Mechanism of corrosion of zirconium hydride and

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May 22, 2015 - Environnement, CEA Saclay, 91191 Gif-sur-Yvette – France .... on the corrosion of Zircaloy-4 in trying to answer to the following questions : ..... First of all, we have to remember elements from the literature date that we have to ...
Accepted Manuscript Title: Mechanism of corrosion of zirconium hydride and impact of precipitated hydrides on the Zircaloy-4 corrosion behaviour Author: M. Tupin C. Bisor P. Bossis J. Chˆene J.L. Bechade F. Jomard PII: DOI: Reference:

S0010-938X(15)00253-X http://dx.doi.org/doi:10.1016/j.corsci.2015.05.058 CS 6345

To appear in: Received date: Revised date: Accepted date:

6-1-2015 22-5-2015 23-5-2015

Please cite this article as: M. Tupin, C. Bisor, P. Bossis, J. Chˆene, J.L. Bechade, F. Jomard, Mechanism of corrosion of zirconium hydride and impact of precipitated hydrides on the Zircaloy-4 corrosion behaviour, Corrosion Science (2015), http://dx.doi.org/10.1016/j.corsci.2015.05.058 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

MECHANISM OF CORROSION OF ZIRCONIUM HYDRIDE AND IMPACT OF PRECIPITATED HYDRIDES ON THE ZIRCALOY-4 CORROSION BEHAVIOUR

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CEA/DEN/Service d'Etude des Matériaux Irradiés, 91191 Gif-sur-Yvette – France

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M. TUPIN(1), C. BISOR(1), P. BOSSIS(1), J. CHÊNE(2), J. L. BECHADE(3), AND F. JOMARD(4)

CEA/DEN/Service de la Corrosion et du Comprtement des Matériaux dans leur

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Environnement, CEA Saclay, 91191 Gif-sur-Yvette – France

CEA/DEN/Service de Recherche en Métallurgie Appliquée, 91191 Gif-sur-Yvette – France

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GEMaC, UMR 8635 / CNRS Université de Versailles-Saint-Quentin 1, place Aristide

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Briand F-92195 Meudon Cedex, FRANCE.

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Corresponding Author : Marc Tupin _ Tel/Fax : +33-(0)1 69 08 88 69/ 90 73 [email protected]

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ABSTRACT

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In Pressurized Water Reactors, zirconium hydrides precipitate in the matrix and could increase the oxidation rate of the claddings. To understand their effect, corrosion tests, TEM and µ-XRD analyses have been performed. This work showed that the oxidation rate and the oxygen diffusion coefficient in the oxide formed on massive hydride are much greater than those of Zircaloy-4. Moreover, oxide characterizations indicated an additional phase indexed as the sub-oxide Zr3O between the oxide film and the massive hydride. Finally, the hydrogen of the hydrides is not incorporated in the oxide during the corrosion process.

Keywords : zirconium (A) reactor conditions (C) hydrogen absorption (C) SIMS (B)

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Introduction

The cladding tubes is the first containement barrier to fission products, their mechanical integrity is so essential for nuclear safety. Due to corrosion and hydriding, Zircaloy-4 cladding is the limiting factor of the lifetime of the fuel assembly. In Pressurized Water Reactors, zirconium alloys are exposed to aggressive aqueous environment (285-325°C, 155 bars, 0.7 to 2.2 ppm Li and 10 to 1200 ppm B, [H2]=25cm3/kg in normal temperature and pressure conditions). In these conditions, the corrosion kinetics of Zircaloy-4 (Zr-1.3Sn-0.2Fe-0.1Cr) shows a drastic acceleration, named “high burnup or phase III acceleration” which occurs for fuel burnups above 40 GWd/tU [1]. To explain this acceleration, three major hypotheses were proposed in the literature. The first cause advanced could be the tin content in the alloy. Indeed, the reduction of tin content in Zircaloy-4 seems to be

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correlated with a lowered corrosion rate of the alloy in PWRs [1, 2, 3]. Moreover, a delayed onset of the high burnup corrosion acceleration has actually been observed on low tin containing samples of Zircaloy-4 (often called Zy4 in the text). A second assumption assigned this accelerated corrosion to the impact of Zr(Fe,Cr)2 precipitates dissolution under irradiation [4, 5]. Finally, several authors suggested that this enhanced corrosion rate would be rather due to the precipitation of massive hydrides at the metal/oxide interface in the metallic part of the cladding [6, 7, 8, 9] and some of them tried to understand the relationship between the precipitation of massive hydrides in the cladding and its high burn-up oxidation rate. Several studies focused actually on the oxidation rate of hydrides compared with the oxidation rate of the -Zr matrix [8, 10, 11]. On Zircaloy-4 claddings oxidized 2 cycles in PWR, a correlation was pointed out between healthy claddings without any rim of hydrides under the metal/oxide interface and defective claddings with a five time thicker oxide associated with rims containing a hydrogen concentration of 1295 to 1520 wt ppm(12). In out-of-pile experiments based on (homogeneously) pre-hydrided and reference samples of Zircaloy-4 oxidized in autoclaves at 360°C, a twice higher weight gain was obtained on the pre-hydrided samples after long time exposures [6, 7]. Other studies propose an impact of hydride precipitates on the -Zr crystalline lattice inducing a loss of coherency with the zirconia lattice, and thereby a less protective oxide barrier at the interface. The oxidizing species in this layer would diffuse more easily through the scale to reach the metal/oxide interface which would lead to accelerated corrosion [6, 7]. During oxidation of cathodically pre-hydrided Zr-2.5Nb samples in water at 350°C, an enhanced corrosion rate was observed in the crystallographic planes previously disturbed by hydride precipitation in metal [13]. Finally, a change in the oxide structure and/or the microstructure induced by the hydride corrosion has also been considered to explain the “phase III” acceleration. Though several studies were carried out to validate this assumption, none of them clearly showed a modification in the corrosion film. An alteration of the oxide microstructure at a finer scale than the one of Transmission Electron Microscopy has been suggested [9]. Moreover, in this study, a higher density and periodicity of cracks in the oxide scales formed on Zircaloy-4 prehydrided samples was evidenced, but it was not stated whether it was a cause or a consequence of the accelerated corrosion rate. Leaning on observations of Zircaloy-4 claddings corroded in reactor, a relieve of the compressive stresses at the metal/oxide interface due to the precipitation of Zr hydrides was envisaged, associated with the quadratic to monoclinic phase transformation in the oxide film, creating more pores and defaults and thus leading to an easier penetration of oxidizing species through the scale [12]. Finally, the potential impact of a hydride rim at the metal/oxide interface on the corrosion kinetics and the related mechanisms is still under debate. Within this framework, this study investigates the effect of massive Zr hydrides on the corrosion of Zircaloy-4 in trying to answer to the following questions : 1. What is the effect of hydrides from a kinetic point of view ? Is their oxidation rate higher than that of an hydrogen free matrix ? 2. What are the transport properties of oxygen in an oxide formed on a massive hydride in comparison to these in a layer formed on a reference Zircaloy-4 ? 3. How do the hydrides change the oxide microstructure and what are the phases formed during their corrosion process ? 4. What does the hydrogen from the hydrides become during the corrosion process ? Are they incorporated in the oxide layer or do they stay in the hydrided matrix ?

Experiment and materials Materials

In this study, sheet materials of recrystallised, fully annealed Zircaloy-4 (called Zy4 in the text) supplied by CEZUS were used. The chemical composition of the alloy is given in Table I. The specimens were cut from a 0.425 mm thick cold-worked sheet which received a final heat treatment of recrystallization at 700°C. The alloy microstructure has been presented in a previous paper [14] and is characterized by an average grain size of about 10 µm.

