- Email: [email protected]
Electrochemical properties and growth mechanism of passive films on Alloy 690 in high-temperature alkaline environments
Corrosion Science 52 (2010) 3444–3452
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Electrochemical properties and growth mechanism of passive films on Alloy 690 in high-temperature alkaline environments Junbo Huang, Xinqiang Wu *, En-Hou Han State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 62 Wencui Road, 110016 Shenyang, PR China
a r t i c l e
i n f o
Article history: Received 29 January 2010 Accepted 20 June 2010 Available online 27 June 2010 Keywords: A. Alloy B. XPS B. TEM C. Passive films C. High temperature corrosion
a b s t r a c t Passive films formed on Alloy 690 in high-temperature alkaline environments were investigated by potentiodynamic polarization, X-ray photoelectron spectroscopy, transmission electron microscopy and Mott–Schottky approach. Passive current density and donor density of the passive films increase with increasing temperature, due to increased diffusion rates of metallic ions and dehydration of hydroxide phases. The passive films show a duplex structure including an inner layer of fine-grained Cr oxide or spinel oxide and an outer layer of Ni–Fe spinel oxide and Ni hydroxide. A growth model of the passive films on Alloy 690 in high-temperature alkaline environments is proposed and discussed. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Alloy 600, a standard steam generators (SG) tube material for pressurized water reactor (PWR) nuclear power plants, was developed in the 1960s. However, during long-term operation, many failures occurred on Alloy 600 and were attributed by various forms of environmental degradation such as pitting, stress corrosion cracking (SCC) and intergranular attack (IGA) [1,2]. For these reasons, Alloy 690, which possesses superior SCC resistance in pure water, is proposed as an alternative to Alloy 600. Alloy 690 is a Ni-based alloy similar to Alloy 600 except for approximately double Cr content (30 wt.%). Such a high Cr content of Alloy 690 results in excellent corrosion resistance in aqueous environments. However, it has been reported that Alloy 690 is susceptible to SCC, IGA [3,4] and pitting corrosion in caustic aqueous environments containing chloride and/or thiosulfate [5,6]. Moreover, attacks in the SG tube have been found mostly at the heat transfer crevices at the tube supports. Some impurities in the SG feedwater are concentrated within the crevices and the local concentrations of impurities, such as sodium, sulfate, etc., can be increased up to six orders of magnitude [7]. It results in a localized highly aggressive environment within the crevices. It is generally recognized that nucleation and growth of corrosion cracking is strongly related to protective properties of passive films formed on metal surface [8]. Therefore, an understanding of the composition and structure of the passive films formed on Alloy 690 can provide important information on its corrosion resistance in high-temperature aqueous environments. * Corresponding author. Tel.: +86 24 2384 1883; fax: +86 24 2389 4149. E-mail address: [email protected] (X. Wu). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.016
It should be noted that the corrosion of metals in high-temperature aqueous environments is considered as an electrochemical process in nature [9]. Although the studies of oxides and passive films on stainless steels and Ni-based alloys have been extensively investigated [10–19], the relationship between the chemical composition and structure of the films and their protective properties is still a matter of considerable debate. Some articles [11,20] claimed that the surface film formed on Alloy 600 and Alloy 690 at the early stages of the oxidation process consisted of Cr2O3, Ni in a transition state of Ni2+ and Ni0 but without any Fe3+/Fe2+. According to some other articles [14,21], a Cr–Fe or Ni–Fe spinel would be formed in the passive films formed on Alloy 690. Machet et al. [11] considered that Cr2O3 would first nucleate and grow on the surface, while Ziemniak et al. [22] noted that the Cr was associated with OH predominantly to form Cr hydroxide at the alloy/ solution interface due to lower standard free energy of Cr(OH)3 than Cr2O3. Moreover, the system composed of the metallic substrate, the passive films and the electrolyte will be quite complex in high-temperature solutions. The electrochemical properties of Alloy 690, which is widely used in PWR, should be studied thoroughly in high-temperature aqueous environments. In our previous work, the influence of pH on electrochemical properties of the passive films formed on Alloy 690 in hightemperature aqueous environments was investigated [14]. In the present work, electrochemical properties of the passive films formed on Alloy 690 in alkaline solutions were further studied in the temperature range of 25–300 °C using potentiodynamic polarization and Mott–Schottky approach. The chemical composition, the element state and the thickness of the passive films were investigated by X-ray photoelectron spectroscopy (XPS) and the
3445
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
285 eV was used as a reference to correct the charging shifts. The depth profile information was obtained by sputtering the specimens with a scanning argon–ion gun operating at ion energy of 2 keV. The sputtering rate was estimated to be about 0.2 nm s1 (vs. Ta2O5). The quantification of the species in the oxide films was performed via XPSpeak4.1 peak fitting software. The specimen for TEM observation was passivated at a potential of 500 mV for 6 h at 300 °C to obtain a stable passive film. Then the thin-foil specimen was prepared using a focused ion beam (FEI Quanta 200 3D) with Ga ion sputtering. A protective tungsten strap (20 lm 1 lm, 2 lm thickness) was first deposited on the surface of the oxide scale to protect it. The specimen was finally thinned to 50–150 nm. A Tecnai G2 F20 TEM instrument equipped with an energy-dispersive spectroscopy system operating at 200 kV was used for selected area electron diffraction analysis.
structure of the passive films was analyzed by transmission electron microscopy (TEM). 2. Experimental procedures 2.1. Material and solution The chemical composition of Alloy 690 used in the present work is given in Table 1. The alloy was solution annealed at 1060 °C for 0.5 h and isothermally treated at 715 °C for 16 h to obtain intragranular and nearly continuous intergranular Cr-carbide precipitation (M23C6) [23]. Specimens of 1 cm 1 cm area were gradually ground with silicon papers up to #1500 grit. Then the specimens were welded to 316 stainless steel wires, which were subsequently shielded by heat-shrinkable polytetrafluoroethylene tubes. The testing electrolyte was chosen as 0.15 M Na2SO4 and 0.04 M NaOH solution to simulate abnormal SG crevice chemistry in the secondary circuit of PWR [14]. The testing temperature range was from 25 to 300 °C. Before the experiments, the solution was deaerated by continuously bubbling with pure nitrogen gas for 2 h.
3. Results 3.1. Potentiodynamic polarization curves Fig. 1a shows potentiodynamic polarization curves for Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at different temperatures. All the potentiodynamic polarization curves showed a primary passivation and a distinct secondary passivation at a given potential. The passivation took place obviously at temperatures above 100 °C. At 25 °C and 50 °C, the current densities were very small (in the magnitude of 106 A cm2). It is generally believed that metal will not corrode when concentrations of dissolved species are less than a value of 106 M [25,26]. Therefore they were also in the region of non-corrosion and passivation and a mean value was taken as the passive current density. A similar passive current was adopted in some articles [27,28]. The passive potentials and the passive current densities at different temperatures are listed in Table 2. Both the primary passive current density (Ip1) and the secondary passive current density (Ip2) increase with increasing temperature. Ep2, indicating the electrochemical potential for the oxidation reaction Cr2O3/CrO2 4 [14,28], descends with increasing temperature. Fig. 1b shows the Arrhenius plot of the passive current density vs. reciprocal temperature. It is the linearity relationship in the Arrhenius plot, indicating that the passivation on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution can be considered a thermally activated process. The activation energy for liquid phase diffusion generally lies below a value of 41.84 kJ mol1 [29,30]. For example, Bazan et al. [29] gave a value of 17.15 kJ mol1 for the diffusion of ferro- and ferricyanide ions in aqueous solutions of sodium hydroxide. In the present work, the corresponding activation energies are 7.61 kJ mol1 and 8.67 kJ mol1, respectively, which are in the region of the activation energy for liquid phase diffusion (below a value of 41.84 kJ mol1). Therefore, the passive process of Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution is controlled by ion diffusion in the liquid phase. Here Ip1 and Ip2 are mainly responsible for a film formation and a corrosion/film dissolution.
2.2. Electrochemical measurements Potentiodynamic polarization data were obtained using a threeelectrode cell with a platinum counter electrode and an external pressure balanced Ag/AgCl reference electrode. The reference electrode was housed in separate compartment that was maintained at ambient temperature and system pressure via a solution bridge. The reference solution was 0.1 M KCl. All electrode potentials in the present work have been converted to standard hydrogen electrode (SHE) according to the following relationship [24],
ESHE ¼ Eobs þ 0:2866 0:001ðT T 0 Þ þ 1:745 107 ðT T 0 Þ2 3:03 109 ðT T 0 Þ3
ð1Þ
where ESHE represents the electrode potential vs. SHE, Eobs the measured electrode potential, T the experimental temperature and T0 the room temperature (25 °C). Then potentiodynamic polarization experiments were performed at a scan rate of 0.5 mV s1 after pretreatment for 10 min at 200 mV below the open-circuit potential. The passive current densities (Ip) in polarization curves were obtained in the passive region of 200–300 mV above the primary passive potential (Ep1) and 150–200 mV above the secondary passive potential (Ep2). The specimens for capacitance measurements were prepared by passivating for 2 h at a constant potential in the passive range (300 mV above Ep1). After a passive film was formed, the capacitance measurement was performed at a frequency of 1 kHz and made at 50 mV intervals in the potential range from 1.0 to 0.4 V. The potentiodynamic experiments and the capacitance measurements were carried out by a EG&G Model 273 and a double phase synchronous detector (Model 5210 lock-in amplifier). 2.3. Analytical study
3.2. XPS results XPS analyses were performed using an ESCALAB 250 X-ray photoelectron spectrometer. Photoelectron emission was excited by monochromatic Al Ka source operated at 150 W with initial photo energy 1486.6 eV. The C 1s peak from contaminative carbon at
Fig. 2 shows the XPS depth profiles of the passive films formed on Alloy 690 at different temperatures. As the solution temperature increased, the thickness of the passive films increased. If the
Table 1 Chemical compositions of Alloy 690 used in the present work (wt.%). C
N
Cr
Fe
Mn
P
Si
Al
Ti
Cu
Nb
Co
Ni
0.013
0.01
29.15
9.19
0.21
0.01
0.02
0.26
0.305
0.01
0.01
0.01
Bal.
