Journal of Nuclear Materials 428 (2012) 110–116
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Impact of the use of the ferritic/martensitic ODS steels cladding on the fuel reprocessing PUREX process B. Gwinner ⇑, M. Auroy, D. Mas, A. Saint-Jevin, S. Pasquier-Tilliette CEA, DEN, DPC, SCCME, Laboratoire d’Etude de la Corrosion Non Aqueuse, F-91191 Gif-sur-Yvette, France
a r t i c l e
i n f o
Article history: Available online 12 November 2011
a b s t r a c t Some ferritic/martensitic oxide dispersed strengthened (F/M ODS) steels are presently developed at CEA for the fuel cladding of the next generation of sodium fast nuclear reactors. The objective of this work is to study if this change of cladding could have any consequences on the spent fuel reprocessing PUREX process. During the fuel dissolution stage the cladding can actually be corroded by nitric acid. But some process specifications impose not to exceed a limit concentration of the corrosion products such as iron and chromium in the dissolution medium. For that purpose the corrosion behavior of these F/M ODS steels is studied in hot and concentrated nitric acid. The influence of some metallurgical parameters such as the chromium content, the elaboration process and the presence of the yttrium oxides is first discussed. The influence of environmental parameters such as the nitric acid concentration, the temperature and the presence of oxidizing species coming from the fuel is then analyzed. The corrosion rate is characterized by mass loss measurements and electrochemical tests. Analyses of the corroded surface are carried out by X-ray photoelectron spectroscopy. Ó 2011 Elsevier B.V. All rights reserved.
1. Introduction The ferritic/martensitic oxide dispersed strengthened (F/M ODS) steels are an interesting alternative to the standard austenitic steels for the cladding materials of the future sodium fast nuclear reactors [1]. An advantage of the F/M steels is that they exhibit a low swelling under irradiation due to the body-center cubic crystal structure. This should allow reaching higher burn-up. To counterbalance the relatively poor mechanical properties at high temperature, the F/M steels are oxide dispersed strengthened (ODS) to improve the mechanical resistance in these conditions. These F/ M ODS steels are presently under study at CEA in order to increase the knowledge on several aspects: the fabrication process, the mechanical properties, the behavior under irradiation, the welding, the spent fuel reprocessing, etc. Concerning this last point, the potential impact of the use of these F/M ODS steels cladding on the spent fuel reprocessing has to be analyzed. The cladding presently used in the French pressurized water reactors (made of zirconium alloys) has actually no interaction with the dissolution medium (nitric acid). In the case of the F/M ODS steels, the cladding can be more or less corroded by nitric acid and the corrosion products (mainly iron and chromium) can be released in the dissolution medium. But some process specifications impose not to exceed a limit concentration of such corrosion products. The main objective ⇑ Corresponding author. Tel.: +33 1 69 08 16 40; fax: +33 1 69 08 15 86. E-mail address:
[email protected] (B. Gwinner). 0022-3115/$ - see front matter Ó 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2011.11.005
of this study is to verify if the cladding corrosion by the dissolution media remains acceptable regarding the process specifications. Three main options are currently envisaged at CEA: one martensitic ODS steel with 9 wt% of chromium and two ferritic ODS steels with respectively 14 or 18 wt% of chromium. The corrosion of the steels by hot and concentrated nitric acid is thermodynamically governed by the following reactions [2]: The global reduction of the nitric acid:
HNO3 þ 2Hþ þ 2e () HNO2 þ H2 O ðE0 ¼ 0:934 V=ESH at 25 C½2Þ
ð1Þ
It is worth mentioning that this reaction is in reality more complex. It has indeed been shown that other species than HNO2 can form and that this is not HNO3 (but other species with a lower oxidation degree) which directly exchanges the electrons with the steel [3–5]. But for the present work we consider as sufficient the use of reaction (1). The oxidation of the major elements of the steel (iron and chromium). In concentrated nitric acid (pH < 0), the Pourbaix diagrams shows that no thermodynamically stable solid oxides of chromium and iron should form [2]. Fig. 1 compares the standard potential at 25 °C of the HNO3/HNO2 couple and those of the most stable species of chromium and iron at this pH. It is obvious that chromium and iron should be mainly oxidized at the valence III by these reactions:
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HNO2
Current density (a.u.)