-Table 1-

2.2

Hydriding treatment

Before corrosion tests, half of the samples was cathodically charged with hydrogen in order to precipitate the δZrH1.66 phase in the outer surface of the specimens. The hydriding protocol was borrowed from EDF R&D laboratory [8]. The samples were previously pickled with a fluoronitric solution (4% HF, 39% HNO3, and water)

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before being cleaned in an ultrasonic-tray. Then, they were immersed 98 hours in a 0.5M sulfuric acid solution under a cathodic current density of -7.5 mA/cm². The resulting thickness of the hydride phase extended from around 5 to 10 µm depending on the global hydrogen concentration. The SEM micrographs exposed in the figure 1a and 1b show a cross section and a fractography of cathodically charged samples, respectively with a 400 ppm and a 170 ppm global hydrogen content.

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-Figure 1-

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The micrograph shown in the Figure 1a indicates that during the cathodic charging, the hydride phase precipitates only at the outer surface of the sample and appears to be massive. For this specimen, an hydrogen content of 400 ppm (wt) corresponds approximately to a 10 µm thick massive hydride phase according to XRD analyses (not presented in this paper). Moreover, the fractography presented in Figure 1b points out a different failure facies of the hydride phase which is brittle compared with the ductile Zircaloy-4 matrix. In this Figure 1b, the average thickness of the hydride phase is about 7 µm and corresponds to a global hydrogen content of 170 ppm. At a finer scale, the transmission electron micrograph presented in Figure 2a shows a microstructure of the hydride phase quite heterogeneous and composed of a mixture of hydrides grains and hydrides platelets. The figure 2a shows a dispersive grain size from 0.2 µm to 1 µm while the average grain size of initial Zircaloy-4 matrix is about 10 µm. The indexation of -ZrH1.66 phase is presented in Figures 2b to 2d and this phase corresponds to that met in fuel claddings in core.

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Corrosion tests

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All the pre-hydrided and reference Zircaloy-4 samples were corroded in static autoclaves at 360°C and 18.7 MPa in light primary water. The water chemistry was composed of 2 ppm of lithium and 1000 ppm of boron prepared using lithium hydroxide and boric acid. To follow the kinetic behavior of the pre-hydrided and reference Zy4, the weight gains were frequently measured after each 7-day exposure time before being replaced in the autoclave for a new oxidation cycle. To understand the oxidation mechanisms of hydrided samples compared with reference specimens, isotopic exchanges of oxygen were carried out. In a first step, all samples were corroded with specific times of exposure in order to obtain nearly the same oxide thickness for the pre-hydrided and the reference samples in the pretransition or in the post-transition stage of Zircaloy-4 kinetics. In a second step, the samples were corroded in H218O/H216O (20% vol. of H218O) during short or long time exposures, from 6 hours to 28 days.

Characterization of samples

1. SIMS analyses The elementary penetration profiles in the oxide scales were analyzed by SIMS technique using an IMS 4F CAMECA ionic analyzer system with a 10 KeV energetic Cs+ primary ion beam. Negative secondary ions were collected from an analyzed surface of 33 µm diameter with a low weight resolution mode. 2. TEM and SEM characterization Cross-sectional thin foils were prepared via Focus Ion Beam technique (thinning and cutting out with a Ga+ ion beam) in order to be observed with a Transmission Electron Microscope FEI Technai 30 G2. Characterizations of the samples were also made using a high resolution FEG SEM ZEISS ULTRA 55. 3. X-Rays Micro-Diffraction (µ-XRD) The µ-XRD technique using synchrotron radiation is now widely used to determine locally the phase formed on the metallic substrate and is clearly described in the publication of A. Motta et al. [15]. Two samples were analysed with this technique at Advanced Photon Source 2ID-D beamline, one hydrided material and the reference one, oxidized until the post-transition stage. Whatever the studied material, the thin foils used for these analyses was performed in the plan formed by the normal direction and the rolling direction (Figure 3). Thus, taking into account the recristallized texture of sheet, the direction analyzed by the X-ray incident beam is perpendicular to the axes (normal to the basal plane).

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Figure 3 represents schematically the crystallographic orientation of the thin foils compared to the X-ray incident beam. -Figure 3-

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The Zy4 reference sample analyzed by µ-XRD has an 3.6 µm oxide thickness. A transverse thin foil was produced in order to characterize the metal/oxide interface and also the external interface : oxide/coating (Sn+Bi layer). Figure 4 shows a SEM micrography of this thin foil. -Figure 4-

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The yellow arrow indicates the displacement of analysed location during the experiment. During the µ-XRD analyses, it was also possible to follow-up thanks to the fluorescence signal of zirconium the beam position on the sample as well as the various interfaces. 31 scans were carried out, with a step of 0.2 µm. Lastly, 19 scans correspond to the analysis of the oxide film, which are equivalent to a thickness close to 3.8 µm and consistent with the mass gain and the SEM observations.

Results Results of corrosion tests

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-Figure 5a-Figure 5b-

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In this part, the aim is to evidence or not the difference of corrosion kinetics between hydrided materials and the reference Zy4. The corrosion behaviour at 360°C in PWR conditions of this latter was first of all studied : the kinetic curves of two samples are presented in figure 5a. The mass gain of the samples measured at each step of the corrosion process is converted to oxide thickness considering the classical relationship between these two parameters [16], supposing that all oxygen is used for oxide growth. It is well established that the pre-transition stage of Zircaloy-4 corrosion is characterized by the growth of a dense and protective oxide layer limited by the oxygen diffusion in the scale. The experimental points on both samples exposed on the figure 5a are reproducible and were fitted according to the usual power law (x = ktn). According to the fitting parameters given in the table 3, the kinetic curve follows a sub-parabolic law with an n parameter of 0.3, value in agreement with the literature data [15, 17]. The gap regarding the classical parabolic law has been widely discussed in the past and could be due to the presence of cracks [16], the compressive stresses [18, 19], the electric field built in the layer [20], the variations of the grain size [21] and so on…

When the alloy reaches the kinetic transition, around 2 µm, the oxide cracks [22, 23] and probably becomes porous [24, 25, 26] whereas a new protective film is forming at the inner interface. Hence, in the post-transition stage, the oxide is divided into two sub-layers: the external one, which presents a lot of defects as pores and cracks, and a quite dense internal layer near the metal/oxide interface. In this post-transition region, a new corrosion cycle occurs that can be approximatly fitted with the same power law. The hydrogen contents absorbed by the alloy during corrosion has been measured, before and after the kinetic transition (results not presented in this paper). The results seem to be in agreement with bibliography with an average hydrogen pick-up fraction of 12%. The corrosion behavior at 360°C in PWR simulating conditions of pre-hydrided Zircaloy-4 samples were also studied and compared to reference specimens. Their characteristics and their whole exposure time are summarized in Table 2. Table 2 The kinetic curves of these materials are presented in Figure 5b. First of all, three groups of specimens must be distinguished: the specimens without hydrogen, the hydrided samples with an average content of 150 ppm, and the (about) 300 ppm containing ones. For each group of specimens, the reproducibility of the kinetic curves was quite satisfactory and these curves show clearly an impact of cathodic charging on the corrosion rate of Zircaloy4. This influence is obviously dependent on the level of the initial hydrogen content in the sample since the corrosion acceleration of hydrided samples compared with reference Zircaloy-4 was more significant on 300 ppm containing samples. This result is quite unexpected because the samples are theoretically covered by a thickness of around few micron hydride phase, whatever the hydrogen content charged in the sample. It tends to prove that either the hydride phase is not completely pure for the sample having the lowest global hydrogen content or a significant fraction of the massive hydride is partially dissolved in the -Zr matrix during the corrosion test.

