Corrosion Science 55 (2012) 126–132
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Application of the electrochemical microcapillary technique to study intergranular stress corrosion cracking of austenitic stainless steel on the micrometre scale Mathias Breimesser a,b,⇑, Stefan Ritter b, Hans-Peter Seifert b, Sannakaisa Virtanen c, Thomas Suter a a b c
EMPA – Swiss Federal Laboratories for Materials Testing and Research, Lab for Corrosion and Materials Integrity, CH-8600 Dübendorf, Switzerland Paul Scherrer Institute PSI, Nuclear Energy and Safety Research Department, Lab for Nuclear Materials, 5232 CH-Villigen PSI, Switzerland Friedrich-Alexander University, Department of Materials Science, Institute for Surface Science and Corrosion (LKO), Martensstrasse 7, D-91058 Erlangen, Germany
a r t i c l e
i n f o
Article history: Received 25 July 2011 Accepted 10 October 2011 Available online 15 October 2011 Keywords: A. Stainless steel B. Potentiostatic B. SEM C. Intergranular corrosion C. Stress corrosion
a b s t r a c t The electrochemical microcapillary technique was applied for the ﬁrst time to study the electrochemical dissolution signals from single growing cracks due to stress corrosion cracking of thermally sensitised AISI 304 stainless steel in potassium tetrathionate solution. Potentiostatic current measurements on initiating cracks along grain boundaries were performed. Typical current signals of potentiostatic and potentiodynamic measurements consisted of a series of current peaks showing fast rise and exponential decay, sometimes interrupted by passive phases. The results indicate that stress corrosion crack growth is a discontinuous process of passive and active phases, which might be explained by the ﬁlm rupture model. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction Intergranular (IG) stress corrosion cracking (SCC) is one of the major causes of material degradation. It is a common phenomenon in the petroleum, chemical, and nuclear sectors [1–6], where austenitic steels and nickel-based alloys are widely used as construction materials for pipes and reactor vessels. In these applications, SCC can lead to catastrophic failures of critical structural components, resulting in signiﬁcant economic losses and potential safety hazards. It is well known that thermally sensitised austenitic stainless steels are prone to IG SCC: The formation of chromium carbides along the grain boundaries during heat treatment leads to Cr depletion, and renders the grain boundaries vulnerable towards corrosion attack [7–9]. A variety of effects on SCC of austenitic stainless steels are still the subject of research : The inﬂuence of applied  and local stress values , surface preparation , inhibiting and accelerating effects of different metal cations [11,14,15] and anions on IG SCC [11,15–17], cold working , or the degree of sensitisation (DOS) [9,11], to name a few. Despite the large amount of data, the mechanism of IG SCC is still under discussion. Several mechanistic models have been ⇑ Corresponding author at: Paul Scherrer Institute PSI, Nuclear Energy and Safety Research Department, Lab for Nuclear Materials, 5232 CH-Villigen PSI, Switzerland. Tel.: +41 56 310 27 93; fax: +41 56 310 21 99. E-mail address: [email protected]
(M. Breimesser). 0010-938X/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2011.10.011
developed, and are used to describe IG SCC in different corrosion systems. Among them are the slip-dissolution model , the hydrogen embrittlement model , the surface mobility model , and the ﬁlm induced cleavage model . The ﬁlm induced cleavage model is typically applied to noble alloy systems, like Au–Ag or Au–Cu, where SCC proceeds through brittle fracture of de-alloyed surfaces. The surface mobility model explains SCC as a surface diffusion process where atoms diffuse away from the crack tip. It was applied to a wide range of systems. However it is stated in , that for systems with pre-existing crack paths, like sensitised austenitic stainless steels, the surface mobility model alone is not ideal, and an anodic dissolution process is responsible for SCC as well. Gutman critically discusses the slip-dissolution model (also called ‘‘ﬁlm rupture model’’) for SCC in  and points out inconsistencies in its theory. He concludes that the range of SCC processes cannot be explained by a single mechanism, but should be treated as multiple separate phenomena which have to be studied individually. Gomez-Duran and Macdonald postulate a hydrogen-embrittlement mechanism for SCC of 304 stainless steel in aqueous thiosulphate solution [25,26]. They state that the discrete crack propagation events they found are of the size of grain boundaries and, therefore, are better explained by hydrogen-embrittlement than by a slip-dissolution mechanism. In an ongoing project at PSI and EMPA, an attempt is made to gain new insights into the underlying mechanisms of SCC and to improve the detection and monitoring of SCC by the electrochemical noise (EN) technique. The application of EN to detect SCC has
M. Breimesser et al. / Corrosion Science 55 (2012) 126–132 Table 1 Chemical composition of the studied austenitic stainless steel (in wt.%). C
2. Experimental 2.1. Material and sample preparation
Fig. 1. Experimental setup of the electrochemical microcapillary and bending cell.
