mardon, Hydrogen Content, Pre Oxidation and Cooling Scenario Influences on Post-Quench Mechanical Properties of Zy-4 and M5™ Alloys in LOCA Conditions - Relationship with the Post-Quench Microstructure, J. ASTM. Int. 5 (4) (2008). Paper JAI101116.
[18] JC. Brachet, C. Toffolon-Masclet, D. Hamon, T. Guilbert, G. Trego, J. Jourdan, A. Stern, C. Raepsaet, « Oxygen, hydrogen and main alloying chemical elements partitioning upon alpha→beta phase transformation in zirconium alloys », in: Solid State Phenomena, 172-174, Trans Tech Publications, Switzerland, 2011, pp. 753-759. Vols.

[19] J.C. Brachet, D. Hamon, J.L. Béchade, P. Forget, C. Toffolon-Masclet, C. Raepsaet, J.P. Mardon, B. Sebbari, Quantification of the chemical elements partitioning within pre-hydried Zircaloy-4 after high temperature steam oxidation as a function of the final cooling scenario (LOCA conditions) and consequences on the (local) materials hardening, IAEA-TECDOC-CD-1709 (2013) 253-265.

[20] JC. Brachet, D. Hamon, M.Le Saux, V. Vandenberghe, C. Toffolon-Masclet, E. Rouesne, S. Urvoy, JL. Béchade, C. Raepsaet, JL. Lacour, G. Bayon, F. Ott, Study of secondary hydriding at high temperature in zirconium based nuclear fuel cladding tubes by coupling information from neutron radiography/tomography, electron probe micro analysis, micro elastic recoil detection analysis and laser induced breakdown spectroscopy microprobe, J. Nucl. Mater. 488 (2017) 267-286.

[21] T.Le Hong, J.-C. Brachet, J. Crépin, M.Le Saux, Combined effects of temperature and of high hydrogen and oxygen contents on the mechanical behavior of a zirconium alloy upon cooling from the βₓ₂ phase temperature range, J. Nucl. Mater. 554 (2021) 153069.

[22] H.M. Chung, T.F. Kassner, Pseudo binary Zircaloy-oxygen diagram, J. Nucl. Mater. 84 (1979) 327.

[23] E.L. Hunt, P. Niessen, The effect of oxygen on the equilibrium Beta / Alpha + Beta transformation temperature of Zirconium-Niobium alloys, J. Nucl. Mater. 325 (1970) 134-136.

[24] A. Stern J.-C. Brachet, V. Maillot, D. Hamon, F. Barcelo, S. Poissonnet, A. Pineau, J.-P. Mardon, A. Lesbros, Investigations of the microstructure and mechanical properties of prior-βₓ₂ structure as function of the oxygen content (up to 0.9 wt %) in two zirconium alloys, in: ASTM 15th International Symposium on Zirconium in the Nuclear Industry, June 24-28, 2007, Sunriver, Oregon, 5, Journal of ASTM International, USA, 2008. Paper ID JAI101119.

[25] D. Kim, Y. Lee, Diffusion of chromium of Cr-coated Zircaloy accident tolerant fuel cladding: Model development and experimental validation, Surf. Coat. Technol. 468 (2023) 129698.

[26] H.K. Yueh, R.J. Comstock, B. Dunn, M.Le Saux, Y.P. Lin, D. Lutz, D.J. Park, E. Perez-Fero, Y. Yan, Loss of coolant accident testing round robin, in: Proceedings of TopFuel 2013, Charlotte, North Carolina, USA, September 15-19, 2013, pp. 550-556.

[27] J.-C. Brachet, S. Urvoy, E. Rouesne, G. Nony, M. Dumerval, M.Le Saux, F. Ott, A. Michau, F. Schuster, F. Maury, DLI-MOCVD Cr₂C₃ coating to prevent Zr-based cladding from inner oxidation and secondary hydriding upon LOCA conditions, J. Nucl. Mater. (2021), https://doi.org/10.1016/j.jnucmat.2021.152953.

[28] L. Baker and L.C. Just, "Studies of metal-water reactions at high temperatures. III. Experimental and theoretical studies of the zirconium-water reaction", ANL-6548, ML050550198, (May 1962).

[29] K. Taneichi, T. Narushima, Y. Iguchi, C. Ouchi, Oxidation or nitridation behavior of pure chromium and chromium alloys containing 10 mass%Ni or Fe in atmospheric heating, Mater. Trans. 47 (10) (2006) 2540, to 2546.

[30] L. Royer, X. Ledoux, S. Mathieu, P. Steinmetz, On the oxidation and nitridation of chromium at 1300°C, Oxid Met 74 (2010) 79-92.

[31] M. Steinbrück, N. Vér, in: High-Temperature Oxidation of Zircaloy-4 in Mixed Steam-Air and Steam-Nitrogen Atmospheres, Proceedings of ICAPP '10, San Diego, CA, USA, 2010, p. 10168. June 13-17Paper.

[32] M. Steinbrück, M. Böttcher, Air oxidation of Zircaloy-4, M5 and ZIRLO™ cladding alloys at high temperatures, J. Nucl. Mater. 414 (2) (2011) 276-285, 15 July.

[33] M. Steinbrück, High-temperature reaction of oxygen-stabilized α-Zr(O) with nitrogen.

[34] M. Steinbrueck, F. Oliveira da Silva, M. Grosse, Oxidation of Zircaloy-4 in steam-nitrogen mixtures at 600-1200°C.

