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In situ and ex situ characterisation of oxide films formed on strained stainless steel surfaces in high-temperature water
Applied Surface Science 252 (2006) 8580–8588 www.elsevier.com/locate/apsusc
In situ and ex situ characterisation of oxide films formed on strained stainless steel surfaces in high-temperature water Yoichi Takeda a, Tetsuo Shoji a, Martin Bojinov b,*, Petri Kinnunen c, Timo Saario c a
Fracture and Reliability Research Institute, Tohoku University, 6-6-01 Aoba Aramaki Aobaku, 980-8579 Sendai, Japan b Department of Physical Chemistry, University of Chemical Technology and Metallurgy, 1756 Sofia, Bulgaria c VTT Industrial Systems, VTT Technical Research Centre of Finland, P.O. Box 1704, FIN-02044 VTT, Espoo, Finland Received 1 October 2005; received in revised form 25 November 2005; accepted 25 November 2005 Available online 4 January 2006
Abstract A possible approach in describing the role of the environment in the phenomena behind initiation and propagation of a stress corrosion crack is to assume that the transport of matter and charge through the oxide film on the material is one of the rate-controlling factors. Straining of the bulk material may affect the transport rates of ionic defects, such as vacancies and interstitials, through the oxide film. The aim of the present work has been to verify the applicability of combined slow strain rate tests (SSRT) and contact electric resistance (CER) measurements to assess the influence of strain on the electric properties of oxide films on AISI 316L stainless steel with or without prior cold work in simulated boiling water reactor (BWR) coolant conditions. The SSRT-CER measurements have been combined with ex situ characterisation of the oxide films after experiments using electron spectroscopy for chemical analysis (ESCA) and scanning electron microscopy (SEM). The results suggest that the effect of strain on the resistance of the oxide films seems to correlate with the effect of the same parameter on the Cr(III) concentration in the inner layer of the oxide. In addition, important differences between the concentration of Ni and Fe in the outer layer formed on stressed and unstressed surface have been observed. Based on the mixed-conduction model for oxide films, an attempt is made to evaluate the effect of straining on the electric properties of the oxide films and to correlate these effects with the changes in film composition and structure. # 2005 Elsevier B.V. All rights reserved. Keywords: Stainless steel; Simulated nuclear reactor coolant; Contact electric resistance; Slow strain rate test; Oxide film; Kinetic model
1. Introduction When determining the susceptibility of a material to environmentally assisted cracking (EAC), it is important to consider the effect of the environment and stress/strain field on the properties of the surface oxide films on the material. Generally, oxidation and/or dissolution take place at the interface of the material and the environment around the crack tip, where stress/strain field exists due to external and/or residual stress, as suggested in the slip-oxidation, coupled environment fracture, and film-induced cleavage models for EAC [1–3]. Once the crack advances or external loading varies, distribution of stress/strain at the crack front changes. This change in the distribution around the crack tip results in the generation of dynamic strain. Consequently, the interfacial
* Corresponding author. Tel.: +359 889 298 679; fax: +359 2 8682 036. E-mail address: [email protected] (M. Bojinov). 0169-4332/$ – see front matter # 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2005.11.073
reactions and mass transport involved in oxidation and dissolution processes in the crack tip region take place on the dynamically strained surface. In a plausible working hypothesis, the straining of the bulk material is thought to affect the transport of ionic defects within the oxide film in the vicinity of the oxide film/metal interface, resulting in a vacancy flux into the bulk metal. Such an influence may determine the rate at which the vacancies react with dislocations at the interface and in the bulk metal. So far, a few attempts have been made to investigate the influence of the stress or strain on the oxide film during its growth and restructuring mainly in ambient temperature electrolytic media [4–12]. However, practically no literature references that treat the influence of dynamic straining are available, especially for simulated boiling water reactor (BWR) coolant conditions. A possible approach to understand how dynamic straining affects film growth is to measure in situ the contact electric resistance (CER) of the oxide formed on straining surface. Such measurements have been already applied for the investigation
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of passive film formation and stability in high-temperature water [13–19]. As a result, the mobilities of electronic and ionic current carriers in the films have been evaluated for different combination of materials and environments. The electric resistance of the formed oxide film, which is measured by CER technique, is one of the important properties of that film that can be correlated to its composition and structure [13–19]. Therefore, the CER technique is thought to be a powerful tool to gain in situ information of material stability in the environment. The general aim of further work in that direction [20–22] has been to test the applicability of combined CER and slow strain rate tests (SSRT) to assess the influence of active straining and strain rate on the electric properties, composition and morphology of oxide films on austenitic stainless steels in simulated BWR conditions. A specific goal of the present paper is to evaluate the effect of prior cold work applied to the steel substrate on the electric properties, composition and morphology of such oxide films. For the purpose, the CER of the oxide film on actively deforming surfaces of stainless steels formed during SSRT was measured continuously in situ by means of the SSRT-CER technique. In addition, the composition and morphology of the formed oxides have been characterised ex situ using scanning electron microscopy (SEM) and electron spectroscopy for chemical analysis (ESCA). A preliminary hypothesis on the relationship between the oxide film composition, its electrical and electrochemical properties and the processes occurring at the alloy/film interface leading eventually to crack initiation is proposed. 2. Experimental The materials used were commercial mill-annealed AISI 316L stainless steel and 20% cold-worked steel of the same grade, the cold work procedure being cross-rolling. The chemical composition of the material is shown in Table 1. The SSRT-CER measurement set-up shown in Fig. 1 consists of a plate-type SSRT specimen and a metal probe electrode. In this study, a cylinder-shaped probe made of pure iridium (99.7%, provided by Goodfellow) was used in order to exclude the influence of oxidation of the probe on measured film resistance [21,22]. The probe contacting the plate-type Table 1 Composition of the studied material (wt%) AISI 316L C Si Mn P S Ni Cr Mo N Fe
0.019 0.44 1.66 0.026 0.008 11.18 16.85 2.07 0.065 Bal.
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Fig. 1. Scheme of the principle of the SSRT-CER technique: (left) during the formation of oxide under straining and (right) during the measurement of film resistance.
specimen was 2 mm in diameter, that can be considered to be a large enough area to give average information on the electrical properties of the oxide. The SSRT-CER technique is based on a four point resistance measurement arrangement, as described in greater detail earlier [20–22]. The SSRT-CER system consisted of a uniaxial SSRT system and a CER system. The SSRT-CER system was connected to a recirculating BWR water loop in order to maintain the desirable water chemistry conditions. The dissolved oxygen (DO) concentration was maintained by purging nitrogen/oxygen mixed gas into the water tank of the loop. The inlet oxygen level was 1.0 ppm. Both inlet and outlet DO and conductivity were monitored using a Model 3600 Analyser (Orbisphere Laboratories) and a Kemotron 2902 conductivity meter. The internal volume of the autoclave containing the SSRT-CER measuring device was about 0.7 dm3 and the water flow rate was 0.3–0.5 dm3 min 1. The SSRT-CER specimens had a gauge length of 18 mm and were 6 mm in width and 3 mm in thickness. The surface of the specimens were finished by emery paper up to #1200 grid and then rinsed with acetone and pure water, before subjecting to the high-temperature water environment. The SSRT-CER measurements were performed at the strain rate of 3 10 7 s 1 at 288 8C and a pressure of 10 MPa. After stabilization of the environment, the specimen was preoxidized for 96 h without straining. The preoxidation was performed in order to minimise the effect of the initial transient oxidation occurring after reaching the target experimental conditions. Once the straining was started, the contact resistance of the gauge part of the respective specimen was measured every 5–20 h during the straining. Three independent runs have been performed for each type of specimen. In one and the same run, the values of the resistance have been measured seven times at each point in time and the result has been averaged. The straining and measurement of the film resistance were interrupted at predetermined strain level. The formed surface oxide on the interrupted SSRT specimens were analysed by using SEM and ESCA.
