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Characterisation of oxygen and hydrogen migration through oxide scales formed on nickel-base alloys in PWR primary medium conditions
Solid State Ionics 231 (2013) 69–73
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Characterisation of oxygen and hydrogen migration through oxide scales formed on nickel-base alloys in PWR primary medium conditions Fanny Jambon a,⁎, Loïc Marchetti a, François Jomard b, Jacques Chêne c a b c
CEA, DEN, DPC, SCCME, Laboratoire d'Étude de la Corrosion Aqueuse, F-91191 Gif-Sur-Yvette Cedex, France Groupe d'Etude de la Matière Condensée (CNRS and University of Versailles Saint Quentin), 45 avenue des Etats-Unis, 78035 Versailles cedex, France UMR 8587 CEA/CNRS, Équipe Hydrogène/Matériaux de structure, CEA, DEN, DPC, SCCME, Laboratoire d'Étude de la Corrosion Aqueuse, F-91191 Gif-Sur-Yvette Cedex, France
a r t i c l e
i n f o
Article history: Received 16 April 2012 Received in revised form 29 August 2012 Accepted 16 October 2012 Available online 4 December 2012 Keywords: Alloy 600 Hydrogen Oxygen Diffusion Isotope labelling Stress corrosion cracking
a b s t r a c t This study aims at providing quantitative data concerning the oxygen and hydrogen diffusion through the oxide scale formed on alloy 600 in PWR primary medium conditions (325 °C) using 2D and 18O isotopes as markers. The diffusion coefficients were calculated according to a short-circuit diffusion of the tracers along the oxide grain boundaries (so-called “C-regime”). The values obtained are respectively Dsc =4.6+/− 0.9× 10−17 cm2/s for 18O, and Dsc =5.2+/− 1.2×10−17 cm2/s for 2D. These experiments confirm that the transport mechanism of hydrogen through the oxide scale as hydroxide ions on the anionic sublattice is highly probable. The effect of hydrogen subsequent absorption by the alloy substrate, notably in stress corrosion cracking phenomenon is discussed. © 2012 Published by Elsevier B.V.
1. Introduction In a former study [1], single nickel-base alloy crystals, with a composition close to alloy 600s one, were exposed to simulated PWR primary medium, where the gaseous hydrogen or the hydrogen in water had been replaced by deuterium. SIMS analysis of the deuterium profiles in those samples allowed the authors to identify the origin of the hydrogen absorbed in nickel-base alloys exposed to PWR primary medium as coming from the water molecule dissociation associated with the oxide scale building. A mechanism was also proposed in order to model the hydrogen transport associated with the oxide growth during the alloy passivation. This mechanism is based on the diffusion of hydrogen as an interstitial proton through the oxide lattice; or as hydroxide ion, towards the oxide anionic sub-lattice. The latter hypothesis implies the oxygen and hydrogen diffusivities through the oxide scale layer to be the same. The aim of this study is to test the validity of this statement. 2. Experimental 2.1. Sample preparation In order to improve the depth resolution of the SIMS analysis used in this study, single-crystal samples of a model alloy, with a chemical ⁎ Corresponding author. Tel.: +33 169081261. E-mail address: [email protected] (F. Jambon). 0167-2738/$ – see front matter © 2012 Published by Elsevier B.V. http://dx.doi.org/10.1016/j.ssi.2012.10.012
composition close to alloy 600 were used. The single-crystal was provided by the Ecole Nationale Supérieure des Mines de Saint-Etienne. Its growing axis is b111> and its elementary composition is given in Table 1. The microstructure of this alloy is quite homogeneous; with the presence of some small micrometric precipitates, mainly chromium carbides. The samples were prepared according to a procedure described elsewhere [1]; the final surface state, corresponding to a mechanical polishing with a 1/10 μm diamond paste, exhibited a typical mirror aspect, characteristic of low surface roughness (Ra b 6.5 nm), as measured via microrugosimetry experiments.
2.2. Corrosion experiments The samples were first passivated during 516 h at PWR nominal conditions (Table 2). The oxide scale usually formed in these conditions, according to literature, is duplex, with a protective mixed spinel internal layer, mainly nickel chromite, compact and continuous; and an external layer mainly composed of micron-sized nickel ferrite crystallites associated to the precipitation of released cations, as extensively studied in [2]. The corrosion tests were thus performed in a recirculation autoclave where the primary medium cations were removed thanks to a set of ions exchanging resins. The oxide scale formed during this first exposure was observed with SEM. As shown in Fig. 1, it is composed of small, scattered sub-micron-sized compact crystallites (30–80 nm on average) that seem to correspond
70
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Table 1 Chemical composition (wt.%) of alloy 600-like single crystals. Ni
Cr
Fe
C
Co
Mn
Ti
Cu
Si
S
P
76.99
16.99
5.57
n.c.