3446
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 1. (a) Potentiodynamic curves for Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution in a temperature range from 25 to 300 °C with a scan rate of 0.5 mV/s and (b) Arrhenius plot of the passive current density for the potentiodynamic curve.
percentage composition was higher (or lower) than the matrix, it was considered as an interface where the composition and structure of the passive films changed. From 100 to 300 °C, Ni–Fe-rich oxide was formed in the outer layer whereas Cr-rich oxide was formed in the inner layer. It was notable that the Ni content reached to 70% (Fig. 2c and d) and the Fe content was 18% (Fig. 2b) at the surface of passive films. The Fe content was slightly low (10%) at 200 and 300 °C, because of the selective dissolution of Fe at high temperatures [11,14]. Both the Ni and Fe contents decreased with increasing sputtering time in the outer layer. Conversely, the Cr content increased up to 40% in the inner layer. The same phenomenon was observed at 50 °C roughly. However, the passive films at 50 °C were too thin (only three experimental points) to identify its layer structure. Fig. 3 shows the Cr 2p core level spectra and their decompositions for passive films formed on Alloy 690 at different temperatures. At
the beginning of sputtering, one peak, which located a binding energy (BE) at 577.3 ± 0.3 eV, was observed in the Cr 2p spectra. A low-intense peak was only observed at a BE of 576.1 ± 0.3 eV at 300 °C. Then the spectra were systematically decomposed into three peaks: one located at a BE of 577.3 ± 0.3 eV, and two other ones located at BEs of 576.1 ± 0.3 eV and 574.0 ± 0.4 eV. With increasing sputtering time, the intensity of the signal at 577.3 eV decreased gradually and disappeared finally. At the same time, the other two components (576.1 and 574.0 eV) increased rapidly and became predominant gradually, in particular the component at 576.1 eV. By comparison with published Refs. [31,32], the signal at a BE of 574.0 ± 0.4 eV is assigned to metallic Cr in the Ni-based alloy, the signal at 576.1 eV to Cr3+ in Cr2O3 and the signal at 577.3 eV to Cr3+ in Cr(OH)3. Fig. 4 shows the Ni 2p spectra and their decompositions for passive films formed on Alloy 690 at different temperatures. Only one component and an associated satellite were detected at the beginning of sputtering. It located at a BE of 856.3 ± 0.3 eV with a satellite at + 5.8 eV (±0.2 eV). With increasing sputtering time, the intensity of the signal at 856.3 eV decreased rapidly and disappeared. The other two components and their associated satellites were detected: a signal at a BE of 852.8 ± 0.1 eV (satellite at + 5.7 (±0.4 eV) with respect to the main signal) and another signal at a BE of 854.0 ± 0.4 eV. By comparison with published Refs. [11,33,34], the signal at the BE of 856.3 ± 0.3 eV (satellite at 862.1 ± 0.2 eV) is attributed to Ni(OH)2 on the surface, two other components (and the associated satellites) are defined: a signal from metallic Ni in the alloy at a BE of 852.8 ± 0.1 eV (satellite at + 5.7 ± 0.4 eV with respect to the main signal) and another Ni oxide (NiO) at a BE of 854.0 ± 0.4 eV. Although a small shift of the BEs of Cr 2p or Ni 2p in the oxide and hydroxide is observed, the shift is within the uncertainty of the reported measurements. It should be noted that the energy difference between O1s and Cr 2p or Ni2p is kept almost constant in the peak fitting, indicating an oxide feature at 530.1 ± 0.2 eV and a hydroxide feature at 531.5 ± 0.1 eV (Fig. 5). It is believed that hydroxides will be formed prior on the surface of passive films in 0.15 M Na2SO4 and 0.04 M NaOH solution at temperatures above 100 °C. From the O 1s, Ni 2p and Cr 2p intensity ratios (Figs. 3–5), there is considerable evidence that a gradual dehydration of the hydroxide film occurs with aging. The oldest portion of the passive film, nearest the alloy interface, is mostly the oxide while the hydroxide becomes so little to be ignored. Ziemniak et al. [22] and Mclntyre et al. [35] also considered that the hydroxide would first form on the surface of the alloy and then dehydrated at high temperatures. The Fe 2p spectra (not shown here) show low-intensity signals on the surface at high temperatures, indicating that Fe is dissolved chemically or electrochemically at high temperatures [14,36]. The chemical composition of passive films seems to be independent of solution temperature investigated presently. 3.3. TEM observations Fig. 6 shows a cross-sectional TEM image of the passive films on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at 300 °C.
Table 2 Passive potentials Ep, transpassive potential Etr and passive current densities Ip at different temperatures. T (°C)
Ep1 (SHE) (V)
Ep2 (SHE) (V)
Etr (SHE) (V)
logIp1 (A cm2)
Error
logIp2 (A cm2)
25 50 100 150 200 250 300
0.5 0.5 0.95 0.80 0.99 0.80 0.80
0.16 0.14 0.07 0.02 0.04 0.13 0.18
0.33 0.42 0.39 0.35 0.36 0.23 0.21
5.14 5.21 5.00 4.31 4.03 4.07 3.87
0.30 0.29 0.12 0.04 0.10 0.02 0.05
4.82 4.61 4.38 3.74 3.55 3.38 3.21
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
3447
Fig. 2. Composition depth profiles obtained from XPS analysis for the passive films formed on Alloy 690 (a) 50 °C, (b) 100 °C, (c) 200 °C, (d) 300 °C.
Fig. 3. Cr 2p3/2 core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 200 °C, (c) 300 °C.
There was a very thin passive film formed on Alloy 690, approximate 200 nm in depth. Examples of the electron diffraction patterns from the oxide film are shown in Fig. 6a and b. The patterns from both the inner and outer layers are in agreement with the spinel structure. Since the diffraction pattern shows a spot image, the outer layer of the passive film is considered to be a single crystal oxide. In contrast, the structure of the inner layer is identified as having a ring diffraction pattern, confirming that the inner layer consists of fine-grained spinel oxide. The similar elec-
tron diffraction patterns were obtained from the oxide film of SUS 316 stainless steel at 320 °C [8]. 3.4. Mott–Schottky plots Fig. 7 shows the Mott–Schottky plot for the passive films formed on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at 200 °C. According to Eq. (2), the primary passive film on Alloy 690 is an n-type semiconductor due to a straight line with a posi-
3448
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 4. Ni 2p3/2 core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 200 °C, (c) 300 °C.
tive slope. The electronic structure of the passive films seems to be independent of solution temperature investigated presently.
1 C2
Fig. 5. O 1s core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 300 °C.
Fig. 6. Cross-sectional TEM image of the passive film after passivation at a passive potential of 500 mV for 6 h at 300 °C. (a) Diffraction pattern with spot from spinel oxide, (b) diffraction pattern with rings from fine-grained spinel oxide.
¼
2 kT E EFB e ee0 qNd
ð2Þ
where C represents the oxide capacitance, e the dielectric constant of passive film, e0 the vacuum permittivity, q the elementary charge (+e for electrons and e for holes), Nd the carrier density, k the Boltzmann constant, T the absolute temperature, E the applied potential and EFB the flat band potential. From the slope of the Mott–Schottky plot, the product eNd can be calculated. If the value 15.6 of the dielectric constant e [37,38] is taken, the carrier density can be obtained. Table 3 shows the slopes of the Mott–Schottky plots and the donor densities of the passive films at different temperatures. The values of Nd at all temperatures are very high (at a magnitude of 1020 cm3) and increase with increasing temperature. The values of Nd at high temperatures (150–300 °C) are almost one order of magnitude higher than that at low temperatures (25–100 °C).
Fig. 7. Mott–Schottky plot for passive films in 0.15 M Na2SO4 and 0.04 M NaOH solution at 200 °C.