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The secondary electron images were obtained by using HITACHI S-4700 cold filed emission scanning electron microscope in which the electron beam accelerated at 15 kV as a light sources was utilised. The SEM system was equipped with two secondary electron detectors. The upper secondary electron detector was placed above the objective lens and secondary electrons were detected through the magnetic field of the lens. The lower secondary electron detector was placed in the specimen chamber. The signals from both detectors were corrected together for the observations. The vacuum level of the specimen chamber (<1.0 10 3 Pa) was maintained by an optionally installed turbomolecular pump. Subsequently the SEM analyses, a scanning ESCA microprobe (PHI Quantum 2000) was employed for the analysis of formed oxide film using Al Ka radiation and an acceleration voltage of 15 keV. The take-off angle of photoelectrons was selected to 908 in order to maintain higher sensitivity of satellite peak and chemical shift. The incident beam diameter used was 200 mm. The average information of the oxide composition within the beam area was obtained. The obtained ESCA spectra were evaluated by using PHI Multipak software. For the depth profiling, Ar-ion milling was used with 2 keV acceleration voltage and 2 mm 2 mm raster. The vacuum level of the specimen chamber of the ESCA system was maintained less than 1.5 10 6 Pa during sputtering by utilising a differential pump. Sputter time was converted to sputter depth by using an experimentally derived sputter rate of 4.03 nm min 1. This rate was determined using a SiO2 sputter standard known to have an oxide thickness of 104 nm. Both the SEM and ESCA analyses were carried out on the oxide films formed on both grip part and gauge part (i.e. unstressed and stressed part) of the specimens. 3. Results 3.1. Influence of stress on the electrical properties of the oxides Fig. 2 shows the stress/strain curve, as well as the dependence of the average contact resistance of the oxide film
Fig. 2. Film resistance normalised to the value in the absence of stress as depending on the strain level for AISI 316L without cold work. Experiment ended at a strain level of 0.18.
on mill-annealed AISI 316L on the strain level. The resistance was normalised to the initial value in the absence of strain in order to isolate the influence of strain on that parameter from the respective influence of other experimental conditions that could lead to a change in the contact area in different experiments (e.g. geometrical and/or energetical surface inhomogeneity [21]). Prior to straining, the specimens were oxidised in simulated BWR water for 96 h and the resistance of the film has been continuously monitored. A logarithmic kinetics of the increase of the resistance has been observed and constant values have been reached after 20–30 h of oxidation. During straining, occasionally, the resistance was measured on the unstressed part of the specimens and was found to be constant within the limits of experimental error. The values of the resistivity of the film on the unstressed part (used as a reference value to normalise the resistance during straining) calculated by using the estimate of the whole film thickness on the unstressed surfaces obtained by ESCA (see below) has been found to be (1.04 0.2) 107 V cm. The increase of the resistance observed up to a strain level of ca. 0.03 has been discussed by us previously [21,22] and has been correlated to the increase of the Cr(III) concentration in the inner layer of the oxide films formed on separate specimens subjected to interrupted SSRT tests as evidenced by ESCA depth profiling [21]. Such a correlation was explained on the basis of the fact that the conductivity of the inner layer (which determines the transport of matter and charge through the whole oxide film) has been reported to decrease with increasing Cr content in that layer because of the decrease of the concentration of ionic current carriers that act as electron donors [23]. On the other hand, the subsequent decrease of resistance with straining up to ca. 0.06 could be explained in terms of the decrease of the concentration of Cr(III) in both the inner and outer layers due to an enhancement of the transpassive dissolution of Cr by straining [22]. The increase of the resistance beyond strain levels of 0.08 is most probably due to a decrease in the contact surface area via formation of cracks (see below). Fig. 3 shows the stress/strain curves and the dependence of the resistance of the oxide film on AISI 316L with 20% cold
Fig. 3. Film resistance normalised to the value in the absence of stress as depending on the additional strain for AISI 316L with 20% cold work from two sets of experiments ended at different additional strain levels.
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Fig. 4. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel without prior cold work. Experiment ended at a strain level of 0.18.
Fig. 5. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel with 20% cold work. Experiment ended at a strain level of 0.037.
work (normalised to its value before the start of the straining for the particular experiment) as depending on the additional strain applied to the samples in two separate experiments that were terminated at different additional strain levels. The differences between the stress/strain curves shown in Figs. 2 and 3 can be explained by the fact that the surface to which dynamic straining was applied during the experiments shown in Fig. 3 has been already subjected to a strain of ca. 0.2 (corresponding to the level of prior cold work of 20%). The normalised resistance increases sharply for low levels of additional strain and subsequently decreases gradually, reaching at the end of the experiments values essentially similar to those before the additional strain was applied. This behaviour is markedly different from the dependence of the resistance on strain shown in Fig. 2 and demonstrates that the proposed measurement method can be used to differentiate
between the electrical properties of films formed on dynamically strained surfaces of stainless steels with or without prior cold work. An explanation for the different behaviour exhibited by the specimens in Figs. 2 and 3 has been sought by ex situ analysis of the morphology, thickness and indepth composition of the formed oxides. 3.2. Surface morphology The SEM micrographs of the oxide film formed on the stressed and unstressed parts of the steel specimen without cold work are shown in Fig. 4. The differences between the morphology of the oxide on the stressed and unstressed parts are relatively small, except for the higher number of large crystals seen on the top of the oxide layer formed on the unstressed part. The morphology of the oxide is that of a typical
Fig. 6. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel with 20% cold work. Experiment ended at an additional strain level of 0.065.