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The oxidation curves were fitted (dotted lines) with a power law in order to determine the n and k parameters of the kinetic law in the pre transition domain. The results are gathered in Table II. For samples charged with 150 ppm (ZH1 and ZH3), the influence of hydrogen charging on the corrosion rate is characterized by a higher k parameter. For samples charged with 300 ppm hydrogen content, the hydride effect results in higher k and n parameters. On the 300 ppm-pre-hydrided curve, 2 kinetic transitions can be pointed out around 1.5 µm and 2.8 µm, whereas the reference curve is still in the pre transition stage. The kinetic curves of the hydrided specimens on the figure 5b show actually an earlier kinetic transition and shorter corrosion cycles with transition thicknesses close to 1.5 µm compared to 2 µm for the reference samples, as already mentioned in the literature [24, 25]. Quantitatively, one can estimate by derivation of the power law the average corrosion rate at a given thickness, typically, 1.2 µm for each kind of material. And, as indicated in the last column in the previous table, the rate ratios between a pre-hydrided material and a reference one increase with the hydrogen content and can reach a factor of approximately 4 for the highest charged samples.

Results of alternated oxidation tests in H218O enriched environment.

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In conclusion, these results confirm that -ZrH1.66 hydrides show, firstly, a corrosion rate significantly higher than that observed on the reference Zy4 in a given kinetic domain and, on the other hand, a transition oxide thickness much lower [8, 9]. To complete this kinetic study, the impact of zirconium hydrides on oxygen diffusion through the oxide scale was investigated using the 18O isotopic marker.

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The techniques of isotopic markers were largely used during these last decades for the study of the oxidation mechanisms. The principle is based on the alternated oxidations : the first oxidation stage is carried out under 16 O2, C16O2 or H216O environment and the second one is performed in similar environments enriched in oxygen 18 (18O2, C18O2, H218O). The 18O penetration profile in the material obtained by SIMS after the last annealing gives actually precious information about the diffusion mechanisms of oxygen when the times of isotopic exchange are sufficiently long [27] and allows to estimate under certain conditions the apparent diffusion coefficient of oxygen in the oxides for short times of annealing [14]. The pre-transition and the post-transition stages were therefore explored by this way with short and long time exposures in H218O enriched enviromnent. The experimental procedure, the samples and their exposure times in the different environments are detailed in Table 3. -Table 3 -

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As indicated in this table, the short times of isotopic exchange in H218O/H216O mixing (20% vol. of H218O) were 6 and 24 hours and the long exposures, 14 and 28 days. Next, samples were analyzed using SIMS technique. 1. Short exposure times

First of all, the 18O SIMS profiles obtained in these conditions were treated in order to know the distribution of this element that we should get if the isotopic exchange had been carried out in a water solution enriched with 100% of oxygen 18. The solution being initially made up of 20 % from H218O and 80 % of H216O, the intensities of the profiles in 18O and 16O presented thereafter were reconstituted in the following way (Eq.1a and Eq.1b) : It

 O   5I  O   I  O 

It

 O   4 I  O   4I  O   I  O 

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18

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16

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18

b

18

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(Eq. 1a) (Eq. 1b)

where It is the intensity of the curves reconstituted by considering an isotopic exchange with water enriched with 100 % of oxygen 18, Ib, the intensity of the rough curves, Inat, the intensity related to the natural percentage of oxygen 18. From these expressions, the 18O atomic fraction in the oxides can be deduced from the following ratio (Eq. 2) :

x18O = It(18O)/(Ib(16O)+ Ib(18O))

(Eq. 2)

Secondly, the pulverization times are converted into abraded depths knowing the oxide thickness deduced from the weight gain. The interface between the metal and the oxide is then positioned at the mid-height of the sum of

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O and 18O signals in the oxide. The results obtained after 24 hours of isotopic exchange on reference and prehydrided samples oxidized in light primary water in the pre transition stage are shown in Figure 6.

-Figure 6 -

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Each curve describes a diffusion profile from the outer surface into the bulk of the oxide. The 18O fraction in the bulk of the film is clearly higher in the scale formed on the prehydrided samples, typically, around 2 % of oxygen 18 O. This result may indicate a higher 18O content in the short-circuits (dislocations, grain boundaries, cracks, pores…) of the oxide scale grown on the hydrogen containing samples. A higher site density or volume fraction of short-circuits in the films formed on pre-hydrided specimens could explain such a difference with reference Zircaloy-4 samples. An estimation of the apparent diffusion coefficient of 18O was calculated by fitting the 18O SIMS profile (dotted lines in Figure 6) with the second Fick’s law analytical solution (Eq. 3) adapted for a semiinfinite domain:   x  x18O  x s  (xs  x0 ).erf   2 Da t   

(Eq. 3)

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with the boundary conditions x18O = xs for x = 0 (outer surface), xO18 = x0 for x = ∞ (away from the diffusion region), t, the time of isotopic exchange and X, the oxide depth. Using the previous equation, the 18O apparent diffusion coefficients have been calculated for reference and 150 ppm pre-hydrided samples corroded during the first stage in H216O up to around 1.2-1.3 µm thicknesses. The fitting parameters, xs, xO and D are given in Table 4.

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After 24 hours of isotopic exchange, the apparent diffusion coefficients of 18O is around twice for the prehydrided samples compared to the reference Zircaloy-4. The value obtained on the reference material is consistent to those found in the literature data [28]. The ratio of 18O diffusion coefficients between the prehydrided and the reference samples can be compared with that of the corrosion rates calculated in the table 3 by derivating the power law obtained by fittings. For the same oxide thickness of 1.2 µm, the ratio of the corrosion rates between the pre-hydrided and the reference samples is also increased by a factor 2, approximately, which is very close to the ratio of 18O apparent diffusion coefficients. As the corrosion rate is proportional to the oxygen diffusion flux through the oxide scale, the accelerated corrosion rate observed on the pre-hydrided samples is likely due to a higher oxygen diffusion coefficient rather than a higher oxygen concentration gradient. Finally, this result supports the hypothesis that, on pre-hydrided samples, the oxide growth is limited by the oxygen diffusion through the scale, as for the reference Zircaloy-4 [14, 16]. 2. Long exposure times

As mentioned before, long time exposures enable us to determine the predominant diffusion mechanism in the oxide layer considering the typical profiles schematized in the publication of S. N. Basu [27]. The figure 7 shows, for each material oxidized during the pre-transition stage, the SIMS profiles of 18O atomic fraction obtained after long times of isotopic exchange, .

-Figure 7-

After 14 days, the SIMS profile of oxygen 18 atomic fraction obtained on the reference material is nearly linear. In agreement with the Wagner's theory, the oxygen concentration gradient is thus quite constant in the oxide. Beyond this exposure time, the 18O atomic fraction increases close to the metal/oxide interface meaning a mixed diffusion process through the bulk and grain boundaries during the oxidation of the reference material. On the other hand, a huge accumulation area of oxygen 18 is observed near the hydride/oxide interface. According to the work of S. N. Basu et al. [27], this 18O profile is a typical one of a predominant diffusion process through the short-circuits (like grain boundaries, dislocations and so on…) of the scale formed on the hydride phase. As a consequence, the new oxide formed on the massive hydride is only composed of oxygen 18 after 28 days of exchange.