Two types of samples were prepared from an AISI 304 stainless steel rod (UNS S30400). The chemical composition is shown in Table 1. The steel was solution annealed at 1050 °C and water quenched. One sample type was thermally sensitised at 620 °C for 8 h and the other type was sensitised for 24 h to create samples with different susceptibilities to IG SCC. Double loop electrochemical potentiodynamic reactivation tests, according to JIS G 05801986, were done on both sample types. Degree of sensitisation (DOS) values of 15.2% (8 h) and 22.4% (24 h) were found. Solution annealed material without any further heat treatment showed a DOS value below 0.2%. Typical grain sizes were in the range of 50–150 lm. The heat treated material was cut into coupon-shaped samples with the dimension 25 10 0.9 mm. Prior to electrochemical measurement, the samples were ground with SiC paper up to 4000 grit and cleaned with ethanol. Three to four measurements were performed in a row on a sample.
2.2. Experimental setup already been described in several studies [25–33], but depending on the observed system and the measuring technique, noise signals varied and were interpreted differently. In one study  SCC was indicated by current pulses of low frequency and high intensity, while samples not prone to SCC exhibited current signals of higher frequency and low intensity. In other works [29,30,33], two types of transients were found in strain tests on stainless steel in thiocyanate solution: sharp peaks, interpreted as the result of slip dissolution, and broader peaks, probably caused by the dissolution of metal along the intergranular faces. Moreover, AISI 304 stainless steel at high temperature has been examined with EN . Sequences of transients, separated by low noise intervals, were detected, which were interpreted as different stages of fracture (IG SCC, transgranular (TG) SCC, ductile fracture). However, the noise signals alone cannot distinguish IG SCC from TG SCC. An earlier study showed that SCC of stainless steel in thiosulphate solutions was indicated by a rise in baseline current and fast current transients . The combination of EN with other detection techniques, like acoustic emission and elongation measurements, to detect IG SCC is presented in several studies [34,35]. In the present study, the characteristic potential and current signals during the initiation and propagation of single stress corrosion cracks are measured on the microscale at EMPA, using the electrochemical microcell technique [36–39]. In a future study, the results of the microcapillary experiments will be compared to large scale, macroscopic EN measurements, and it may become possible to identify characteristic signal parameters in macroscopic noise signals that can be correlated with the onset of SCC. This paper reports the ﬁrst known instance where electrochemical signal measurement tracked the growth of single cracks due to SCC. A sensitised 304 austenitic stainless steel was investigated in potassium tetrathionate solution. The used microcapillary technique [40–43] was already applied to study the inﬂuence of surface stress on the initiation of pitting in different systems [43–45]. However, this technique has not been used before to track the electrochemical signals of a growing stress corrosion crack.