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452 Table 3 Slopes of Mott–Schottky plots and donor densities Nd of passive films at different temperatures. T (°C)
Slope of Mott–Schottky plot (106 F2 cm4 V1)
Donor density Nd (1020 cm3)
25 50 100 150 200 300
1328 1013 317 46 33 32
1.2 1.5 4.8 33.3 46.5 47.9
4. Discussion 4.1. Duplex oxide layer Due to lack of the experimental data, the passive films formed on the Fe–Cr–Ni alloys and their protective properties are still a matter of considerable debate, especially for Alloy 690 at high temperatures. Castle and Masterson [30] recognized that the corrosion process was limited by the solution phase transport of metallic ions through pores in the oxide film, while Robertson [39] proposed that the corrosion rate was limited by the solid state diffusion of metallic ions along grain boundaries of the oxide layer. The key divergence of their opinions is the value of the activation energy of corrosion which is mainly from diffusion coefficients of metallic ions. The activation energy in the Castle–Masterson model (15–20 kJ mol1) is rather lower than Robertson’s (120 kJ mol1). In the present work, the activation energies are so low (Fig. 2) that the transport of metallic ions through micropores in the oxide film is mainly considered. The inner layer grows by access of OH to the oxide/alloy interface through the micropores in the oxide (only 1 nm in diameter [39]). The outer layer grows by the diffusion of metallic ions through micropores in the oxide film. The controlling step of the corrosion rate is the diffusion of metallic ions through micropores in the oxide film. According to the Cr 2p core level spectra and their decompositions from XPS (Fig. 3), a Cr(OH)3-based passive film grows on Alloy 690 in high-temperature aqueous solutions by the selective dissolution of the other majority components (Ni, Fe). The film growth rate now exceeds the Fe/Ni dissolution rate. With increasing passive time, the Cr(OH)3 will dehydrate into Cr2O3 on the alloy surface, which plays a main role in the inner layer of passive films. A similar behavior on the passivation process of stainless steels (at room temperature) was observed by Maurice et al. [40], who considered the conversion of Cr(OH)3 into Cr2O3 and then the further growth of Cr2O3. Asami et al. [41] noted that the films changed from a Fe-based oxide to nearly pure Cr2O3 when the alloy Cr content exceeded about 13% on a range of Fe–Cr alloys. This shows that the Cr2O3-based film do not grow by the depletion of Cr in the alloy, but by the selective release of the excess Fe to the solution. The location of each alloying component depends on whether or not it diffuses faster than the base alloy in its oxide. Faster diffusing component will pass through to the outer layer, while slower diffusing component like Cr will be oxidized without movement and remain in the inner layer. The diffusion rates of these alloy ions in the oxides have been measured [42,43] and show a well defined order 2þ
Fe2þ > Ni 2+
Cr3þ
ð3Þ 3+
Ni moves faster than the other component Cr . Moreover, the transport of NiO is faster than NiCr2O4 and Cr2O3 thermodynamically. Ni(OH)2/NiO as the main product is predominant in the outer layer (Fig. 4), containing the balance of Ni and Cr as a Ni, Cr spinel.
3449
Therefore, the outer layer is Ni-rich and the inner one retains Crrich oxide. A model was proposed earlier by Mclntyre et al. [35] and Machet et al. [11] on Alloy 600, consisting of an outer layer of Ni(OH)2/NiO and an inner of Cr2O3. However, both of them did not consider the transport of Fe ions in the oxide in the long-term oxidation. If Ni has diffused through the Cr2O3 oxide layer, it is quite reasonable to consider that Fe is also transported through the oxide layer, forming some spinel oxides with the oxide and/ or releasing Fe ions into the solution. The transport of Fe through Cr oxide at high temperatures has been observed in dry oxidation experiments [44]. Therefore a Cr or Ni spinel (Ni, Fe) (Cr, Fe)2O4 will be formed in the passive films after the long-term passivation. Fig. 8 shows the potential-pH diagram for ternary Fe–Cr–Ni system at 300 °C, which is calculated from the available thermodynamic data as described in Appendix A. The full lines on the diagram separate the stable regions for Fe–Cr–Ni species. The dash lines separate the stable regions for Cr species in the absence of Fe and Ni. In the passive potential region in the present work (the vertical dot line in Fig. 8), Cr2O3 can easily formed on the Fe–Cr–Ni alloy, even at very low potentials. Some spinel oxides (Ni, Fe) Cr2O4 and NiFe2O4 are formed with increasing passive potential, due to the transport of Fe. The results of electron diffractions (Fig. 6) also indicate that some spinel structure oxide exist in the passive films. Beverskog et al. [45], Lister et al. [46] and Lemire et al. [21] considered that the most stable product for the Fe–Cr–Ni alloy was FeCr2O4/NiFe2O4 through the thermodynamic calculation, which was more stable than a mixture of Cr2O3/NiO with a Fe oxide or hydroxide.
4.2. Influence of temperature Although the composition and the electronic structure of the passive films are the similar, the passive current density and the
Fig. 8. Potential-pH diagram for the Fe–Cr–Ni–H2O system at 300 °C. Full lines separate the stable regions for Fe–Cr–Ni species. Dash lines separate the stable regions for Cr species in the absence of Fe and Ni. The oxygen and hydrogen equilibria at 1 atm are represented by dash dot line a and b, respectively. The vertical dot line reveals the passive region in the high-temperature alkaline solution in the present work.
3450
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 9. Schematic diagrams of the growth mechanism of the passive films on Alloy 690 in high-temperature alkaline solutions.
donor density increase with increasing temperature. On the one hand, diffusion rates of metallic ions increase strongly with temperature, so that the diffusion flux tends to dominate above 100– 150 °C and the film growth rate becomes diffusion controlled [47]. Paola et al. [48] and Hakiki et al. [49] considered Fe2+ or Ni2+ as the doping species in the passive films. The high Nd values (Table 3) attribute to the highly disordered amorphous nature of the passive films, indicating that the passive films become instable with increasing temperature though the thickness of the passive films increased. This is in good agreement with the results of the polarization curves (Fig. 1a). On the other hand, the current density arises from the dehydration of hydroxide phases at higher temperature. The higher temperature means the more intension of dehydration. This is why a low-intense Cr oxide peak was observed at 300 °C (Fig. 3c). A similar behavior was suggested by Douglas et al. [50] who considered that corrosion rates were controlled by dehydration at intermediate temperature, 100–240 °C. All of these indicate that compactness and stability of the passive films are more important than the thicknesses.
alkaline solution and then dehydrating into NiO (Fig. 9b). The absence of the Fe enrichment under the oxide layer (Fig. 2) indicates that Fe also dissolves. 2þ
Ni
þ 2OH ! NiðOHÞ2 #
NiðOHÞ2 ! NiO þ H2 O
2þ
þ Fe2þ þ 2OH þ 2e ! ðNi; FeÞCr2 O4 þ H2 O
Cr2 O3 þ Ni
Based on the analyses above, a growth mechanism is proposed for the passive films of Alloy 690 in high-temperature alkaline environments. Fig. 9 shows the schematic diagrams of the growth mechanism. Because of selective dissolution of Ni and Fe from the alloy surface, the nucleation and growth of Cr(OH)3 islands prefer on the surface of Alloy 690 (surface enrichment of Cr3+) (Reaction (4)). Then the Cr(OH)3 is converted into Cr2O3 (Reaction (5)) and the nucleation of Cr2O3 islands starts (Fig. 9a).
NiO þ 2Fe3þ þ 6OH ! NiFe2 O4 þ 3H2 O
CrðOHÞ3 þ Cr þ 3OH ! Cr2 O3 þ 3H2 O þ 3e
ð4Þ
ð5Þ
The coalescence of the Cr2O3 islands leads to the formation of a continuous Cr2O3 layer. Such a layer will block the transport of Cr and further slow down the growth of Cr2O3 layer. Instead, Ni will transport through the Cr2O3 oxide layer, forming Ni(OH)2 in the
ð7Þ
With increasing passivation time, the coalescence of NiO islands leads to the formation of a continuous layer with a very thin outer layer of Ni(OH)2 (Fig. 9c). The Ni hydroxide is predominant on the surface of the passive films while the Cr hydroxide becomes very few. Fe ions do not play any important role at the early stages of passivation. However, with the transport of Fe through the oxide layer, the Cr2O3 in the inner layer will be converted into FeCr2O4 or a non-stoichiometric spinel oxide (Ni, Fe) Cr2O4. On the other hand, NiO in the outer layer will be combined with Fe and become NiFe2O4 spinel oxide. At the same time, OH in the solution transports into the oxide interface through mircopores in the oxide. The probable reactions are as follows.
4.3. Growth mechanism
Cr3þ þ 3OH ! CrðOHÞ3 #
ð6Þ
ð8Þ ð9Þ
As a result, the passive films are constructed by a Cr oxide or spinel oxide inner layer (grown from alloy matrix) and a Ni–Fe spinel oxide and Ni hydroxide outer layer (formed by a dissolution and precipitation process) (Fig. 9d). 5. Conclusions The passive films on Alloy 690 in alkaline solutions have been investigated in the temperature range from 25 to 300 °C using potentiodynamic curves, XPS, TEM and Mott–Schottky plots. The passive films in high-temperature alkaline solutions are composed of a duplex film structure. The films are very thin and consist of an inner layer of fine-grained Cr oxide or spinel oxide and an outer layer of Ni–Fe spinel oxide and Ni hydroxide. The chemical compo-
3451
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
sition and the electronic structure of the passive films seem to be independent of solution temperature investigated presently. With increasing solution temperature, the thickness of the passive films increases. However, the donor density in the passive films increases and in turn results in an increase of the passive current density. This is believed to be due to increased diffusion rates of metallic ions and dehydration of hydroxide phases with increasing temperature. It indicates that compactness and stability of the passive films are more important than the thickness. A mechanism is proposed to explain the growth process of the passive films on Alloy 690 in high-temperature alkaline environments, in which the inner layer of the passive film is believed to grow directly from the alloy matrix and the outer layer is formed by a dissolution and precipitation process.
To obtain these data, the Gibbs energies are calculated from:
DG ¼ DH T DS
ðA:1Þ
The thermodynamic values at 298 K are given in some Refs. [25,26,36,51,52]. Compound or species considered and their thermodynamic values are listed in Table A.1. For the higher temperatures, the DG and DS are modified by the integral of Cp and Cp/T, respectively. Therefore:
G0T ¼ G0298 ðT 298Þ S0298 þ
Z
Z
T
C p dT T
298
T
298
Cp dT T
For solids, gases and liquids, the heat capacity equation used is:
C p ¼ A þ BT þ CT 2
Acknowledgements This study was jointly supported by the Science and Technology Foundation of China (50871113), the Special Funds for the Major State Basic Research Projects (2006CB605001) and the Innovation Fund of Institute of Metal Research (IMR), Chinese Academy of Sciences (CAS).