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As shown on the figure 8 presenting all the SIMS profiles of 18O atomic fraction, the evolutions of the posttransition diffusion profiles are also very different between the two materials. First of all, it appears clearly that, for a given exposure time, the 18O amount integrated in the oxide layer formed on the hydride phase is much greater than that in the reference material. These post-transition profiles independently of the initial hydrogen concentration in the specimen present oxygen accumulation zones located close to the thicknesses corresponding to kinetic transitions: about 2 µm for the reference samples [14, 17] and about 1.3 and 2.6 µm for the prehydrided specimens (Figure 8a and 8b). For the reference Zy4 (Figure 8a), the oxide layer can be clearly divided into two parts, the outer one and the inner one. Contrary to what is generally assumed concerning the external part of the oxide, this one seems obviously to keep a protective character taking into account the oxygen diffusion profile evolution in this part as a function of the exposure times. In the inner part, oxygen 18 is accumulated at the entrance of the inner layer and progressively diffuses through this scale. Moreover, it seems that a small amount of oxygen 18 has a relatively high penetration depth and arrives not so far from the metal/oxide interface after a short isotopic exchange (6 hours). This part of oxygen 18 diffuses probably through the short-circuits existing in the whole scale.

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Concerning the hydrided material (Figure 8b), we observe that a relatively high quantity of oxygen 18 (~5%) reaches the internal interface even for short exposure times. The relatively uniform level of oxygen 18 inside the scale grows progressively with the time of exchange which means that the oxide layer is not very protective and is probably not a scale limiting the oxygen transport. The oxygen 18 accumulation close to the hydride/oxide interface increases with exposure times (up to 70% after 28 days isotopic exchange). As previously, the predominant oxygen diffusion way occurs obviously through the short-circuits in the scale formed on the hydride phase. These results obtained in the post-transition stage confirm the interpretation, mentioned above, concerning the higher volume fraction of short-circuits and/or a significantly greater diffusion coefficient through the shortcircuits of the corrosion films formed on pre-hydrided specimens.

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In conclusion, this study evidenced that the oxygen diffusion flux in the oxide formed on zirconium hydrides is much higher than that estimated in reference Zy4 oxide. Apparent oxygen diffusion coefficient in the oxide formed on pre-hydrided materials is much greater compared to that in the scales formed on Zy4. The oxidation rate increase deduced from the corrosion tests seems to be due to the variations of the oxygen diffusion coefficient in the oxides. Finally, the preferential diffusion way of oxygen occurs through the short-circuits of the oxide scale formed on hydrided materials while a mixed diffusion process controls the corrosion kinetics of Zy4.

Characterization of the materials by SEM and TEMs

After the corrosion tests, the pre-hydrided and reference samples were observed using SEM and TEM in order to characterize their oxide scale. To be comparative, the samples were oxidized, respectively, 7 days and 14 days so that their oxide thickness was close to 1.3 µm and in the pre transition stage for both. The description of the samples is presented in Table 5.

-Table 5-

The micrographs of oxide films formed on pre-hydrided and reference samples after 7 and 14 days of corrosion at 360°C are presented in Figures 9, 10 and 11. First of all, it was checked that the hydride phase was still present under the oxide film after the corrosion test in order to be sure that the scale was only formed on the hydride phase during the corrosion exposure (Figure 11). Secondly, according to the Figures 9a and 9b no change in microstructure is evidenced at the TEM scale. The oxide actually remains columnar with similar grain orientation and average size of the columns (about 200 nm in length x 30 nm in diameter). These observations lead to the same conclusion as in the literature [9]. -Figure 9a,b,c,d On the contrary, a small difference was observed according to the SEM micrographs. The outer part of the oxide scale formed on the pre-hydrided sample shows a significant porosity adjacent to small grains on a 100 to 200 nm thickness region as evidenced on the Figure 9c and 9d.

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Finally, the presence of initial hydrogen at the surface of the specimen does not involve a change of the "bulk" oxide microstructure (inner part of the oxide but induce a microstructure modification of the 200 nanometer outer oxide film compared with reference Zircaloy-4 corroded in the same conditions. The interface between the oxide film and the underlying metal (hydrided or not) was also observed by TEM. As evidenced by the figure 10e, dispersed hydrides appeared in α-Zr grains of the reference Zircaloy-4 sample. These hydrides result from the precipitation of hydrogen coming from the oxidation reaction of the metallic matrix with water.

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-Figure 10a, b,c,d,e –

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When the metal is previously hydrided, it seems that there is a higher density of micro-cracks in the oxide but it could be due to the sample preparation. The most important feature is that an additional phase obviously appears between the ZrO2 oxide film and the rim of massive hydrides because of a darker contrast compared with adjacent phases observable on STEM image of the Figure 11a. As shown in the Figure 11b, this phase is composed of grains with dimensions varying from 200 to 300 nm and their crystallographic structure corresponds to the Zr3O sub-oxide according to the electron diffraction pattern. In literature data, several authors studied this complex region of metal/oxide interface and Zr3O sub-oxides were sometimes but not systematically observed on Zircaloy-4 without hydrogen [29, 30, 31]. Nevertheless, even if it is impossible to conclude on whether this phase is always present in this region during corrosion of Zircaloy-4, the following hypothesis can be advanced : in the experimental conditions of this study, the corrosion mechanism of massive hydrides is apparently not the same as that of Zircaloy-4. -Figure 11a, b,c,d,e -

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In conclusion, for a same oxide thickness in the pre-transition stage, the characterization of pre-hydrided and reference samples revealed no impact of hydrides on the grain morphology in the inner part of the oxide scale, as previouly observed [9]. However, the microstructure of the 200 nanometer outer oxide part of the hydrided material seems to be more disturbed (presence of pores and small grains) and less columnar compared to the reference Zircaloy-4. A higher density of micro-cracks has also been observed inside the oxide film formed on the hydride phase and a new phase indexed as the sub-oxide Zr3O located at the interface between zirconia and the massive hydride phase -ZrH1.66 has been identified.

3.4

Results of XRD analyses

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To reinforce the previous conclusions, the corrosion films formed on two pre-hydrided samples and one reference material in the post-transition regime were analysed using micro X-Rays Diffraction technique with synchrotron radiation. The oxide thicknesses for the hydrided samples and the reference one are 3.4, 7.4 and 3.6 µm and only the highest oxide thicknesses on both types of materials are presented here. The local X-ray analysis allowed us to build the profile of the present phases from the matrix towards the external interface (oxide/water). In order to identify each diffraction peak, the following ICDD files of the phases likely to be present in the samples were used, the 037-1484 file corresponding to the monoclinic zirconia, the 042-1164 describing the quadratic form of zirconia, the 005-0665 for the α-Zr hcp phase and the 034-0649 corresponding to the cubic hydride centered faces δ-ZrH1.66.

3.4.1

Hydrogen free Zy4

Figure 10 shows the superposition of the diffractograms as well as the indexing of the peaks related to the phases previously mentioned. -Figure 12 Some comments can be made about the figure 12. First of all, we can notice that the (10.0) and (00.2) peaks of zirconium α are present only on the three first scans, which enables us to position the metal/oxide interface in a rather satisfactory way. Whereas the (00.2) line intensity corresponding to the basal plane of α-Zr is very weak thanks to the texture effect (see Figure 3), the one due to the prismatic plane (10.0) is a little more intense. In addition, it is worth noting that the peak corresponding to the pyramidal plane (10.1) is always present, whatever the beam position. Two factors can explain their presence : its higher factor of structure and its position with a larger angle 2θ, for which the penetration of the beam under the surface is deeper. Lastly, one can emphasize that the simultaneous indexing of the three peaks of metal by means of an adjustment of the cell parameters does not

8

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make possible to suitably position the (10.1) peak. We think that the distortion of the network due to the oxygen enrichment of the matrix under the interface could be the cause for that. Concerning the hydride phase, the scans raised in the metal part of the reference Zy4 sample show fortunatly the presence of δ-ZrH1.66 phase. While the (111) diffraction peak of this phase appears very clearly, we suppose that the one corresponding to the (200) plane results in the visible shoulder observed on the (10.1) diffraction peak of the metal. Once the metal/oxide interface is exceeded, the main peaks of monoclinic zirconia appear and their intensities grow with the distance from the metal/oxide interface. On the other hand, whereas the (101) diffraction peak of quadratic zirconia is always present, the (002) peak appears only on the first scans raised on the oxide and in a rather high quantity cause of its strong intensity. This observation seems right now to indicate that the quadratic phase is, as expected [32], concentrated close to the metal/oxide interface and that the oxide texture is very marked there. The map presented in figure 11 gathers together all the diffratograms previously exposed and evidenced, in a 3D map, the progressive disappearance of the metal to the detriment of monoclinic zirconia, as well as the presence of hydrides and intense quadratic zirconia lines located near the metal/oxide interface. -Figure 13-

Hydrided material

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3.4.2

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However, this representation raises a new remark. A cyclic evolution of the monoclinic phase peaks is actually highlighted in all the thickness of the studied zone. Knowing that the depth analyzed by the beam is about 3 µm under surface, we can make the assumption that this undulation is related to the roughness of the interface.