The electrochemical microcapillary technique was applied for all experiments. The technique uses silicon coated glass capillaries with diameters between 1 and 1000 lm as electrochemical cells. The precise positioning of the capillary on a speciﬁc surface spot allows the electrochemical investigation of single surface features, such as grain boundaries or inclusions. Contrary to scanning techniques , where macroscopic samples are immersed in an electrolyte and scanned with a probing microelectrode, the microcapillary technique allows microscopic spots on the surface to be exposed to electrolyte and polarised [43,47]. The capillary was mounted in a modiﬁed socket of a light microscope objective carousel, and the sample was ﬁxed on the microscope stage. This setup allowed the microscope to be used for searching suitable positions on the sample and then, once such a place was found, the capillary could be switched in for the measurement. A saturated calomel electrode, connected through a Haber-Luggin capillary was used as reference and a platinum wire was used as counter electrode (see Fig. 1). A high resolution potentiostat/galvanostat (Jaissle IMP 83 PC T-BC) with a current detection limit of 1 fA was used. A scanning frequency of 9 Hz was applied. A Faraday cage shielded the experimental setup from external electromagnetic and acoustic interference. All measurements were performed under atmospheric pressure and at 23 °C. A three point bending cell was used to stress the specimen by constant deformation  (see Fig. 1). A screw was used to deform the sample before the measurement. The applied nominal load was calculated from the displacement of the screw. No further calculations of the real stress values at the sample surfaces were made. Macroscopic immersion tests of stressed samples showed that SCC initiation occurred on the sample below the yield stress. (Rp0.225 °C = 291 MPa, measured on solution annealed material). For the ﬁrst microelectrochemical measurements, different stress values above and below the yield stress were applied. It was found that higher stress values generally induced more active spots and larger resulting cracks. These experiments were therefore easier
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to analyse. For the systematic measurement of potentiostatic current curves, two degrees of plastic deformation were chosen: a nominal stress level of 360 MPa and 720 MPa, which could be set easily by turning the screw of the deformation cell by 45° or 90°. A capillary with a diameter of 60 lm was used for all experiments, exposing a 2800 lm2 area of the working electrode. The exposed area was similar in size to the grains. Therefore, most measurements hit at least one grain boundary. Smaller capillaries made it increasingly difﬁcult to ﬁnd active spots on the surface. Larger capillaries often covered more than one growing crack and the measured current signals became more difﬁcult to interpret. The choice of a suitable electrolyte was crucial for the successful measurement of IG SCC. An electrolyte had to be found, which speciﬁcally attacks the sensitised grain boundaries of stressed samples and does not cause other types of localised corrosion. Furthermore, the initiation of cracks should occur in a reasonable time frame. Prior to measurement, the sample surface had to be scanned for an active crack initiation point. Therefore, crack initiation should happen within seconds or minutes to avoid too many measurements on passive surface spots. Sodium and magnesium chloride solutions at various pH values, which are widely used to study localised corrosion, were found to initiate IG SCC with the samples immersed and under a static deformation in immersion tests, but severe pitting occurred as well. Experiments with sodium thiosulphate at pH values between 3 and 7 did not show these problems. Thiosulphate solutions are also known to initiate IG SCC in sensitised stainless steel . However, using the applied load conditions, no cracks occurred in a reasonable time frame. An acidic solution of potassium tetrathionate was found to be a suitable electrolyte. Tetrathionate solutions have already been used to induce IG SCC in sensitised austenitic steels [48,49]. A 10 mM tetrathionate solution with a pH value of 2.2 induced IG SCC under the applied loading conditions with very little pitting corrosion. Unstressed samples showed usually minor IG corrosion. This same electrolyte was used for all experiments.
2.3. Experimental procedure After grinding, the samples have been scanned for IG SCC initiation sites. Only few speciﬁc places on a sample surface act as crack initiation sites. The systematic scanning of a surface with the electrochemical microcapillary has already been explained in detail
Fig. 2. OCP measurements on a passive and an active surface spot. The potential on the passive spot stabilised around 300 mV/SHE, while the active spot showed a ﬂuctuating potential around 50 mV/SHE. SEM images of the active surface spot showed an intergranular crack of 140 lm.
[50,51]. A similar approach was chosen in this work. Measurements were performed point by point in an array pattern over a rectangular area. Open circuit potentials (OCP) were measured on each spot for 10 to 30 s. Values around 250 mV/SHE indicated passivity (see Section 3.1.1), active sites showed values below 40 mV/ SHE. These active spots usually initiated SCC. Depending on the applied load, 10 to 50 surface spots had to be checked to ﬁnd an active one. Initiation sites on partially sensitised samples were often large MnS-inclusions. On fully sensitised samples, additional active sites on bulk material were found as well. It might be that not all active sites could be found with the applied technique. Some sites might not initiate within 30 s after electrolyte contact. On active and passive surface spots, both potentiodynamic (scan rate of 0.5 mV/s) and potentiostatic current measurements were performed. Potentiostatic current measurements are comparable with potentiostatic electrochemical current noise measurements.
3. Results and discussion 3.1. Microcapillary measurements 3.1.1. OCP and potentiodynamic current measurements OCP measurements on a passive and an active spot are shown in Fig. 2. Both surface spots were initially active. Most tested spots showed a passive behaviour after a few seconds or minutes, indicated by an OCP value around +300 mV/SHE. On active surface spots metal dissolution occurs immediately, indicated by an OCP value between 0 to 50 mV/SHE. Polarization curves demonstrate 100 times higher current densities on active spots (Fig. 3). Once the capillary is positioned on the surface, OCP values allow active spots to be identiﬁed.