0.01
0.01
b0.01
0.02
0.08
0.01
0.01
to the precipitates described in [2], covering less than 50% of a homogeneous layer with nano-sized grains irresolvable with this SEM (b20 nm), the internal part of the oxide layer. In a second sequence, the pre-oxidized samples were exposed to PWR primary water in a 300 mL titanium alloy autoclave (in order to minimize nickel and iron oxide and hydroxide precipitation phenomena at the surface of the samples). The relatively small size of this autoclave also allows decreasing the usual temperature transient to 45 min. In order to compare oxygen and hydrogen diffusivities, two series of experiments were performed, where these species were isotopically marked with 18O (20%) and deuterium (99.97%) respectively. The samples were exposed to this modified PWR medium during 9 h. At the end of the corrosion experiments, the samples were stored into liquid nitrogen until their characterisation. 2.3. Sample analyses The isotopic tracer profiles through the oxide layer were obtained with Secondary Ion Mass Spectroscopy (SIMS) analyses at the GEMaC CNRS laboratory (Meudon, France), as described in reference [1], but using the SIMS high mass resolution slits in order to differentiate 18O from the molecular ion 2D 16O at mass 18 (m/Δm = 1800). Then the oxide scale thickness was measured with Nuclear Reaction Analysis (NRA), done at JANNuS (Joint Accelerators for Nanoscience and Nuclear Simulation), Saclay, France and supported by the French Network EMIR. The experiment was performed with 900 keV 2D + ions using a 2 MV Van de Graaff accelerator. The quantitative determination of 16O content was made collecting the 1258 keV protons produced with the 16O(d,p1) 17O reaction, under an analysing angle of 150°. A Mylar stopping foil of 12 μm thickness was used to filter out the backscattered deuterons. The oxygen content calibration was made using some references composed of a tantalum substrate covered by a Ta2O5 oxide of well-known thickness. 3. Results The oxide scale thickness determined by NRA after the 525 h exposure to PWR primary medium was 29 +/− 3 nm. SIMS profiles of the two tracers through the oxide scale, representative of the results obtained on the different samples, are presented in Fig. 2. The initial raw intensity-sputtering time data have been converted to a depth-scale display after rescaling the sputtering time scale to the oxide scale thickness measured by NRA so that this thickness would coincide with the oxide/alloy interface on each sample. The oxide/alloy interface is arbitrarily set at the total oxygen signal (that is, [18O + 16O] for the sample exposed to the 18O marked medium) inflexion point. For this rescaling, a mean sputtering velocity is thus calculated from the oxide thickness measured by NRA and the sputtering time at which the oxygen inflexion point occurs; and the abscise axis is multiplied by this rescaling factor, which leads to the depth scale (only valid for the oxide part of the graph), of Fig. 2.
1 µm
Fig. 1. Oxide scale observation with SEM; magnification × 10 k.
Finally, the naturally present 18O contribution, associated to the 16O of the first sequence and second sequences has been subtracted to the total 18O signal. If some 18O associated with the 16O from the second sequence also contribute to the total 18O signal, they are negligible compared with the other contributions, since the second sequence is a really short one, performed at a low oxidation rate. In any case, this correction mainly changed the signal amplitude, which was to be expected since, except at the interfaces, the oxygen signal is constant through the oxide scale. This procedure permits to have an estimation of the total oxygen amount which penetrated in the scale during the second sequence, and to make sure that the growth was limited and that this amount is compatible with the C-regime assumption (see below). Since the tracer profiles are measured in a homogenous oxide scale, it is not necessary to correct the 2D intensity as was shown in reference [1] because of the need to compare 2D intensity in the oxide and in the metal, which have quite different ionisation and abrasion rates. One can distinguish three main domains in these profiles. On the left hand side of the graph Fig. 2a, at the beginning of the 18O profile appears a tip, corresponding to the fact that the 18O concentration increases faster than the 16O at the external interface. This phenomenon is linked to the presence of an inflexion point at P1 =3.5 nm on the total oxygen signal at the beginning of the oxide scale, and can be explained by a small but significant part of external growth during the second sequence (H218O medium). This part of the profile will thus have to be spared for the calculation of the tracer diffusion coefficient [3]. The second domain (the internal oxide scale), starts at the abovementioned first inflexion point, and ends at the internal oxide/ metal interface, that is, at the second inflexion point; and finally the third part corresponds to the metallic substrate. According to the work of [4], where the accumulation of oxygen isotopic tracers at the internal interface on a similar system was evidenced, demonstrating that the growth of the oxide proceeds by short-circuit diffusion, the tracer diffusion coefficients were determined assuming a short-circuit diffusion regime (or “C-regime” [3]), through the second domain (internal oxide growth part only). The 18O and 2D profiles were fitted with a C-regime equation modified to exclude the first x−P 1 ffiffiffiffiffiffiffi , where x is the recalculated domain interval: cH ðxÞ ¼ k:erf c p 2
Table 2 Experimental PWR primary medium conditions used in this study. Temperature Pressure (static autoclave) Pressure (recirculating autoclave)
325 °C 135 bar 155 bar
Boron (H3BO3) Lithium (LiOH) H2 (pressurisation)
1000 appm 2 appm 0.28 bar at 325 °C
Dsc :t
depth, k a coefficient depending on the tracer concentration at the oxide surface and on the short-circuits density, t is the second sequence duration, and Dsc the short-circuit diffusion coefficient. Finally, the main uncertainty for the determination of the tracer diffusion coefficients is the oxide scale thickness, as determined by NRA. The tracer diffusion coefficients were thus calculated for the mean oxide thickness, and for the minimum and maximum oxide
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Ext. growth
16
16
1.5 105
1200
6 104
600
3 104
0
0 0
1
10
20
30
40
50
60
70
80
Intensity [counts/s]
Intensity [counts/s] 18O
9 104
D [counts/s]
1.2 105
400
M/Ox interface
300
9 104
200
6 104
100
3 104
0
Intensity [counts/s]
M/Ox interface 1800
Intensity [counts/s]
1.2 105
1.5 105 2
16O
O [counts/s]
2400
O [counts/s]
500 16O
18
1.8 105
600
O [counts/s]
3000
(16O+18O)"
b
1.8 105
3600
2D
a
71
0 0
10
20
30
Depth (NRA) [nm]
40
50
60
70
80
Depth (NRA) [nm]
0.5 0 -0.5 -1
Fig. 2. SIMS profiles of the isotopic tracers through the oxide scale as a function of the recalculated depth (the oxide scale, represented by the 16O profile is displayed as a guide to the eyes): (a) 18O tracer (subplot: total oxygen signal's second derivative used to identify the domains); (b) 2D tracer.