-Figure 14 -

Comparison of the quadratic zirconia volume fraction

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3.4.3

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The diffractogramms obtained on the pre-hydrided sample corroded in primary water are presented in a 3D view exposed on the Figure 14. As previously, the persistence of the (10.0) diffraction peak of -Zr all along the scan indicates that residual zirconium is also present in small proportion within the hydrided matrix which potentially underlines the partial dissolution of the hydride during the corrosion test. In addition, a new extra hydride phase is identified in a strong proportion and seems to correspond to the ε-ZrH2 hydride with a centered quadratic structure (ICDD 017-0314). On the other hand, the (002) diffraction peak of the Zr3O suboxide appears clearly on several scans and this phase is located between zirconia and the hydride phase as indicated on the 3D map of the figure 14. The thickness of this Zr3O layer is around 1 µm according to the number of scans showing this phase. Thus, the µ-XRD analyses confirm the presence of this sub-oxide at the hydride/oxide interface which is absent on the reference sample.

A qualitative determination of the polymorphic oxide phases of zirconia using the Garvie-Nicholson formula (so without correction of texture [33]) was performed in order to estimate the evolution of the quadratic phase towards the corrosion film but the results obtained are not presented here. We just want to evidence that, considering the (101) and (002) diffraction peaks of the quadratic zirconia indicated in the figure 15, the oxides formed on the pre-hydrided samples show obviously lower proportions of quadratic phase and marked differences of texture in the scale near the metal/oxide interface compared to those in the reference material. -Figure 15 -

In conclusion, the diffraction peaks confirmed the presence of the sub-oxide Zr3O between the zirconia and massive hydride phase, evidenced significant differences of cristalline phase proportion and texture close to the metal (or hydride)/interface and revealed the appearance of a new hydride phase indexed as ε-ZrH2 located beneath this suboxide.

3.5

Distribution of the hydrogen from hydride phase after corrosion

During the corrosion process, once the limit of solubility in temperature is reached and the hydrogen is precipitated as zirconium hydride, one issue of interest is what become the hydrogen coming from hydrides when the oxidation front moves forward by tranforming the hydrided phase in oxide. Three assumptions can be formulated in this case : 1. dissolution/diffusion/re-precipitation of the hydrogen in the hydrided matrix ahead of the oxidation front ;

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2. Trapping in zirconia; 3. Release in the enviromnent. To answer this issue, corrosion tests in D2O (pure to 99,9%) on hydrided and reference Zircaloy-4 samples were carried out. Table 6 shows the experimental matrix performed during this study. -Table 6-

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The hydrogen profile obtained on the reference sample after the corrosion test in D2O is compared to that in the hydrided matrix on the figures 16a and 16b. -Figure 16a-

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-Figure 16b-

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On the figure 16a, a very strong hydrogen intensity in the metal part of the hydrided Zircaloy-4 sample shows the reminiscence of the hydride phase after the corrosion test. Typcally, it is around 100 times higher than the initial contents of hydrogen in the Zy4 matrix. In addition, the hydrogen profile in the oxide formed on the hydrided sample presented in the figure 16b does not show any accumulation of this specy and is aprroximately the same as that obtained on the reference material. The hydrogen concentrations in the oxides seem to be apparently independent on the material analysed which means that the hydrogen from the hydrides is not integrated in the scale. Knowing that the hydrogen content in the metal is around 20 ppm (mass) and supposing that the ionization yield is comparable in the oxide and the metal, that means that the hydrogen contents in the oxides are rather weak. Deuterium profiles in the scale were also analysed by SIMS, as shown on the figure 17.

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-Figure 17-

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Deuterium concentration in the oxide formed on pre-hydrided matrix is clearly higher and the global amount of deuterium absorbed by the scale is around twice greater than in the reference material. It would seem, according to this observation, that there would be more deuterium trapping sites in an oxide formed on an hydrided matrix.

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In conclusion, the hydrogen initially present in the hydride phase is not incorporated in the oxide and mainly pushed ahead of the oxidation front. Thus hydrogen stays in the hydrided matrix.

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The table 7 sums up the results and the conclusions obtained during this work Table 7

4

Discussion

First of all, we have to remember elements from the literature date that we have to take into account for the description of the corrosion mechanism [15, 34]. Especially, it has been reported that the intermetallic precipitates such as Zr(Fe, Cr) 2 play a role in the hydrogen pick-up process [35]. These second phase particles in a metallic state close to the metal/oxide interface would be cathodic sites for hydrogen according to Ramasubramanian et al. [36] and "windows" or "bridges" for hydrogen to cross the metal/oxide interface [37]. Moreover, L. Aufore showed during her PhD [38] that the oxide layer can be divided into two parts, an outer one, quite permeable to hydrogen and an inner part where it would be more difficult for hydrogen to diffuse. In the mechanism exposed on the figure 18, which describes both oxidation and hydrogen uptake, we consider an adsorption step of water (1) and a reaction of the adsorbed water molecule with a vacancy on the surface (2). The protons produced during the last step can either be reduced and desorbed (3-4) or be incorporated in the oxide (5). Oxygens and protons can diffuse either in the bulk or in the grain boundaries of the oxide layer [39]. The metal is oxidized at the metal/oxide interface and hydrogen is reduced and incorporated in the Laves phases (metallic state) (7). Finally, hydrogen is accumulated in the zirconium matrix (8) and precipitates as zirconium hydrides (-ZrH1.66) when the solubility limit is reached (9). -Figure 18 -

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The experimental results obtained during this work showed that the corrosion mechanism of hydrided Zy4 was undoubtedly different because of the presence of the Zr3O phase between zirconium hydride and oxide and a lower quadratic zirconia volume fraction in the oxide scale. Moreover we proved that the corrosion rate increase is mainly caused by the growth of the apparent oxygen diffusion coefficient in the oxide layer formed on prehydrided material compared to reference Zy4. Secondly, the preferential way of oxygen transport is clearly through the grain boundaries (short-circuits) and the diffusion flux by this way is very high compared to that in the oxide formed on hydrogen free Zy4. Finally we have to take into account the fact that, during the corrosion process, the hydrogen of the hydride phase stays inside this phase while the oxidation front moves forward. Figure 19 shows a scheme of the corrosion mechanism of the -hydride phase considering the elements indicated previously. Next to it, the 18O diffusion profiles obtained after 14 days of isotopic exchange respectively for hydrogen free Zy4 and hydride phase are exposed on this figure and evidence the huge 18O accumulation located around the hydride/oxide interface. As explained previously, this accumulation is due to the very fast diffusion process of oxygen through the strong density of short-circuits, indicated by a broad arrow on the scheme of the corrosion mechanism. This scheme also describes the dissolution process of hydrogen from the -hydrides into the underlying Zy4 metallic matrix as well as the hydrogen accumulation into the hydride phase when the oxidation front progresses. This process leads to the formation of -ZrH2 hydride. -Figure 19 -