ðA:2Þ
ðA:3Þ
With T in K and the values of A, B, C are constants (Table A.1). For the free energies of dissolved ions in solutions, the equation is:
T a þ ðb 1ÞS0298 C p 298 ¼ T ln 298
ðA:4Þ
Appendix A A.1. Method used for calculating potential-pH diagram Constructing potential-pH diagrams requires calculating freeenergy values, electrode potential and pH of testing solutions. One of the major obstacles is the lack of the high temperature thermodynamic data for solids and ions in solutions.
where a and b are constants at different temperatures (Table A.2). Constructing potential-pH diagrams further requires calculating a large number of reactions and determining the equilibrium phases based on the resulting line equations. Because of the indeterminacy and complexity of the possible reactions and their associated equations, the assured solids and ions are chosen such as Cr2O3, Fe2O3, CrO2 4 , (Ni, Fe) Cr2O4 and NiFe2O4.
Table A.1 Compound or species considered and thermodynamic values used. Compound or species
Cr Cr2O3 Cr2+ Cr3+ CrO2 4 Fe Fe2O3 Fe2+ Fe3+ FeCr2O4 Ni NiO Ni2O3 Ni2+ NiFe2O4 NiCr2O4 H2 O2 H2O H+ OH a
Gibbs energy at 298 K (kJ mol1)
Entropy at 298 K (J mol1 K1)
Heat capacity (J mol1 K1) 3
References 5
A
B 10
C 10
0 1058.1 176.1 215.5 727.8
23.77 81.17 18 307 50.2
24.4 119.4
9.9 9.2
3.68 15.6
[26] [26,36] [25,26,36,52] [25,26,36,52] [25,26,36,52]
0 742.2 78.9 4.6 1343 0 211.17 469.7 45.6 974.6 1275.7 0 0 285.8 0 230.1
27.3 87.3 137.7 315.9 146 29.87 38 94 128.8 125.9 124.26 130.6 205 70 0 10.9
12.7 98.3
31.7 78
2.5 14.8
22.3 29.5 157 0
31.9 0 16.3 0
235 9.2 3.3 4.2 0 0 1180
1.42 15.6 0.5 1.7 0 0 246
[26,36] [26,36] [25,26,36,52] [25,26,36,52] [26,36] [26] [26] [26] [25,26,52] [26] [26] [36,51] [36,51] [36,51] [36,51] [36]
a a a
a a
163 16.98 20.88 92 a
77.4 177.4 27.28 30 75.44 0 506.4
Heat capacities of ions in solution are calculated by Eq. (A.4).
Table A.2 Heat capacity constants of ions in solution for Eq. (A.4) (data from Refs. [25,52], in J mol1 K1). T (K)
298 573
Simple cations
Oxyanion, XOm
Simple anion
Acid oxyanion
a
b
a
b
a
b
a
b
0 80.1
1 0.348
0 107.7
1 0.972
0 232.1
1 2.618
0 240.9
1 4.1
3452
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
References [1] [2] [3] [4] [5] [6] [7] [8] [9]
[10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
D. Gómez-Briceno, M.L. Castano, M.S. Carcía, Nucl. Eng. Des. 165 (1996) 161–169. R.S. Dutta, J. Nucl. Mater. 393 (2009) 343–349. C.M. Brown, W.J. Mills, Corrosion 55 (1999) 173–186. R.S. Dutta, R. Tewari, P.K. De, Corros. Sci. 49 (2007) 303–318. R.C. Newman, W.P. Wong, H. Ezuber, A. Garner, Corrosion 45 (1989) 282–287. Y.Y. Chen, L.B. Chou, H.C. Shih, Mater. Chem. Phys. 97 (2006) 37–49. B.T. Lu, J.L. Luo, Y.C. Lu, J. Electrochem. Soc. 154 (2007) C379–C389. T. Terachi, K. Fujii, K. Arioka, J. Nucl. Sci. Technol. 42 (2005) 225–232. R.W. Bosch, D. Feron, J.P. Celis, Electrochemistry in Light Water Reactors, Reference Electrodes, Measurement, Corrosion and Tribocorrosion Issues, EFC Publication No. 49, Woodhead Publishing in Materials, Cambridge, UK, 2007. M.F. Montemor, M.G.S. Ferreira, M. Walls, B. Rondot, M.C. Belo, Corrosion 59 (2003) 11–21. A. Machet, A. Galtayries, S. Zanna, L. Klein, V. Maurice, P. Jolivet, M. Foucault, P. Combrade, P. Scott, P. Marcus, Electrochim. Acta 49 (2004) 3957–3964. X. Gao, X.Q. Wu, Z.E. Zhang, H. Guan, E.H. Han, J. Supercrit, Fluids 42 (2007) 157–163. M.C. Sun, X.Q. Wu, Z.E. Zhang, E.H. Han, J. Supercrit, Fluids 47 (2008) 309–317. J.B. Huang, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2976–2982. A. Gebert, F. Schneider, K. Mummert, Nucl. Eng. Des. 174 (1997) 327–334. H. Sun, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2565–2572. H. Sun, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2840–2847. M.C. Sun, X.Q. Wu, E.H. Han, J.C. Rao, Scr. Mater. 61 (2009) 996–999. M.C. Sun, X.Q. Wu, Z.E. Zhang, E.H. Han, Corros. Sci. 51 (2009) 1069–1072. R.S. Dutta, A. Lobo, R. Purandare, S.K. Kulkarni, G.K. Dey, Metall. Mater. Trans. A 33A (2002) 1437–1447. R.J. Lemire, G.A. McRae, J. Nucl. Mater. 294 (2001) 141–147. S.E. Ziemniak, M.E. Jones, K.E.S. Combs, J. Solut. Chem. 27 (1998) 33–66. S.Y. Qiu, X.W. Su, Y. Wen, F.G. Yan, Y.H. Yu, Y.C. He, Nucl. Power Eng. 4 (1995) 336–341. D.D. Macdonald, A.C. Scott, P. Wentrcek, J. Electrochem. Soc. 126 (1979) 908–911.
[25] C.M. Chen, K. Aral, G.J. Theus, Computer Calculated Potential-pH Diagrams to 300 °C, EPRI NP-3137, vols. 1–3, Electric Power Research Institute, Palo Alto, CA, 1983. [26] D. Cubicciotti, J. Nucl. Mater. 201 (1993) 176–183. [27] D.J. Kim, H.C. Kwon, H.P. kim, Corros. Sci. 50 (2008) 1221–1227. [28] R.M. Carranza, M.G. Alvarez, Corros. Sci. 38 (1996) 909–925. [29] J.C. Bazan, A.J. Arvia, Eletrochim. Acta 10 (1965) 1025–1032. [30] J.E. Castle, H.G. Masterson, Corros. Sci. 6 (1966) 93–104. [31] B. Elsenser, A. Rossi, Electrochim. Acta 37 (1992) 2269–2276. [32] A.A. Hermas, Corros. Sci. 50 (2008) 2498–2505. [33] D. Shintani, T. Ishida, H. Izumi, T. Fukutsuka, Y. Matsuo, Y. Sugie, Corros. Sci. 50 (2008) 2840–2845. [34] I. Czekaj, F. Loviat, F. Raimondi, J. Wambach, S. Biollaz, A. Wokaun, Appl. Catal. A: General 329 (2007) 68–78. [35] N.S. Mclntyre, D.G. Zetaruk, D. Owen, J. Electrochem. Soc. 126 (1979) 750–760. [36] D. Cubicciotti, J. Nucl. Mater. 152 (1988) 259–264. [37] G. Okamoto, T. Shibata, Corros. Sci. 10 (1970) 371–378. [38] M.E. Curley-Fiorino, M.G. Schmid, Corros. Sci. 20 (1980) 313–329. [39] J. Robertson, Corros. Sci. 29 (1989) 1275–1291. [40] V. Maurice, W. Yang, P. Marcus, J. Electrochem. Soc. 145 (1998) 909–920. [41] K. Asami, K. Hashimoto, S. Shirodaira, Corros. Sci. 18 (1978) 151–160. [42] R. Dieckmann, T.O. Mason, J.D. Hodge, H. Schmalzried, Ber. Bunsen. Phys. Chem. 82 (1978) 778–783. [43] R. Dieckmann, Solid State Ionics 12 (1984) 1–22. [44] H.J. Mathieu, D. Landolt, Corros. Sci. 26 (1986) 547–559. [45] B. Beverskog, I. Pujgdomenech, Corrosion 55 (1999) 1077–1087. [46] D.H. Lister, R.D. Davidson, E. McAlpine, Corros. Sci. 27 (1987) 113–140. [47] J. Robertson, Corros. Sci. 32 (1991) 443–465. [48] A.D. Paola, Electrochim. Acta 34 (1989) 203–210. [49] N.E. Hakiki, M.F. Montemor, M.G.S. Ferreira, M.D.C. Belo, Corros. Sci. 42 (2000) 687–702. [50] D.L. Douglas, F.C. Zyzes, Corrosion 13 (1957) 361t–374t. [51] D. Cubicciotti, J. Nucl. Mater. 182 (1991) 287–290. [52] C.M. Criss, J.W. Cobble, J. Am. Chem. Soc. 86 (1964) 5385–5390.