thicknesses, as given by the standard deviation obtained during NRA experiments. This allows for the determination of a mean diffusion coefficient, and for a standard deviation associated. The diffusion coefficients obtained after this treatment are respectively Dsc = 4.6 +/− 0.9 × 10−17 cm2/s for 18O, and Dsc = 5.2 +/− 1.2 × 10−17 cm2/s for 2D. The two values can be considered as equal, considering the overlapping of their standard error range. The 18O diffusion coefficient is coherent to the one found in another study performed on alloy 690 [5], corroded in similar conditions, which yielded to a “C-regime” diffusion coefficient equal to 2.10 −18 cm2/s, which is somewhat smaller. This difference could possibly be explained by the slight composition difference of the inner spinel oxide growing on alloy 690. Indeed, comparing the elementary profiles given in reference [2] on alloy 690, and in reference [5] for alloy 600, it seems that, in the later case, the spinel oxide scale is richer in iron. This difference in the diffusion coefficients could also be linked to the 18O accumulation at the internal interface reported in [4]. The deuterium diffusion coefficient in the oxide is ten orders of magnitude lower compared with the deuterium coefficient in the alloy given in the literature [6,7] at the same temperature: the transport of hydrogen through the oxide scale could possibly be the rate limiting step in the process of hydrogen absorption by the passivated alloy in service. Indeed, the diffusion time necessary for the hydrogen transport through a membrane of thickness d where its diffusion coefficient is D, d 2/D, is slightly larger in the oxide scale than in the alloy, indicating that the rate limiting step during permeation experiments is possibly the hydrogen transport through the oxide.
migration of oxygen as a hydroxide ion is largely encouraged by steric effects due to the lesser ionic volume of the hydroxide ion (95 pm) compared with oxygen ion (140 pm) [8]. As a consequence, the validated mechanism for the hydrogen transport through the oxide scale (supposed here to be only chromium and nickel mixed spinel), leading to its uptake in the material is the following one, according to Kröger–Vink notations [9] (the notation 〈〈X〉〉 refers to the element X in solid solution in the metal phase), in three steps: - The adsorption of water molecules at the surface of the sample, - Oxide scale edification giving rise to the OHO• species - The growth of the oxide film, supported by the transport OHO• through the gradient of oxygen vacancies along the oxide. Writing the equations at the external interface (oxide film/surface interface) yields to: - water adsorption: H2 O þ s⇔H2 O s
ð1Þ
- oxide growth: ••
X
þ
H2 O s þ VO ⇔OO þ 2H þ s þ
X
•
H þ OO ⇔OHO
ð2aÞ ð2bÞ
- proton reduction: 4. Discussion
4.1. Oxide growth mechanism The diffusivity compatibility between oxygen and hydrogen allows a collaborative migration of the two species on the anionic sublattice. This
þ ′ 1H : H þ e ⇔ 2 2
ð2cÞ
Then, through the oxide scale, OHO• diffusion occurs, transporting H-species to the internal interface ðoxide scale=alloy substrate interfaceÞ:
ð3Þ
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F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
4.2. Effect of hydrogen partial pressure on the oxide growth
Finally, at the internal interface the reactions are: - oxide growth:
••
′
2〈〈Cr〉〉 þ 〈〈Ni〉〉⇔2CrCr þ NiNi þ 4VO þ 8e
ð4aÞ
- hydrogen absorption (by alloy): •
′
Also, the hydrogen partial pressure can have an effect on the OH-species production rate, according to this mechanism. In effect, combining Eqs. (2a) and (2c), and Eqs. (2b) and (2c), one can formulate the second step in an alternative way: ••
OHO þ e ⇔OO þ 〈〈H〉〉:
ð4bÞ
It is interesting to note that, for Zirconium alloys corroded in high temperature water (>300 °C) and in steam, a similar mechanism for the hydrogen uptake has recently been proposed [10]. Like Ni-base alloys [11], Zirconium alloys form an n-type semi-conducting oxide at low oxidising powers, such as those encountered in PWR primary medium conditions, with an internal growth based on oxygen, anionic diffusion [12]. This mechanism was deduced from the observation that, on the one hand, hydrogen uptake during the corrosion of these alloys in high temperature water in PWR conditions [12] or in steam [13] was strongly linked to the oxidation kinetics. On the other hand, on the fact that the amount of hydrogen uptake in these conditions is indicative of a higher solubility than predicted by H2-exposure experiments, which can be explained by different H-species solubilities in the oxide, depending on the charge carried by the ingressing hydrogen, that is, neutral charge in the case of H2-originating hydrogen (low solubility), or a single positive charge for water-originating hydrogen (higher solubility). It has been demonstrated in both this study and its first part [1] that the hydrogen uptake in nickel-base alloys in PWR primary medium conditions is not directly linked to the dissolved H2 present in the medium. This could, at first sight, seem contradictory with some studies in which hydrogen permeation tests were performed, showing an influence of the H2 fugacity with the permeating flux through nickelbase alloy membranes exposed to PWR primary medium. For instance, Morton et al. [14] showed, for low oxidising powers imposed by H2, a dependence of the flux with the hydrogen fugacity; concluding that the predominant source of hydrogen in non-strained specimens was the hydrogen dissolved in the water. However, the authors also mentioned that the flux actually was found to decrease with exposure times in these conditions. These observations are actually fully compatible with the mechanism proposed in this article. Of course, since the oxidation rate tends to decrease with exposure time, accordingly, OH-flux also has to decrease.