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The difference of corrosion mechanism between hydrogen free Zy4 and zirconium hydrides brings up many questions. Especially how can be explained the marked difference of microstructure (density of cracks) and crystal phases between the oxides formed on these materials ? The second interesting issue is why the oxygen diffusion flux is so high in the oxide formed on hydrided Zy4 ? 1. The TEM observations as well as the µ-XRD analyses attested the presence of a Zr3O suboxide between the massive hydride and zirconia. This phase is actually an intermediate phase formed during the corrosion process of the hydrides but it is not a final product of the oxidation of hydrides. Its presence shows that the formation of this phase on the surface of zirconium hydride is more favourable than that of zirconia. This can be explained by considering the Pilling Bedworth Ratios. Indeed, using the ICDD file datas of the -ZrH1,66, ZrO2, and Zr3O, phases the PBR is equal to 0.96 between the hydride and the suboxide against 1.45 in the case of ZrO2ZrH1.66 interface. This means that the compressive stresses generated at the -ZrH1.66/Zr3O interface are lower and due to the less ductile behaviour of the massive hydride (compared to the Zy4 matrix) and its weak aptitude to adapt the constraints, the growth of the Zr3O phase on the hydride phase is favoured. However, the presence in small proportion of suboxide is not excluded on the alloy reference (this phase was already observed by other authors [29]) but it remains negligible compared with that formed on hydrided Zy4. This oxide layer also absorbs a greater amount of hydrogen coming from the oxidizing medium according to the corrosion tests in pure D2O. By referring to several publications found in the literature [38], this accumulation could promote the allotropic transformation from quadratic to monoclinic zirconia and justify the lower proportion of quadratic phase on these hydrided samples [12]. This polymorphic transformation could also explain the density of micro-cracks inside the oxide. 2. To explain the difference of oxygen diffusion process, some hypotheses can be suggested, specially the grain size in the oxide, the cracks and the structure of the grain boundaries. The grain size in the oxides are similar in both cases as evidenced by SEM and TEM analyses. The oxide cracks can act either as impassable barrier or short-circuits if they are connected. As the oxide thickness is relatively thin, the porosity is likely not opened to the medium. So these hypotheses cannot explain the difference of the oxygen diffusion flux. The second hypothesis that we can put forward is that intergranular spaces in oxide formed on this hydrided material is broader than that of hydrogen free Zy4. This idea can be supported by the following example coming from the literature. Figure 20 shows actually atomic order in a symmetrical flexion grain boundary around [001] for various orientations between two adjacent grains [41]. -Figure 20 The preceding figures clearly highlight the variations of interatomic spaces between two grains for various types of grain boundaries. We can therefore speculate that the grain boundaries in oxide formed on hydrided material can be regarded as general grain boundaries, presenting broad intergranular spaces and so numerous available sites allowing a fast migration of oxygen in this space. For reference Zy4, interatomic space in the grain boundaries could be more narrow and the number of available sites more reduced because of either a weaker disorientation or a better coherence from a crystallographic point of view between two adjacent grains as in a singular grain boundary.

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Conclusions

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Finally, to estimate the hydride effect on the high burn up acceleration, we need to specify that the hydrogen concentration of hydride rims in the fuel rod is 4000 to 5000 wt ppm [42]. This content is around the third of the hydrogen concentration obtained by cathodic charging. So we can approximately consider that the matrix at high burn-up is composed of around 30 % of delta hydrides. Assuming that the global corrosion rate of the cladding is the sum of corrosion rates of each phase balanced by the volume fraction, we obtain only a corrosion rate ratio due to the presence of hydride phase comprised between 1.3 and 2, value clearly lower compared to the phase III accelerating factor which is around 10. Consequently, the hydriding process of the matrix results in an increase of the corrosion rate and contributes to the huge acceleration observed in PWR. It is however difficult to estimate the part of this contribution because we do not accurately know the temperature increase at the metal/oxide interface during the in-core lifetime of the claddings and its effect on the corrosion rate.

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The impact of hydrides on Zircaloy-4 corrosion at 360°C in primary water has been evaluated through corrosion tests and characterizations on samples catholically charged with hydrogen compared to reference specimens. The results of these investigations are the following ones: 1. The presence of a massive hydride phase at the surface of Zircaloy-4 samples clearly increases the corrosion rate in simuated PWR conditions. As for the reference Zircaloy-4, the oxide growth on prehydrided specimens seems to be limited by the oxygen diffusion through the scale. 2. A slight difference has been observed with SEM in the structure of the oxide films. Indeed, on the pre-hydrided samples, the external part (around 200 nm under the surface) seems to be more porous and associated with small equiaxed grains. Beyond this thickness, the oxide microstructure observed with TEM (columnar grains of about 200 nm in length and 30 nm in diameter) is similar for both types of samples. 3. When the oxide forms on a hydrided matrix, an additional phase appears at the interface between ZrO2 and δ-ZrH1.66, which has been identified as the sub-oxide Zr3O. Its length extends from 200 to 1 µm depending on the kinetic regime (pre- or post-transition). This phase was not observed on reference samples oxidized and observed in the same conditions. The corrosion mechanism of pre-hydrided samples is actually not the same as that of reference Zy4. 4. Isotopic exchanges carried out with oxygen 18 on these materials revealed an acceleration of oxygen diffusion in the scales formed on pre-hydrided samples. The average apparent diffusion coefficient at 360°C reached 2.8x10-15 cm²/s instead of 1.6x10-15 cm²/s for reference Zircaloy-4. The diffusion coefficient ratio equal to 1.8 is moreover quite close to the oxidation rate ratio. 5. According to SIMS profiles obtained after short and long re-oxidations of the oxide scales formed on pre-hydrided and reference samples, the faster enrichment of oxygen 18 at the (pre-hydrided) metal/oxide interface leads to suppose a greater contribution of oxygen diffusion via the grain boundaries for the oxide scales grown on the pre-hydrided samples. This observation could be due to a higher disorientation of the oxide grains formed on pre-hydrided specimens compared with reference Zircaloy-4. But this interpretation has to be confirmed by TEM analyses. The aim of this study was to explore the role of hydride precipitation on the corrosion behavior of Zircaloy-4. Based on the results mentioned above, the hydriding process of the matrix results in an increase of the corrosion rate and contributes to the huge acceleration observed in PWR. But we do not know if this factor is the only one responsible for this acceleration.

ACKNOWLEDGMENTS

The authors want to thank AREVA NP/CEZUS for supplying material, M. Blat from EDF for sharing the hydrogen cathodic charging protocol and the hydrogen fusion measurements, Pr. A. Motta to have enable us to use the APS synchrotron and Barry Lai, Zhonghou Cai, Advanced Photon Source, Argonne National Laboratory, USA for the µ-XRD analyses.

REFERENCES 1.

Garzarolli, F., Stehle, H., and Steinberg, E., Zirconium in the Nuclear Industry: 11th International Symposium, ASTM STP 1295 (1996) 12-32.

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Bossis, P., Pêcheur, D., Hanifi, K., Thomazet, J., and Blat, M., Zirconium in the Nuclear Industry: 14th International Symposium, ASTM STP 1467 (2005) 494-525.

3.

Garde, A. M., Pati, A. R., Krammen, M. A., Smith, G. P., and Endter, R. K., Zirconium in the Nuclear Industry: 10th International Symposium, ASTM STP 1245 (1994) 760-778.

4.

Iltis, X., Lefebvre F., and Lemaignan, C., Zirconium in the Nuclear Industry: 11th International Symposium, ASTM STP 1295 (1996) 242-264.