[35] M. Steinbrück, S. Prestel, U. Gerhards, High-temperature interaction of oxygen preloaded Zr1Nb alloy with nitrogen, Nucl. Eng. Technol. 50 (2) (2018) 237-245. March.

[36] C. Duriez, T. Dupont, B. Schmet, F. Enoch, Zircaloy-4 and M5™ high temperature oxidation and nitriding in air, J. Nucl. Mater. 380 (2008) 30-45.

[37] O. Coindreau, C. Duriez, S. Ederli, Air oxidation of Zircaloy-4 in the 600-1000°C temperature range: Modeling for ASTEC code application, J. Nucl. Mater. 405 (3) (2010) 207-215, 30 October.

[38] I. Idarraga, M. Mermoux, C. Duriez, A. Crisci, J.P. Mardon, Raman investigation of pre- and post-breakaway oxide scales formed on Zircaloy-4 and M5® in air at high temperature, J. Nucl. Mater. 421 (1-3) (2012) 160-171. IssuesFebruary.

[39] C. Duriez, D. Drouan, G. Pouzadoux, Reaction in air and in nitrogen of pre-oxidised Zircaloy-4 and M5™ claddings, Journal of Nuclear Materials 441 (2013) 84-95.

[40] M. Lasserre, V. Peres, M. Pijolat, O. Coindreau, C. Duriez, J.-P. Mardon, Modelling of Zircaloy-4 accelerated degradation kinetics in nitrogen-oxygen mixtures at 850°C, J. Nucl. Mater. 462 (2015) 221-229. July.

[41] C. Duriez, O. Coindreau, M. Gestin, A. Kasperski, V. Peres, M. Pijolat, H. Buscail, C. Issartel, R. Rolland, M. Mermoux, Zircaloy-4 high temperature oxidation in atmospheres representative of SFP-LOCA: Investigation of the influence of a low temperature pre-oxidation scale, J. Nucl. Mater. 513 (2019) 152e174.

[42] J. Ribis, A. Wu, J.-C. Brachet, F. Barcelo, B. Arnal, Atomic-scale interface structure of a Cr-coated Zircaloy-4 material, J. Mater. Sci. 53 (14) (July 2018) 9879-9895.

[43] A. Wu, J. Ribis, J.-C. Brachet, E. Clouet, F. Leprêtre, E. Bordas, B. Arnal, HRTEM and chemical study of an ion-irradiated chromium/zircaloy-4 Interface, J. Nucl. Mater. 504 (2018) 289-299.

[44] R. Chen, et al., Microstructure evolution and failure mechanism of Cr-coated Zry-4 in air and steam at 1200°C, J. Mater. Res. Technol. 30 (2024) 4412-4423.

[45] T. Furuta, S. Kawasaki, « Reaction Behavior of Zircaloy-4 in Steam-Hydrogen Mixtures at High Temperatures », JNM 105 (1982) 119-131.

[46] M.Le Saux, J.-C. Brachet, V. Vandenberghe, A. Ambard, R. Chosson, Breakaway oxidation of zirconium alloys exposed to steam around 1000°C, Corros. Sci. 176 (2020) 108936.

[47] J.-C. Brachet, V. Vandenberghe, ... on the hydrogen uptake of Zircaloy-4 and M5™ alloys subjected to steam oxidation in the 1100-1250°C temperature range, Letter to the Editor, J. Nucl. Mater. 395 (2009) 169-172.

[48] J. Stuckert, M. Grosse, M. Steinbrueck, M. Walter, A. Wensauer, Results of the QUENCH-LOCA experimental program at KIT, J. Nucl. Mater. 534 (2020) 152143.

[49] M.C. Billone, Y. Yan, T. Burtseva, R. Daum, "Cladding embrittlement during postulated loss-of-coolant accidents", NUREG/CR-6967, ML082130389, (July 2008).

[50] L. Portier, T. Bredel, J.C. Brachet, V. Maillot, J.P. Mardon, A. Lesbros, Influence of long service exposures on the thermal-mechanical behavior of Zy-4 and M5™ Alloys in LOCA Conditions, J. ASTM. Int. 2 (2) (2005). Feb.N° paper ID JAI12468.

[51] J.-C. Brachet, T. Guilbert, D. Hamon, G. Nony, Paul Gokelaere, A new "ECR" parameter to support HT one-sided steam oxidation behavior of Cr-coated Zr-based E-ATF claddings..., in: Proceedings of International Conference, Grenoble, France, 2024 "Topfuel- 2024".

[52] J.-C. Brachet, P. Gokelaere, T. Guilbert, C. Toffolon-Masclet, S. Urvoy, M. Dumerval, G. Nony, et al., Effect of chromium content on the on-cooling phase transformations and induced prior-βₓ₂ mechanical hardening and failure mode (in relation to enhanced accident-tolerant fuel chromium-coated zirconium-based cladding behavior upon and after high-temperature Transients), in: S.K. Yagnik, M. Preuss (Eds.), Zirconium in the Nuclear Industry: 20th International Symposium, ASTM International, West Conshohocken, PA, 2023, pp. 433-458. http://doi.org/10.1520/STP164520220028.

[53] W.M. Rumball, « The effect of chromium, oxygen and microstructure on the hardness of zirconium-chromium alloys », J. Common Met. 39 (1) (1975) 35-42, https://doi.org/10.1016/0022-5088(75)90215-5, janv.