−
X
H2 O s þ VO þ 2e ⇔OO þ H2 þ s
ð5Þ
1 X • ′ H þ OO ⇔OHO þ e : 2 2
ð6Þ
Considering that the steps described by Eqs. (1), (5) and (6) are reversible, and that the electroneutrality is respected locally all through the oxide scale, the concentration of OHO• can be determined after resolving the coupled system composed of the laws of mass action equations of the aforementioned equations and the electroneutrality equation, e.g. ••
•
½e′ ¼ 2½VO þ ½OHO :
ð7Þ
The thermodynamic constant K associated with reactions (1), (5) and (6) are respectively:
K1 ¼
θeq ð1−θ Þ aðH2 OÞ
K5 ¼
θ aðH2 Þ ð1−θ Þ e′ eq 2 ½VO •• eq aðH2 OÞ
eq
ð8Þ
eq
eq
• eq h ′ ieq OHO e pffiffiffiffiffiffiffiffiffiffiffiffi K6 ¼ aðH2 Þ
ð9Þ
ð10Þ
where θ eq refers to the equilibrium occupation of adsorption sites for water molecules at the surface of the oxide (no hypothesis is made about the nature of these sites, only considered here as different from regular oxide sites), symbols in brackets refer to the element concentration, a(H2) refers to hydrogen fugacity as imposed by the H2 partial pressure, and a(H2O) to the activity of water.
This yields to:
• OHO ¼
1=2
3
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi1=3 1=4 1=2 1=4 aðH2 Þ 3 aðH2 Þ þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 K 6 ðK 1 K 5 aðH2 OÞÞ2=3 þ 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 1=3
K 6 ðK 1 K 5 aðH2 OÞÞ
evidencing the link between the gaseous hydrogen fugacity and the point defect concentration: increasing the hydrogen fugacity yields to an increase of the point defect concentration, thus permitting a higher hydroxide ion flux, and a susceptibly higher hydrogen uptake (originating from the water molecules) by the alloy, which could explain the results published by Morton.
0
½VO ∘∘ext ¼ eq
With a similar treatment, one can extract the oxygen vacancy concentration at the external interface and express its dependence with the hydrogen partial pressure and thermodynamic constants of the previously mentioned reactions:
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 3aðH2 Þ1=2 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 ðK 1 K 5 aðH2 OÞÞ
1=3
ð11Þ
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 2 K 6 aðH2 OÞ2=3 þ 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3
ð12Þ
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Eqs. (11) and (12) show an increase of both defects with increasing hydrogen partial pressure: this could be an alternative way of expressing the oxidising power with hydrogen partial pressure, since, here, an increase of hydrogen pressure leads to a decrease in the oxygen vacancy gradient, thus reducing the oxidation kinetics.
73
compatibility between hydrogen and oxygen diffusivities confirms the possibility for the major part of hydrogen produced during the corrosion process to be transported through the oxide scale in the hydroxide ion form. According to the general mechanism proposed for the hydrogen uptake by the alloy, hydrogen may play a significant role in crack propagation of stress corrosion cracking defects.
4.3. Consequences in SCC phenomenon Acknowledgement As shown previously, the hydrogen uptake is linked to the oxidation kinetics. A consequence of the above exposed mechanism is that, a significant hydrogen uptake by the bulk material has to be expected, in case of local break in the passive film, in the vicinity of the oxide film flaw. For instance, in case of a crack opening, an important quantity of hydrogen will be produced at the crack tip during its repassivation. Furthermore, just after the alloy repassivation, the oxide scale will locally be thinner than elsewhere; and thus the hydrogen transport through the oxide will be lessened, keeping a higher flux of hydrogen locally even after the repassivation. The subsequent, temporally high concentration of hydrogen in the alloy around the crack tip may then have significant importance on the crack propagation behaviour. On the other hand, this mechanism does not seem, at first sight, to be able to play a direct, significant role in crack initiation. 5. Conclusion Hydrogen and oxygen diffusion coefficients through a mixed chromium and nickel spinel oxide scale grown on alloy 600-like single crystals corroded in simulated PWR primary medium have been determined using isotopically labelled species. The tracer diffusion coefficients calculated assuming a short-circuit diffusion regime are respectively Dsc =4.6×10−17 cm2/s for 18O, and Dsc =5.2×10−17 cm2/s for 2D. The
The authors would like to thank Anna Fraczkiewicz from the Claude Goux Laboratory at Ecole Nationale Supérieure des Mines de SaintEtienne for providing the alloy 600 single crystal. References [1] F. Jambon, L. Marchetti, F. Jomard, J. Chêne, J. Nucl. Mater. 414 (3) (2011) 386–392. [2] M. Sennour, L. Marchetti, F. Martin, S. Perrin, R. Molins, M. Pijolat, J. Nucl. Mater. 402 (2010) 147–156. [3] Y.M. Mishin, G. Borchardt, J. Phys. III 3 (1993) 863–881. [4] L. Marchetti, S. Perrin, O. Raquet, M. Pijolat, Mater. Sci. Forum 595–598 (2008) 529–537. [5] J. Panter, B. Viguier, J.-M. Cloué, M. Foucault, P. Combrade, E. Andrieu, J. Nucl. Mater. 348 (1–2) (2006) 213–221. [6] K. Sakamoto, M. Sukisagi, Fusion Sci. Technol. 41 (2002) 912–914. [7] E. Rota, F. Waelbroeck, P. Wienhold, J. Winter, J. Nucl. Mater. 111&112 (1982) 233–239. [8] Y. Wouters, A. Galerie, J.-P. Petit, Solid State Ionics 104 (1997) 89–96. [9] F.A. Kröger, H.J. Vink, Solid State Phys. 3 (1956) 273–301. [10] M.S. Veshchunov, A.V. Berdyshev, J. Nucl. Mater. 255 (1998) 250–262. [11] L. Marchetti, S. Perrin, Y. Wouters, F. Martin, M. Pijolat, Electrochim. Acta 55 (2010) 5384–5392. [12] B. Cox, Adv. Corros. Sci. Technol. 5 (1976) 173. [13] B. Cox, J. Nucl. Mater. 264 (1999) 283–294. [14] D. Morton, S.A. Attanasio, G.A. Young, P.L. Andresen, T.M. Angeliu, in: Conference paper n° 01117 at NACE CORROSION 2001, March 11–16, 2001, Houston, USA, 2001.
