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New insights in structural characterization of zirconium alloys oxidation at high temperature
Journal of Nuclear Materials 429 (2012) 19–24
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Journal of Nuclear Materials journal homepage: www.elsevier.com/locate/jnucmat
New insights in structural characterization of zirconium alloys oxidation at high temperature Dominique Gosset a,⇑, Matthieu Le Saux b,1, David Simeone a,2, Didier Gilbon b,3 a b
CEA, DEN, LRC CARMEN, CEA Saclay, 91191 Gif/Yvette cedex, France CEA, DEN, CEA Saclay, 91191 Gif/Yvette cedex, France
a r t i c l e
i n f o
Article history: Received 16 December 2011 Accepted 3 May 2012 Available online 23 May 2012
a b s t r a c t Zircaloy-4 samples have been analyzed by grazing incidence X-ray diffraction during oxidation at high temperature (900–1100 °C) in a He–O2 mixture simulating a LOCA atmosphere. We clearly show that the structure of the oxide layer strongly depends on the oxidation temperature, i.e. tetragonal at 1100 °C, mixture of tetragonal and monoclinic at 1000 °C, mainly monoclinic at 900 °C. Moreover, the high temperature phases are not stable during cooling; at room temperature, the oxide is mainly monoclinic. Furthermore, re-heating the samples to the oxidation temperature in neutral atmosphere does not lead to the pristine structural composition. This behavior is explained by considering the martensitic character of the tetragonal to monoclinic transition of zirconia. As a conclusion, the structural and thermo-mechanical properties of the oxide layer as it forms during oxidation at high temperature cannot be deduced from post-mortem analysis. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction The zirconium alloys fuel claddings used in nuclear water reactors oxidize during normal operation or under low probability accidental situations [1,2]. The resulting oxide layer notably affects the thermal exchanges and the mechanical properties of the cladding. More specifically, during a Loss of Coolant Accident (LOCA), the cladding is exposed to steam at high temperature (typically between 800 and 1200 °C), which leads to a relatively fast growth of the oxide layer and a diffusion of oxygen into the metallic substrate [2–5]. Even higher temperatures may be reached during reactivity initiated or severe accidents. The oxidation kinetics of zirconium alloys is logically related to the structural and microstructural properties of the oxide layer, which depends mainly on the oxidation temperature, oxidizing atmosphere or type of alloy. For example, it is known that between approximately 600 °C and 1050 °C, ‘‘breakaway’’ oxidation phenomena can occur during steam oxidation [3–6]. Shortest critical times, at which this transition takes place, are observed around 1000 °C. This phenomenon is mainly characterized by a rapid increase of the oxidation rate (transition from a (sub-) parabolic to a nearly linear kinetics) and a significant hydrogen uptake into the metal when critical oxida-
⇑ Corresponding author. Tel.: +33 1 6908 5857; fax: +33 1 6907 7130. E-mail addresses: [email protected] (D. Gosset), [email protected] (M. Le Saux), [email protected] (D. Simeone), [email protected] (D. Gilbon). 1 Tel.: +33 1 6908 1228. 2 Tel.: +33 1 6908 2920. 3 Tel.: +33 1 6908 2200. 0022-3115/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnucmat.2012.05.003
tion time or oxide layer thickness (which depend mainly on temperature) is reached. One common explanation to this transition is related to the structural properties of the zirconia layer in this temperature range: transformation of the initially formed metastable tetragonal zirconia to the stable monoclinic phase during oxide growth [2–5]. Thus, taking into account the potential consequences on the cladding integrity, it is important to be able to characterize the oxides formed during oxidation at high temperature of zirconium alloys. The phase diagram of the zirconium–oxygen system is well documented [7]. In particular, the zirconia ZrO2 undergoes two structural transitions leading to drastic symmetry changes [8]. The high fluorite like), with a wide stoitemperature phase is cubic (Fm3m chiometry range allowed by high oxygen vacancy concentration. It undergoes a second order transition around 2300 °C to a tetragonal structure (P42/nmc). The second structural change occurs around 1100 °C, from tetragonal to monoclinic (P21/c). The latter is a first order, martensitic-like transition. The structure undergoes strong distortions and volume changes (about 3.5%). A marked hysteresis, around 200 °C wide, is observed (Fig. 1). It is worth noticing that this transition occurs in the temperature range the fuel claddings are supposed to undergo during a LOCA. Complex behaviors are then to be expected. Moreover, the tetragonal phase can be observed down to room temperature in pure zirconia either in nanometric grains or in irradiated materials. In both cases, it has been shown that this stabilization is induced by high internal stresses [9,10]. Such stresses are also responsible for the stabilization of the tetragonal phase, e.g. by point defects [11], doping elements [12] or high pressure [13].
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D. Gosset et al. / Journal of Nuclear Materials 429 (2012) 19–24
10000
analyzed depth (nm)
90% 1000
αc
10
1
2. Experimental procedures The experiment have been performed on a laboratory goniometer equipped with a fixed curved position sensitive detector (INEL CPS-120) allowing recording the X-ray diffraction diagrams simultaneously on a 120° range with a 0.015° resolution. The angular calibration of the detector is performed with a Y2O3 secondary standard. The X-ray source is a classical Cu tube. A Ge (1 1 1) plane monochromator allows first to select the Cu Ka1 reflection and second to obtain a narrow (80 lm width), very parallel beam. Such a setting is then used to perform the analyses in an asymmetric configuration [25]. In this case, modifying the impinging angle of the incident X-ray beam allows an accurate control of the analysed thickness of the material (Fig. 2). The simultaneous recording of a wide 2h angular range allows fast analyses: here, identification of the zirconia phases and semi-quantitative analyses can be obtained within 300 s records. The goniometer is equipped with a HTK-1200 Anton-Paar furnace allowing analysis up to 1200 °C in neutral or oxidizing conditions. A 190°-wide modified window allows sample alignment and analyses on the whole angular range of the CPS-120 detector in grazing incidence beam conditions (GIXRD). The window of the furnace is closed with a 130 lm kapton foil protected with a 20 lm nickel foil to resist to the inner high temperature oxidizing
0.1
1
10
incidence (°)
Fig. 1. Monoclinic-tetragonal transition of pure zirconia (tetragonal volumic fraction) from [8]. : on cooling. j: on heating.