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BWR bilayer film (see e.g. Ref. [25]), the grain size of the outer layer being of the order of 200–300 nm, whereas that of the underlayer (seen e.g. in the left micrograph) is an order of magnitude smaller. In addition, several cracks and protrusions are observed in the micrograph of the film formed on the stressed part of the specimen. They are most probably due to processes of cracking and repair of the oxide layer under the influence of applied strain. Figs. 5 and 6 show the morphology of the oxides formed on the stressed and unstressed parts of the two AISI 316L specimens subjected to a 20% prior cold work before straining. In these cases, the difference between the morphology of the oxide formed on the stressed and unstressed parts is somewhat more significant than that in Fig. 4, i.e. the number of crystallites and crystallite size of the outer layer formed on the stressed part is significantly lower. This indicates that on a substrate that is subjected to cold work, an oxide with different properties is formed during dynamic straining. Further proof for the difference between the oxides was sought in their ESCA indepth compositional profiles. The second type of experiments on CW samples has been terminated at quite high strain levels, which most probably resulted in changes the mechanism of repassivation due to the significant extent of cracking of the oxide. Thus, the ESCA results for the second type of experiments on CW samples have been omitted because they were not considered to be representative for the whole area of the stressed region. 3.3. In-depth composition of the oxides Fig. 7 shows the ESCA depth profiles of the concentrations of oxygen in the films formed on the stressed and unstressed surfaces of the AISI 316L steel with and without prior cold work. An estimate of the thickness of the oxide was obtained by approximating the curves with sigmoidal functions (solid lines in Fig. 7) and taking the value of depth at the inflection point of the curves as the position of the metal/film interface The values for the film thicknesses formed on the stressed and unstressed surfaces on the sample without prior cold work are rather close to each other (960 and 860 nm, respectively), i.e. the straining does not seem to influence the thickness of the film significantly (ca. with 10%) in accordance with our earlier results [21,22]. On the other hand, the estimated thickness of the oxide formed on the stressed part of the specimen is ca. 20% smaller (700 nm versus 900 nm) in the case of the experiment stopped at an additional strain of 0.037, i.e. when prior cold work is applied to the substrate, the differences in thickness due to straining appear to be larger. Fig. 8 shows the depth profiles of Cr(III) and metallic Cr in the films on the specimens with and without cold work. It is important to mention that in the present paper, the concentrations of the individual metallic constituents have been normalised to the total concentration of metallic elements. This has been done in order to exclude the influence of oxygen on the in-depth profiles of the metallic elements in the oxide. In addition, the depth profiles of the concentrations of metallic Cr and Cr(III) have been evaluated
Fig. 7. ESCA depth profiles of the atomic concentration of oxygen in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
separately from the respective spectra in analogy to earlier work [21,22]. Another estimate of the film thickness was obtained from these profiles by approximating them with sigmoidal functions and averaging the depth at which both Cr(III) and Cr(0) profiles have inflection points. These estimates agree very well with those resulting from the evaluation of the oxygen profiles, which lends credibility to the values of the film thickness derived from ESCA measurements. The maximum concentrations of Cr(III) in the inner layer of the film formed on the unstressed surface are significantly (more than two times in the case of the specimen without cold work and ca. 60% for the cold worked specimen) higher than those in the inner layer formed on the stressed surface. This probably demonstrates that straining enhances the rate of transpassive dissolution of Cr, as concluded also earlier [22]. It is worth mentioning that the maximum content of Cr(III) in the inner layer on the unstressed part of the cold-worked specimen is ca. 33% higher than that for the specimen without cold work. This could be explained by a higher mobile dislocation density in the cold-worked specimen affecting the transport of Cr atoms that are preferentially oxidised at the metal/oxide interface. However, a more detailed description of the oxidation process
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Fig. 8. ESCA depth profiles of the Cr(III) and metallic Cr contents normalised to the total cation content in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
at the alloy/film interface remains out of the scope of the present work. Fig. 9 depicts the depth profiles of Fe and Ni in the oxides on the specimens with and without cold work. The Fe content in the outer layer of the oxide formed on the stressed surface is significantly higher (and correspondingly, the Ni content significantly lower) than that formed on the unstressed surface, which may indicate that the portion of the spinel (NiFe2O4) in the outer layer with respect to hematite (Fe2O3) is smaller in the case of the oxide formed on the stressed surface. This might be in turn connected to an enhanced dissolution of Fe from the substrate through the inner layer with subsequent deposition which creates an excess iron in the outer layer. The differences between the Fe/Ni ratios in the outer layer formed on the stressed and unstressed part of the specimen are more significant for the oxide formed on the specimen with 20% cold work, which may once again indicate a change in the oxidation mechanism of cold-worked specimens in comparison to mill-annealed ones. In that connection, it is worth mentioning that the importance of the Ni content in the oxides formed on strained AISI 316L surfaces in BWR conditions for the overall resistance of the material to cracking has been emphasised by earlier authors [24].
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Fig. 9. ESCA depth profiles of the Fe and Ni contents normalised to the total cation content in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
4. Discussion 4.1. Influence of straining on the electrochemical and transport properties of the oxide film In order to explain the results reported earlier [21,22] and those obtained in the present paper, a mechano-electrochemical model of the growing oxide layer during straining has to be advanced. A plausible approach towards such a model is to examine the effect of applied stress on oxidation by considering its interaction with the oxide growth induced stress and stress relaxation modes in growing oxides. A number of recent papers treat the effect of stress on the behaviour of passive films and thermal oxide films [8–10], the different forms of stress generated by the primary oxide growth and further oxidation, as well as the stress relaxation modes [26–28]. On the other hand, the electrochemical aspects of the growth and dissolution of passivating films can be understood in terms of the point defect model (PDM) by Macdonald and co-workers [29–31]. The PDM describes the metal/oxide/ solution system at steady state in terms of generation, transport and consumption of point defects (anion and cation vacancies and interstitial cations) under the influence of both the concentration and potential gradients in the oxide and at its
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Fig. 10. A simplified scheme of the process during oxidation/dissolution of stainless steel in high-temperature water according to MCM.
boundaries with the metal and the solution. The model has been further developed by some of us as mixed-conduction model (MCM) [13–16]. A possible explanation of the experimental results can be given in terms of the MCM. A simplified scheme of the processes considered during growth and dissolution of the oxide on a stainless steel in high-temperature water is shown in Fig. 10. The film thickness versus strain data reported earlier indicate that the overall kinetics of the film growth (reaction sequence k2–k4 in Fig. 10) is not significantly influenced by the strain level. However, as near steady state the growth of the film is balanced by its dissolution, the effect of strain level on film thickness may appear small. Based on the experimental evidence for that from earlier work [21], in the following the reactions involving cation transport through an oxide with constant thickness (reaction sequences k1–k3 and k1i–k3i–kr) are assumed to predominate over the processes involving oxygen vacancies. As indicated in a range of papers [11,12,26,28], the growth of an oxide film via vacancy generation and transport initiates compressive stresses at the metal/film interface regardless of the exact growth mechanism. In addition, compressive surface stresses at the 316L/passive film interface have been measured in situ even under an externally applied tensile stress up to a certain critical value [9]. Thus, only internal mechanisms of stress relief are believed to control the influence of stress on oxidation up to the maximum in resistance shown in Fig. 2. Two important stress relief mechanisms are considered. First, it has been recently demonstrated by first-principle calculations [12] that the oxidation-induced strain can be compensated by a relatively large flow of cations from the metal into the oxide. Such an emission which is accompanied by the generation of vacancies in the metal can be favoured by the tensile stress applied to the substrate (the self-diffusion through a vacancy mechanism is accelerated by tensile and sloweddown by compressive stress [10,28]). In the case of the inner layer formed on stainless steel (a (CrxFe1 x)3 dO4 spinel [25]), taking into account the processes shown in Fig. 10, it is