Contents lists available at SciVerse ScienceDirect
Solid State Ionics journal homepage: www.elsevier.com/locate/ssi
Characterisation of oxygen and hydrogen migration through oxide scales formed on nickel-base alloys in PWR primary medium conditions Fanny Jambon a,⁎, Loïc Marchetti a, François Jomard b, Jacques Chêne c a b c
CEA, DEN, DPC, SCCME, Laboratoire d'Étude de la Corrosion Aqueuse, F-91191 Gif-Sur-Yvette Cedex, France Groupe d'Etude de la Matière Condensée (CNRS and University of Versailles Saint Quentin), 45 avenue des Etats-Unis, 78035 Versailles cedex, France UMR 8587 CEA/CNRS, Équipe Hydrogène/Matériaux de structure, CEA, DEN, DPC, SCCME, Laboratoire d'Étude de la Corrosion Aqueuse, F-91191 Gif-Sur-Yvette Cedex, France
a r t i c l e
i n f o
Article history: Received 16 April 2012 Received in revised form 29 August 2012 Accepted 16 October 2012 Available online 4 December 2012 Keywords: Alloy 600 Hydrogen Oxygen Diffusion Isotope labelling Stress corrosion cracking
a b s t r a c t This study aims at providing quantitative data concerning the oxygen and hydrogen diffusion through the oxide scale formed on alloy 600 in PWR primary medium conditions (325 °C) using 2D and 18O isotopes as markers. The diffusion coefficients were calculated according to a short-circuit diffusion of the tracers along the oxide grain boundaries (so-called “C-regime”). The values obtained are respectively Dsc =4.6+/− 0.9× 10−17 cm2/s for 18O, and Dsc =5.2+/− 1.2×10−17 cm2/s for 2D. These experiments confirm that the transport mechanism of hydrogen through the oxide scale as hydroxide ions on the anionic sublattice is highly probable. The effect of hydrogen subsequent absorption by the alloy substrate, notably in stress corrosion cracking phenomenon is discussed. © 2012 Published by Elsevier B.V.
1. Introduction In a former study [1], single nickel-base alloy crystals, with a composition close to alloy 600s one, were exposed to simulated PWR primary medium, where the gaseous hydrogen or the hydrogen in water had been replaced by deuterium. SIMS analysis of the deuterium profiles in those samples allowed the authors to identify the origin of the hydrogen absorbed in nickel-base alloys exposed to PWR primary medium as coming from the water molecule dissociation associated with the oxide scale building. A mechanism was also proposed in order to model the hydrogen transport associated with the oxide growth during the alloy passivation. This mechanism is based on the diffusion of hydrogen as an interstitial proton through the oxide lattice; or as hydroxide ion, towards the oxide anionic sub-lattice. The latter hypothesis implies the oxygen and hydrogen diffusivities through the oxide scale layer to be the same. The aim of this study is to test the validity of this statement. 2. Experimental 2.1. Sample preparation In order to improve the depth resolution of the SIMS analysis used in this study, single-crystal samples of a model alloy, with a chemical ⁎ Corresponding author. Tel.: +33 169081261. E-mail address: [email protected] (F. Jambon). 0167-2738/$ – see front matter © 2012 Published by Elsevier B.V. http://dx.doi.org/10.1016/j.ssi.2012.10.012
composition close to alloy 600 were used. The single-crystal was provided by the Ecole Nationale Supérieure des Mines de Saint-Etienne. Its growing axis is b111> and its elementary composition is given in Table 1. The microstructure of this alloy is quite homogeneous; with the presence of some small micrometric precipitates, mainly chromium carbides. The samples were prepared according to a procedure described elsewhere [1]; the final surface state, corresponding to a mechanical polishing with a 1/10 μm diamond paste, exhibited a typical mirror aspect, characteristic of low surface roughness (Ra b 6.5 nm), as measured via microrugosimetry experiments.
2.2. Corrosion experiments The samples were first passivated during 516 h at PWR nominal conditions (Table 2). The oxide scale usually formed in these conditions, according to literature, is duplex, with a protective mixed spinel internal layer, mainly nickel chromite, compact and continuous; and an external layer mainly composed of micron-sized nickel ferrite crystallites associated to the precipitation of released cations, as extensively studied in [2]. The corrosion tests were thus performed in a recirculation autoclave where the primary medium cations were removed thanks to a set of ions exchanging resins. The oxide scale formed during this first exposure was observed with SEM. As shown in Fig. 1, it is composed of small, scattered sub-micron-sized compact crystallites (30–80 nm on average) that seem to correspond
70
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Table 1 Chemical composition (wt.%) of alloy 600-like single crystals. Ni
Cr
Fe
C
Co
Mn
Ti
Cu
Si
S
P
76.99
16.99
5.57
n.c.