Cr
Table 1 Crystalline structure and chemical composition of the steels (wt.%).
HNO3
Cr 3+
Cr 2+
H2CrO4
Fe 2+
Fe
Fe3+
Cr
Steel
Structure
Fe
Cr
W
Ti
Y
Si
Ni
Mn
9Cr–1W–Ti ODS 14Cr–1W–Ti ODS 18Cr–1W–Ti ODS
Martensitic Ferritic Ferritic
Bal. Bal. Bal.
9 14 18
1 1 1
0.3 0.3 0.3
0.3 0.3 0.4
0.2 0.3 0.3
0.4 0.2 0.2
0.3 0.3 0.3
Fe Steel
-1
-0,5
0
0,5
1
1,5
2
Potential (V/ENH) Fig. 1. Standard potentials at 25 °C for different electrochemical couples [2] and steady anodic curves for iron, chromium and steel (in H2SO4 0.5 mol/L) [6].
Fe () Fe2þ þ 2e ðE0 ¼ 0:44 V=ESH at 25 C½2Þ and Fe2þ () Fe3þ þ e ðE0 ¼ 0:771 V=ESH at 25 C½2Þ
ð2aÞ
Cr () Cr2þ þ 2e ðE0 ¼ 0:913 V=ESH at 25 C½2Þ and Cr2þ () Cr3þ þ e ðE0 ¼ 0:407 V=ESH at 25 C½2Þ
ð2bÞ
To these thermodynamic considerations should be added some kinetic elements. For that purpose, the qualitative steady anodic curves of iron, chromium and steel in H2SO4 0.5 mol/L are given in Fig. 1 [6]. It is supposed that there is no major difference of the anodic behavior of iron, chromium and steel in nitric acid and in sulfuric acid as previously shown by Schosger for a 304L stainless steel [7]. In the range of the HNO3/HNO2 standard potential, the steel is corroded in its passive range, where the current density is nearly independent of the potential. This passivity is mainly due to the presence of chromium which shows an excellent resistance to the oxidation, as illustrated by the relatively low oxidation current density of chromium in this range of potential. This good corrosion resistance in acidic media is often attributed to the formation of a metastable chromium oxide layer on the surface [8] [9], which slows down the exchanges between the bulk metal and the solution. This good behavior of the steel is limited at higher potential (around 1.2 V/ESH) as illustrated by the sharp increase of the current density (transpassive range). This corresponds to the oxidation of chromium from the valence III to the valence VI, this later being highly soluble in nitric acid. Moreover it has to be noted that the HNO3/HNO2 standard potential (E0 = 0.934 V/ESH at 25 °C) is not far from the transition potential (around 1.2 V/ESH) between the passive and transpassive ranges. As a consequence, only the part of the anodic curve from the end of the passive range to the transpassive range is measurable by electrochemical measurement. Indeed, the anodic curve of the steels is only electrochemically measurable above the HNO3/HNO2 standard potential. Under this potential, the anodic curve is ‘‘masked’’ by the nitric acid reduction. The aim of the following paragraphs is to discuss on the influence of the steel (chromium content, etc.) and the media properties (temperature, nitric acid concentration, etc.) on the kinetics of these redox reactions (1), (2a), and (2b).