presumed that the rate of reaction k1i is increased above the unstressed state. A second mechanism of stress generation and relief can be proposed by considering the fact that the stoichiometry of the inner layer varies from the metal/film to the film/solution interface [12,26]. Due to the higher oxidation state of iron at the oxide layer/water interface when compared to that at the steel/inner layer interface, the internal part of the layer is probably subjected to larger compressive stresses than the external part. As a result, the rate of cation transport through the film in its outer part will be higher [26], and so will be the rate of the reaction sequence k3i–kr that control iron dissolution and subsequent reprecipitation to form an outer layer (Fig. 10). Thus, a similar line of reasoning could explain qualitatively the changes in the Fe/Ni ratio in the outer layer under the influence of strain by taking into account the enhancement of the transport of interstitial cations at the inner layer/water interface and an unchanged rate constant for outer layer redeposition (kr) which is mainly determined by nickel and iron solubility in high-temperature water. Accordingly, it is hypothesized that for strain levels up to around 0.02–0.03, the emission of interstitial cations, their diffusion through the inner layer and the outer layer deposition are enhanced. This most probably results in an increase of the concentration of Cr(III) in the inner layer, and hence the electronic resistance of that layer, in accordance with experimental observations [21,22]. Above a strain of 0.02– 0.03, it can be assumed that tensile stresses start to dominate also at the steel/inner layer interface resulting in an increase of both the fluxes of cation vacancies and interstitial cations. Thus, also the reactions of Cr dissolution is accelerated, and therefore, Cr(III) concentration is decreased, as observed experimentally. Such a hypothesis is in agreement with recent observations that tensile stresses are measured upon potential steps in oxides with predominating cation vacancy conduction mechanism [11]. Summarising, the inner layer formed on 316SS in hightemperature water is considered to be a normal spinel of the FeCr2O4 type which is a p-type semiconductor, i.e. cation vacancies are predominant ionic current carriers in it. However,
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interstitial conduction through it is also significant, and Fe dissolves through the inner layer via that route. At lower strains, the rate of the processes involving interstitial cations are increased, and cation vacancies are accumulated in the inner layer due to their slower rate of transport, wheras at higher strains, also the vacancy transport is accelerated. In other words, no qualitative change in the conduction mechanism is envisaged in the present treatment, only a quantitative change in the rate constants and diffusion coefficients and hence of the ratio between concentrations of different defects in the film. 4.2. Effect of prior cold work applied to the substrate In principle, the results obtained in the present work for the specimens subjected to 20% cold work could be qualitatively explained on the basis of the model framework presented above. Indeed, it can be assumed that due to the residual stress (associated with the residual strain accumulated in the metal substrate during cold work), the tensile stresses start to dominate at the metal/film interface almost immediately after the additional strain is applied, therefore the resistance of the oxide decreases during most part of the experiments (Fig. 3). This decrease in the resistance could be once more correlated to the decrease of the Cr(III) content in the inner layer of the oxide on the stressed part of the specimens in comparison to the unstressed part, as shown by the ESCA profiles (see Fig. 8). This in turn shows that straining accelerates markedly the transpassive dissolution of Cr (and probably also the dissolution of Ni) from the substrate. The acceleration of the transpassive dissolution will enhance the flux of cation vacancies in the inner layer of the oxide and also at the metal/film interface until a critical vacancy concentration at this interface is reached which cannot be accomodated at the metal side, and as a result a crack could be initiated [32]. During machining of the specimens, a significant portion of the residual stress due to cross rolling would have diminished and stress contour would have changed. The dislocation density and morphology (such as cellular, tangled, etc.) would be different in CW specimens as compared to that of mill annealed. The higher mobile dislocation density in CW specimens is expected to affect the transport of metal atoms at the metal/oxide interface. In this connection, a more complete understanding of the differences between the oxidation behaviour of mill-annealed and cold-worked AISI 316L under dynamic straining would be reached if a comprehensive investigation of the differences between the dislocation and grain boundary structure, as well as the energetic inhomogeneity of the surface of both substrates is carried out. Such studies represent an important trend in our ongoing research and will be reported in the near future. 5. Conclusions The present paper has demonstrated the applicability of a combination between an in situ measurement of the contact resistance of an oxide film and ex situ analysis of its composition and morphology for the characterisation of oxides
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formed on dynamically strained stainless steel surfaces in simulated nuclear power plant coolant conditions. The obtained results permitted the establishment of a qualitative relationship between the electrical properties of the film, its in-depth composition and structure. This relationship has been qualitatively interpreted in terms of the mixed-conduction model for surface films. Using the framework of this model, a mechanism for enhancing of cation vacancy and interstitial cation flux under the stress/strain field of the alloy substrate was proposed. It can be assumed that when these elemental fluxes are accelerated by straining, periodic formation of new oxides will be expected and repeated Cr(III) enrichment/depletion in the oxide will take place. Within this model approach, a tentative explanation of the enrichment of Fe in the outer deposited layer of the oxide as a result of the applied strain could be advanced. In addition, the differences between the electric properties, the composition and the structure of the oxide film formed on strained surfaces with or without prior cold work could be also rationalised on the basis of the proposed mechanism. Summarising, as a result of the present work, the concentration and mobility of cationic defects in the inner layer of the oxide film and at the alloy/film interface are shown to be one of the important factors that dominate the protectiveness of the oxide film. However, it is well known that especially in BWR conditions, the crack tip water chemistry is markedly different from the bulk water chemistry. In addition, the structure and distribution of grain boundaries in the steel substrate is expected to play a significant role in both stress corrosion crack initiation and propagation processes. Thus, an important direction for further investigation would be to characterise the oxide films formed in simulated crack tip chemistry conditions on substrate steel materials with a markedly different grain boundary structure and distribution. The work on this task is in progress and will be reported in the near future. Acknowledgements This work has been conducted within the integrity and Life Time of Reactor Circuits (INTELI) project as a part of SAFIR (Safety of nuclear power plants—Finnish national research programme 2003–2006) funded by the Finish Ministry of Trade and Industry (KTM) and coordinated by the Finnish Nuclear Safety Authority (STUK). This work has been supported by the grand-in aid for the 21st Century COE Program, ‘‘The Exploration of the Frontiers of Mechanical Science Based on Nanotechnology’’, The Japanese Ministry of Education, Culture, Sports, Science and Technology (MEXT). References [1] F.P. Ford, Corrosion 52 (1996) 375. [2] D.D. Macdonald, P.C. Lu, M. Urquidi-Macdonald, T.-K. Yeh, Corrosion 52 (1996) 768. [3] K. Sieradzki, R.C. Newman, Phil. Mag. A 51 (1985) 95. [4] F. Navaı¨, J. Mater. Sci. 30 (1995) 1166. [5] F. Navaı¨, O. Debbouz, J. Mater. Sci. 34 (1999) 1073. [6] F. Navaı¨, J. Mater. Sci. 35 (2000) 5921.
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[7] R.W. Bosch, F. Moons, J.H. Zheng, W.F. Bogaerts, Corrosion 57 (2001) 532. [8] V. Vignal, R. Oltra, M. Verneau, L. Coudreuse, Mater. Sci. Eng. A 303 (2001) 173. [9] V. Vignal, C. Valot, R. Oltra, M. Verneau, L. Coudreuse, Corros. Sci. 44 (2002) 1477. [10] Q. Yang, J.G. Yu, J.L. Luo, J. Electrochem. Soc. 150 (2003) B389. [11] J. Kim, M. Seo, J. Electrochem. Soc. 150 (2003) B193. [12] R. Krishnamurthy, D.J. Srolovitz, Acta Mater. 51 (2003) 2171. [13] B. Beverskog, M. Bojinov, A. Englund, P. Kinnunen, T. Laitinen, K. Ma¨kela¨, T. Saario, P. Sirkia¨, Corros. Sci. 44 (2002) 1901. [14] B. Beverskog, M. Bojinov, P. Kinnunen, T. Laitinen, K. Ma¨kela¨, T. Saario, Corros. Sci. 44 (2002) 1923. [15] M. Bojinov, P. Kinnunen, G. Sundholm, Corrosion 59 (2003) 91. [16] I. Betova, M. Bojinov, P. Kinnunen, K. Ma¨kela¨, T. Saario, J. Electroanal. Chem. 572 (2004) 211. [17] K.S. Raja, T. Shoji, J. Mater. Sci. Lett. 21 (2002) 435. [18] K.S. Raja, T. Shoji, J. Mater. Sci. Lett. 22 (2003) 1347. [19] K.S. Raja, T. Shoji, J. Mater. Sci. 39 (2004) 1033. [20] Y. Takeda, T. Masuda, T. Shoji, in: Proceedings of Asian Pacific Conference on Fracture and Strength ’01 and International Conference on Advanced Technology in Experimental Mechanics ’01 (APCFS&ATEM’01), Sendai, Japan, The Japan Society of Mechanical Engineers, (2001), pp. 114–119.