0.01
0.01
b0.01
0.02
0.08
0.01
0.01
to the precipitates described in [2], covering less than 50% of a homogeneous layer with nano-sized grains irresolvable with this SEM (b20 nm), the internal part of the oxide layer. In a second sequence, the pre-oxidized samples were exposed to PWR primary water in a 300 mL titanium alloy autoclave (in order to minimize nickel and iron oxide and hydroxide precipitation phenomena at the surface of the samples). The relatively small size of this autoclave also allows decreasing the usual temperature transient to 45 min. In order to compare oxygen and hydrogen diffusivities, two series of experiments were performed, where these species were isotopically marked with 18O (20%) and deuterium (99.97%) respectively. The samples were exposed to this modified PWR medium during 9 h. At the end of the corrosion experiments, the samples were stored into liquid nitrogen until their characterisation. 2.3. Sample analyses The isotopic tracer profiles through the oxide layer were obtained with Secondary Ion Mass Spectroscopy (SIMS) analyses at the GEMaC CNRS laboratory (Meudon, France), as described in reference [1], but using the SIMS high mass resolution slits in order to differentiate 18O from the molecular ion 2D 16O at mass 18 (m/Δm = 1800). Then the oxide scale thickness was measured with Nuclear Reaction Analysis (NRA), done at JANNuS (Joint Accelerators for Nanoscience and Nuclear Simulation), Saclay, France and supported by the French Network EMIR. The experiment was performed with 900 keV 2D + ions using a 2 MV Van de Graaff accelerator. The quantitative determination of 16O content was made collecting the 1258 keV protons produced with the 16O(d,p1) 17O reaction, under an analysing angle of 150°. A Mylar stopping foil of 12 μm thickness was used to filter out the backscattered deuterons. The oxygen content calibration was made using some references composed of a tantalum substrate covered by a Ta2O5 oxide of well-known thickness. 3. Results The oxide scale thickness determined by NRA after the 525 h exposure to PWR primary medium was 29 +/− 3 nm. SIMS profiles of the two tracers through the oxide scale, representative of the results obtained on the different samples, are presented in Fig. 2. The initial raw intensity-sputtering time data have been converted to a depth-scale display after rescaling the sputtering time scale to the oxide scale thickness measured by NRA so that this thickness would coincide with the oxide/alloy interface on each sample. The oxide/alloy interface is arbitrarily set at the total oxygen signal (that is, [18O + 16O] for the sample exposed to the 18O marked medium) inflexion point. For this rescaling, a mean sputtering velocity is thus calculated from the oxide thickness measured by NRA and the sputtering time at which the oxygen inflexion point occurs; and the abscise axis is multiplied by this rescaling factor, which leads to the depth scale (only valid for the oxide part of the graph), of Fig. 2.
1 µm
Fig. 1. Oxide scale observation with SEM; magnification × 10 k.
Finally, the naturally present 18O contribution, associated to the 16O of the first sequence and second sequences has been subtracted to the total 18O signal. If some 18O associated with the 16O from the second sequence also contribute to the total 18O signal, they are negligible compared with the other contributions, since the second sequence is a really short one, performed at a low oxidation rate. In any case, this correction mainly changed the signal amplitude, which was to be expected since, except at the interfaces, the oxygen signal is constant through the oxide scale. This procedure permits to have an estimation of the total oxygen amount which penetrated in the scale during the second sequence, and to make sure that the growth was limited and that this amount is compatible with the C-regime assumption (see below). Since the tracer profiles are measured in a homogenous oxide scale, it is not necessary to correct the 2D intensity as was shown in reference [1] because of the need to compare 2D intensity in the oxide and in the metal, which have quite different ionisation and abrasion rates. One can distinguish three main domains in these profiles. On the left hand side of the graph Fig. 2a, at the beginning of the 18O profile appears a tip, corresponding to the fact that the 18O concentration increases faster than the 16O at the external interface. This phenomenon is linked to the presence of an inflexion point at P1 =3.5 nm on the total oxygen signal at the beginning of the oxide scale, and can be explained by a small but significant part of external growth during the second sequence (H218O medium). This part of the profile will thus have to be spared for the calculation of the tracer diffusion coefficient [3]. The second domain (the internal oxide scale), starts at the abovementioned first inflexion point, and ends at the internal oxide/ metal interface, that is, at the second inflexion point; and finally the third part corresponds to the metallic substrate. According to the work of [4], where the accumulation of oxygen isotopic tracers at the internal interface on a similar system was evidenced, demonstrating that the growth of the oxide proceeds by short-circuit diffusion, the tracer diffusion coefficients were determined assuming a short-circuit diffusion regime (or “C-regime” [3]), through the second domain (internal oxide growth part only). The 18O and 2D profiles were fitted with a C-regime equation modified to exclude the first x−P 1 ffiffiffiffiffiffiffi , where x is the recalculated domain interval: cH ðxÞ ¼ k:erf c p 2
Table 2 Experimental PWR primary medium conditions used in this study. Temperature Pressure (static autoclave) Pressure (recirculating autoclave)
325 °C 135 bar 155 bar
Boron (H3BO3) Lithium (LiOH) H2 (pressurisation)
1000 appm 2 appm 0.28 bar at 325 °C
Dsc :t
depth, k a coefficient depending on the tracer concentration at the oxide surface and on the short-circuits density, t is the second sequence duration, and Dsc the short-circuit diffusion coefficient. Finally, the main uncertainty for the determination of the tracer diffusion coefficients is the oxide scale thickness, as determined by NRA. The tracer diffusion coefficients were thus calculated for the mean oxide thickness, and for the minimum and maximum oxide
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Ext. growth
16
16
1.5 105
1200
6 104
600
3 104
0
0 0
1
10
20
30
40
50
60
70
80
Intensity [counts/s]
Intensity [counts/s] 18O
9 104
D [counts/s]
1.2 105
400
M/Ox interface
300
9 104
200
6 104
100
3 104
0
Intensity [counts/s]
M/Ox interface 1800
Intensity [counts/s]
1.2 105
1.5 105 2
16O
O [counts/s]
2400
O [counts/s]
500 16O
18
1.8 105
600
O [counts/s]
3000
(16O+18O)"
b
1.8 105
3600
2D
a
71
0 0
10
20
30
Depth (NRA) [nm]
40
50
60
70
80
Depth (NRA) [nm]
0.5 0 -0.5 -1
Fig. 2. SIMS profiles of the isotopic tracers through the oxide scale as a function of the recalculated depth (the oxide scale, represented by the 16O profile is displayed as a guide to the eyes): (a) 18O tracer (subplot: total oxygen signal's second derivative used to identify the domains); (b) 2D tracer.