3.1.2. Monitoring of crack initiation and propagation Crack initiation and propagation of single cracks were recorded during potentiostatic current measurements at 100 mV/SHE. The potential was chosen to be higher than the OCP of an active spot but lower than the OCP of a passive spot. Passive spots showed usually quite stable current densities. Values varied between surface spots in the range of 0.2 to 0 mA/cm2 (Fig. 4a and Fig. 5a, dotted curves). These values correspond well with those values of
Fig. 3. Potentiodynamic polarization curves of an active (thin line) and a passive (bold line) surface spot. A lower OCP value and a higher anodic current density were measured on the active spot. Visual inspection of the active spots after the measurements showed intergranular cracks, while passive spots showed no corrosion attack.
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Fig. 4a. Exp. A: Potentiostatic current time curve measured over 2 h on an active surface spot (2800 lm2) of a sensitised sample. The current indicates an extended passive phase between 6000 and 7000 s. For comparison, the current time curve of a passive spot on an unstressed sample is indicated by the dotted line.
Fig. 4b. Detail view of Fig. 4a. The steep rise in current at 3950 s is caused by three overlapping peaks of 5–10 nA.
polarization curves (Fig. 3). No corrosion attack was observed after the measurements. Current curves of active surface spots on the highly and on the partially sensitised material are shown in Figs. 4a and 4b and Figs. 5a and 5b. For comparison, current curves measured on passive spots on the two materials are plotted as well. Crack growth is observed immediately after the measurement was started, indicated by high initial currents. Crack growth is indicated by a series of current peaks. The typical peak shape includes a fast rise and slow decay, indicating sudden metal dissolution followed by a repassivation process. High current signals were more often found on spots on the highly sensitised material. But generally, a large variation between single surface spots was found, and the difference between the two materials was small. Current peak shapes were similar on both materials. Some measurements showed passive phases between active peaks which can last several minutes. Current signals measured during these periods were similar to those measured on passive surface spots, showing stable signals in the region of 0.2–0 mA/cm2. These ﬁndings suggest that the growth of intergranular cracks is a discontinuous process, even on the scale of single cracks. Since the
Fig. 5a. Exp. B: Potentiostatic current time curve of an active surface spot on the partially sensitised sample. The current signal reaches more than 100 nA in the ﬁrst seconds of the measurements, followed by a fast drop. Single transients appear after 800 s and active dissolution occurs at 1200 s. The current time curve of a passive spot on an unstressed sample is indicated by the dotted line.
Fig. 5b. Detail view of Fig. 5a. Due to the applied potential a cathodic current is measured. Three peaks can be identiﬁed, reaching 6 nA and exhibiting fast rise and exponential decay.
measurements were performed under static deformation conditions, no dynamic load was necessary to reinitiate crack growth after the passive phase. At potentials below 0 mV/SHE, repassivation of formerly active spots was observed. This implies that crack propagation is governed by anodic dissolution. Single current peaks (Fig. 5b) show an exponential decay. According to the ﬁlm rupture model, as schematically shown in Fig. 6, such transients are interpreted as individual local events of ﬁlm rupture, anodic metal dissolution depending on material sensitisation, and repassivation. When plotted on a double logarithmic scale, the slopes of such peaks appear linear. If several events occur simultaneously, the current peaks overlap and the overall current signal increases. Large current peaks (Fig. 4a) deviate from this shape. Their slopes indicate additional continuous metal dissolution along the grain boundaries. These ﬁndings of discontinuous crack propagation are in good agreement with the studies from Newman and Healey . They concluded that along a crack front, even within one grain, all possible stages of the ﬁlm rupture mechanism are present simultaneously.
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Fig. 8. SEM picture of a crack on a partially sensitised sample. The crack initiation occurred most likely at the partially dissolved MnS-inclusion in the centre of the picture. The inner rim of the capillary is indicated by the dotted line. Fig. 6. Stages of a single event as explained by the ﬁlm rupture model: 0. Passivity. The surface is protected by an oxide layer (black). ? Low current. (1) Rupture of the oxide ﬁlm and start of metal dissolution. ? Current peak. (2) Metal dissolution and ongoing repassivation. ? Decreasing current. The repassivation rate depends on the DOS. 3. Surface is passive again. ? Low current.