[21] Y. Takeda, M. Bojinov, H. Ha¨nninen, P. Kinnunen, T. Laitinen, K. Ma¨kela¨, T. Saario, K. Sakaguchi, T. Shoji, P. Sirkia¨, A. Toivonen, in: G.S. Was, L. Nelson, P. King (Eds.), Proceedings of the 11th International Conference on Environmental Degradation of Materials in Nuclear Power Systems— Water Reactors, American Nuclear Society, (2003), pp. 611–616. [22] Y. Takeda, M. Bojinov, H. Hanninen, P. Kinnunen, T. Laitinen, K. Ma¨kela¨, T. Saario, T. Sakaguchi, T. Shoji, P. Sirkia¨, A. Toivonen, Key Eng. Mater. 261–263 (2004) 925. [23] J. To¨pfer, S. Aggarwal, R. Dieckmann, Solid State Ionics 81 (1995) 251. [24] J. Congleton, W. Yang, Corros. Sci. 37 (1995) 429. [25] Y.J. Kim, Corrosion 51 (1995) 849. [26] P. Periat, B. Domenichini, B. Gillot, J. Phys. Chem. Solids 57 (1996) 1641. [27] M. Uematsu, H. Kageshima, K. Shirashi, Comp. Mater. Sci. 24 (2002) 229. [28] K. Takahashi, N. Inoue, J. Mater. Sci.: Mater. Electron. 10 (1999) 359. [29] D.D. Macdonald, J. Electrochem. Soc. 139 (1992) 3434. [30] D.D. Macdonald, Pure Appl. Chem. 71 (1999) 951. [31] D.D. Macdonald, M. Al Rifaie, G.R. Engelhardt, J. Electrochem. Soc. 148 (2001) B343. [32] P. Aaltonen, T. Saario, P. Karjalainen-Roikonen, J. Piippo, S. Ta¨htinen, M. Ita¨aho, H. Ha¨nninen, in: Corrosion’96, NACE International, Houston, 1996, paper no. 81.
In situ and ex situ characterisation of oxide films formed on strained stainless steel surfaces in high-temperature water Yoichi Takeda a, Tetsuo Shoji a, Martin Bojinov b,*, Petri Kinnunen c, Timo Saario c a
Fracture and Reliability Research Institute, Tohoku University, 6-6-01 Aoba Aramaki Aobaku, 980-8579 Sendai, Japan b Department of Physical Chemistry, University of Chemical Technology and Metallurgy, 1756 Sofia, Bulgaria c VTT Industrial Systems, VTT Technical Research Centre of Finland, P.O. Box 1704, FIN-02044 VTT, Espoo, Finland Received 1 October 2005; received in revised form 25 November 2005; accepted 25 November 2005 Available online 4 January 2006
Abstract A possible approach in describing the role of the environment in the phenomena behind initiation and propagation of a stress corrosion crack is to assume that the transport of matter and charge through the oxide film on the material is one of the rate-controlling factors. Straining of the bulk material may affect the transport rates of ionic defects, such as vacancies and interstitials, through the oxide film. The aim of the present work has been to verify the applicability of combined slow strain rate tests (SSRT) and contact electric resistance (CER) measurements to assess the influence of strain on the electric properties of oxide films on AISI 316L stainless steel with or without prior cold work in simulated boiling water reactor (BWR) coolant conditions. The SSRT-CER measurements have been combined with ex situ characterisation of the oxide films after experiments using electron spectroscopy for chemical analysis (ESCA) and scanning electron microscopy (SEM). The results suggest that the effect of strain on the resistance of the oxide films seems to correlate with the effect of the same parameter on the Cr(III) concentration in the inner layer of the oxide. In addition, important differences between the concentration of Ni and Fe in the outer layer formed on stressed and unstressed surface have been observed. Based on the mixed-conduction model for oxide films, an attempt is made to evaluate the effect of straining on the electric properties of the oxide films and to correlate these effects with the changes in film composition and structure. # 2005 Elsevier B.V. All rights reserved. Keywords: Stainless steel; Simulated nuclear reactor coolant; Contact electric resistance; Slow strain rate test; Oxide film; Kinetic model
1. Introduction When determining the susceptibility of a material to environmentally assisted cracking (EAC), it is important to consider the effect of the environment and stress/strain field on the properties of the surface oxide films on the material. Generally, oxidation and/or dissolution take place at the interface of the material and the environment around the crack tip, where stress/strain field exists due to external and/or residual stress, as suggested in the slip-oxidation, coupled environment fracture, and film-induced cleavage models for EAC [1–3]. Once the crack advances or external loading varies, distribution of stress/strain at the crack front changes. This change in the distribution around the crack tip results in the generation of dynamic strain. Consequently, the interfacial
* Corresponding author. Tel.: +359 889 298 679; fax: +359 2 8682 036. E-mail address: [email protected] (M. Bojinov). 0169-4332/$ – see front matter # 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.apsusc.2005.11.073
reactions and mass transport involved in oxidation and dissolution processes in the crack tip region take place on the dynamically strained surface. In a plausible working hypothesis, the straining of the bulk material is thought to affect the transport of ionic defects within the oxide film in the vicinity of the oxide film/metal interface, resulting in a vacancy flux into the bulk metal. Such an influence may determine the rate at which the vacancies react with dislocations at the interface and in the bulk metal. So far, a few attempts have been made to investigate the influence of the stress or strain on the oxide film during its growth and restructuring mainly in ambient temperature electrolytic media [4–12]. However, practically no literature references that treat the influence of dynamic straining are available, especially for simulated boiling water reactor (BWR) coolant conditions. A possible approach to understand how dynamic straining affects film growth is to measure in situ the contact electric resistance (CER) of the oxide formed on straining surface. Such measurements have been already applied for the investigation
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of passive film formation and stability in high-temperature water [13–19]. As a result, the mobilities of electronic and ionic current carriers in the films have been evaluated for different combination of materials and environments. The electric resistance of the formed oxide film, which is measured by CER technique, is one of the important properties of that film that can be correlated to its composition and structure [13–19]. Therefore, the CER technique is thought to be a powerful tool to gain in situ information of material stability in the environment. The general aim of further work in that direction [20–22] has been to test the applicability of combined CER and slow strain rate tests (SSRT) to assess the influence of active straining and strain rate on the electric properties, composition and morphology of oxide films on austenitic stainless steels in simulated BWR conditions. A specific goal of the present paper is to evaluate the effect of prior cold work applied to the steel substrate on the electric properties, composition and morphology of such oxide films. For the purpose, the CER of the oxide film on actively deforming surfaces of stainless steels formed during SSRT was measured continuously in situ by means of the SSRT-CER technique. In addition, the composition and morphology of the formed oxides have been characterised ex situ using scanning electron microscopy (SEM) and electron spectroscopy for chemical analysis (ESCA). A preliminary hypothesis on the relationship between the oxide film composition, its electrical and electrochemical properties and the processes occurring at the alloy/film interface leading eventually to crack initiation is proposed. 2. Experimental The materials used were commercial mill-annealed AISI 316L stainless steel and 20% cold-worked steel of the same grade, the cold work procedure being cross-rolling. The chemical composition of the material is shown in Table 1. The SSRT-CER measurement set-up shown in Fig. 1 consists of a plate-type SSRT specimen and a metal probe electrode. In this study, a cylinder-shaped probe made of pure iridium (99.7%, provided by Goodfellow) was used in order to exclude the influence of oxidation of the probe on measured film resistance [21,22]. The probe contacting the plate-type Table 1 Composition of the studied material (wt%) AISI 316L C Si Mn P S Ni Cr Mo N Fe
0.019 0.44 1.66 0.026 0.008 11.18 16.85 2.07 0.065 Bal.