However, observing the actual structure of zirconia as it forms in the conditions of a LOCA remains a quite difficult task. This is why most of the structural characterizations are performed ‘‘post-mortem’’, i.e. either at room temperature [14,15] or after re-heating in neutral atmosphere at representative temperatures [16–18]. It is however expected that, owing to the martensitic features of the monoclinic to tetragonal transition, such an approach is unable to lead to a correct description of the actual phase structure of the oxide as it forms [19]. Furthermore, the very few ‘‘in situ’’ structural analysis reported in the literature have been performed under air (which leads to an accelerated degradation compared to steam due to the formation of oxynitrides [20]) and/or at temperatures lower than 900 °C [21–24]. In this frame, we have launched a program aiming at an accurate structural analysis of the oxide ‘‘in situ’’, i.e. remaining at the temperature of oxidation. In this paper, we present preliminary results obtained through high temperature X-ray diffraction analysis of Zircaloy-4 maintained at high temperature (900–1100 °C) in controlled atmosphere.
1-1/e
100
Fig. 2. Analyzed depth in grazing incidence X-ray diffraction as a function of the impinging X-ray incidence in ZrO2, for two absorption yields (90% and (1 1/e), i.e. the classical absorption length); ac: critical angle for total reflection.
atmosphere. The parasitic diffraction lines generated by those foils do not interfere with the Zr or ZrO2 main lines. The furnace-holder is motorized in order to tune the height (±5 lm) and the incidence (±0.02°) of the sample with regard to the X-ray beam and compensate for the sample holder dilatation at high temperature. A gas circuitry allows measurements under neutral (5.5 N He) or strongly oxidizing (40%He–60%O2) atmospheres. According to the oxidation kinetics reported in [26,27], this He–O2 mixture is expected to be a rather good surrogate to steam (LOCA conditions) and is easier to handle. Although the hydrogen production (and absorption into the metal for instance in the post-breakaway regime) and the enthalpy of the Zr–O reaction are not the same during oxidation under O2 and steam, the structural properties of the growing oxide are expected to be similar. Thanks to the small volume of the furnace chamber, the gas circuitry allows fast gas change (less than 10 s): however, this would induce turbulent flows in the furnace chamber and a strong transient cooling of the sample, up to 100 K. Here, such a transient cooling is prohibited since abrupt, irreversible structural changes are expected. As a consequence, the gas change is made slower and practically obtained in around 1 min. The temperature is measured with a Pt 10%Rh–Pt thermocouple close to the sample. The temperature difference between the sample and the thermocouple has been previously determined placing a secondary K-thermocouple on the back face of a standard sample through the sample holder axis. The difference (less than 15 °C in the whole temperature range in gaseous environments) is then automatically corrected. 3. Results The measurements have been performed on Zircaloy-4 material (Zr–1.3%Sn–0.2%Fe–0.1%Cr–0.13%O, in wt.%). The samples are 12 12 mm2 flat squares with a thickness of about 1 mm. The size of the samples allows grazing incidence X-ray diffraction analysis at incidence down to 0.5° without the X-ray beam to propagate over the sample. The flatness is required to make sure the incident angle (and then the analyzed depth) is constant on the whole illuminated surface. The analyses have been performed with a constant 5° incidence angle for the impinging X-ray beam, for which most of the diffracted intensity comes from the first 3 lm (Fig. 2), whatever the diffracted angle. In such conditions, the fingerprint of the X-ray beam on the sample surface is about
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0.9 mm wide, leading to a moderately wide instrumental component FWinstr to the width of the diffracted lines:
Table 1 Theoretical intensities of monoclinic and tetragonal zirconia and zirconium in the (2h) range of interest for phase identification.
FWinstr ¼ ðe=R sin aÞ sinð2h aÞ where e and a are the width and the incident angle of the impinging beam, R is the radius of the goniometer and 2h is the angle of diffraction. The following measurement sequence has been performed on fresh samples. First, the sample is heated under pure helium (heating rate 1 K/s). A diffraction analysis is performed in order to check there is no significant oxidation during heating. Then, the oxygen– helium mixing is introduced. Short (30 s) diffraction analyses are performed in order to observe the oxide formation. When the oxide layer is thick enough to induce the disappearance of the metal diffraction diagram, pure helium is reintroduced. There is no need here to continue oxidizing since the subsequent oxidation would no longer be visible by diffraction analysis, unless phase changes occur in the outer part of the oxide layer due in particular to relaxation effects: this point is not intended to be studied within the scope of this preliminary study. Accurate diffraction analyses (500 s) are then performed from the oxidation temperature down to room temperature (cooling rate 1 K/s then plateaus) by 100 °C steps in order to analyze the oxide composition during cooling. At last, the sample is re-heated under pure helium and analyzed again at the oxidation temperature. The observed diffraction patterns correspond to well crystallized but highly oriented materials (Fig. 3 to be compared to Table 1): for example, the (1 1 1) monoclinic ZrO2 line (2h 31.5°) is nearly invisible. Quantitative analyses are thus impossible. Here, the oxide phase ratios are then roughly estimated by the following expression:
T ð1 0 1ÞT ¼ T þ M ð1 1 1ÞM þ ð1 0 1ÞT þ ð1 1 1ÞM where the (h k l)P are the intensities of the (h k l) lines (when visible) for the P phase (T for tetragonal, M for monoclinic). All other lines are too strongly convoluted to be easily considered. This expression is obviously far from the actual phase ratios but can be used for relative estimates. At the end of the experiments, all the samples show a black, smooth, slightly rounded surface with no spalling or desquamation. Oxide layer thicknesses measured by optical microscopy are close to those obtained elsewhere for cladding tube specimens oxidized under pure steam at nearly atmospheric pressure [2,28]: the
Fig. 3. Diffraction patterns after oxidation recorded at the temperature of oxidation. Extra lines (e.g. at 56°) come from the kapton–Ni window.
