Effect of nitrogen content in sensitised austenitic stainless steel on the crack growth rate in simulated BWR environment

Journal of Nuclear Materials 410 (2011) 59–68

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Effect of nitrogen content in sensitised austenitic stainless steel on the crack growth rate in simulated BWR environment S. Roychowdhury a,⇑, V. Kain a, R.C. Prasad b a b

Bhabha Atomic Research Centre, Mumbai, India Indian Institute of Technology, Powai, Mumbai, India

a r t i c l e

i n f o

Article history: Received 25 November 2010 Accepted 29 December 2010 Available online 6 January 2011

a b s t r a c t The focus of the study is to establish the role of nitrogen addition to type 304L stainless steel, in sensitised condition, on the crack growth rate (CGR) by intergranular stress corrosion cracking (IGSCC) in the simulated boiling water reactor (BWR) environment. The CGR studies were carried out for two heats of type 304L stainless steel with 0.10 and 0.16 wt.% nitrogen and at different dissolved oxygen (DO) levels in high temperature demineralised water whose chemistry was maintained in a recirculating loop. The degree of sensitisation (DOS) was characterised quantitatively by double loop electrochemical potentiodynamic reactivation (DL-EPR) technique. The results clearly show that the susceptibility to IGSCC was substantially lower in the stainless steel with a higher level of nitrogen as reflected by the CGR values. This was attributed to the beneficial role of nitrogen addition against sensitization i.e. lesser coverage of the chromium depleted regions and higher level of chromium in the depleted regions in the stainless steels with higher nitrogen content. Ó 2011 Elsevier B.V. All rights reserved.

1. Introduction Intergranular stress corrosion cracking (IGSCC) of austenitic stainless steel (SS) components is a generic problem in boiling water reactor (BWR) operating conditions [1–5]. Sensitization in the weld heat affected zone (HAZ) is an important material condition which makes SS susceptible to IGSCC. Chromium carbides form at the grain boundaries of austenitic SS when exposed in the temperature range 500–850 °C. This is accompanied by chromium depletion in adjacent areas of the grain boundary carbides and is known as sensitization. The extent of sensitization in type 304 stainless steel is determined by a combination of carbon, nickel and chromium content [6–11]. Low Temperature Sensitisation (LTS) is another phenomenon which results in an increase in the degree of sensitization (DOS) below 500 °C due to the growth of pre-existing carbides, without precipitation of any new carbide [12–15]. The SS becomes prone to IGSCC by an increase in the DOS due to LTS at reactor operation temperature (typically around 300 °C) for durations of the order of 10 years or longer. In operating BWRs oxidizing species from the in-core radiolytic dissociation of water equivalent to around 200–300 ppb of oxygen are present which brings the electrochemical potential (ECP) of the SS in a range in which IGSCC can easily initiate and propagate in a sensi⇑ Corresponding author. Address: Materials Science Division, 2nd Floor, Mod-Lab, D-Block, Bhabha Atomic Research Centre, Mumbai 400 085, India. Tel.: +91 22 25593456; fax: +91 22 25505151. E-mail address: [email protected] (S. Roychowdhury). 0022-3115/$ - see front matter Ó 2011 Elsevier B.V. All rights reserved. doi:10.1016/j.jnucmat.2010.12.313

tised material. The IGSCC susceptibility of structural materials in BWRs is strongly influenced by the ECP and it is affected by the concentration of the oxidizing and reducing species. Compared to a typical ECP of 100 to +100 mV Vs standard hydrogen electrode (SHE) for SS and nickel based alloy components, an increase in the ECP by as much as 250 mV is observed for components in the core of BWR [1,16,17]. The cracking of sensitised SS in BWR environment has been found to occur at potentials more oxidizing than a threshold potential (230 mVSHE at 288 °C) [18]. Several kinds of ionic impurities, like sulphate, chloride, carbonate, fluoride, chromate, nitrate and phosphate ions, etc., have also been reported to be able to accelerate the initiation and propagation of intergranular stress corrosion cracks in sensitised type 304 stainless steel and alloy 600 in high-temperature water [19–21]. The sensitization induced IGSCC has been a dominant reason for stress corrosion cracking (SCC) of components fabricated from austenitic SS in BWRs. Hence, for mitigation of this problem, type 316NG (nuclear grade) stainless steel containing molybdenum and nitrogen was developed and has been used in BWRs extensively as it is highly resistant to sensitization during welding (and to LTS) and hence resistant to IGSCC [18,22]. In presence of molybdenum, nitrogen has an enhanced effect in imparting resistance to sensitization but is detrimental if present in excess of 0.16 wt.% [23]. Hence, type 316NG has nitrogen content restricted to 0.10 wt.% [24]. However, the weld fusion zone of the molybdenum bearing SS have delta ferrite that is richer in chromium (and contains molybdenum) hence is more prone to low temperature embrittlement (LTE) during long term operation of the