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Fig. 1. Scheme of the principle of the SSRT-CER technique: (left) during the formation of oxide under straining and (right) during the measurement of film resistance.
specimen was 2 mm in diameter, that can be considered to be a large enough area to give average information on the electrical properties of the oxide. The SSRT-CER technique is based on a four point resistance measurement arrangement, as described in greater detail earlier [20–22]. The SSRT-CER system consisted of a uniaxial SSRT system and a CER system. The SSRT-CER system was connected to a recirculating BWR water loop in order to maintain the desirable water chemistry conditions. The dissolved oxygen (DO) concentration was maintained by purging nitrogen/oxygen mixed gas into the water tank of the loop. The inlet oxygen level was 1.0 ppm. Both inlet and outlet DO and conductivity were monitored using a Model 3600 Analyser (Orbisphere Laboratories) and a Kemotron 2902 conductivity meter. The internal volume of the autoclave containing the SSRT-CER measuring device was about 0.7 dm3 and the water flow rate was 0.3–0.5 dm3 min 1. The SSRT-CER specimens had a gauge length of 18 mm and were 6 mm in width and 3 mm in thickness. The surface of the specimens were finished by emery paper up to #1200 grid and then rinsed with acetone and pure water, before subjecting to the high-temperature water environment. The SSRT-CER measurements were performed at the strain rate of 3 10 7 s 1 at 288 8C and a pressure of 10 MPa. After stabilization of the environment, the specimen was preoxidized for 96 h without straining. The preoxidation was performed in order to minimise the effect of the initial transient oxidation occurring after reaching the target experimental conditions. Once the straining was started, the contact resistance of the gauge part of the respective specimen was measured every 5–20 h during the straining. Three independent runs have been performed for each type of specimen. In one and the same run, the values of the resistance have been measured seven times at each point in time and the result has been averaged. The straining and measurement of the film resistance were interrupted at predetermined strain level. The formed surface oxide on the interrupted SSRT specimens were analysed by using SEM and ESCA.
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The secondary electron images were obtained by using HITACHI S-4700 cold filed emission scanning electron microscope in which the electron beam accelerated at 15 kV as a light sources was utilised. The SEM system was equipped with two secondary electron detectors. The upper secondary electron detector was placed above the objective lens and secondary electrons were detected through the magnetic field of the lens. The lower secondary electron detector was placed in the specimen chamber. The signals from both detectors were corrected together for the observations. The vacuum level of the specimen chamber (<1.0 10 3 Pa) was maintained by an optionally installed turbomolecular pump. Subsequently the SEM analyses, a scanning ESCA microprobe (PHI Quantum 2000) was employed for the analysis of formed oxide film using Al Ka radiation and an acceleration voltage of 15 keV. The take-off angle of photoelectrons was selected to 908 in order to maintain higher sensitivity of satellite peak and chemical shift. The incident beam diameter used was 200 mm. The average information of the oxide composition within the beam area was obtained. The obtained ESCA spectra were evaluated by using PHI Multipak software. For the depth profiling, Ar-ion milling was used with 2 keV acceleration voltage and 2 mm 2 mm raster. The vacuum level of the specimen chamber of the ESCA system was maintained less than 1.5 10 6 Pa during sputtering by utilising a differential pump. Sputter time was converted to sputter depth by using an experimentally derived sputter rate of 4.03 nm min 1. This rate was determined using a SiO2 sputter standard known to have an oxide thickness of 104 nm. Both the SEM and ESCA analyses were carried out on the oxide films formed on both grip part and gauge part (i.e. unstressed and stressed part) of the specimens. 3. Results 3.1. Influence of stress on the electrical properties of the oxides Fig. 2 shows the stress/strain curve, as well as the dependence of the average contact resistance of the oxide film
Fig. 2. Film resistance normalised to the value in the absence of stress as depending on the strain level for AISI 316L without cold work. Experiment ended at a strain level of 0.18.
on mill-annealed AISI 316L on the strain level. The resistance was normalised to the initial value in the absence of strain in order to isolate the influence of strain on that parameter from the respective influence of other experimental conditions that could lead to a change in the contact area in different experiments (e.g. geometrical and/or energetical surface inhomogeneity [21]). Prior to straining, the specimens were oxidised in simulated BWR water for 96 h and the resistance of the film has been continuously monitored. A logarithmic kinetics of the increase of the resistance has been observed and constant values have been reached after 20–30 h of oxidation. During straining, occasionally, the resistance was measured on the unstressed part of the specimens and was found to be constant within the limits of experimental error. The values of the resistivity of the film on the unstressed part (used as a reference value to normalise the resistance during straining) calculated by using the estimate of the whole film thickness on the unstressed surfaces obtained by ESCA (see below) has been found to be (1.04 0.2) 107 V cm. The increase of the resistance observed up to a strain level of ca. 0.03 has been discussed by us previously [21,22] and has been correlated to the increase of the Cr(III) concentration in the inner layer of the oxide films formed on separate specimens subjected to interrupted SSRT tests as evidenced by ESCA depth profiling [21]. Such a correlation was explained on the basis of the fact that the conductivity of the inner layer (which determines the transport of matter and charge through the whole oxide film) has been reported to decrease with increasing Cr content in that layer because of the decrease of the concentration of ionic current carriers that act as electron donors [23]. On the other hand, the subsequent decrease of resistance with straining up to ca. 0.06 could be explained in terms of the decrease of the concentration of Cr(III) in both the inner and outer layers due to an enhancement of the transpassive dissolution of Cr by straining [22]. The increase of the resistance beyond strain levels of 0.08 is most probably due to a decrease in the contact surface area via formation of cracks (see below). Fig. 3 shows the stress/strain curves and the dependence of the resistance of the oxide film on AISI 316L with 20% cold
Fig. 3. Film resistance normalised to the value in the absence of stress as depending on the additional strain for AISI 316L with 20% cold work from two sets of experiments ended at different additional strain levels.
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Fig. 4. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel without prior cold work. Experiment ended at a strain level of 0.18.
Fig. 5. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel with 20% cold work. Experiment ended at a strain level of 0.037.
work (normalised to its value before the start of the straining for the particular experiment) as depending on the additional strain applied to the samples in two separate experiments that were terminated at different additional strain levels. The differences between the stress/strain curves shown in Figs. 2 and 3 can be explained by the fact that the surface to which dynamic straining was applied during the experiments shown in Fig. 3 has been already subjected to a strain of ca. 0.2 (corresponding to the level of prior cold work of 20%). The normalised resistance increases sharply for low levels of additional strain and subsequently decreases gradually, reaching at the end of the experiments values essentially similar to those before the additional strain was applied. This behaviour is markedly different from the dependence of the resistance on strain shown in Fig. 2 and demonstrates that the proposed measurement method can be used to differentiate
between the electrical properties of films formed on dynamically strained surfaces of stainless steels with or without prior cold work. An explanation for the different behaviour exhibited by the specimens in Figs. 2 and 3 has been sought by ex situ analysis of the morphology, thickness and indepth composition of the formed oxides. 3.2. Surface morphology The SEM micrographs of the oxide film formed on the stressed and unstressed parts of the steel specimen without cold work are shown in Fig. 4. The differences between the morphology of the oxide on the stressed and unstressed parts are relatively small, except for the higher number of large crystals seen on the top of the oxide layer formed on the unstressed part. The morphology of the oxide is that of a typical
Fig. 6. Typical morphologies of the oxide film at the stressed (left) and unstressed (right) surfaces of the AISI 316L steel with 20% cold work. Experiment ended at an additional strain level of 0.065.