(h k l)
2h (°)
I/Imax (%)
ZrO2 M
1 1) (1 (1 1 1) (0 0 2) (0 2 0) (2 0 0)
28.2 31.5 33.9 34.7 35.4
100 68 21 13 15
ZrO2 T
(1 0 1) (0 0 2) (1 1 0)
30.1 34.5 35.1
100 9 12
Zirconium a
(1 0 0) (0 0 2) (1 0 1)
32.0 34.9 36.1
24 25 100
Zirconium b
(1 1 0)
35.8
100
actual geometry (flat squares versus tubes) and atmosphere (He–O2 mixing versus steam) have thus no significant effect here. A peak to peak analysis of the diagrams has been performed limited to the angular range 20–40° which allows phase identification and semi-quantitative phase ratio estimates according to the previous equation. The main results can be summarized as follows (Figs. 3 and 4): As mentioned above, the patterns correspond to well-crystallized compounds, with narrow lines and low background levels. They can be identified to highly textured monoclinic and/or tetragonal zirconia: the cubic phase is not observed (systematic splitting of the observable lines, clearly corresponding to the tetragonal phase). The structure composition of the oxide layer at the oxidation temperature drastically depends on this temperature. At 900 °C, nearly pure monoclinic ZrO2 formed. At 1000 °C a monoclinic-tetragonal mixing is observed. At 1100 °C, pure tetragonal ZrO2 formed. When cooled down to room temperature, all the samples show mainly monoclinic, oriented zirconia. The phase transformation during cooling depends on the oxidation temperature. From 1100 °C, the tetragonal phase abruptly decreases when cooled from 900 °C to 800 °C, then smoothly decreases down to ca. 300 °C. From 1000 °C and 900 °C, a slow decrease is observed from the oxidation temperature down to ca. 300 °C.
Fig. 4. Relative intensity of the tetragonal component of the oxide layer (T/(T + M)) observed on Zircaloy-4 samples after oxidation under O2-He at high temperature then cooling to room temperature then re-heating to the oxidation temperature under pure He.
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When re-heated at the oxidation temperature under pure helium, the tetragonal component remains much lower than observed just after oxidation. Complementary tests have then been performed. A Zircaloy-4 sample has been oxidized in the DEZIROX 1 facility [6,28] in LOCA representative conditions (steam at nearly atmospheric pressure at 1000 °C during 534 s), leading to the formation of a 24 lm thick oxide layer. At room temperature, the X-ray diffraction diagram shows monoclinic oriented zirconia with a minor tetragonal component. After re-heating at 1000 °C under neutral atmosphere, the only modification – excepting lines shift due to thermal expansion – of the diagram was a narrowing of the diffraction lines, most probably due to decreasing in residual stresses since grain growth could hardly occur in such a low temperature range. No noticeable modification of the relative intensities of the former lines or new lines was observed. This is very similar to the observations we made during the re-heating step performed after the 1000 °C oxidation treatment under He–O2 mixture. Another sample has been pre-oxidized in autoclave in Pressurized Water Reactor (PWR) representative conditions (43 days at 360 °C, water at 190 bars with 1000 w.ppm B and 2 w.ppm Li) leading to the formation of a 1.2 lm thick oxide layer. The oxide we observed is mainly highly oriented monoclinic zirconia and does not change when heated at 1000 °C under pure He (Fig. 5). This sample has then been oxidized at 1000 °C then cooled, following the same sequence as above. As previously observed at 1000 °C, a mixture of tetragonal and monoclinic zirconia formed, the ratio of which cannot be precisely determined due to the pre-existing oxide layer. However, it is worth noticing that the ‘‘fresh’’ high temperature oxide layer begins to form only after 1200 s under oxygen (as identified by the disappearance of the zirconium substrate). This is in good agreement with the observations made on pre-oxidized samples subsequently oxidized at 1000 or 1200 °C in a steam environment [29]: in the presence of a pre-oxide layer, the formation of a ‘‘fresh’’ oxide at the metal/oxide interface is delayed (time necessary for oxygen anions coming from the oxidizing environment to diffuse across the pre-oxide layer). Furthermore, a different texture formed after 1200 s at 1000 °C since the (1 1 1) monoclinic line clearly appears (Fig. 5). An analysis at a lower Xray beam incidence (1°, corresponding to a 500 nm analyzed thickness) has been performed, showing that the structure of the upper oxide layer remained unchanged. During cooling, the tetragonal
phase and the (1 1 1) monoclinic line disappear below 750 °C and do not reappear when re-heated under pure He at 1000 °C.
4. Discussion 4.1. X-ray diffraction analysis The method we use here allows a qualitative analysis of the phase change in the materials, as evidenced in Fig. 4. However, accurate structural description with methods such as Rietveld analyses cannot be obtained: first, such a derivation would require longer records with more accurate diagrams and, above all, texture analyses would be previously required since the relative intensities of the lines are far from those of the isotropic phases. Moreover, it is worth noticing that one of the main advantages of the asymmetric configuration we use here, when compared to the classical symmetrical ones, is that the material thickness which is analyzed does not depend on the given diffraction lines: no volume correction are then to be performed even for the first steps of oxidation when the metal substrate is still visible. However, performing structural refinements would require specific geometrical corrections of the texture effects (Fig. 6). The reason is that on a classical Bragg– Brentano configuration, the diffraction vector (bisecting direction between the incident and the diffracted beams) is fixed (OA in Fig. 6, perpendicular to the sample surface) and the texture correction depends only on the angle between the diffraction planes normal and this vector. But, on the asymmetric configuration, first, a given (h k l) line corresponds to diffraction domains which have a different orientation to the sample surface than in the symmetrical configuration, and second, this direction turns according to h (OA then OA0 in Fig. 6), and then the texture correction would have to be accordingly modified for each line. For instance, different corrections would have to be applied to the harmonic lines of the same diffracting planes (such as (h, k, l), (2h, 2k, 2l), and (3h, 3k, 3l)) since different populations of diffracting domains, with different directions in the sample, have to be considered. Such a correction is not implemented in the classical structural refinement programs (e.g. Rietveld method).