mechanical alloying was performed to add the yttrium in an attritor under hydrogen. After the mechanical alloying powders are sealed into cans and hot extruded at around 1150 °C. To obtain cylindrical bars or sheet bars different dies are used for the hot extruded products. During the mechanical alloying there is dissolution of the yttria (Y2O3). Yttrium, oxygen and titanium are put in solid solution in the powders and the precipitation of the nanoclusters (Y, Ti, O) occurs during the firing of the cans before the hot extrusion. The final microstructure depends on the alloy and the final metallurgical state. After hot extrusion the microstructure is made of very fine grains or ‘‘fibers’’ which are elongated along the hot extrusion direction. The corrosion experiments consist in immersion and electrochemical tests on the steels in nitric media. The immersion tests are performed with parallelepiped samples of the steels (10 30 1 mm) immersed in nitric acid (solutions prepared by dilution of HNO3 52.5%, AnalaR NORMAPUR, ProlaboÒ) at different temperatures (70–90 °C-boiling point) and at different concentrations (3–6–9 mol/L). The ratio S/V (the steel area on the volume of nitric acid) and the test duration are identical for each test (S/ V = 4.3 dm2/L and T = 48 h). The corrosion rate is estimated by mass loss measurement. XPS analyses (VG ESCALAB 220i XL) are carried out to determine the chemical composition of the oxide layer after corrosion. The electrochemical experiments consist in linear sweep voltammetry tests (Biologic VM). The exposed sample surface is 1.25 cm2 in contact with 300 mL of solution. The sample is first immerged in the nitric acid during 24 h to ensure that the steel surface is in equilibrium before starting the linear sweep voltammetry test. During this period the corrosion potential is measured to verify that the stabilized state is well reached. After stabilization the sample is slowly linearly polarized from the corrosion potential to more anodic potentials. The scan rate (0.167 mV/s) is chosen sufficiently low to consider that the system stays at stationary state during the whole polarization test. The experiment is stopped when the current density reaches 10 mA/ cm2. The anodic polarization curve is thus obtained. With a new sample the cathodic curve is obtained by a similar way (decreasing potential). The entire polarization curve is obtained by the superposition of both curves. In some cases, the oxidizing specie vanadium(V) is added to the nitric acid solution. Vanadium(V) is introduced as vanadium oxide V2O5 (purity 99.5%, Rectapur™ Prolabo). The relative concentration of vanadium(V)/vanadium(IV) is controlled by in situ measuring the redox potential on a platinum electrode and by ex situ measuring the absorbance of the solution (350 nm for the vanadium(V) and 750 nm for the vanadium (IV)) by UV–visible spectrophotometry (Carry 50 scan, Varian).
3. Results and discussion 3.1. Corrosion morphology
2. Materials and experimental methods The crystalline structure and the chemical composition of the three steels made at CEA are given in Table 1. The F/M ODS steels were prepared from pre-alloyed powders atomized [1]. The
At millimeter scale the corrosion can be considered as mainly homogeneous. Fig. 2 illustrates the typical appearance of the F/M ODS steels after corrosion. But this global homogeneous behavior should be completed by two observations.
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Fig. 2. Metallographic sections of the ODS steels after 48 h corrosion in nitric acid 9 mol/L at 90 °C.
Fig. 3. SEM observations of the ODS steels surfaces corroded 48 h in nitric acid 9 mol/L at 90 °C.
3.2. Influence of the metallurgical properties 3.2.1. Influence of the chromium content in the steel The chromium content is the main difference between the three ODS steels. It is then important to study its impact on the corrosion resistance in nitric acid. It is well known that chromium could increase the corrosion resistance in acidic media [11–14], even if it should thermodynamically be oxidized in Cr3+. This has to be
confirmed for the present system F/M ODS steels in nitric acid media. Fig. 4 presents the variation of the corrosion rate obtained from the immersion tests (48 h in nitric acid 9 mol/L at 109 ± 1 °C) as a function of the chromium content for different ferritic/martensitic steels. The lower the chromium content is, the faster the steel is corroded. The corrosion rate difference is about one order of magnitude between each steel. The beneficial role of chromium is in this way confirmed for the F/M ODS steels. This positive effect of chromium has also been observed during electrochemical test performed in the same conditions as these immersion tests [15]. It has been shown that all the three F/M ODS steels are corroded in their passive domain but not far from the transpassive one (the measured corrosion potentials are 0.94 ± 0.20, 1.07, and 1.10 V/ESH for the 9Cr-, 14Cr- and 18Cr–1W–Ti ODS steels, respectively), but the oxidation current in the passive domain depends on the chromium content in the steel. The higher the chromium content in the steel is, the lower the oxidation current is. The presence of chromium in the steel seems to improve the oxidizing resistance in the passive domain. To go further in the explanation of the beneficial role of chromium, some ex situ X-ray photoelectron spectroscopy (XPS)
25
Corrosion rate (µm/d)