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BWR bilayer film (see e.g. Ref. [25]), the grain size of the outer layer being of the order of 200–300 nm, whereas that of the underlayer (seen e.g. in the left micrograph) is an order of magnitude smaller. In addition, several cracks and protrusions are observed in the micrograph of the film formed on the stressed part of the specimen. They are most probably due to processes of cracking and repair of the oxide layer under the influence of applied strain. Figs. 5 and 6 show the morphology of the oxides formed on the stressed and unstressed parts of the two AISI 316L specimens subjected to a 20% prior cold work before straining. In these cases, the difference between the morphology of the oxide formed on the stressed and unstressed parts is somewhat more significant than that in Fig. 4, i.e. the number of crystallites and crystallite size of the outer layer formed on the stressed part is significantly lower. This indicates that on a substrate that is subjected to cold work, an oxide with different properties is formed during dynamic straining. Further proof for the difference between the oxides was sought in their ESCA indepth compositional profiles. The second type of experiments on CW samples has been terminated at quite high strain levels, which most probably resulted in changes the mechanism of repassivation due to the significant extent of cracking of the oxide. Thus, the ESCA results for the second type of experiments on CW samples have been omitted because they were not considered to be representative for the whole area of the stressed region. 3.3. In-depth composition of the oxides Fig. 7 shows the ESCA depth profiles of the concentrations of oxygen in the films formed on the stressed and unstressed surfaces of the AISI 316L steel with and without prior cold work. An estimate of the thickness of the oxide was obtained by approximating the curves with sigmoidal functions (solid lines in Fig. 7) and taking the value of depth at the inflection point of the curves as the position of the metal/film interface The values for the film thicknesses formed on the stressed and unstressed surfaces on the sample without prior cold work are rather close to each other (960 and 860 nm, respectively), i.e. the straining does not seem to influence the thickness of the film significantly (ca. with 10%) in accordance with our earlier results [21,22]. On the other hand, the estimated thickness of the oxide formed on the stressed part of the specimen is ca. 20% smaller (700 nm versus 900 nm) in the case of the experiment stopped at an additional strain of 0.037, i.e. when prior cold work is applied to the substrate, the differences in thickness due to straining appear to be larger. Fig. 8 shows the depth profiles of Cr(III) and metallic Cr in the films on the specimens with and without cold work. It is important to mention that in the present paper, the concentrations of the individual metallic constituents have been normalised to the total concentration of metallic elements. This has been done in order to exclude the influence of oxygen on the in-depth profiles of the metallic elements in the oxide. In addition, the depth profiles of the concentrations of metallic Cr and Cr(III) have been evaluated
Fig. 7. ESCA depth profiles of the atomic concentration of oxygen in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
separately from the respective spectra in analogy to earlier work [21,22]. Another estimate of the film thickness was obtained from these profiles by approximating them with sigmoidal functions and averaging the depth at which both Cr(III) and Cr(0) profiles have inflection points. These estimates agree very well with those resulting from the evaluation of the oxygen profiles, which lends credibility to the values of the film thickness derived from ESCA measurements. The maximum concentrations of Cr(III) in the inner layer of the film formed on the unstressed surface are significantly (more than two times in the case of the specimen without cold work and ca. 60% for the cold worked specimen) higher than those in the inner layer formed on the stressed surface. This probably demonstrates that straining enhances the rate of transpassive dissolution of Cr, as concluded also earlier [22]. It is worth mentioning that the maximum content of Cr(III) in the inner layer on the unstressed part of the cold-worked specimen is ca. 33% higher than that for the specimen without cold work. This could be explained by a higher mobile dislocation density in the cold-worked specimen affecting the transport of Cr atoms that are preferentially oxidised at the metal/oxide interface. However, a more detailed description of the oxidation process
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Fig. 8. ESCA depth profiles of the Cr(III) and metallic Cr contents normalised to the total cation content in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
at the alloy/film interface remains out of the scope of the present work. Fig. 9 depicts the depth profiles of Fe and Ni in the oxides on the specimens with and without cold work. The Fe content in the outer layer of the oxide formed on the stressed surface is significantly higher (and correspondingly, the Ni content significantly lower) than that formed on the unstressed surface, which may indicate that the portion of the spinel (NiFe2O4) in the outer layer with respect to hematite (Fe2O3) is smaller in the case of the oxide formed on the stressed surface. This might be in turn connected to an enhanced dissolution of Fe from the substrate through the inner layer with subsequent deposition which creates an excess iron in the outer layer. The differences between the Fe/Ni ratios in the outer layer formed on the stressed and unstressed part of the specimen are more significant for the oxide formed on the specimen with 20% cold work, which may once again indicate a change in the oxidation mechanism of cold-worked specimens in comparison to mill-annealed ones. In that connection, it is worth mentioning that the importance of the Ni content in the oxides formed on strained AISI 316L surfaces in BWR conditions for the overall resistance of the material to cracking has been emphasised by earlier authors [24].
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Fig. 9. ESCA depth profiles of the Fe and Ni contents normalised to the total cation content in the oxide on the stressed and unstressed regions of the AISI 316L specimens with 0% and 20% prior cold work experiment for the specimen without cold work ended at a strain level of 0.18, and for the specimen with 20% prior cold work at an additional strain level of 0.037.
4. Discussion 4.1. Influence of straining on the electrochemical and transport properties of the oxide film In order to explain the results reported earlier [21,22] and those obtained in the present paper, a mechano-electrochemical model of the growing oxide layer during straining has to be advanced. A plausible approach towards such a model is to examine the effect of applied stress on oxidation by considering its interaction with the oxide growth induced stress and stress relaxation modes in growing oxides. A number of recent papers treat the effect of stress on the behaviour of passive films and thermal oxide films [8–10], the different forms of stress generated by the primary oxide growth and further oxidation, as well as the stress relaxation modes [26–28]. On the other hand, the electrochemical aspects of the growth and dissolution of passivating films can be understood in terms of the point defect model (PDM) by Macdonald and co-workers [29–31]. The PDM describes the metal/oxide/ solution system at steady state in terms of generation, transport and consumption of point defects (anion and cation vacancies and interstitial cations) under the influence of both the concentration and potential gradients in the oxide and at its
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Fig. 10. A simplified scheme of the process during oxidation/dissolution of stainless steel in high-temperature water according to MCM.