θ θ
θ
θ
Fig. 5. Diffraction diagrams after heating at 1000 °C under He (the line at 35° and the shoulder at 37° correspond to the zirconium substrate), after 3000 s at 1000 °C under O2–He and after cooling at 750 °C under He for a sample pre-oxidized in autoclave at 360 °C under representative PWR conditions.
Fig. 6. Texture correction in Bragg–Brentano and asymmetric configurations. In the asymmetric setup, a given (h k l) line corresponds to different grains with different directions as compared to the symmetric setup and different texture corrections would have to be applied to the harmonics of the same (h k l) family: the SOD and SOD0 X-rays correspond to different populations of diffraction domains.
D. Gosset et al. / Journal of Nuclear Materials 429 (2012) 19–24
The resolution of the diagrams would allow cell parameters estimations as long as unambiguous line indexation is possible: this condition is certainly fulfilled for the tetragonal phase, but a careful observation of the diagrams would have to be performed for the monoclinic phase. The main reason of this limitation is here again the high orientation of the diffracting grains leading to ambiguous line indexation in the large h range. Now, the cell parameters values that could be derived may have complex evolutions since they depend on both the temperature (thermal expansion) and the stresses induced by the metal substrate and the relaxation of the two zirconia phases. It is also well known that analyzing the evolution of line widths along a diffraction diagram can lead to estimations of structural parameters of the material, size of diffraction domains and microstrains. However, the line widths can here be determined only on the first part of the diagrams, up to about 60° (2h): on the high angle part of the diagrams, the lines superposition and the low signal over noise ratio make such an analysis inaccurate. We then observe a significant broadening of the lines as the samples are cooled (then a narrowing when re-heated), specifically for the tetragonal phase. However, it is here impossible to attribute this broadening to a decrease of the size of the diffraction domains or to an increase of the internal microstrains. Such an analysis could be tentatively performed by the well-known Hall–Williamson method but it requires the use of lines over a larger angular range than the one available here. As a result, it is impossible to discriminate between grain fragmentation (as observed in ion-irradiated zirconia [30]) and higher microstrains. It is however observed that the materials are relatively isotropic: for example, the (2 0 0), (0 2 0), (0 0 2) lines of the monoclinic phase of zirconia have close widths, which means that the micro-strains or the dimensions of diffracting domains do not significantly depend on the crystal direction. 4.2. Zirconia formation and phase transitions As mentioned above, the oxidation steps have been performed until the diffraction lines of the metallic substrate disappear, in order to avoid a possible oxide desquamation. This corresponds here to an oxide thickness of a few micrometers (Fig. 2), which is certainly lower that the critical thickness at which the so-called breakaway oxidation phenomenon occurs. By continuing the oxidation, initial existing constraints (compressive stress, misorientations, etc.) may be relaxed and a partial tetragonal to monoclinic conversion may occur in the growing oxide layer. These mechanisms are often postulated to explain the breakaway oxidation. Although the thickness analyzed by the ‘‘in situ’’ XRD method presented in this paper is only a few micrometers, one can expect that these transformations in the oxide may be visible within the outer part of the oxide layer. This point is to be analyzed in a further study. Those results clearly show that in all investigated cases, the zirconia phases formed at high temperature can be observed neither at room temperature nor when re-heated at high temperature. This can be qualitatively explained by the properties of the zirconia phase diagram. First, the higher the oxidation temperature, the higher the tetragonal phase content: the compositions we estimate are similar to those observed on the cooling branch of the hysteresis in Fig. 1 at the corresponding temperatures. Then, cooling the material allows the normal tetragonal to monoclinic transition to take place, by coarsely following this cooling branch. The differences between the curves in Figs. 4 and 1 may be attributed to stresses induced by the pre-existing oxide layer and the zirconium substrate or by different microstructures. The differences in residual tetragonal contents we observe at room temperature in Fig. 4 could also be attributed to different microstructures (grain size, micro-strains) that allow the stabilization of different tetragonal proportions or different preferred orientations. Then, re-heating
23
the cooled samples to the oxidation temperatures cannot lead to the reverse monoclinic to tetragonal transition since the temperature remains lower than the heating branch of the hysteresis (around 1200 °C). And even if such a reverse transition would occur, the obtained microstructure could not correspond to the pristine one. It then clearly appears that the transitions we observe are irreversible. This explains why the composition of the oxide when reheated at the oxidation temperature is strongly different from its composition during oxidation. Strong inferences can then be drawn from this study: The structural observations performed at room temperature after high temperature oxidation are never representative. Re-heating a sample to the previous oxidation temperature does not lead to the pristine microstructures, even if the same phases are observed. 5. Conclusions We have performed X-ray diffraction analysis during the oxidation under a mixture of He and O2 of Zircaloy-4 at high temperature (900–1100 °C) then during cooling and subsequent reheating in neutral atmosphere. We clearly show that any study based on post-mortem observations, whatever at room temperature or after re-heating, cannot be representative of the actual state of the material during the high temperature oxidation: the phase composition and microstructure of the oxide are irreversibly changed during cooling. This has strong inferences on the supposed structure of the oxide, but overall on its thermo-mechanical and microstructural properties as estimated from such post-mortem observations. Indeed, the tetragonal to monoclinic transition is associated with large distortions and volume changes. During cooling, this could for example induce (in addition to the effects of thermal stresses due to differences between thermal expansion coefficients of the oxide and the underlying metallic layers) an irreversible and non representative micro-fragmentation or deformation of the materials under consideration. Further experiments are in progress in order to determine the microstructural properties of the high temperature oxide and the effect of such structural transitions first on the oxidation kinetics of zirconium alloys and second on their consequences on the thermo-mechanics of the oxidized claddings. Acknowledgments The authors thank AREVA-NP for providing the materials. They are also highly indebted to Marc Tupin (CEA) for preparing the samples and performing the pre-oxidation treatments and Valérie Vandenberghe (CEA) for fruitful discussions. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