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Electrochemical properties and growth mechanism of passive films on Alloy 690 in high-temperature alkaline environments Junbo Huang, Xinqiang Wu *, En-Hou Han State Key Laboratory for Corrosion and Protection, Institute of Metal Research, Chinese Academy of Sciences, 62 Wencui Road, 110016 Shenyang, PR China
a r t i c l e
i n f o
Article history: Received 29 January 2010 Accepted 20 June 2010 Available online 27 June 2010 Keywords: A. Alloy B. XPS B. TEM C. Passive films C. High temperature corrosion
a b s t r a c t Passive films formed on Alloy 690 in high-temperature alkaline environments were investigated by potentiodynamic polarization, X-ray photoelectron spectroscopy, transmission electron microscopy and Mott–Schottky approach. Passive current density and donor density of the passive films increase with increasing temperature, due to increased diffusion rates of metallic ions and dehydration of hydroxide phases. The passive films show a duplex structure including an inner layer of fine-grained Cr oxide or spinel oxide and an outer layer of Ni–Fe spinel oxide and Ni hydroxide. A growth model of the passive films on Alloy 690 in high-temperature alkaline environments is proposed and discussed. Ó 2010 Elsevier Ltd. All rights reserved.
1. Introduction Alloy 600, a standard steam generators (SG) tube material for pressurized water reactor (PWR) nuclear power plants, was developed in the 1960s. However, during long-term operation, many failures occurred on Alloy 600 and were attributed by various forms of environmental degradation such as pitting, stress corrosion cracking (SCC) and intergranular attack (IGA) [1,2]. For these reasons, Alloy 690, which possesses superior SCC resistance in pure water, is proposed as an alternative to Alloy 600. Alloy 690 is a Ni-based alloy similar to Alloy 600 except for approximately double Cr content (30 wt.%). Such a high Cr content of Alloy 690 results in excellent corrosion resistance in aqueous environments. However, it has been reported that Alloy 690 is susceptible to SCC, IGA [3,4] and pitting corrosion in caustic aqueous environments containing chloride and/or thiosulfate [5,6]. Moreover, attacks in the SG tube have been found mostly at the heat transfer crevices at the tube supports. Some impurities in the SG feedwater are concentrated within the crevices and the local concentrations of impurities, such as sodium, sulfate, etc., can be increased up to six orders of magnitude [7]. It results in a localized highly aggressive environment within the crevices. It is generally recognized that nucleation and growth of corrosion cracking is strongly related to protective properties of passive films formed on metal surface [8]. Therefore, an understanding of the composition and structure of the passive films formed on Alloy 690 can provide important information on its corrosion resistance in high-temperature aqueous environments. * Corresponding author. Tel.: +86 24 2384 1883; fax: +86 24 2389 4149. E-mail address: [email protected] (X. Wu). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.016
It should be noted that the corrosion of metals in high-temperature aqueous environments is considered as an electrochemical process in nature [9]. Although the studies of oxides and passive films on stainless steels and Ni-based alloys have been extensively investigated [10–19], the relationship between the chemical composition and structure of the films and their protective properties is still a matter of considerable debate. Some articles [11,20] claimed that the surface film formed on Alloy 600 and Alloy 690 at the early stages of the oxidation process consisted of Cr2O3, Ni in a transition state of Ni2+ and Ni0 but without any Fe3+/Fe2+. According to some other articles [14,21], a Cr–Fe or Ni–Fe spinel would be formed in the passive films formed on Alloy 690. Machet et al. [11] considered that Cr2O3 would first nucleate and grow on the surface, while Ziemniak et al. [22] noted that the Cr was associated with OH predominantly to form Cr hydroxide at the alloy/ solution interface due to lower standard free energy of Cr(OH)3 than Cr2O3. Moreover, the system composed of the metallic substrate, the passive films and the electrolyte will be quite complex in high-temperature solutions. The electrochemical properties of Alloy 690, which is widely used in PWR, should be studied thoroughly in high-temperature aqueous environments. In our previous work, the influence of pH on electrochemical properties of the passive films formed on Alloy 690 in hightemperature aqueous environments was investigated [14]. In the present work, electrochemical properties of the passive films formed on Alloy 690 in alkaline solutions were further studied in the temperature range of 25–300 °C using potentiodynamic polarization and Mott–Schottky approach. The chemical composition, the element state and the thickness of the passive films were investigated by X-ray photoelectron spectroscopy (XPS) and the
3445
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
285 eV was used as a reference to correct the charging shifts. The depth profile information was obtained by sputtering the specimens with a scanning argon–ion gun operating at ion energy of 2 keV. The sputtering rate was estimated to be about 0.2 nm s1 (vs. Ta2O5). The quantification of the species in the oxide films was performed via XPSpeak4.1 peak fitting software. The specimen for TEM observation was passivated at a potential of 500 mV for 6 h at 300 °C to obtain a stable passive film. Then the thin-foil specimen was prepared using a focused ion beam (FEI Quanta 200 3D) with Ga ion sputtering. A protective tungsten strap (20 lm 1 lm, 2 lm thickness) was first deposited on the surface of the oxide scale to protect it. The specimen was finally thinned to 50–150 nm. A Tecnai G2 F20 TEM instrument equipped with an energy-dispersive spectroscopy system operating at 200 kV was used for selected area electron diffraction analysis.
structure of the passive films was analyzed by transmission electron microscopy (TEM). 2. Experimental procedures 2.1. Material and solution The chemical composition of Alloy 690 used in the present work is given in Table 1. The alloy was solution annealed at 1060 °C for 0.5 h and isothermally treated at 715 °C for 16 h to obtain intragranular and nearly continuous intergranular Cr-carbide precipitation (M23C6) [23]. Specimens of 1 cm 1 cm area were gradually ground with silicon papers up to #1500 grit. Then the specimens were welded to 316 stainless steel wires, which were subsequently shielded by heat-shrinkable polytetrafluoroethylene tubes. The testing electrolyte was chosen as 0.15 M Na2SO4 and 0.04 M NaOH solution to simulate abnormal SG crevice chemistry in the secondary circuit of PWR [14]. The testing temperature range was from 25 to 300 °C. Before the experiments, the solution was deaerated by continuously bubbling with pure nitrogen gas for 2 h.
3. Results 3.1. Potentiodynamic polarization curves Fig. 1a shows potentiodynamic polarization curves for Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at different temperatures. All the potentiodynamic polarization curves showed a primary passivation and a distinct secondary passivation at a given potential. The passivation took place obviously at temperatures above 100 °C. At 25 °C and 50 °C, the current densities were very small (in the magnitude of 106 A cm2). It is generally believed that metal will not corrode when concentrations of dissolved species are less than a value of 106 M [25,26]. Therefore they were also in the region of non-corrosion and passivation and a mean value was taken as the passive current density. A similar passive current was adopted in some articles [27,28]. The passive potentials and the passive current densities at different temperatures are listed in Table 2. Both the primary passive current density (Ip1) and the secondary passive current density (Ip2) increase with increasing temperature. Ep2, indicating the electrochemical potential for the oxidation reaction Cr2O3/CrO2 4 [14,28], descends with increasing temperature. Fig. 1b shows the Arrhenius plot of the passive current density vs. reciprocal temperature. It is the linearity relationship in the Arrhenius plot, indicating that the passivation on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution can be considered a thermally activated process. The activation energy for liquid phase diffusion generally lies below a value of 41.84 kJ mol1 [29,30]. For example, Bazan et al. [29] gave a value of 17.15 kJ mol1 for the diffusion of ferro- and ferricyanide ions in aqueous solutions of sodium hydroxide. In the present work, the corresponding activation energies are 7.61 kJ mol1 and 8.67 kJ mol1, respectively, which are in the region of the activation energy for liquid phase diffusion (below a value of 41.84 kJ mol1). Therefore, the passive process of Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution is controlled by ion diffusion in the liquid phase. Here Ip1 and Ip2 are mainly responsible for a film formation and a corrosion/film dissolution.
2.2. Electrochemical measurements Potentiodynamic polarization data were obtained using a threeelectrode cell with a platinum counter electrode and an external pressure balanced Ag/AgCl reference electrode. The reference electrode was housed in separate compartment that was maintained at ambient temperature and system pressure via a solution bridge. The reference solution was 0.1 M KCl. All electrode potentials in the present work have been converted to standard hydrogen electrode (SHE) according to the following relationship [24],
ESHE ¼ Eobs þ 0:2866 0:001ðT T 0 Þ þ 1:745 107 ðT T 0 Þ2 3:03 109 ðT T 0 Þ3
ð1Þ
where ESHE represents the electrode potential vs. SHE, Eobs the measured electrode potential, T the experimental temperature and T0 the room temperature (25 °C). Then potentiodynamic polarization experiments were performed at a scan rate of 0.5 mV s1 after pretreatment for 10 min at 200 mV below the open-circuit potential. The passive current densities (Ip) in polarization curves were obtained in the passive region of 200–300 mV above the primary passive potential (Ep1) and 150–200 mV above the secondary passive potential (Ep2). The specimens for capacitance measurements were prepared by passivating for 2 h at a constant potential in the passive range (300 mV above Ep1). After a passive film was formed, the capacitance measurement was performed at a frequency of 1 kHz and made at 50 mV intervals in the potential range from 1.0 to 0.4 V. The potentiodynamic experiments and the capacitance measurements were carried out by a EG&G Model 273 and a double phase synchronous detector (Model 5210 lock-in amplifier). 2.3. Analytical study
3.2. XPS results XPS analyses were performed using an ESCALAB 250 X-ray photoelectron spectrometer. Photoelectron emission was excited by monochromatic Al Ka source operated at 150 W with initial photo energy 1486.6 eV. The C 1s peak from contaminative carbon at
Fig. 2 shows the XPS depth profiles of the passive films formed on Alloy 690 at different temperatures. As the solution temperature increased, the thickness of the passive films increased. If the
Table 1 Chemical compositions of Alloy 690 used in the present work (wt.%). C
N
Cr
Fe
Mn
P
Si
Al
Ti
Cu
Nb
Co
Ni
0.013
0.01
29.15
9.19
0.21
0.01
0.02
0.26
0.305
0.01
0.01
0.01
Bal.