Firstly it has already been shown that the corrosion morphology of these ODS steels is slightly different between the longitudinal and the transverse extrusion directions [10]. If the corrosion can be considered as mainly homogeneous in the extrusion direction, some pits appear in the transverse direction. It has been shown that this specific corrosion morphology could result from the preferential dissolution of inclusions aligned along the extrusion direction. Thermodynamic calculations and energy dispersive X-ray analyses tend to prove that these inclusions should be titanium carbides. On a practical point of view the transverse surface represents only a small part of the total corroded surface. This heterogeneous corrosion observed in the transverse direction should not have a significant influence on the global corrosion rate. For this reason this specific of corrosion morphology is no more considered in the rest of this study. Secondly some SEM observations allow characterizing more precisely the corrosion morphology. Fig. 3 shows the microscopic appearance of the surface for the different F/M ODS steels after corrosion in nitric acid. The 18Cr–1W–Ti ODS steel shows no corrosion as the surface presents the initial appearance (polishing). On the contrary the 9Cr–1W–Ti and 14Cr–1W–Ti ODS steels reveal a corroded structure. The grain morphology seems to be revealed. The 9Cr–1W–Ti ODS steel presents a quite isomorphous grain structure as the 14Cr–1W–Ti ODS steel shows a more elongated morphology (‘‘fibers’’ structure). This grains revelation should imply that the grain boundaries are more or less sensitive to corrosion in comparison with the grains themselves. The exact mechanism responsible of these grains revelation is not yet well explained. It is not the objective of this study to explain this specific behavior. On the first order it is considered that the ODS steels are corroded on a homogeneous way.
20 ODS sheet bars
15
ODS cyl. Bars Non ODS cyl. Bars
10
5
0 9Cr-1W-Ti ODS
14Cr-1W-Ti ODS
18Cr-1W-Ti ODS
Fig. 4. Evolution of the corrosion rate (estimated from the immersion tests) as a function of the elaboration geometry and the presence of yttrium oxide (48 h in nitric acid 9 mol/L at 109 ± 1 °C).
B. Gwinner et al. / Journal of Nuclear Materials 428 (2012) 110–116
analyses have been performed on the samples surface before and after the above-mentioned corrosion tests (samples prepared from a cylindrical bar and corroded in nitric acid 9 mol/L at 90 °C during 48 h). For the ones performed before the corrosion test, the XPS analyses show the presence of chromium and iron in both oxidized and metallic states. This means that a native oxide layer is present on the surface (due to the oxidation by the ambient air) whose thickness is thinner than the analysis depth (around 10–15 nm). Additionally the XPS analyses allowed to precise the chemical state of oxidized chromium (Cr2O3) and iron (Fe2O3 or FeOOH). The relative chromium content [Cr]/([Cr]+[Fe]) (where [Cr] and [Fe] are expressed in at.%) in the native oxide layer is compared in Fig. 5 with the relative chromium content in the steel. Both results are sufficiently similar to consider that the native oxide layer has the same composition as the steel. Concerning the XPS analyses performed on the F/M ODS steels after corrosion, Fig. 5 gives the relative chromium content in the oxide layer of the corroded samples. The XPS spectra show no metallic state for chromium or iron, which means that the layer thickness is thicker than the analysis depth (around 10–15 nm). It can be deduced that the analyses results concern only the external part of the oxide layer. Some complementary analyses are presently in progress to estimate the exact thickness and chemical composition of these oxide layers. The chemical state of chromium and iron in the external part of the oxide layer is the same as the one observed before corrosion. It is worth mentioning that XPS analyses are ex situ measurements, which can slightly modify the chemical state reached in solution. These XPS results before and after corrosion are compared in Fig. 5. For every ODS steels the relative chromium content in the oxide layer after corrosion is largely higher than before corrosion. From these observations it can be deduced that chromium seems to be kinetically less dissolved than iron. This has been confirmed by elementary analyses of the solution (by ICP-AES) during the corrosion test, which show that chromium is proportionally less detected in solution than iron [10]. Moreover the comparison of the results obtained for the three F/ M ODS steels indicates that the more the steel contains chromium, the more chromium is concentrated in the oxide layer. This could at least partially explain why a steel containing more chromium is more efficiently protected. The nature of the oxide layer on ODS steels has been extensively discussed in Ref. [10]. It has been shown that the chemical composition of the oxide layer does not depend a lot on the temperature (in the range 70 °C to boiling point) and on the nitric acid concentration (in the range 3–9 mol/L). Moreover it appears that the oxide layer contains other elements such as tungsten (from the steel) or
Steel composition XPS analysis of the oxide layer (before corrosion)
Ratio [Cr]/([Cr]+[Fe])
100%
XPS analysis of the oxide layer (after corrosion)
90% 80% 70% 60% 50% 40% 30% 20% 10% 0% 9Cr-1W-Ti ODS
14Cr-1W-Ti ODS
18Cr-1W-Ti ODS
Fig. 5. Relative chromium content in the F/M ODS steels, in their native oxide layer (air) and in their external oxide layer after corrosion in nitric acid 9 mol/L at 90 °C during 48 h.