Fig. 7. SEM picture of a large crack formed on a fully sensitised sample. Selective dissolution of the metal occurred along grain boundaries. Crack growth apparently continued underneath and even beyond the capillary rim. The inner rim of the capillary is indicated by the dotted line. The outer edge of the capillary can be seen as a circular imprint on the metal.
3.2. Crack characterisation 3.2.1. SEM investigation on active and passive sites Figs. 7 and 8 show SEM images of the surface spots corresponding to the experiments shown in Figs. 4a and 4b and Figs. 5a and 5b. After the measurements, formed cracks appeared as very ﬁne lines on the surface. In order to improve their visibility, they were slightly opened by bending the specimens. Crack lengths ranged from 20 to 400 lm. Crack growth was not limited to the surface area exposed to the electrolyte, but could continue under the capillary rim, especially on highly stressed samples (see Fig. 7). It is assumed that the electrolyte continuously ﬁlled the crack void and allowed crack growth well beyond the capillary dimension. No signiﬁcant leakage of the capillary was observed, even when a crack grew past the capillary rim. The bulk material did not show signiﬁcant corrosion damage, and no pitting was observed. Anodic dissolution was limited to cracks and MnS inclusions. Intergranular cracks usually grew perpendicular to the applied tensile stress direction, but deviations from this direction were observed, e.g. in Fig. 7. This can be explained by the geometry of the stressing cell. It used a single screw to apply stress to the sample, and the intensity and direction of stress on the surface varied around the centre of highest deformation. Therefore, the exact direction of crack growth depended on the position of a measure-
Fig. 9. Mean surface crack lengths after a 2 h measurement on differently sensitised materials at high (720 MPa) and low (360 MPa) load. A total of 21 cracks were analysed. A higher degree of deformation and sensitisation induces larger cracks.
ment relative to the centre of highest deformation, as well as the orientation of the grain boundaries present. Crack initiation often occurred at MnS-inclusions. In some cases, the MnS-inclusions were still intact and the crack followed the metal–inclusion interface. In other cases, a partial or complete break out of an inclusion was observed (e.g. shown in Fig. 8). Crack initiation also occurred at surface defects created during grinding. These observations are in good agreement with ﬁndings reported in the literature . However, some cracks did not show a clear initiation point (e.g. in Fig. 7). In these cases, cracks initiated probably at microscopic defects or weak spots of the oxide layer, which formed during surface preparation. As it was not known in advance, which surface spots would show crack initiation, no investigation of such defects before the measurement was possible. The inﬂuence of sensitisation and stress on mean crack lengths after 2 h of potentiostatic measurements is shown in Fig. 9. With increasing stress longer cracks were observed on both materials. However, the difference is small for the partially sensitised material. On average, longer cracks were measured on the fully sensitised material, but the difference with partially sensitised material was small. 3.2.2. Comparison of charge values and corresponding cracks Charge transfer values from current data were compared to the resulting crack sizes, as seen by SEM, using Faraday’s law. For the calculation, the cracks were assumed to have the shape of a semi-elliptical disk with constant width and walls into the grain
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Measurements which showed signs of pitting or crevice corrosion are indicated as open points in Fig. 11. They typically show higher charge values. They were excluded for the best ﬁt calculation. Despite the limited number of measurements, the calculations might indicate that the formed cracks are caused by IG crack growth via metal dissolution. Considering the fact that crack propagation is a discontinuous process, it is suggested that the ﬁlm rupture model ﬁts best. 4. Conclusions
Fig. 10. Schematic illustration for the calculation of crack wall areas and the width of the dissolution zone: Crack shapes were approximated as semi-elliptical discs (dark grey) along the grain boundaries (light grey) with visible surface length a, depth b and a constant width c. The area A of a crack wall was derived from a and b. The width c was calculated by dividing the volume of dissolved metal by the wall area A (dissolved metal volumes were calculated from the current curves using Faraday’s law).