3446
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 1. (a) Potentiodynamic curves for Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution in a temperature range from 25 to 300 °C with a scan rate of 0.5 mV/s and (b) Arrhenius plot of the passive current density for the potentiodynamic curve.
percentage composition was higher (or lower) than the matrix, it was considered as an interface where the composition and structure of the passive films changed. From 100 to 300 °C, Ni–Fe-rich oxide was formed in the outer layer whereas Cr-rich oxide was formed in the inner layer. It was notable that the Ni content reached to 70% (Fig. 2c and d) and the Fe content was 18% (Fig. 2b) at the surface of passive films. The Fe content was slightly low (10%) at 200 and 300 °C, because of the selective dissolution of Fe at high temperatures [11,14]. Both the Ni and Fe contents decreased with increasing sputtering time in the outer layer. Conversely, the Cr content increased up to 40% in the inner layer. The same phenomenon was observed at 50 °C roughly. However, the passive films at 50 °C were too thin (only three experimental points) to identify its layer structure. Fig. 3 shows the Cr 2p core level spectra and their decompositions for passive films formed on Alloy 690 at different temperatures. At
the beginning of sputtering, one peak, which located a binding energy (BE) at 577.3 ± 0.3 eV, was observed in the Cr 2p spectra. A low-intense peak was only observed at a BE of 576.1 ± 0.3 eV at 300 °C. Then the spectra were systematically decomposed into three peaks: one located at a BE of 577.3 ± 0.3 eV, and two other ones located at BEs of 576.1 ± 0.3 eV and 574.0 ± 0.4 eV. With increasing sputtering time, the intensity of the signal at 577.3 eV decreased gradually and disappeared finally. At the same time, the other two components (576.1 and 574.0 eV) increased rapidly and became predominant gradually, in particular the component at 576.1 eV. By comparison with published Refs. [31,32], the signal at a BE of 574.0 ± 0.4 eV is assigned to metallic Cr in the Ni-based alloy, the signal at 576.1 eV to Cr3+ in Cr2O3 and the signal at 577.3 eV to Cr3+ in Cr(OH)3. Fig. 4 shows the Ni 2p spectra and their decompositions for passive films formed on Alloy 690 at different temperatures. Only one component and an associated satellite were detected at the beginning of sputtering. It located at a BE of 856.3 ± 0.3 eV with a satellite at + 5.8 eV (±0.2 eV). With increasing sputtering time, the intensity of the signal at 856.3 eV decreased rapidly and disappeared. The other two components and their associated satellites were detected: a signal at a BE of 852.8 ± 0.1 eV (satellite at + 5.7 (±0.4 eV) with respect to the main signal) and another signal at a BE of 854.0 ± 0.4 eV. By comparison with published Refs. [11,33,34], the signal at the BE of 856.3 ± 0.3 eV (satellite at 862.1 ± 0.2 eV) is attributed to Ni(OH)2 on the surface, two other components (and the associated satellites) are defined: a signal from metallic Ni in the alloy at a BE of 852.8 ± 0.1 eV (satellite at + 5.7 ± 0.4 eV with respect to the main signal) and another Ni oxide (NiO) at a BE of 854.0 ± 0.4 eV. Although a small shift of the BEs of Cr 2p or Ni 2p in the oxide and hydroxide is observed, the shift is within the uncertainty of the reported measurements. It should be noted that the energy difference between O1s and Cr 2p or Ni2p is kept almost constant in the peak fitting, indicating an oxide feature at 530.1 ± 0.2 eV and a hydroxide feature at 531.5 ± 0.1 eV (Fig. 5). It is believed that hydroxides will be formed prior on the surface of passive films in 0.15 M Na2SO4 and 0.04 M NaOH solution at temperatures above 100 °C. From the O 1s, Ni 2p and Cr 2p intensity ratios (Figs. 3–5), there is considerable evidence that a gradual dehydration of the hydroxide film occurs with aging. The oldest portion of the passive film, nearest the alloy interface, is mostly the oxide while the hydroxide becomes so little to be ignored. Ziemniak et al. [22] and Mclntyre et al. [35] also considered that the hydroxide would first form on the surface of the alloy and then dehydrated at high temperatures. The Fe 2p spectra (not shown here) show low-intensity signals on the surface at high temperatures, indicating that Fe is dissolved chemically or electrochemically at high temperatures [14,36]. The chemical composition of passive films seems to be independent of solution temperature investigated presently. 3.3. TEM observations Fig. 6 shows a cross-sectional TEM image of the passive films on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at 300 °C.
Table 2 Passive potentials Ep, transpassive potential Etr and passive current densities Ip at different temperatures. T (°C)
Ep1 (SHE) (V)
Ep2 (SHE) (V)
Etr (SHE) (V)
logIp1 (A cm2)
Error
logIp2 (A cm2)
25 50 100 150 200 250 300
0.5 0.5 0.95 0.80 0.99 0.80 0.80
0.16 0.14 0.07 0.02 0.04 0.13 0.18
0.33 0.42 0.39 0.35 0.36 0.23 0.21
5.14 5.21 5.00 4.31 4.03 4.07 3.87
0.30 0.29 0.12 0.04 0.10 0.02 0.05
4.82 4.61 4.38 3.74 3.55 3.38 3.21
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
3447
Fig. 2. Composition depth profiles obtained from XPS analysis for the passive films formed on Alloy 690 (a) 50 °C, (b) 100 °C, (c) 200 °C, (d) 300 °C.
Fig. 3. Cr 2p3/2 core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 200 °C, (c) 300 °C.
There was a very thin passive film formed on Alloy 690, approximate 200 nm in depth. Examples of the electron diffraction patterns from the oxide film are shown in Fig. 6a and b. The patterns from both the inner and outer layers are in agreement with the spinel structure. Since the diffraction pattern shows a spot image, the outer layer of the passive film is considered to be a single crystal oxide. In contrast, the structure of the inner layer is identified as having a ring diffraction pattern, confirming that the inner layer consists of fine-grained spinel oxide. The similar elec-
tron diffraction patterns were obtained from the oxide film of SUS 316 stainless steel at 320 °C [8]. 3.4. Mott–Schottky plots Fig. 7 shows the Mott–Schottky plot for the passive films formed on Alloy 690 in 0.15 M Na2SO4 and 0.04 M NaOH solution at 200 °C. According to Eq. (2), the primary passive film on Alloy 690 is an n-type semiconductor due to a straight line with a posi-
3448
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 4. Ni 2p3/2 core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 200 °C, (c) 300 °C.
tive slope. The electronic structure of the passive films seems to be independent of solution temperature investigated presently.
1 C2
Fig. 5. O 1s core level spectra and their decompositions of the passive films formed on Alloy 690 at (a) 100 °C, (b) 300 °C.
Fig. 6. Cross-sectional TEM image of the passive film after passivation at a passive potential of 500 mV for 6 h at 300 °C. (a) Diffraction pattern with spot from spinel oxide, (b) diffraction pattern with rings from fine-grained spinel oxide.
¼
2 kT E EFB e ee0 qNd
ð2Þ
where C represents the oxide capacitance, e the dielectric constant of passive film, e0 the vacuum permittivity, q the elementary charge (+e for electrons and e for holes), Nd the carrier density, k the Boltzmann constant, T the absolute temperature, E the applied potential and EFB the flat band potential. From the slope of the Mott–Schottky plot, the product eNd can be calculated. If the value 15.6 of the dielectric constant e [37,38] is taken, the carrier density can be obtained. Table 3 shows the slopes of the Mott–Schottky plots and the donor densities of the passive films at different temperatures. The values of Nd at all temperatures are very high (at a magnitude of 1020 cm3) and increase with increasing temperature. The values of Nd at high temperatures (150–300 °C) are almost one order of magnitude higher than that at low temperatures (25–100 °C).
Fig. 7. Mott–Schottky plot for passive films in 0.15 M Na2SO4 and 0.04 M NaOH solution at 200 °C.