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weldment at reactor operation temperature [12,25–27]. Thus LTE is a matter of concern for molybdenum bearing SS (type 316NG) when considering its use in the future generation nuclear reactors with a design life of up to 100 years [12]. Problems due to LTE are expected to be less in type 304L stainless steel making it suitable for use in reactors with a long design life [25,27]. Nitrogen addition to SS has been shown to increase its resistance to sensitization but the extent of beneficial effect varies in type 304L and 316L stainless steels [23,28,29]. The sensitisation behaviour as a function of nitrogen content and its effect on IGSCC crack growth rate (CGR) in simulated BWR environment for type 304LN stainless steel thus needs to be established. Most of the reported studies related to the effect of nitrogen on the sensitisation behaviour and susceptibility to IGSCC has been on type 316L stainless steel or on type 304 stainless steel. There are no reported results on the effect of nitrogen addition in type 304L stainless steel on the sensitisation behaviour affecting the CGR in simulated BWR environment. Dutta et al. [30] have reported that higher nitrogen levels (0.19 wt.%) in type 316L stainless steel lead to a greater DOS due to chromium nitride precipitation as compared to the SS with 0.16 wt.% nitrogen. In some of the studies [31,32] it has been reported that presence of nitrogen in type 304 austenitic SS results in earlier crack initiation and higher CGR in boiling magnesium chloride solution. It has also been reported that nitrogen addition up to 0.16 wt.% reduces susceptibility to IGSCC in sensitised type 304 stainless steel in sodium sulphate added water at 250 °C [29]. This behaviour was attributed to the retarding effect of nitrogen on the sensitization kinetics when nitrogen is present in the range of 0.04–0.16 wt.% [28,29]. Bruemmer and Was [33] reported that the IGSCC of type 304 stainless steel in oxygenated, high-temperature water was controlled by the grain boundary chromium concentration. Intergranular cracking was initiated in slow strain rate tests (SSRT) when local chromium levels dropped below a threshold level. Cracking severity and ductility loss increased as the grain boundary chromium content decreased [33]. It has been shown by laboratory studies and also from the data collected from plants that DOS above 2 C/cm2 in the single loop electrochemical potentiodynamic reactivation – EPR (or double loop – DL-EPR value of above 1) makes SS prone to IGSCC in BWR environment [12,34,35]. The IGSCC CGR for sensitised SS has also been successfully modelled and the CGR can be fairly accurately predicted for different material and environmental conditions [36–38]. Such type of cracking was also observed in austenitic SS stabilized by titanium or niobium which are inherently more resistant to sensitization. Cracking in the stabilized SS was primarily attributed to the decomposition of the stabilized carbides during processing and the subsequent local sensitization of the material during reactor operation [19,39,40]. Presence of nitrogen may also affect IGSCC susceptibility of austenitic SS in BWR environment as it promotes planar dislocation motion leading to strain localization [41], increases the strain hardening rate and the work hardening potential [41,42] and improves resistance to sensitization, pitting and crevice corrosion [43,44]. The effect of nitrogen in type 304 L stainless steel on the CGR in the strain hardened non-sensitised SS has been reported earlier [45] and it was shown that higher nitrogen levels cause a three fold increase in the IGSCC CGR values in simulated BWR conditions at dissolved oxygen (DO) levels in the range of <10– 760 ppb. The effect of various parameters on the IGSCC CGR in SS has been studied. The effect of test temperature [21,46,47], environmental variables like DO and ECP [39,48–50] and ionic impurities [19,21,39] have been widely studied. Nitrogen containing type 304L stainless steel thus becomes an attractive alternative material to type 316 NG with lower tendency for LTE for use in reactors with a long design life. However, the nitrogen levels in type 304L stainless steel which would impart