113
silicon (which can come from the steel or be an experimental artifact coming from the corrosion of the glass reactor). 3.2.2. Influence of the elaboration conditions Two kinds of bars of ODS steels have been made: cylindrical and sheet bars. The latter are prepared on a same way as the cylindrical bars with an additional stage of hot rolling at around 700 °C. The corrosion rates measured for both geometries are given in Fig. 4. Whatever the chromium content is, the corrosion rate of sheet bars is on the same order but systematically inferior to than those of cylindrical bars. This small difference is at present time not explained but allows illustrating the variability of results which could be induced by the elaboration process. 3.2.3. Influence of the presence of nanoparticles of yttrium oxides The presence of nanoparticles of yttrium oxides improves the mechanical properties of steels at high temperature. An additional question could be if their presence has also an impact on the corrosion resistance. From Pourbaix [2], the thermodynamic stability of Y2O3 in aqueous solution is governed by the following equilibrium:
2Y3þ þ 3H2 O () Y2 O3 þ 6Hþ
ð3Þ
In an acidic medium such as in nitric acid solution this equilibrium is highly displaced to the left side and Y2O3 highly dissolved. Indeed the concentration of Y3+ can be estimated as follows [2]:
Log½Y3þ ¼ 35; 32 3pH
ð4Þ
It can be concluded that Y2O3 can be considered as thermodynamically not stable. If the dissolution kinetic is sufficiently fast, the presence of yttrium in the steels should not have a significant role in the protection of the material. To confirm this point, cylindrical bars have been made with and without yttrium oxides. The corrosion rates are given in Fig. 4. In the absence of yttrium, the corrosion rate is lower for the 9Cr– 1W–Ti ODS (40%) and the 14Cr–1W–Ti ODS (50%) steels, but higher for the 18Cr–1W–Ti ODS (+20%) steels. The variability of these results is considered to be on the same order as the variability observed for the elaboration process (see Section 3.2.2). As a consequence the presence of yttrium is considered to have no major influence on the corrosion process. 3.3. Influence of the medium properties 3.3.1. Influence of the temperature The temperature is a process parameter that could be adjusted to control the dissolution kinetic of the spent nuclear fuel. It is essential to quantify the exact impact of the temperature on the corrosion by nitric acid of the ODS steel cladding. Corrosion tests (for the three F/M ODS steels and three nitric acid concentrations 3–6–9 mol/L) have been performed at various temperatures namely 70, 90 and boiling point (113 °C for the nitric acid 9 mol/L). For all these conditions the measured corrosion rates have been traced in a semi-logarithmic scale as a function of the inverse of the temperature in Fig. 6. The corrosion process appears to be thermally activated and the corrosion rate depends on the temperature according to an Arrhenius law Vcorr = Exp (Ea/RT) where Vcorr is the corrosion rate (lm/d), A the preexponenial factor (lm/d), Ea the apparent activation energy (kJ/mol), R the universal gas constant (8.314 10-3 kJ/mol/K) and T the temperature (K). The apparent activation energies calculated by fitting the experimental results are listed in Table 2. It appears that the energy values are slightly influenced by the chromium content in the steel and the nitric acid concentration. Nevertheless the apparent acti-
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B. Gwinner et al. / Journal of Nuclear Materials 428 (2012) 110–116 Table 3 Values of the activity coefficient of HNO3 cHNO3 , the mean activity coefficient c±, the H+ concentration for different nitric acid concentrations and the dissociation constant K [16].