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452 Table 3 Slopes of Mott–Schottky plots and donor densities Nd of passive films at different temperatures. T (°C)
Slope of Mott–Schottky plot (106 F2 cm4 V1)
Donor density Nd (1020 cm3)
25 50 100 150 200 300
1328 1013 317 46 33 32
1.2 1.5 4.8 33.3 46.5 47.9
4. Discussion 4.1. Duplex oxide layer Due to lack of the experimental data, the passive films formed on the Fe–Cr–Ni alloys and their protective properties are still a matter of considerable debate, especially for Alloy 690 at high temperatures. Castle and Masterson [30] recognized that the corrosion process was limited by the solution phase transport of metallic ions through pores in the oxide film, while Robertson [39] proposed that the corrosion rate was limited by the solid state diffusion of metallic ions along grain boundaries of the oxide layer. The key divergence of their opinions is the value of the activation energy of corrosion which is mainly from diffusion coefficients of metallic ions. The activation energy in the Castle–Masterson model (15–20 kJ mol1) is rather lower than Robertson’s (120 kJ mol1). In the present work, the activation energies are so low (Fig. 2) that the transport of metallic ions through micropores in the oxide film is mainly considered. The inner layer grows by access of OH to the oxide/alloy interface through the micropores in the oxide (only 1 nm in diameter [39]). The outer layer grows by the diffusion of metallic ions through micropores in the oxide film. The controlling step of the corrosion rate is the diffusion of metallic ions through micropores in the oxide film. According to the Cr 2p core level spectra and their decompositions from XPS (Fig. 3), a Cr(OH)3-based passive film grows on Alloy 690 in high-temperature aqueous solutions by the selective dissolution of the other majority components (Ni, Fe). The film growth rate now exceeds the Fe/Ni dissolution rate. With increasing passive time, the Cr(OH)3 will dehydrate into Cr2O3 on the alloy surface, which plays a main role in the inner layer of passive films. A similar behavior on the passivation process of stainless steels (at room temperature) was observed by Maurice et al. [40], who considered the conversion of Cr(OH)3 into Cr2O3 and then the further growth of Cr2O3. Asami et al. [41] noted that the films changed from a Fe-based oxide to nearly pure Cr2O3 when the alloy Cr content exceeded about 13% on a range of Fe–Cr alloys. This shows that the Cr2O3-based film do not grow by the depletion of Cr in the alloy, but by the selective release of the excess Fe to the solution. The location of each alloying component depends on whether or not it diffuses faster than the base alloy in its oxide. Faster diffusing component will pass through to the outer layer, while slower diffusing component like Cr will be oxidized without movement and remain in the inner layer. The diffusion rates of these alloy ions in the oxides have been measured [42,43] and show a well defined order 2þ
Fe2þ > Ni 2+
Cr3þ
ð3Þ 3+
Ni moves faster than the other component Cr . Moreover, the transport of NiO is faster than NiCr2O4 and Cr2O3 thermodynamically. Ni(OH)2/NiO as the main product is predominant in the outer layer (Fig. 4), containing the balance of Ni and Cr as a Ni, Cr spinel.
3449
Therefore, the outer layer is Ni-rich and the inner one retains Crrich oxide. A model was proposed earlier by Mclntyre et al. [35] and Machet et al. [11] on Alloy 600, consisting of an outer layer of Ni(OH)2/NiO and an inner of Cr2O3. However, both of them did not consider the transport of Fe ions in the oxide in the long-term oxidation. If Ni has diffused through the Cr2O3 oxide layer, it is quite reasonable to consider that Fe is also transported through the oxide layer, forming some spinel oxides with the oxide and/ or releasing Fe ions into the solution. The transport of Fe through Cr oxide at high temperatures has been observed in dry oxidation experiments [44]. Therefore a Cr or Ni spinel (Ni, Fe) (Cr, Fe)2O4 will be formed in the passive films after the long-term passivation. Fig. 8 shows the potential-pH diagram for ternary Fe–Cr–Ni system at 300 °C, which is calculated from the available thermodynamic data as described in Appendix A. The full lines on the diagram separate the stable regions for Fe–Cr–Ni species. The dash lines separate the stable regions for Cr species in the absence of Fe and Ni. In the passive potential region in the present work (the vertical dot line in Fig. 8), Cr2O3 can easily formed on the Fe–Cr–Ni alloy, even at very low potentials. Some spinel oxides (Ni, Fe) Cr2O4 and NiFe2O4 are formed with increasing passive potential, due to the transport of Fe. The results of electron diffractions (Fig. 6) also indicate that some spinel structure oxide exist in the passive films. Beverskog et al. [45], Lister et al. [46] and Lemire et al. [21] considered that the most stable product for the Fe–Cr–Ni alloy was FeCr2O4/NiFe2O4 through the thermodynamic calculation, which was more stable than a mixture of Cr2O3/NiO with a Fe oxide or hydroxide.
4.2. Influence of temperature Although the composition and the electronic structure of the passive films are the similar, the passive current density and the
Fig. 8. Potential-pH diagram for the Fe–Cr–Ni–H2O system at 300 °C. Full lines separate the stable regions for Fe–Cr–Ni species. Dash lines separate the stable regions for Cr species in the absence of Fe and Ni. The oxygen and hydrogen equilibria at 1 atm are represented by dash dot line a and b, respectively. The vertical dot line reveals the passive region in the high-temperature alkaline solution in the present work.
3450
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
Fig. 9. Schematic diagrams of the growth mechanism of the passive films on Alloy 690 in high-temperature alkaline solutions.
donor density increase with increasing temperature. On the one hand, diffusion rates of metallic ions increase strongly with temperature, so that the diffusion flux tends to dominate above 100– 150 °C and the film growth rate becomes diffusion controlled [47]. Paola et al. [48] and Hakiki et al. [49] considered Fe2+ or Ni2+ as the doping species in the passive films. The high Nd values (Table 3) attribute to the highly disordered amorphous nature of the passive films, indicating that the passive films become instable with increasing temperature though the thickness of the passive films increased. This is in good agreement with the results of the polarization curves (Fig. 1a). On the other hand, the current density arises from the dehydration of hydroxide phases at higher temperature. The higher temperature means the more intension of dehydration. This is why a low-intense Cr oxide peak was observed at 300 °C (Fig. 3c). A similar behavior was suggested by Douglas et al. [50] who considered that corrosion rates were controlled by dehydration at intermediate temperature, 100–240 °C. All of these indicate that compactness and stability of the passive films are more important than the thicknesses.
alkaline solution and then dehydrating into NiO (Fig. 9b). The absence of the Fe enrichment under the oxide layer (Fig. 2) indicates that Fe also dissolves. 2þ
Ni
þ 2OH ! NiðOHÞ2 #
NiðOHÞ2 ! NiO þ H2 O
2þ
þ Fe2þ þ 2OH þ 2e ! ðNi; FeÞCr2 O4 þ H2 O
Cr2 O3 þ Ni
Based on the analyses above, a growth mechanism is proposed for the passive films of Alloy 690 in high-temperature alkaline environments. Fig. 9 shows the schematic diagrams of the growth mechanism. Because of selective dissolution of Ni and Fe from the alloy surface, the nucleation and growth of Cr(OH)3 islands prefer on the surface of Alloy 690 (surface enrichment of Cr3+) (Reaction (4)). Then the Cr(OH)3 is converted into Cr2O3 (Reaction (5)) and the nucleation of Cr2O3 islands starts (Fig. 9a).
NiO þ 2Fe3þ þ 6OH ! NiFe2 O4 þ 3H2 O
CrðOHÞ3 þ Cr þ 3OH ! Cr2 O3 þ 3H2 O þ 3e
ð4Þ
ð5Þ
The coalescence of the Cr2O3 islands leads to the formation of a continuous Cr2O3 layer. Such a layer will block the transport of Cr and further slow down the growth of Cr2O3 layer. Instead, Ni will transport through the Cr2O3 oxide layer, forming Ni(OH)2 in the
ð7Þ
With increasing passivation time, the coalescence of NiO islands leads to the formation of a continuous layer with a very thin outer layer of Ni(OH)2 (Fig. 9c). The Ni hydroxide is predominant on the surface of the passive films while the Cr hydroxide becomes very few. Fe ions do not play any important role at the early stages of passivation. However, with the transport of Fe through the oxide layer, the Cr2O3 in the inner layer will be converted into FeCr2O4 or a non-stoichiometric spinel oxide (Ni, Fe) Cr2O4. On the other hand, NiO in the outer layer will be combined with Fe and become NiFe2O4 spinel oxide. At the same time, OH in the solution transports into the oxide interface through mircopores in the oxide. The probable reactions are as follows.
4.3. Growth mechanism
Cr3þ þ 3OH ! CrðOHÞ3 #
ð6Þ
ð8Þ ð9Þ
As a result, the passive films are constructed by a Cr oxide or spinel oxide inner layer (grown from alloy matrix) and a Ni–Fe spinel oxide and Ni hydroxide outer layer (formed by a dissolution and precipitation process) (Fig. 9d). 5. Conclusions The passive films on Alloy 690 in alkaline solutions have been investigated in the temperature range from 25 to 300 °C using potentiodynamic curves, XPS, TEM and Mott–Schottky plots. The passive films in high-temperature alkaline solutions are composed of a duplex film structure. The films are very thin and consist of an inner layer of fine-grained Cr oxide or spinel oxide and an outer layer of Ni–Fe spinel oxide and Ni hydroxide. The chemical compo-
3451
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
sition and the electronic structure of the passive films seem to be independent of solution temperature investigated presently. With increasing solution temperature, the thickness of the passive films increases. However, the donor density in the passive films increases and in turn results in an increase of the passive current density. This is believed to be due to increased diffusion rates of metallic ions and dehydration of hydroxide phases with increasing temperature. It indicates that compactness and stability of the passive films are more important than the thickness. A mechanism is proposed to explain the growth process of the passive films on Alloy 690 in high-temperature alkaline environments, in which the inner layer of the passive film is believed to grow directly from the alloy matrix and the outer layer is formed by a dissolution and precipitation process.
To obtain these data, the Gibbs energies are calculated from:
DG ¼ DH T DS
ðA:1Þ
The thermodynamic values at 298 K are given in some Refs. [25,26,36,51,52]. Compound or species considered and their thermodynamic values are listed in Table A.1. For the higher temperatures, the DG and DS are modified by the integral of Cp and Cp/T, respectively. Therefore:
G0T ¼ G0298 ðT 298Þ S0298 þ
Z
Z
T
C p dT T
298
T
298
Cp dT T
For solids, gases and liquids, the heat capacity equation used is:
C p ¼ A þ BT þ CT 2
Acknowledgements This study was jointly supported by the Science and Technology Foundation of China (50871113), the Special Funds for the Major State Basic Research Projects (2006CB605001) and the Innovation Fund of Institute of Metal Research (IMR), Chinese Academy of Sciences (CAS).