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New insights in structural characterization of zirconium alloys oxidation at high temperature Dominique Gosset a,⇑, Matthieu Le Saux b,1, David Simeone a,2, Didier Gilbon b,3 a b
CEA, DEN, LRC CARMEN, CEA Saclay, 91191 Gif/Yvette cedex, France CEA, DEN, CEA Saclay, 91191 Gif/Yvette cedex, France
a r t i c l e
i n f o
Article history: Received 16 December 2011 Accepted 3 May 2012 Available online 23 May 2012
a b s t r a c t Zircaloy-4 samples have been analyzed by grazing incidence X-ray diffraction during oxidation at high temperature (900–1100 °C) in a He–O2 mixture simulating a LOCA atmosphere. We clearly show that the structure of the oxide layer strongly depends on the oxidation temperature, i.e. tetragonal at 1100 °C, mixture of tetragonal and monoclinic at 1000 °C, mainly monoclinic at 900 °C. Moreover, the high temperature phases are not stable during cooling; at room temperature, the oxide is mainly monoclinic. Furthermore, re-heating the samples to the oxidation temperature in neutral atmosphere does not lead to the pristine structural composition. This behavior is explained by considering the martensitic character of the tetragonal to monoclinic transition of zirconia. As a conclusion, the structural and thermo-mechanical properties of the oxide layer as it forms during oxidation at high temperature cannot be deduced from post-mortem analysis. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction The zirconium alloys fuel claddings used in nuclear water reactors oxidize during normal operation or under low probability accidental situations [1,2]. The resulting oxide layer notably affects the thermal exchanges and the mechanical properties of the cladding. More specifically, during a Loss of Coolant Accident (LOCA), the cladding is exposed to steam at high temperature (typically between 800 and 1200 °C), which leads to a relatively fast growth of the oxide layer and a diffusion of oxygen into the metallic substrate [2–5]. Even higher temperatures may be reached during reactivity initiated or severe accidents. The oxidation kinetics of zirconium alloys is logically related to the structural and microstructural properties of the oxide layer, which depends mainly on the oxidation temperature, oxidizing atmosphere or type of alloy. For example, it is known that between approximately 600 °C and 1050 °C, ‘‘breakaway’’ oxidation phenomena can occur during steam oxidation [3–6]. Shortest critical times, at which this transition takes place, are observed around 1000 °C. This phenomenon is mainly characterized by a rapid increase of the oxidation rate (transition from a (sub-) parabolic to a nearly linear kinetics) and a significant hydrogen uptake into the metal when critical oxida-
⇑ Corresponding author. Tel.: +33 1 6908 5857; fax: +33 1 6907 7130. E-mail addresses: [email protected] (D. Gosset), [email protected] (M. Le Saux), [email protected] (D. Simeone), [email protected] (D. Gilbon). 1 Tel.: +33 1 6908 1228. 2 Tel.: +33 1 6908 2920. 3 Tel.: +33 1 6908 2200. 0022-3115/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jnucmat.2012.05.003
tion time or oxide layer thickness (which depend mainly on temperature) is reached. One common explanation to this transition is related to the structural properties of the zirconia layer in this temperature range: transformation of the initially formed metastable tetragonal zirconia to the stable monoclinic phase during oxide growth [2–5]. Thus, taking into account the potential consequences on the cladding integrity, it is important to be able to characterize the oxides formed during oxidation at high temperature of zirconium alloys. The phase diagram of the zirconium–oxygen system is well documented [7]. In particular, the zirconia ZrO2 undergoes two structural transitions leading to drastic symmetry changes [8]. The high fluorite like), with a wide stoitemperature phase is cubic (Fm3m chiometry range allowed by high oxygen vacancy concentration. It undergoes a second order transition around 2300 °C to a tetragonal structure (P42/nmc). The second structural change occurs around 1100 °C, from tetragonal to monoclinic (P21/c). The latter is a first order, martensitic-like transition. The structure undergoes strong distortions and volume changes (about 3.5%). A marked hysteresis, around 200 °C wide, is observed (Fig. 1). It is worth noticing that this transition occurs in the temperature range the fuel claddings are supposed to undergo during a LOCA. Complex behaviors are then to be expected. Moreover, the tetragonal phase can be observed down to room temperature in pure zirconia either in nanometric grains or in irradiated materials. In both cases, it has been shown that this stabilization is induced by high internal stresses [9,10]. Such stresses are also responsible for the stabilization of the tetragonal phase, e.g. by point defects [11], doping elements [12] or high pressure [13].
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10000
analyzed depth (nm)
90% 1000
αc
10
1
2. Experimental procedures The experiment have been performed on a laboratory goniometer equipped with a fixed curved position sensitive detector (INEL CPS-120) allowing recording the X-ray diffraction diagrams simultaneously on a 120° range with a 0.015° resolution. The angular calibration of the detector is performed with a Y2O3 secondary standard. The X-ray source is a classical Cu tube. A Ge (1 1 1) plane monochromator allows first to select the Cu Ka1 reflection and second to obtain a narrow (80 lm width), very parallel beam. Such a setting is then used to perform the analyses in an asymmetric configuration [25]. In this case, modifying the impinging angle of the incident X-ray beam allows an accurate control of the analysed thickness of the material (Fig. 2). The simultaneous recording of a wide 2h angular range allows fast analyses: here, identification of the zirconia phases and semi-quantitative analyses can be obtained within 300 s records. The goniometer is equipped with a HTK-1200 Anton-Paar furnace allowing analysis up to 1200 °C in neutral or oxidizing conditions. A 190°-wide modified window allows sample alignment and analyses on the whole angular range of the CPS-120 detector in grazing incidence beam conditions (GIXRD). The window of the furnace is closed with a 130 lm kapton foil protected with a 20 lm nickel foil to resist to the inner high temperature oxidizing
0.1
1
10
incidence (°)
Fig. 1. Monoclinic-tetragonal transition of pure zirconia (tetragonal volumic fraction) from [8]. : on cooling. j: on heating.