100
Corrosion rate (µm/d)
113 C (B.P.) 90 °C
10 9Cr-1W-Ti ODS
70°C
1
HNO3 9 mol/L HNO3 6 mol/L
14Cr-1W-Ti ODS
HNO3 3 mol/L HNO3 9 mol/L HNO3 6 mol/L HNO3 3 mol/L
0.1 18Cr-1W-Ti ODS
HNO3 9 mol/L HNO3 6 mol/L
0.01
HNO3 3 mol/L
0.001 0.0025
0.0026
0.0027
0.0028
0.0029
0.003
Inverse of temperature (1/K) Fig. 6. Variation of the corrosion rate (estimated from the immersion tests) as a function of the inverse of the temperature for the three F/M ODS steels and for different nitric acid concentrations.
Table 2 Apparent activation energies (kJ/mol) calculated from experimental results for the three F/M ODS steels and for various nitric acid concentrations. Steel
Concentration of nitric acid
9Cr–1W–Ti ODS 14Cr–1W–Ti ODS 18Cr–1W–Ti ODS
3 mol/L
6 mol/L
9 mol/L
69 ± 6 70 ± 5 62
73 ± 1 79 ± 5 82 ± 13
74 ± 2 95 ± 5 91 ± 1
[HNO3]0
cHNO3
c±
[H+]
K
3 mol/L 6 mol/L 9 mol/L
2.49 5.99 13.9
1.149 2.63 5.79
2.7 mol/L 4.5 mol/L 5.0 mol/L
15.4
cathodic currents are translated to higher values when the nitric acid concentration is increases. It means that the reduction process (reaction (1)) is accelerated by the presence of nitric acid. It confirms that nitric acid can be considered as the main oxidizing specie responsible of the steel corrosion. On the other hand the anodic currents (resulting from reactions (2)) are also influenced by the nitric acid concentration, in particular in the passive zone. At a fixed anodic potential, the more concentrated the nitric acid is, the higher the mean passive current is. Considering that the passive current results from the state the passive layer on the surface (mainly composed by Cr2O3 as shown in Section 3.2.1), it can be deduced that the passive layer is less efficient in a more concentrated acid, even the steel is polarized at the same potential. The following thermodynamic considerations could be proposed to explain this trend. The thermodynamical stability of the main component of the oxide layer Cr2O3 is controlled by the following equilibrium [2]:
Cr2 O3 þ 6Hþ () 2Cr3þ þ 3H2 O
ð5Þ
vation energy seems to increase slightly when the nitric acid is more concentrated.
It appears that the Cr2O3 solubility is highly influence by the concentration of H+, which is mainly controlled by the following reaction [16]:
3.3.2. Influence of the nitric acid concentration The nitric acid concentration is another parameter that could be adjusted for process reasons. The more concentrated the nitric acid is, the faster the spent nuclear fuel dissolution is. It is then important to study the impact of a nitric acid concentration variation on the corrosion of the cladding. Immersion tests (for the three ODS steels and at 70 °C, 90 °C and boiling point) have been performed at various concentrations of nitric acid (3, 6 and 9 mol/L). The corrosion rates obtained at 90 °C are given in Fig. 7 in a semi-logarithmic scale. The whole results are given in Fig. 6. For a given ODS steel, the more concentrated the nitric acid is, the larger the corrosion rate is. To explain this trend, some linear sweep voltammetry tests have been carried out in similar conditions [15]. It has been observed that the corrosion potential increases when the nitric acid is more concentrated. This corrosion potential variation has been explained by two complementary ways. On the one hand the
HNO3 () Hþ þ NO3
10
Corrosion rate (µm/d)
9Cr-1W-Ti ODS 14Cr-1W-Ti ODS
1
18Cr-1W-Ti ODS
0.1
0.01
ð6Þ
+
+
The H concentration [H ] (mol/L) can be estimated from the relationship K ¼ c
c2 ½Hþ 2
HNO3 ð½HNO3 0 ½H
þ 0
Þ
where K is the dissociation con-
stant of the reaction (6), c± the mean activity coefficient, cHNO3 the activity coefficient of HNO3 and [HNO3]0 the global nitric acid concentration (the sum of the dissociated NO 3 and non dissociated HNO3 forms). Some values of these parameters at 25 °C are proposed in Table 3 for different nitric acid concentrations [16]. It is then shown that the H+ concentration is increased when the nitric acid is more concentrated (reaction (6)). This increase of H+ concentration should displace the reaction (5) to the right side increasing the dissolution of Cr2O3. These thermodynamically considerations could explain the variation of the film resistance with the nitric acid concentration. 3.3.3. Influence the dissolved fuel In a real dissolution medium, the nitric acid contains a lot of dissolved species coming from the nuclear fuel. Under certain conditions (at high temperature, with low HNO2 concentration), some dissolved species can exist at a high valence state (for example plutonium(VI), neptunium(VI), etc.). Taking into account their high standard redox potentials (compared to the one of the HNO3/ HNO2 couple, E = 934 mV/SHE at 25 °C), they are considered as oxidizing species [5]. þ 4þ PuO2þ þ 2H2 O ðE0 ¼ 1040 mV=SHE at 25 CÞ 2 þ 4H þ 2e ¼ Pu
ð7Þ 0.001 3
6
9
HNO3 concentration (mol/L) Fig. 7. Corrosion rate variation (estimated from the immersion tests) as a function of the nitric acid concentration for the three ODS steels at 90 °C.