ðA:2Þ
ðA:3Þ
With T in K and the values of A, B, C are constants (Table A.1). For the free energies of dissolved ions in solutions, the equation is:
T a þ ðb 1ÞS0298 C p 298 ¼ T ln 298
ðA:4Þ
Appendix A A.1. Method used for calculating potential-pH diagram Constructing potential-pH diagrams requires calculating freeenergy values, electrode potential and pH of testing solutions. One of the major obstacles is the lack of the high temperature thermodynamic data for solids and ions in solutions.
where a and b are constants at different temperatures (Table A.2). Constructing potential-pH diagrams further requires calculating a large number of reactions and determining the equilibrium phases based on the resulting line equations. Because of the indeterminacy and complexity of the possible reactions and their associated equations, the assured solids and ions are chosen such as Cr2O3, Fe2O3, CrO2 4 , (Ni, Fe) Cr2O4 and NiFe2O4.
Table A.1 Compound or species considered and thermodynamic values used. Compound or species
Cr Cr2O3 Cr2+ Cr3+ CrO2 4 Fe Fe2O3 Fe2+ Fe3+ FeCr2O4 Ni NiO Ni2O3 Ni2+ NiFe2O4 NiCr2O4 H2 O2 H2O H+ OH a
Gibbs energy at 298 K (kJ mol1)
Entropy at 298 K (J mol1 K1)
Heat capacity (J mol1 K1) 3
References 5
A
B 10
C 10
0 1058.1 176.1 215.5 727.8
23.77 81.17 18 307 50.2
24.4 119.4
9.9 9.2
3.68 15.6
[26] [26,36] [25,26,36,52] [25,26,36,52] [25,26,36,52]
0 742.2 78.9 4.6 1343 0 211.17 469.7 45.6 974.6 1275.7 0 0 285.8 0 230.1
27.3 87.3 137.7 315.9 146 29.87 38 94 128.8 125.9 124.26 130.6 205 70 0 10.9
12.7 98.3
31.7 78
2.5 14.8
22.3 29.5 157 0
31.9 0 16.3 0
235 9.2 3.3 4.2 0 0 1180
1.42 15.6 0.5 1.7 0 0 246
[26,36] [26,36] [25,26,36,52] [25,26,36,52] [26,36] [26] [26] [26] [25,26,52] [26] [26] [36,51] [36,51] [36,51] [36,51] [36]
a a a
a a
163 16.98 20.88 92 a
77.4 177.4 27.28 30 75.44 0 506.4
Heat capacities of ions in solution are calculated by Eq. (A.4).
Table A.2 Heat capacity constants of ions in solution for Eq. (A.4) (data from Refs. [25,52], in J mol1 K1). T (K)
298 573
Simple cations
Oxyanion, XOm
Simple anion
Acid oxyanion
a
b
a
b
a
b
a
b
0 80.1
1 0.348
0 107.7
1 0.972
0 232.1
1 2.618
0 240.9
1 4.1
3452
J. Huang et al. / Corrosion Science 52 (2010) 3444–3452
References [1] [2] [3] [4] [5] [6] [7] [8] [9]
[10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
D. Gómez-Briceno, M.L. Castano, M.S. Carcía, Nucl. Eng. Des. 165 (1996) 161–169. R.S. Dutta, J. Nucl. Mater. 393 (2009) 343–349. C.M. Brown, W.J. Mills, Corrosion 55 (1999) 173–186. R.S. Dutta, R. Tewari, P.K. De, Corros. Sci. 49 (2007) 303–318. R.C. Newman, W.P. Wong, H. Ezuber, A. Garner, Corrosion 45 (1989) 282–287. Y.Y. Chen, L.B. Chou, H.C. Shih, Mater. Chem. Phys. 97 (2006) 37–49. B.T. Lu, J.L. Luo, Y.C. Lu, J. Electrochem. Soc. 154 (2007) C379–C389. T. Terachi, K. Fujii, K. Arioka, J. Nucl. Sci. Technol. 42 (2005) 225–232. R.W. Bosch, D. Feron, J.P. Celis, Electrochemistry in Light Water Reactors, Reference Electrodes, Measurement, Corrosion and Tribocorrosion Issues, EFC Publication No. 49, Woodhead Publishing in Materials, Cambridge, UK, 2007. M.F. Montemor, M.G.S. Ferreira, M. Walls, B. Rondot, M.C. Belo, Corrosion 59 (2003) 11–21. A. Machet, A. Galtayries, S. Zanna, L. Klein, V. Maurice, P. Jolivet, M. Foucault, P. Combrade, P. Scott, P. Marcus, Electrochim. Acta 49 (2004) 3957–3964. X. Gao, X.Q. Wu, Z.E. Zhang, H. Guan, E.H. Han, J. Supercrit, Fluids 42 (2007) 157–163. M.C. Sun, X.Q. Wu, Z.E. Zhang, E.H. Han, J. Supercrit, Fluids 47 (2008) 309–317. J.B. Huang, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2976–2982. A. Gebert, F. Schneider, K. Mummert, Nucl. Eng. Des. 174 (1997) 327–334. H. Sun, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2565–2572. H. Sun, X.Q. Wu, E.H. Han, Corros. Sci. 51 (2009) 2840–2847. M.C. Sun, X.Q. Wu, E.H. Han, J.C. Rao, Scr. Mater. 61 (2009) 996–999. M.C. Sun, X.Q. Wu, Z.E. Zhang, E.H. Han, Corros. Sci. 51 (2009) 1069–1072. R.S. Dutta, A. Lobo, R. Purandare, S.K. Kulkarni, G.K. Dey, Metall. Mater. Trans. A 33A (2002) 1437–1447. R.J. Lemire, G.A. McRae, J. Nucl. Mater. 294 (2001) 141–147. S.E. Ziemniak, M.E. Jones, K.E.S. Combs, J. Solut. Chem. 27 (1998) 33–66. S.Y. Qiu, X.W. Su, Y. Wen, F.G. Yan, Y.H. Yu, Y.C. He, Nucl. Power Eng. 4 (1995) 336–341. D.D. Macdonald, A.C. Scott, P. Wentrcek, J. Electrochem. Soc. 126 (1979) 908–911.
[25] C.M. Chen, K. Aral, G.J. Theus, Computer Calculated Potential-pH Diagrams to 300 °C, EPRI NP-3137, vols. 1–3, Electric Power Research Institute, Palo Alto, CA, 1983. [26] D. Cubicciotti, J. Nucl. Mater. 201 (1993) 176–183. [27] D.J. Kim, H.C. Kwon, H.P. kim, Corros. Sci. 50 (2008) 1221–1227. [28] R.M. Carranza, M.G. Alvarez, Corros. Sci. 38 (1996) 909–925. [29] J.C. Bazan, A.J. Arvia, Eletrochim. Acta 10 (1965) 1025–1032. [30] J.E. Castle, H.G. Masterson, Corros. Sci. 6 (1966) 93–104. [31] B. Elsenser, A. Rossi, Electrochim. Acta 37 (1992) 2269–2276. [32] A.A. Hermas, Corros. Sci. 50 (2008) 2498–2505. [33] D. Shintani, T. Ishida, H. Izumi, T. Fukutsuka, Y. Matsuo, Y. Sugie, Corros. Sci. 50 (2008) 2840–2845. [34] I. Czekaj, F. Loviat, F. Raimondi, J. Wambach, S. Biollaz, A. Wokaun, Appl. Catal. A: General 329 (2007) 68–78. [35] N.S. Mclntyre, D.G. Zetaruk, D. Owen, J. Electrochem. Soc. 126 (1979) 750–760. [36] D. Cubicciotti, J. Nucl. Mater. 152 (1988) 259–264. [37] G. Okamoto, T. Shibata, Corros. Sci. 10 (1970) 371–378. [38] M.E. Curley-Fiorino, M.G. Schmid, Corros. Sci. 20 (1980) 313–329. [39] J. Robertson, Corros. Sci. 29 (1989) 1275–1291. [40] V. Maurice, W. Yang, P. Marcus, J. Electrochem. Soc. 145 (1998) 909–920. [41] K. Asami, K. Hashimoto, S. Shirodaira, Corros. Sci. 18 (1978) 151–160. [42] R. Dieckmann, T.O. Mason, J.D. Hodge, H. Schmalzried, Ber. Bunsen. Phys. Chem. 82 (1978) 778–783. [43] R. Dieckmann, Solid State Ionics 12 (1984) 1–22. [44] H.J. Mathieu, D. Landolt, Corros. Sci. 26 (1986) 547–559. [45] B. Beverskog, I. Pujgdomenech, Corrosion 55 (1999) 1077–1087. [46] D.H. Lister, R.D. Davidson, E. McAlpine, Corros. Sci. 27 (1987) 113–140. [47] J. Robertson, Corros. Sci. 32 (1991) 443–465. [48] A.D. Paola, Electrochim. Acta 34 (1989) 203–210. [49] N.E. Hakiki, M.F. Montemor, M.G.S. Ferreira, M.D.C. Belo, Corros. Sci. 42 (2000) 687–702. [50] D.L. Douglas, F.C. Zyzes, Corrosion 13 (1957) 361t–374t. [51] D. Cubicciotti, J. Nucl. Mater. 182 (1991) 287–290. [52] C.M. Criss, J.W. Cobble, J. Am. Chem. Soc. 86 (1964) 5385–5390.