boundaries with the metal and the solution. The model has been further developed by some of us as mixed-conduction model (MCM) [13–16]. A possible explanation of the experimental results can be given in terms of the MCM. A simplified scheme of the processes considered during growth and dissolution of the oxide on a stainless steel in high-temperature water is shown in Fig. 10. The film thickness versus strain data reported earlier indicate that the overall kinetics of the film growth (reaction sequence k2–k4 in Fig. 10) is not significantly influenced by the strain level. However, as near steady state the growth of the film is balanced by its dissolution, the effect of strain level on film thickness may appear small. Based on the experimental evidence for that from earlier work [21], in the following the reactions involving cation transport through an oxide with constant thickness (reaction sequences k1–k3 and k1i–k3i–kr) are assumed to predominate over the processes involving oxygen vacancies. As indicated in a range of papers [11,12,26,28], the growth of an oxide film via vacancy generation and transport initiates compressive stresses at the metal/film interface regardless of the exact growth mechanism. In addition, compressive surface stresses at the 316L/passive film interface have been measured in situ even under an externally applied tensile stress up to a certain critical value [9]. Thus, only internal mechanisms of stress relief are believed to control the influence of stress on oxidation up to the maximum in resistance shown in Fig. 2. Two important stress relief mechanisms are considered. First, it has been recently demonstrated by first-principle calculations [12] that the oxidation-induced strain can be compensated by a relatively large flow of cations from the metal into the oxide. Such an emission which is accompanied by the generation of vacancies in the metal can be favoured by the tensile stress applied to the substrate (the self-diffusion through a vacancy mechanism is accelerated by tensile and sloweddown by compressive stress [10,28]). In the case of the inner layer formed on stainless steel (a (CrxFe1 x)3 dO4 spinel [25]), taking into account the processes shown in Fig. 10, it is
presumed that the rate of reaction k1i is increased above the unstressed state. A second mechanism of stress generation and relief can be proposed by considering the fact that the stoichiometry of the inner layer varies from the metal/film to the film/solution interface [12,26]. Due to the higher oxidation state of iron at the oxide layer/water interface when compared to that at the steel/inner layer interface, the internal part of the layer is probably subjected to larger compressive stresses than the external part. As a result, the rate of cation transport through the film in its outer part will be higher [26], and so will be the rate of the reaction sequence k3i–kr that control iron dissolution and subsequent reprecipitation to form an outer layer (Fig. 10). Thus, a similar line of reasoning could explain qualitatively the changes in the Fe/Ni ratio in the outer layer under the influence of strain by taking into account the enhancement of the transport of interstitial cations at the inner layer/water interface and an unchanged rate constant for outer layer redeposition (kr) which is mainly determined by nickel and iron solubility in high-temperature water. Accordingly, it is hypothesized that for strain levels up to around 0.02–0.03, the emission of interstitial cations, their diffusion through the inner layer and the outer layer deposition are enhanced. This most probably results in an increase of the concentration of Cr(III) in the inner layer, and hence the electronic resistance of that layer, in accordance with experimental observations [21,22]. Above a strain of 0.02– 0.03, it can be assumed that tensile stresses start to dominate also at the steel/inner layer interface resulting in an increase of both the fluxes of cation vacancies and interstitial cations. Thus, also the reactions of Cr dissolution is accelerated, and therefore, Cr(III) concentration is decreased, as observed experimentally. Such a hypothesis is in agreement with recent observations that tensile stresses are measured upon potential steps in oxides with predominating cation vacancy conduction mechanism [11]. Summarising, the inner layer formed on 316SS in hightemperature water is considered to be a normal spinel of the FeCr2O4 type which is a p-type semiconductor, i.e. cation vacancies are predominant ionic current carriers in it. However,
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interstitial conduction through it is also significant, and Fe dissolves through the inner layer via that route. At lower strains, the rate of the processes involving interstitial cations are increased, and cation vacancies are accumulated in the inner layer due to their slower rate of transport, wheras at higher strains, also the vacancy transport is accelerated. In other words, no qualitative change in the conduction mechanism is envisaged in the present treatment, only a quantitative change in the rate constants and diffusion coefficients and hence of the ratio between concentrations of different defects in the film. 4.2. Effect of prior cold work applied to the substrate In principle, the results obtained in the present work for the specimens subjected to 20% cold work could be qualitatively explained on the basis of the model framework presented above. Indeed, it can be assumed that due to the residual stress (associated with the residual strain accumulated in the metal substrate during cold work), the tensile stresses start to dominate at the metal/film interface almost immediately after the additional strain is applied, therefore the resistance of the oxide decreases during most part of the experiments (Fig. 3). This decrease in the resistance could be once more correlated to the decrease of the Cr(III) content in the inner layer of the oxide on the stressed part of the specimens in comparison to the unstressed part, as shown by the ESCA profiles (see Fig. 8). This in turn shows that straining accelerates markedly the transpassive dissolution of Cr (and probably also the dissolution of Ni) from the substrate. The acceleration of the transpassive dissolution will enhance the flux of cation vacancies in the inner layer of the oxide and also at the metal/film interface until a critical vacancy concentration at this interface is reached which cannot be accomodated at the metal side, and as a result a crack could be initiated [32]. During machining of the specimens, a significant portion of the residual stress due to cross rolling would have diminished and stress contour would have changed. The dislocation density and morphology (such as cellular, tangled, etc.) would be different in CW specimens as compared to that of mill annealed. The higher mobile dislocation density in CW specimens is expected to affect the transport of metal atoms at the metal/oxide interface. In this connection, a more complete understanding of the differences between the oxidation behaviour of mill-annealed and cold-worked AISI 316L under dynamic straining would be reached if a comprehensive investigation of the differences between the dislocation and grain boundary structure, as well as the energetic inhomogeneity of the surface of both substrates is carried out. Such studies represent an important trend in our ongoing research and will be reported in the near future. 5. Conclusions The present paper has demonstrated the applicability of a combination between an in situ measurement of the contact resistance of an oxide film and ex situ analysis of its composition and morphology for the characterisation of oxides
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formed on dynamically strained stainless steel surfaces in simulated nuclear power plant coolant conditions. The obtained results permitted the establishment of a qualitative relationship between the electrical properties of the film, its in-depth composition and structure. This relationship has been qualitatively interpreted in terms of the mixed-conduction model for surface films. Using the framework of this model, a mechanism for enhancing of cation vacancy and interstitial cation flux under the stress/strain field of the alloy substrate was proposed. It can be assumed that when these elemental fluxes are accelerated by straining, periodic formation of new oxides will be expected and repeated Cr(III) enrichment/depletion in the oxide will take place. Within this model approach, a tentative explanation of the enrichment of Fe in the outer deposited layer of the oxide as a result of the applied strain could be advanced. In addition, the differences between the electric properties, the composition and the structure of the oxide film formed on strained surfaces with or without prior cold work could be also rationalised on the basis of the proposed mechanism. Summarising, as a result of the present work, the concentration and mobility of cationic defects in the inner layer of the oxide film and at the alloy/film interface are shown to be one of the important factors that dominate the protectiveness of the oxide film. However, it is well known that especially in BWR conditions, the crack tip water chemistry is markedly different from the bulk water chemistry. In addition, the structure and distribution of grain boundaries in the steel substrate is expected to play a significant role in both stress corrosion crack initiation and propagation processes. Thus, an important direction for further investigation would be to characterise the oxide films formed in simulated crack tip chemistry conditions on substrate steel materials with a markedly different grain boundary structure and distribution. The work on this task is in progress and will be reported in the near future. Acknowledgements This work has been conducted within the integrity and Life Time of Reactor Circuits (INTELI) project as a part of SAFIR (Safety of nuclear power plants—Finnish national research programme 2003–2006) funded by the Finish Ministry of Trade and Industry (KTM) and coordinated by the Finnish Nuclear Safety Authority (STUK). This work has been supported by the grand-in aid for the 21st Century COE Program, ‘‘The Exploration of the Frontiers of Mechanical Science Based on Nanotechnology’’, The Japanese Ministry of Education, Culture, Sports, Science and Technology (MEXT). References [1] F.P. Ford, Corrosion 52 (1996) 375. [2] D.D. Macdonald, P.C. Lu, M. Urquidi-Macdonald, T.-K. Yeh, Corrosion 52 (1996) 768. [3] K. Sieradzki, R.C. Newman, Phil. Mag. A 51 (1985) 95. [4] F. Navaı¨, J. Mater. Sci. 30 (1995) 1166. [5] F. Navaı¨, O. Debbouz, J. Mater. Sci. 34 (1999) 1073. [6] F. Navaı¨, J. Mater. Sci. 35 (2000) 5921.
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