However, observing the actual structure of zirconia as it forms in the conditions of a LOCA remains a quite difficult task. This is why most of the structural characterizations are performed ‘‘post-mortem’’, i.e. either at room temperature [14,15] or after re-heating in neutral atmosphere at representative temperatures [16–18]. It is however expected that, owing to the martensitic features of the monoclinic to tetragonal transition, such an approach is unable to lead to a correct description of the actual phase structure of the oxide as it forms [19]. Furthermore, the very few ‘‘in situ’’ structural analysis reported in the literature have been performed under air (which leads to an accelerated degradation compared to steam due to the formation of oxynitrides [20]) and/or at temperatures lower than 900 °C [21–24]. In this frame, we have launched a program aiming at an accurate structural analysis of the oxide ‘‘in situ’’, i.e. remaining at the temperature of oxidation. In this paper, we present preliminary results obtained through high temperature X-ray diffraction analysis of Zircaloy-4 maintained at high temperature (900–1100 °C) in controlled atmosphere.
1-1/e
100
Fig. 2. Analyzed depth in grazing incidence X-ray diffraction as a function of the impinging X-ray incidence in ZrO2, for two absorption yields (90% and (1 1/e), i.e. the classical absorption length); ac: critical angle for total reflection.
atmosphere. The parasitic diffraction lines generated by those foils do not interfere with the Zr or ZrO2 main lines. The furnace-holder is motorized in order to tune the height (±5 lm) and the incidence (±0.02°) of the sample with regard to the X-ray beam and compensate for the sample holder dilatation at high temperature. A gas circuitry allows measurements under neutral (5.5 N He) or strongly oxidizing (40%He–60%O2) atmospheres. According to the oxidation kinetics reported in [26,27], this He–O2 mixture is expected to be a rather good surrogate to steam (LOCA conditions) and is easier to handle. Although the hydrogen production (and absorption into the metal for instance in the post-breakaway regime) and the enthalpy of the Zr–O reaction are not the same during oxidation under O2 and steam, the structural properties of the growing oxide are expected to be similar. Thanks to the small volume of the furnace chamber, the gas circuitry allows fast gas change (less than 10 s): however, this would induce turbulent flows in the furnace chamber and a strong transient cooling of the sample, up to 100 K. Here, such a transient cooling is prohibited since abrupt, irreversible structural changes are expected. As a consequence, the gas change is made slower and practically obtained in around 1 min. The temperature is measured with a Pt 10%Rh–Pt thermocouple close to the sample. The temperature difference between the sample and the thermocouple has been previously determined placing a secondary K-thermocouple on the back face of a standard sample through the sample holder axis. The difference (less than 15 °C in the whole temperature range in gaseous environments) is then automatically corrected. 3. Results The measurements have been performed on Zircaloy-4 material (Zr–1.3%Sn–0.2%Fe–0.1%Cr–0.13%O, in wt.%). The samples are 12 12 mm2 flat squares with a thickness of about 1 mm. The size of the samples allows grazing incidence X-ray diffraction analysis at incidence down to 0.5° without the X-ray beam to propagate over the sample. The flatness is required to make sure the incident angle (and then the analyzed depth) is constant on the whole illuminated surface. The analyses have been performed with a constant 5° incidence angle for the impinging X-ray beam, for which most of the diffracted intensity comes from the first 3 lm (Fig. 2), whatever the diffracted angle. In such conditions, the fingerprint of the X-ray beam on the sample surface is about
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0.9 mm wide, leading to a moderately wide instrumental component FWinstr to the width of the diffracted lines:
Table 1 Theoretical intensities of monoclinic and tetragonal zirconia and zirconium in the (2h) range of interest for phase identification.
FWinstr ¼ ðe=R sin aÞ sinð2h aÞ where e and a are the width and the incident angle of the impinging beam, R is the radius of the goniometer and 2h is the angle of diffraction. The following measurement sequence has been performed on fresh samples. First, the sample is heated under pure helium (heating rate 1 K/s). A diffraction analysis is performed in order to check there is no significant oxidation during heating. Then, the oxygen– helium mixing is introduced. Short (30 s) diffraction analyses are performed in order to observe the oxide formation. When the oxide layer is thick enough to induce the disappearance of the metal diffraction diagram, pure helium is reintroduced. There is no need here to continue oxidizing since the subsequent oxidation would no longer be visible by diffraction analysis, unless phase changes occur in the outer part of the oxide layer due in particular to relaxation effects: this point is not intended to be studied within the scope of this preliminary study. Accurate diffraction analyses (500 s) are then performed from the oxidation temperature down to room temperature (cooling rate 1 K/s then plateaus) by 100 °C steps in order to analyze the oxide composition during cooling. At last, the sample is re-heated under pure helium and analyzed again at the oxidation temperature. The observed diffraction patterns correspond to well crystallized but highly oriented materials (Fig. 3 to be compared to Table 1): for example, the (1 1 1) monoclinic ZrO2 line (2h 31.5°) is nearly invisible. Quantitative analyses are thus impossible. Here, the oxide phase ratios are then roughly estimated by the following expression:
T ð1 0 1ÞT ¼ T þ M ð1 1 1ÞM þ ð1 0 1ÞT þ ð1 1 1ÞM where the (h k l)P are the intensities of the (h k l) lines (when visible) for the P phase (T for tetragonal, M for monoclinic). All other lines are too strongly convoluted to be easily considered. This expression is obviously far from the actual phase ratios but can be used for relative estimates. At the end of the experiments, all the samples show a black, smooth, slightly rounded surface with no spalling or desquamation. Oxide layer thicknesses measured by optical microscopy are close to those obtained elsewhere for cladding tube specimens oxidized under pure steam at nearly atmospheric pressure [2,28]: the
Fig. 3. Diffraction patterns after oxidation recorded at the temperature of oxidation. Extra lines (e.g. at 56°) come from the kapton–Ni window.