5.

Barberis, P., Zirconium in the Nuclear Industry: 13th International Symposium, ASTM STP 1423 (2002) 3358.

6.

Garde, A. M., Zirconium in the Nuclear Industry: 9th International Symposium, ASTM STP 1132 (1991) 566-594.

7.

Kido T. "A study on enhanced uniform corrosion of Zircaloy-4 cladding during high burnup operation in PWR's" in proceedings of the 6th International Symposium on Environmental Degradation of Materials in Nuclear Power Systems (1993) 449-554.

8.

Blat, M., and Noel, D., Zirconium in the Nuclear Industry: 11th International Symposium, ASTM STP 1295 (1996) 319-337.

9.

Blat, M., Legras, L., Noel, D., and Amanrich, H., Zirconium in the Nuclear Industry: 12th International Symposium, ASTM STP 1354 (2000) 563-591.

us

cr

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2.

an

10. Oskarsson, M., Ahlberg, E., Södervall, U., Andersson, U., and Pettersson, K., Journal of Nuclear Materials 289 (2001) 315-328.

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11. Shemet, V. Zh. Lavrenko, V. A., Teplov, O. A.and Ratushnaya, V. Z., Oxidation of Metals 38 (1992) 8998. 12. Y. S. KIM et al., Zirconium in the Nuclear Industry: 10th International Symposium, ASTM STP 1245 (1994) 745-759.

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13. Kim, Y. S., Rheem, K. S., and Min, D. K., Zirconium in the Nuclear Industry: 13th International Symposium, ASTM STP 1423 (2002) 247-296. 14. Dali, Y., Tupin, M., Bossis, P., Pijolat, M., Wouters, Y. Jomard, F., J. Nucl. Mat. 426 (2012) 148-159.

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15. Motta, A. T., Yilmazbayhan, A., Comstock, R. J., Partezana, J., Sabol, G. P., Lai, B., and Cai, Z., Zirconium in the Nuclear Industry: Fourteenth International Symposium, ASTM STP 1467 (2005) 205-233.

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16. Tupin, M., Pijolat, M., Valdivieso, F., Soustelle, M., Frichet, A., and Barberis, P., Journal of Nuclear Materials 317 (2003) 130-144. 17. Cox, B., Journal of Nuclear Materials 336 (2005) 331-368. 18. Dollins, C. C., and Jursich, M., Journal of Nuclear Materials 113 (1983) 19-24. 19. Guerain, M., Duriez C., Grosseau-Poussard, J. L. and Mermoux M., "Review of stress fields in Zirconium allys corrosion scales ", Corros. Sci. 95 (2015) 11-21. 20. Eloff, G. A., Greyliong, C. J. and Vijoen, P. E., Journal of Nuclear Materials 202 (1993) 239-244. 21. Sabol, G. P., and Dalgaard, S. B., Journal of Electrochemical Society 112 (1975) 316-317. 22. Bryner, J. S., Journal of Nuclear Materials 82 (1979) 84-101. 23. Parise, M., Sicardy, O., and Cailletaud, G., Journal of Nuclear Materials 256 (1998) 35-46. 24. Cox, B., Journal of Nuclear Materials 29 (1969) 50-66. 25. Garzarolli, F., Seidel, H., Tricot, R., and Gros, J. P., Zirconium in the Nuclear Industry: 9th International Symposium, ASTM STP 1132 (1991) 395-415. 26. Sabol, G. P., MacDonald, S. G. and Airey, G. P., Zirconium in Nuclear Applications, ASTM STP 551 (1974) 435-447. 27. Basu, S. N., and Halloran, J.W., Oxidation of Metals 27 (1987) 143-155. 28. Cox, B., and Pemsler, J.P., Journal of Nuclear Materials 28 (1968) 73-78.

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29. Yilmazbayhan, A., Breval, E., Motta, A. T., and Comstock, R. J., Journal of Nuclear Materials, 349 (2006) 265-281. 30. Iltis, X., and Michel, H., Journal of Alloys and Compounds 177 (1991) 71-82. 31. Anada, H., Zirconium in the Nuclear Industry: 11th International Symposium, ASTM STP 1295 (1996) 7493.

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32. Bechade, J-L., Brenner, R., Goudeau, P., Gailhanou, Influence of temperature on X-ray diffraction analysis of ZrO2 oxide layers formed on zirconium based alloys using synchrotron radiation, Materials-ScienceForum 404-407 (2002) 803-808. 33. Garvie, R. C., and Nicholson, P. S., Journal of the American Ceramic Society 55 (1972) 303-305.

34. Tupin, M., Pijolat, M., Valdivieso, F., and Soustelle, M., Journal of Nuclear Materials 342 (2005) 108-118.

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36. Ramasubramanian, N., Journal of Nuclear Materials 55 (1975) 134-154.

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35. Hatano, Y., Hitaka, R., Sugisaki, M., and Hayashi, M., Journal of Radioanalytical and Nuclear Chemistry 239 (1999) 445-448. 37. Lelièvre, G., PhD thesis, Institut National Polytechnique de Grenoble (1998).

38. Aufore, L., PhD Thesis, Université de la Méditerranée Aix-Marseille II (1997).

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39. Sundell, G., Thuvander, M. Yatim, A. K., Nordin, H. Andren, H.-O., Corrosion Science 90 (2015) 1-4. 40. Barberis, P., Journal of Nuclear Materials 226 (1995) 34-43.

41. Sainfort, G., "Le rôle des joints de grain dans les matériaux", Ecole d'été de Métallurgie Physique (1984) 3 (in french) 3-40.

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42. Raepsaet, C., Bossis Ph., Hamon, D., Béchade, J.-L., and Brachet, J.-C., Nuclear Instruments and Methods in Physics Research B 266 (2008) 2424-2428.

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Figure Captions : Figure 1 : Scanning electron micrographs of cathodically charged samples of Zy4. (a) Cross section and (b) fractography of specimens charged with a 400 ppm and a 170 ppm global hydrogen content respectively. Figure 2 : Transmission electron micrographs of pre-hydrided samples of Zy4. (a) Overview of a mixed microstructure composed of hydride platelets (A) and hydrides grains (B) indexed as δ-ZrH1.66 phase. (b) The

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diffraction pattern is associated with the dark and the bright fields presented in (c) and (d).

Figure 3 : Schematic of the crystallographic orientation of the thin foil regarding the incident beam.

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Figure 4 : SEM micrography of the thin foil taken coming from the reference Zircaloy-4 oxidized 171 days in primary medium at 360°C.

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Figure 5 : a) Corrosion kinetics of Zircaloy-4 samples during corrosion in primary water at 360°C and 18.7 MPa. The kinetic transition appears around 2 µm and 125 days of corrosion. b) Corrosion kinetics of pre-hydrided and

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reference Zircaloy-4 samples during corrosion in primary water at 360°C and 18.7 MPa.

Figure 6 : SIMS profiles of oxygen 18 atomic fraction obtained on pre transition oxide films formed on reference (Z2_full squares) and hydrided (ZH2_full circles) samples exchanged during 24 hours.

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Figure 7 : SIMS profiles of oxygen 18 after 14 days (,), and 28 days (,) of isotopic exchange on pre transition oxide films formed on reference (Z_squares) and pre-hydrided (ZH_circles) samples. The metal/oxide interfaces for reference and hydrided samples are, respectively, defined by continuous and dotted vertical lines.