Fig. 11. Plot of estimated crack ﬂank surface vs. charge ﬂown during the potentiostatic current measurements. A trend line is plotted, indicating a linear relationship between crack size and measured charge. Open dots indicate measurements that showed corrosion damage not directly related to crack growth and were omitted for the calculation of the trend line.
boundary (see Fig. 10). The volume of a crack can be derived from the width of the metal dissolution zone (labelled c in Fig. 10) and the area of a crack wall (labelled A in Fig. 10). A similar calculation has already been described . To estimate the mean crack depth (labelled b in Fig. 10), the cross sections of 11 selected cracks were investigated using light microscopy. A mean ratio of maximum crack depth to surface length of 0.47 ± 0.07 was found. This value was used to estimate the depths of all other investigated cracks and to calculate their wall areas. Current values from potentiostatic measurements were used to calculate electrode charge transfer. In order to do this, a baseline correction was applied to each current measurement to account for the cathodic passive current. The resulting wall area and charge transfer values are plotted in Fig. 11. A best ﬁt line for the data has a slope of 2.2 mC/mm2. Passive charge transfer values measured on freshly ground steel samples were found to be more than 20 times less. Therefore, signiﬁcant additional metal dissolution occurred in chromium depleted regions along the grain boundaries during the active phase. Using Faraday’s Law the width of the dissolution front was calculated. A mean width of 80 nm was found, which compares reasonably well to values from other studies on a similar system .
A procedure to measure local IG SCC propagation with the electrochemical microcapillary technique was established. For the ﬁrst time it was possible to track the growth of single cracks under electrochemically controlled conditions. Thermally sensitised AISI 304 stainless steel samples were subjected to static plastic deformation and exposed to 0.01 M potassium tetrathionate solution at pH 2.2. Under these conditions, single cracks were induced and the characteristic current signals of crack propagation were measured. Intergranular crack growth was found to be a stepwise process. Current signals showed characteristic peaks with fast rise and exponential decay. These ﬁndings on the initiation and propagation of cracks can be explained by the ﬁlm rupture model. Cracks formed after 2 h potentiostatic measurements were found to be longer on fully sensitised material than on partially sensitised material. Higher stress on the surface also increased crack length and additionally led to more active sites on the surface. The comparison of estimated crack ﬂank surfaces and measured current charge showed a very good correlation, further supporting the ﬁlm rupture model. Variations in microstructure are likely responsible for the scatter of data, so a larger data set and further studies are necessary to improve the understanding of the correlation between the signal and crack shape. The obtained charge transfer values exceed the expected values necessary for crack wall passivation and indicate signiﬁcant anodic metal dissolution along the grain boundaries. Acknowledgements Financial support of by the Swiss Federal Nuclear Safety Inspectorate (ENSI) is gratefully acknowledged. We would also like to acknowledge J.A. DeRose for useful discussion and help with preparation of the manuscript. References  F. Ford, P. Andresen, Corrosion in nuclear systems, in: Corrosion Mechanisms in Theory and Practice, CRC Press, 2002, pp. 605–642.  Y. Sato, T. Atsumi, T. Shoji, Continuous monitoring of back wall stress corrosion cracking growth in sensitized type 304 stainless steel weldment by means of potential drop techniques, Int. J. Pres. Ves. Pip. 84 (2007) 274–283.  A. Almubarak, M. Belkharchouche, A. Hussain, Stress corrosion cracking of sensitized austenitic stainless steels in Kuwait petroleum reﬁneries, AntiCorros. Methods Mater. 57 (2010) 58–64.  R. Kilian, A. Roth, Corrosion behaviour of reactor coolant system materials in nuclear power plants, Mater. Corros. 53 (2002) 727–739.  P. Scott, An Overview of Materials Degradation by Stress Corrosion in PWRs, in: EUROCORR 2004, Nice, France, 2004.  P. Fauvet, F. Balbaud, R. Robin, Q.T. Tran, A. Mugnier, D. Espinoux, Corrosion mechanisms of austenitic stainless steels in nitric media used in reprocessing plants, J. Nucl. Mater. 375 (2008) 52–64.  V. Cíhal, I. Kasová, Relation between carbide precipitation and intercrystalline corrosion of stainless steels, Corros. Sci. 10 (1970) 875–881.  Y.F. Yin, R.G. Faulkner, P. Moreton, I. Armson, P. Coyle, Grain boundary chromium depletion in austenitic alloys, J. Mater. Sci. 45 (2010) 5872–5882.  A. Abou-Elazm, R. Abdel-Karim, I. Elmahallawi, R. Rashad, Correlation between the degree of sensitization and stress corrosion cracking susceptibility of type 304H stainless steel, Corros. Sci. 51 (2009) 203–208.  K.R. Trethewey, Some observations on the current status in the understanding of stress-corrosion cracking of stainless steels, Mater. Des. 29 (2008) 501–507.
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