(h k l)
2h (°)
I/Imax (%)
ZrO2 M
1 1) (1 (1 1 1) (0 0 2) (0 2 0) (2 0 0)
28.2 31.5 33.9 34.7 35.4
100 68 21 13 15
ZrO2 T
(1 0 1) (0 0 2) (1 1 0)
30.1 34.5 35.1
100 9 12
Zirconium a
(1 0 0) (0 0 2) (1 0 1)
32.0 34.9 36.1
24 25 100
Zirconium b
(1 1 0)
35.8
100
actual geometry (flat squares versus tubes) and atmosphere (He–O2 mixing versus steam) have thus no significant effect here. A peak to peak analysis of the diagrams has been performed limited to the angular range 20–40° which allows phase identification and semi-quantitative phase ratio estimates according to the previous equation. The main results can be summarized as follows (Figs. 3 and 4): As mentioned above, the patterns correspond to well-crystallized compounds, with narrow lines and low background levels. They can be identified to highly textured monoclinic and/or tetragonal zirconia: the cubic phase is not observed (systematic splitting of the observable lines, clearly corresponding to the tetragonal phase). The structure composition of the oxide layer at the oxidation temperature drastically depends on this temperature. At 900 °C, nearly pure monoclinic ZrO2 formed. At 1000 °C a monoclinic-tetragonal mixing is observed. At 1100 °C, pure tetragonal ZrO2 formed. When cooled down to room temperature, all the samples show mainly monoclinic, oriented zirconia. The phase transformation during cooling depends on the oxidation temperature. From 1100 °C, the tetragonal phase abruptly decreases when cooled from 900 °C to 800 °C, then smoothly decreases down to ca. 300 °C. From 1000 °C and 900 °C, a slow decrease is observed from the oxidation temperature down to ca. 300 °C.
Fig. 4. Relative intensity of the tetragonal component of the oxide layer (T/(T + M)) observed on Zircaloy-4 samples after oxidation under O2-He at high temperature then cooling to room temperature then re-heating to the oxidation temperature under pure He.
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When re-heated at the oxidation temperature under pure helium, the tetragonal component remains much lower than observed just after oxidation. Complementary tests have then been performed. A Zircaloy-4 sample has been oxidized in the DEZIROX 1 facility [6,28] in LOCA representative conditions (steam at nearly atmospheric pressure at 1000 °C during 534 s), leading to the formation of a 24 lm thick oxide layer. At room temperature, the X-ray diffraction diagram shows monoclinic oriented zirconia with a minor tetragonal component. After re-heating at 1000 °C under neutral atmosphere, the only modification – excepting lines shift due to thermal expansion – of the diagram was a narrowing of the diffraction lines, most probably due to decreasing in residual stresses since grain growth could hardly occur in such a low temperature range. No noticeable modification of the relative intensities of the former lines or new lines was observed. This is very similar to the observations we made during the re-heating step performed after the 1000 °C oxidation treatment under He–O2 mixture. Another sample has been pre-oxidized in autoclave in Pressurized Water Reactor (PWR) representative conditions (43 days at 360 °C, water at 190 bars with 1000 w.ppm B and 2 w.ppm Li) leading to the formation of a 1.2 lm thick oxide layer. The oxide we observed is mainly highly oriented monoclinic zirconia and does not change when heated at 1000 °C under pure He (Fig. 5). This sample has then been oxidized at 1000 °C then cooled, following the same sequence as above. As previously observed at 1000 °C, a mixture of tetragonal and monoclinic zirconia formed, the ratio of which cannot be precisely determined due to the pre-existing oxide layer. However, it is worth noticing that the ‘‘fresh’’ high temperature oxide layer begins to form only after 1200 s under oxygen (as identified by the disappearance of the zirconium substrate). This is in good agreement with the observations made on pre-oxidized samples subsequently oxidized at 1000 or 1200 °C in a steam environment [29]: in the presence of a pre-oxide layer, the formation of a ‘‘fresh’’ oxide at the metal/oxide interface is delayed (time necessary for oxygen anions coming from the oxidizing environment to diffuse across the pre-oxide layer). Furthermore, a different texture formed after 1200 s at 1000 °C since the (1 1 1) monoclinic line clearly appears (Fig. 5). An analysis at a lower Xray beam incidence (1°, corresponding to a 500 nm analyzed thickness) has been performed, showing that the structure of the upper oxide layer remained unchanged. During cooling, the tetragonal
phase and the (1 1 1) monoclinic line disappear below 750 °C and do not reappear when re-heated under pure He at 1000 °C.
4. Discussion 4.1. X-ray diffraction analysis The method we use here allows a qualitative analysis of the phase change in the materials, as evidenced in Fig. 4. However, accurate structural description with methods such as Rietveld analyses cannot be obtained: first, such a derivation would require longer records with more accurate diagrams and, above all, texture analyses would be previously required since the relative intensities of the lines are far from those of the isotropic phases. Moreover, it is worth noticing that one of the main advantages of the asymmetric configuration we use here, when compared to the classical symmetrical ones, is that the material thickness which is analyzed does not depend on the given diffraction lines: no volume correction are then to be performed even for the first steps of oxidation when the metal substrate is still visible. However, performing structural refinements would require specific geometrical corrections of the texture effects (Fig. 6). The reason is that on a classical Bragg– Brentano configuration, the diffraction vector (bisecting direction between the incident and the diffracted beams) is fixed (OA in Fig. 6, perpendicular to the sample surface) and the texture correction depends only on the angle between the diffraction planes normal and this vector. But, on the asymmetric configuration, first, a given (h k l) line corresponds to diffraction domains which have a different orientation to the sample surface than in the symmetrical configuration, and second, this direction turns according to h (OA then OA0 in Fig. 6), and then the texture correction would have to be accordingly modified for each line. For instance, different corrections would have to be applied to the harmonic lines of the same diffracting planes (such as (h, k, l), (2h, 2k, 2l), and (3h, 3k, 3l)) since different populations of diffracting domains, with different directions in the sample, have to be considered. Such a correction is not implemented in the classical structural refinement programs (e.g. Rietveld method).