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Figure 8 : SIMS profiles of oxygen 18 after (○) 6 hours, (□) 24 hours, (◊) 14 days, and (x) 28 days of isotopic

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exchange on post transition oxide films formed on (a) reference (Z) and (b) pre-hydrided (ZH) samples. Figure 9 : (a) Transmission electron micrographs of the oxide scale formed on reference and (b) pre-hydrided

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samples. The dotted line represents the metal/oxide interface and arrows the oxide growth direction. (c) Scanning electron micrograph of the fractured oxide scale formed on reference and (d) pre-hydrided samples. Arrows point out porosity in the outer part of the oxide scale. Figure 10 : TEM images of the thin foil obtained by FIB coming from a sample of reference Zy4 corroded 28 days, a) STEM image of the thin foil. (b) Bright field and (c) Dark field images of an α-Zr grain just under the metal/oxide interface associated with the electron diffraction pattern (d). (e) Dispersed hydrides in the metallic matrix with another tilt angle. Dotted lines represent the metal/oxide interface. Figure 11 : TEM images of the thin foil obtained by FIB coming from a sample of hydrided Zy4 corroded 14 days (a) STEM image of the thin foil (b) TEM image in bright field of the dark zone between the oxide scale and the hydride phase.c) Bright and e) Dark field of one Zr3O grain associated with the electron diffraction pattern (d). Red dotted lines represent the metal/oxide interface. Figure 12 : Diffraction patterns obtained the sample of Zircaloy-4 oxidized 171 days at 360°C (3.6µm) along the scan from the metal up to the surface of the oxide. The vertical lines indicate the X-ray diffraction peaks of the phases given by the ICDD files. Figure 13 : 3D map of µ-XRD patterns obtained on a reference sample corroded in light primary water at 360°C during 171 days (oxide thickness of 3.6 µm).

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Figure 14 : 3D map of µ-XRD patterns obtained on a pre-hydrided sample corroded in light primary water at 360°C (oxide thickness of 7.4 µm). Figure 15 : Diffraction patterns obtained in the oxide close to the internal interface for the pre-hydrided samples and one reference material . Figure 16 : a) SIMS profiles of hydrogen versus abraded time obtained on reference material (full circles)

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oxidised during 28 days and on hydrided sample (full squares) oxidised during 14 days in deuterium primary medium (b) SIMS profiles of hydrogen versus oxide depth obtained on reference material (full circles) oxidised

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during 28 days and on hydrided sample (full squares) oxidised during 14 days in deuterium primary medium.

Figure 17 : Deuterium SIMS profiles in the oxide formed on hydrided (full squares) and reference (full circles)

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Zircaloy-4, oxidised, respectively, during 28 and 14 days in deuterium primary medium.

Figure 18 : Schematic corrosion mechanism of Zy4 alloy in PWR simulated conditions and reaction steps of the mechanism in the Kröger Vink formalism

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Figure 18 : Schematic of the corrosion mechanism of Zy4 alloy in PWR simulated conditions and reaction steps of the mechanism in the Kröger Vink formalism.

18

O SIMS diffusion

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Figure 19 : Schematic of the corrosion mechanism of an hydride phase forme on Zy4.

profiles obtained after 14 days of isotopic exchange exposure time for hydrogen free Zy4 (continuous line) and hydride phase (empty circles).

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Figure 20 : Atomic structure in symmetrical flexion grain boundaries around [001] a, b : weak disorientation, c :

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high density of sites in coincidence, d : unspecified (39).

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Highlights for Editors, The originality of this work is to study the corrosion behavior of a massive zirconium hydride formed on Zy4. The results obtained are the followings: higher corrosion rate of hydride compared to matrix (Zy4)



higher oxygen diffusion coefficient through oxide formed on hydride



presence of Zr3O phase between hydride and oxide



Hydrogen from the hydride phase is not integrated in the oxide during the corrosion process

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Tables : Table 1 : Chemical composition of Zircaloy-4 (wt %)

Alloying elements

Cr

Sn

O

H, ppm

Zr

0.22

0.11

1.46

0.13

21

Bal.

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Zircaloy-4

Fe

Hydrogen content

Corrosion duration (days)

Final oxide thickness (µm)

50

1.54

-

50

1.52

ZH1

156

14

1.22

ZH3

136

14

1.2

ZH5

315

14

3.37

ZH9

320

14

3.3

ZH10

325

14

3.38

pre transition

n parameter

pre transition

0.47

0.3

0.47

0.3

0.57

0.3

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k parameter

post transition

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Z9 Z10

Kinetic domain

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Name of specimen

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Table 2 : Exposure time of corrosion, final oxide thicknesses, kinetic domains of pre-hydrided and reference Zircaloy-4 samples. Fitting parameters of the power law (x=ktn), oxidation rate at 1.2 µm thickness and corrosion rate ratios compared to Zy4 without hydrogen.

0.55

0.3

0.52

0.4

0.56

0.37

0.54

0.38

dX dt 

X1.2 µm

(µm / d)

dX dt  dX dt 

ZH _ X 1.2 µm

ZNH _ X 1.2 µm

0.015

1.0

0.028

1.8

0.058

3.8

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Table 3 : samples, exposure times, final oxide thicknesses for hydrided and reference materials.

Hydrogen content* (ppm wt)**

H216O exposure

Oxide thickness (µm)*

H218O exposure

Oxide thickness (µm)*

Z1

-

28 d

6h

1.26

-

28 d

1.26

24 h

1.26

-

28 d

1.29

14 d

1.31

-

28 d

1.29

28 d

1.32

kinetic domain

1.26

156

14 d

1.22

6h

1.22

156

14 d

1.22

24 h

1.22

136

14 d

1.2

14 d

1.6

136

14 d

1.2

28 d

1.79

-

171 d

3.59

6h

3.59

Z6

-

171 d

3.59

24 h

3.59

Z7

-

171 d

3.59

14 d

3.75

Z8

-

171 d

3.59

28 d

4.36

Z2 Z3 Z4 ZH1 ZH2 ZH3 ZH4 Z5

Ac

ce pt

Name of specimen

pre transition

post transition ZH5

315

50 d

3.37

6h

3.37

ZH6

315

50 d

3.37

24 h

3.37

ZH7

315

50 d

3.37

14 d

3.64

ZH8

315

50 d

3.37

28 d

4.19

*Hydrogen content and oxide thickness were estimated from weight gain. **This concentration corresponds to a global hydrogen content in the samples which is localized in massive hydrides.

17

Page 18 of 42

Table 4 : Apparent diffusion coefficients of samples in pre transition stage.

specimen

Da (cm²/s) -15

Z2

1.3x10

ZH2

2.8x.10

-15

18

O in the oxide layers formed on reference and pre-hydrided

Xs

X0

0.22

-

0.22

0.02

H content (ppm wt)*

Corrosion duration (days)

Final oxide thickness (µm)*

Z11

-

14

1.26

ZH11

240

7

1.33

Kinetic domain

cr

Name of sample

ip t

Table 5 :. Characteristics of samples observed in TEM and SEM-FEG

pre transition

us

pre transition

*Hydrogen content and oxide thickness were estimated from weight gain.

-

28

~ 200 (erim ~7 µm)

14

oxide thickness (µm)

M

oxidation time (days)

1.2 1.3

ed

[H] (ppm mass.)

an

Table 6 : Materials, oxidation times, oxide thicknesses deduced from mass gain after oxidation in D2O with 2 ppm of lithium and 1000 ppm of boron.

Ac

ce pt

Table 7 : Summary of the results and the conclusions.

18

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Ac ce p

te

d

M

an

us

cr

ip t

Figure 1

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Figure 2

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Figure 3

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Figure 4

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Figure 5a

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Figure 5b

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Figure 6

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Figure 7

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Figure 8a

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Figure 8b

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Figure 9

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Figure 10

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Figure 11

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Figure 12

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Figure 13

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Figure 14

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Figure 15

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i

Figure 16a

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Figure 16b

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Figure 17

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Figure 18

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Figure 19

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Figure 20

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