thicknesses, as given by the standard deviation obtained during NRA experiments. This allows for the determination of a mean diffusion coefficient, and for a standard deviation associated. The diffusion coefficients obtained after this treatment are respectively Dsc = 4.6 +/− 0.9 × 10−17 cm2/s for 18O, and Dsc = 5.2 +/− 1.2 × 10−17 cm2/s for 2D. The two values can be considered as equal, considering the overlapping of their standard error range. The 18O diffusion coefficient is coherent to the one found in another study performed on alloy 690 [5], corroded in similar conditions, which yielded to a “C-regime” diffusion coefficient equal to 2.10 −18 cm2/s, which is somewhat smaller. This difference could possibly be explained by the slight composition difference of the inner spinel oxide growing on alloy 690. Indeed, comparing the elementary profiles given in reference [2] on alloy 690, and in reference [5] for alloy 600, it seems that, in the later case, the spinel oxide scale is richer in iron. This difference in the diffusion coefficients could also be linked to the 18O accumulation at the internal interface reported in [4]. The deuterium diffusion coefficient in the oxide is ten orders of magnitude lower compared with the deuterium coefficient in the alloy given in the literature [6,7] at the same temperature: the transport of hydrogen through the oxide scale could possibly be the rate limiting step in the process of hydrogen absorption by the passivated alloy in service. Indeed, the diffusion time necessary for the hydrogen transport through a membrane of thickness d where its diffusion coefficient is D, d 2/D, is slightly larger in the oxide scale than in the alloy, indicating that the rate limiting step during permeation experiments is possibly the hydrogen transport through the oxide.
migration of oxygen as a hydroxide ion is largely encouraged by steric effects due to the lesser ionic volume of the hydroxide ion (95 pm) compared with oxygen ion (140 pm) [8]. As a consequence, the validated mechanism for the hydrogen transport through the oxide scale (supposed here to be only chromium and nickel mixed spinel), leading to its uptake in the material is the following one, according to Kröger–Vink notations [9] (the notation 〈〈X〉〉 refers to the element X in solid solution in the metal phase), in three steps: - The adsorption of water molecules at the surface of the sample, - Oxide scale edification giving rise to the OHO• species - The growth of the oxide film, supported by the transport OHO• through the gradient of oxygen vacancies along the oxide. Writing the equations at the external interface (oxide film/surface interface) yields to: - water adsorption: H2 O þ s⇔H2 O s
ð1Þ
- oxide growth: ••
X
þ
H2 O s þ VO ⇔OO þ 2H þ s þ
X
•
H þ OO ⇔OHO
ð2aÞ ð2bÞ
- proton reduction: 4. Discussion
4.1. Oxide growth mechanism The diffusivity compatibility between oxygen and hydrogen allows a collaborative migration of the two species on the anionic sublattice. This
þ ′ 1H : H þ e ⇔ 2 2
ð2cÞ
Then, through the oxide scale, OHO• diffusion occurs, transporting H-species to the internal interface ðoxide scale=alloy substrate interfaceÞ:
ð3Þ
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F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
4.2. Effect of hydrogen partial pressure on the oxide growth
Finally, at the internal interface the reactions are: - oxide growth:
••
′
2〈〈Cr〉〉 þ 〈〈Ni〉〉⇔2CrCr þ NiNi þ 4VO þ 8e
ð4aÞ
- hydrogen absorption (by alloy): •
′
Also, the hydrogen partial pressure can have an effect on the OH-species production rate, according to this mechanism. In effect, combining Eqs. (2a) and (2c), and Eqs. (2b) and (2c), one can formulate the second step in an alternative way: ••
OHO þ e ⇔OO þ 〈〈H〉〉:
ð4bÞ
It is interesting to note that, for Zirconium alloys corroded in high temperature water (>300 °C) and in steam, a similar mechanism for the hydrogen uptake has recently been proposed [10]. Like Ni-base alloys [11], Zirconium alloys form an n-type semi-conducting oxide at low oxidising powers, such as those encountered in PWR primary medium conditions, with an internal growth based on oxygen, anionic diffusion [12]. This mechanism was deduced from the observation that, on the one hand, hydrogen uptake during the corrosion of these alloys in high temperature water in PWR conditions [12] or in steam [13] was strongly linked to the oxidation kinetics. On the other hand, on the fact that the amount of hydrogen uptake in these conditions is indicative of a higher solubility than predicted by H2-exposure experiments, which can be explained by different H-species solubilities in the oxide, depending on the charge carried by the ingressing hydrogen, that is, neutral charge in the case of H2-originating hydrogen (low solubility), or a single positive charge for water-originating hydrogen (higher solubility). It has been demonstrated in both this study and its first part [1] that the hydrogen uptake in nickel-base alloys in PWR primary medium conditions is not directly linked to the dissolved H2 present in the medium. This could, at first sight, seem contradictory with some studies in which hydrogen permeation tests were performed, showing an influence of the H2 fugacity with the permeating flux through nickelbase alloy membranes exposed to PWR primary medium. For instance, Morton et al. [14] showed, for low oxidising powers imposed by H2, a dependence of the flux with the hydrogen fugacity; concluding that the predominant source of hydrogen in non-strained specimens was the hydrogen dissolved in the water. However, the authors also mentioned that the flux actually was found to decrease with exposure times in these conditions. These observations are actually fully compatible with the mechanism proposed in this article. Of course, since the oxidation rate tends to decrease with exposure time, accordingly, OH-flux also has to decrease.
−
X
H2 O s þ VO þ 2e ⇔OO þ H2 þ s
ð5Þ
1 X • ′ H þ OO ⇔OHO þ e : 2 2
ð6Þ
Considering that the steps described by Eqs. (1), (5) and (6) are reversible, and that the electroneutrality is respected locally all through the oxide scale, the concentration of OHO• can be determined after resolving the coupled system composed of the laws of mass action equations of the aforementioned equations and the electroneutrality equation, e.g. ••
•
½e′ ¼ 2½VO þ ½OHO :
ð7Þ
The thermodynamic constant K associated with reactions (1), (5) and (6) are respectively:
K1 ¼
θeq ð1−θ Þ aðH2 OÞ
K5 ¼
θ aðH2 Þ ð1−θ Þ e′ eq 2 ½VO •• eq aðH2 OÞ
eq
ð8Þ
eq
eq
• eq h ′ ieq OHO e pffiffiffiffiffiffiffiffiffiffiffiffi K6 ¼ aðH2 Þ
ð9Þ
ð10Þ
where θ eq refers to the equilibrium occupation of adsorption sites for water molecules at the surface of the oxide (no hypothesis is made about the nature of these sites, only considered here as different from regular oxide sites), symbols in brackets refer to the element concentration, a(H2) refers to hydrogen fugacity as imposed by the H2 partial pressure, and a(H2O) to the activity of water.