θ θ
θ
θ
Fig. 5. Diffraction diagrams after heating at 1000 °C under He (the line at 35° and the shoulder at 37° correspond to the zirconium substrate), after 3000 s at 1000 °C under O2–He and after cooling at 750 °C under He for a sample pre-oxidized in autoclave at 360 °C under representative PWR conditions.
Fig. 6. Texture correction in Bragg–Brentano and asymmetric configurations. In the asymmetric setup, a given (h k l) line corresponds to different grains with different directions as compared to the symmetric setup and different texture corrections would have to be applied to the harmonics of the same (h k l) family: the SOD and SOD0 X-rays correspond to different populations of diffraction domains.
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The resolution of the diagrams would allow cell parameters estimations as long as unambiguous line indexation is possible: this condition is certainly fulfilled for the tetragonal phase, but a careful observation of the diagrams would have to be performed for the monoclinic phase. The main reason of this limitation is here again the high orientation of the diffracting grains leading to ambiguous line indexation in the large h range. Now, the cell parameters values that could be derived may have complex evolutions since they depend on both the temperature (thermal expansion) and the stresses induced by the metal substrate and the relaxation of the two zirconia phases. It is also well known that analyzing the evolution of line widths along a diffraction diagram can lead to estimations of structural parameters of the material, size of diffraction domains and microstrains. However, the line widths can here be determined only on the first part of the diagrams, up to about 60° (2h): on the high angle part of the diagrams, the lines superposition and the low signal over noise ratio make such an analysis inaccurate. We then observe a significant broadening of the lines as the samples are cooled (then a narrowing when re-heated), specifically for the tetragonal phase. However, it is here impossible to attribute this broadening to a decrease of the size of the diffraction domains or to an increase of the internal microstrains. Such an analysis could be tentatively performed by the well-known Hall–Williamson method but it requires the use of lines over a larger angular range than the one available here. As a result, it is impossible to discriminate between grain fragmentation (as observed in ion-irradiated zirconia [30]) and higher microstrains. It is however observed that the materials are relatively isotropic: for example, the (2 0 0), (0 2 0), (0 0 2) lines of the monoclinic phase of zirconia have close widths, which means that the micro-strains or the dimensions of diffracting domains do not significantly depend on the crystal direction. 4.2. Zirconia formation and phase transitions As mentioned above, the oxidation steps have been performed until the diffraction lines of the metallic substrate disappear, in order to avoid a possible oxide desquamation. This corresponds here to an oxide thickness of a few micrometers (Fig. 2), which is certainly lower that the critical thickness at which the so-called breakaway oxidation phenomenon occurs. By continuing the oxidation, initial existing constraints (compressive stress, misorientations, etc.) may be relaxed and a partial tetragonal to monoclinic conversion may occur in the growing oxide layer. These mechanisms are often postulated to explain the breakaway oxidation. Although the thickness analyzed by the ‘‘in situ’’ XRD method presented in this paper is only a few micrometers, one can expect that these transformations in the oxide may be visible within the outer part of the oxide layer. This point is to be analyzed in a further study. Those results clearly show that in all investigated cases, the zirconia phases formed at high temperature can be observed neither at room temperature nor when re-heated at high temperature. This can be qualitatively explained by the properties of the zirconia phase diagram. First, the higher the oxidation temperature, the higher the tetragonal phase content: the compositions we estimate are similar to those observed on the cooling branch of the hysteresis in Fig. 1 at the corresponding temperatures. Then, cooling the material allows the normal tetragonal to monoclinic transition to take place, by coarsely following this cooling branch. The differences between the curves in Figs. 4 and 1 may be attributed to stresses induced by the pre-existing oxide layer and the zirconium substrate or by different microstructures. The differences in residual tetragonal contents we observe at room temperature in Fig. 4 could also be attributed to different microstructures (grain size, micro-strains) that allow the stabilization of different tetragonal proportions or different preferred orientations. Then, re-heating
23
the cooled samples to the oxidation temperatures cannot lead to the reverse monoclinic to tetragonal transition since the temperature remains lower than the heating branch of the hysteresis (around 1200 °C). And even if such a reverse transition would occur, the obtained microstructure could not correspond to the pristine one. It then clearly appears that the transitions we observe are irreversible. This explains why the composition of the oxide when reheated at the oxidation temperature is strongly different from its composition during oxidation. Strong inferences can then be drawn from this study: The structural observations performed at room temperature after high temperature oxidation are never representative. Re-heating a sample to the previous oxidation temperature does not lead to the pristine microstructures, even if the same phases are observed. 5. Conclusions We have performed X-ray diffraction analysis during the oxidation under a mixture of He and O2 of Zircaloy-4 at high temperature (900–1100 °C) then during cooling and subsequent reheating in neutral atmosphere. We clearly show that any study based on post-mortem observations, whatever at room temperature or after re-heating, cannot be representative of the actual state of the material during the high temperature oxidation: the phase composition and microstructure of the oxide are irreversibly changed during cooling. This has strong inferences on the supposed structure of the oxide, but overall on its thermo-mechanical and microstructural properties as estimated from such post-mortem observations. Indeed, the tetragonal to monoclinic transition is associated with large distortions and volume changes. During cooling, this could for example induce (in addition to the effects of thermal stresses due to differences between thermal expansion coefficients of the oxide and the underlying metallic layers) an irreversible and non representative micro-fragmentation or deformation of the materials under consideration. Further experiments are in progress in order to determine the microstructural properties of the high temperature oxide and the effect of such structural transitions first on the oxidation kinetics of zirconium alloys and second on their consequences on the thermo-mechanics of the oxidized claddings. Acknowledgments The authors thank AREVA-NP for providing the materials. They are also highly indebted to Marc Tupin (CEA) for preparing the samples and performing the pre-oxidation treatments and Valérie Vandenberghe (CEA) for fruitful discussions. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]
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