þ 0 NpO2þ 2 þ e ¼ NpO2 ðE ¼ 1150 mV=SHE at 25 CÞ
ð8Þ
For some austenitic stainless steels (as 304L), it has been established that in the presence of such oxidizing species, the global cathodic reaction is accelerated [5,17]. The corrosion potential
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B. Gwinner et al. / Journal of Nuclear Materials 428 (2012) 110–116
100
1.E-01
Current density (A/cm2)
Without vanadium(V) With 100 mg/L vanadium(V)
Corrosion rate (µm/d)
With 300 mg/L vanadium(V)
10
1
1.E-02 1.E-03 1.E-04 HNO3 9N + V 300 mg/L
1.E-05
HNO3 9N + V 100 mg/L HNO3 9N
1.E-06 1.E-07 0
0.2
0.4
0 14Cr-1W-Ti ODS
18Cr-1W-Ti ODS
Fig. 8. Corrosion rate variation (estimated from the immersion tests) as a function of the vanadium concentration (for the three ODS steels in HNO3 9 mol/L at 110 ± 1 °C).
0.8
1
1.2
1.4
tends to higher values and the corrosion rate is increased. It is not clear if the cathodic reduction of these oxidizing species replaces or increases (by a catalyst process) the one of nitric acid. It is probably a combination of both effects which the contribution ratio could depends on the nature of the oxidizing specie and the conditions (concentrations, temperature, etc.). Considering these observations established on austenitic stainless steels, it is essential to verify the possibility to transpose these results to the F/M ODS steels. Some corrosion tests have been performed to study the influence of the presence of the oxidizing specie vanadium(V). The later is often used to simulate the corrosive behavior of plutonium(VI) [5,17], especially for corrosion tests on austenitic steels. The use of vanadium(V) as a simulant of plutonium(VI) seems to be justified by the nearly close standard potentials and the similar cationic forms Pu4+ and VOþ 2.
VðOHÞþ4 þ Hþ þ e ¼ VO2þ þ 3H2 O ðE0 ¼ 1000 mV=SHE at 25 CÞ
Fig. 9. Polarization curves of the 9Cr–1W–Ti ODS steel in HNO3 9 mol/L at 110 °C with different vanadium(V) concentrations.
1.E-01
Current density (A/cm2)
9Cr-1W-Ti ODS
0.6
Potential (V/SHE)
1.E-02 1.E-03 1.E-04 1.E-05
HNO3 9 mol/L + V 300 mg/L
1.E-06
HNO3 9 mol/L + V 100 mg/L HNO3 9 mol/L
1.E-07
1a E-08 0
0.2
0.4
0.6
0.8
1
1.2
1.4
Potential (V/SHE) Fig. 10. Polarization curves of the 18Cr–1W–Ti ODS steel in HNO3 9 mol/L at 110 °C with different vanadium(V) concentrations.