This yields to:
• OHO ¼
1=2
3
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi1=3 1=4 1=2 1=4 aðH2 Þ 3 aðH2 Þ þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 K 6 ðK 1 K 5 aðH2 OÞÞ2=3 þ 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 1=3
K 6 ðK 1 K 5 aðH2 OÞÞ
evidencing the link between the gaseous hydrogen fugacity and the point defect concentration: increasing the hydrogen fugacity yields to an increase of the point defect concentration, thus permitting a higher hydroxide ion flux, and a susceptibly higher hydrogen uptake (originating from the water molecules) by the alloy, which could explain the results published by Morton.
0
½VO ∘∘ext ¼ eq
With a similar treatment, one can extract the oxygen vacancy concentration at the external interface and express its dependence with the hydrogen partial pressure and thermodynamic constants of the previously mentioned reactions:
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 3aðH2 Þ1=2 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3 ðK 1 K 5 aðH2 OÞÞ
1=3
ð11Þ
qffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi2=3 2 K 6 aðH2 OÞ2=3 þ 33=2 aðH2 Þ1=4 þ 27aðH2 Þ1=2 −ðK 1 K 5 aðH2 OÞÞ2 K 6 3
ð12Þ
F. Jambon et al. / Solid State Ionics 231 (2013) 69–73
Eqs. (11) and (12) show an increase of both defects with increasing hydrogen partial pressure: this could be an alternative way of expressing the oxidising power with hydrogen partial pressure, since, here, an increase of hydrogen pressure leads to a decrease in the oxygen vacancy gradient, thus reducing the oxidation kinetics.
73
compatibility between hydrogen and oxygen diffusivities confirms the possibility for the major part of hydrogen produced during the corrosion process to be transported through the oxide scale in the hydroxide ion form. According to the general mechanism proposed for the hydrogen uptake by the alloy, hydrogen may play a significant role in crack propagation of stress corrosion cracking defects.
4.3. Consequences in SCC phenomenon Acknowledgement As shown previously, the hydrogen uptake is linked to the oxidation kinetics. A consequence of the above exposed mechanism is that, a significant hydrogen uptake by the bulk material has to be expected, in case of local break in the passive film, in the vicinity of the oxide film flaw. For instance, in case of a crack opening, an important quantity of hydrogen will be produced at the crack tip during its repassivation. Furthermore, just after the alloy repassivation, the oxide scale will locally be thinner than elsewhere; and thus the hydrogen transport through the oxide will be lessened, keeping a higher flux of hydrogen locally even after the repassivation. The subsequent, temporally high concentration of hydrogen in the alloy around the crack tip may then have significant importance on the crack propagation behaviour. On the other hand, this mechanism does not seem, at first sight, to be able to play a direct, significant role in crack initiation. 5. Conclusion Hydrogen and oxygen diffusion coefficients through a mixed chromium and nickel spinel oxide scale grown on alloy 600-like single crystals corroded in simulated PWR primary medium have been determined using isotopically labelled species. The tracer diffusion coefficients calculated assuming a short-circuit diffusion regime are respectively Dsc =4.6×10−17 cm2/s for 18O, and Dsc =5.2×10−17 cm2/s for 2D. The
The authors would like to thank Anna Fraczkiewicz from the Claude Goux Laboratory at Ecole Nationale Supérieure des Mines de SaintEtienne for providing the alloy 600 single crystal. References [1] F. Jambon, L. Marchetti, F. Jomard, J. Chêne, J. Nucl. Mater. 414 (3) (2011) 386–392. [2] M. Sennour, L. Marchetti, F. Martin, S. Perrin, R. Molins, M. Pijolat, J. Nucl. Mater. 402 (2010) 147–156. [3] Y.M. Mishin, G. Borchardt, J. Phys. III 3 (1993) 863–881. [4] L. Marchetti, S. Perrin, O. Raquet, M. Pijolat, Mater. Sci. Forum 595–598 (2008) 529–537. [5] J. Panter, B. Viguier, J.-M. Cloué, M. Foucault, P. Combrade, E. Andrieu, J. Nucl. Mater. 348 (1–2) (2006) 213–221. [6] K. Sakamoto, M. Sukisagi, Fusion Sci. Technol. 41 (2002) 912–914. [7] E. Rota, F. Waelbroeck, P. Wienhold, J. Winter, J. Nucl. Mater. 111&112 (1982) 233–239. [8] Y. Wouters, A. Galerie, J.-P. Petit, Solid State Ionics 104 (1997) 89–96. [9] F.A. Kröger, H.J. Vink, Solid State Phys. 3 (1956) 273–301. [10] M.S. Veshchunov, A.V. Berdyshev, J. Nucl. Mater. 255 (1998) 250–262. [11] L. Marchetti, S. Perrin, Y. Wouters, F. Martin, M. Pijolat, Electrochim. Acta 55 (2010) 5384–5392. [12] B. Cox, Adv. Corros. Sci. Technol. 5 (1976) 173. [13] B. Cox, J. Nucl. Mater. 264 (1999) 283–294. [14] D. Morton, S.A. Attanasio, G.A. Young, P.L. Andresen, T.M. Angeliu, in: Conference paper n° 01117 at NACE CORROSION 2001, March 11–16, 2001, Houston, USA, 2001.