ð9Þ Immersion tests have been performed with various concentrations of vanadium(V). The results are given in Fig. 8. Whatever the steel is, the corrosion rate is increased by the presence of vanadium(V). The more concentrated in vanadium(V) the solution is, the higher the corrosion rate is. Moreover the influence of the presence of vanadium(V) has a relative stronger impact on the steels with higher chromium content. As an important consequence these results seems to indicate that the high differences observed on the corrosion rate (see Section 3.2.1) between three steels in pure nitric acid (about one order of magnitude) decrease strongly in the presence of an added oxidizing specie. To explain this behavior some linear sweep voltammetry tests have been carried out in similar chemical conditions (nitric acid and vanadium(V) concentrations, temperature). The polarization curves are traced in Figs. 9 and 10 for the F/M ODS steels containing respectively 9 and 18 wt%. For every F/M ODS steels, the presence of vanadium has qualitatively the same impact on the polarization curves. At highest potentials the anodic curve is not modified by the presence of such oxidizing specie. It means that the presence of vanadium does not influence the passivity or transpassivity behavior of the steels. This observation is similar to the one made on some austenitic steels corroded by nitric acid in the presence of oxidizing species [5,17,18]. It has been shown that the presence of oxidizing species such as cerium(IV), vanadium(V), iron(III), manganese(VII), and plutonium(VI) does not change the oxidizing processes in the passive and transpassive ranges. On the contrary the cathodic curve is highly modified by the presence of vanadium. As expected the cathodic kinetic is accelerated by the presence of the oxidizing specie. In other words the cathodic
current is increased for a given cathodic potential, when the concentration of vanadium(V) is higher. As a result of the increase of the cathodic process and the stability of the anodic process, the corrosion potential increases when the concentration in vanadium(V) is higher. The potential shift is slightly higher for the 9Cr–1W–Ti ODS steel (0.94 ± 0.02– 0.99 ± 0.01–1.04 V/SHE) than for the 18Cr–1W–Ti ODS steel (1.1– 1.14–1.16 V/SHE). Despite a lower potential shift, it has been observed with the immersion tests that the corrosion rate is stronger increased for the 18Cr–1W–Ti ODS steel than for the 9Cr–1W–Ti ODS steel. This can be explained by the fact that the 9Cr F/M ODS steel stay in its passive zone even in presence of vanadium. In this zone the oxidation current is nearly independent of the potential. As a consequence the corrosion rate does not evolve a lot when the potential increases. On the contrary the 18Cr–1W–Ti ODS steel is corroded near the passive–transpassive transition. In this zone, the slope of the anodic curve is high. A slight increase in potential leads to a stronger current. This explains the higher sensitivity of the corrosion rate of the 18Cr ODS steel to the presence of vanadium(V). 4. Conclusions To quantify the potential impact of the use the F/M ODS steels cladding on the fuel reprocessing PUREX process, the corrosion behavior of these steels has been studied in nitric acid. The corrosion morphology can be considered as homogeneous in spite of the presence of some pits on the transverse direction and the revelation of the grains. It appears that the elaboration
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conditions or the presence of yttrium oxides have no significant influence on the corrosion rate. On the opposite, the chromium content in the steel impacts severely the corrosion behavior. The higher the chromium content is, the lower the corrosion rate is. Higher chromium content seems to allow a higher concentration of chromium in the oxide layer, which leads to a better corrosion resistance. As a consequence, the 9Cr–1W–Ti ODS steel has the worse corrosion resistance comparing to the others. The difference between the three F/M ODS steels seems to decrease for a more representative corrosion medium that includes the presence of oxidizing species from dissolved fuel. The conditions of the fuel dissolution (mainly the nitric acid concentration and the temperature) impact the corrosion the F/M ODS steels. The higher the temperature and the more concentrated the nitric acid solution are, the faster the corrosion is. From all these results it now possible to determine for each F/M ODS steels, the spent fuel dissolution conditions which leads to an acceptable corrosion level regarding to the process specifications for the maximum concentration of the corrosion products iron and chromium. Further efforts will be made to study the potential impact of the cladding aging in the reactor on the further corrosion resistance against nitric acid. Indeed all the present corrosion tests have been performed on non aged steels. In a more realistic point of view, the F/M ODS steel cladding will spend some years under working conditions (at high temperature, under irradiation, in contact with liquid sodium on the external surface, in contact with the fuel on the internal surface). This could leads to an aging of the F/M ODS cladding with some metallurgical evolutions. The corrosion resistance against nitric acid during the following stage of fuel recycling could then be affected. Acknowledgements This work is financially supported by Areva and EdF. The authors would like to sincerely thank M. Tabarant and H. Badji
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