resistance to sensitization to an extent that it has low enough CGR in BWR environment needs to be established. This would allow use of a SS that has sufficiently low DOS and also resist LTE. Thus, the effect of nitrogen content on the CGR of sensitised type 304L stainless steel in simulated BWR environment has been the main focus of this study. Type 304 LN stainless steel was used with two different nitrogen levels to investigate the effect of the nitrogen on sensitization and the SCC susceptibility in the simulated BWR water chemistry. The crack growth kinetics was established using fatigue precracked CT specimens in a recirculating loop that maintained a steady simulated BWR water chemistry. The effect of nitrogen addition on the DOS and its final effect on the CGR are studied and explained. 2. Experimental 2.1. Materials and heat treatment The chemical composition of the SS with two different nitrogen levels used in this investigation is given in Table 1. The material was in the form of plates of 25 mm thickness in the solution annealed condition. The 0.5T CT specimens were fabricated from these plates. The CT specimens were subjected to a sensitization heat treatment at 675 °C for 24 h and subsequently quenched in water. 2.2. Evaluation of sensitization Small specimens of size 10 mm  10 mm  5 mm (thickness) were cut from the CT specimen after completion of the crack growth experiment and used for evaluation of sensitization. The microstructures of the SS prior to (i.e. in the starting solution annealed condition) and after the sensitization heat treatment were examined after electrochemical etching in 10% oxalic acid according to ASTM A262, Practice A [51]. The DL-EPR test was conducted to assess the DOS of SS before as well as after the sensitization heat treatment. The DL-EPR test was carried out at room temperature using a solution of 0.5 M sulphuric acid and 0.01 M potassium thiocyanate [52]. The solution was deaerated for 1 h prior to and also during the test. In this method the potential is swept from the corrosion potential to +300 mV Vs saturated calomel electrode (SCE) and then reversed back to the starting potential. The specimens are polarized at a scan rate of 6 V/h. The maximum current for each scan direction was measured i.e. If for the forward scan and Ir for the reverse scan. The ratio of the Ir to If multiplied by 100 was taken as the DL-EPR value which is a measure of the DOS. The reported values are an average of three tests for each specimen. The microstructure obtained after the DL-EPR test was examined and recorded. 2.3. Fatigue precracking The dimensions of the 0.5T CT specimen used in this study are shown in Fig. 1. Fatigue precracking of the CT specimens was carried out in air to obtain an initial a/W of 0.45 (where a – crack length, W – 25 mm). Fatigue precracking in air was carried out at a frequency of 25 Hz in the decreasing load mode. The maximum p stress intensity factor (K) was maintained at  11 MPa m in the final step of fatigue precracking. Side grooving was done after fatigue precracking to obtain a total thickness reduction of 10% (5% thickness groove on each side). After loading the specimen in the autoclave, fatigue precracking was done in the aqueous environment at the test temperature (288 ± 0.3 °C) and pressure (10 MPa). The loading schedule used for fatigue precracking in the environment and the subsequent

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S. Roychowdhury et al. / Journal of Nuclear Materials 410 (2011) 59–68 Table 1 Chemical composition of the stainless steels (in wt.%) used in the study.

SS 304 LN1 SS 304 LN2

%C

%S

%P

% Mn

% Si

% Cr

% Ni

% Mo

%N

0.028 0.022

0.004 0.005

0.033 0.024

1.63 1.63

0.49 0.42

18.82 18.77

9.20 9.47

0.11 0.30

0.08 0.16

Fig. 1. Dimensions of the CT specimen used for the crack growth rate studies.

SCC testing has been used earlier for SCC testing at high temperature and pressure [45,53,54,47]. Triangular loading was applied at different R values and frequency for fatigue precracking in environment. The fatigue precracking in environment was done in four steps. In the first three steps the R value was increased gradually starting from 0.3 to 0.5 and then to 0.7 at a frequency of 0.01 Hz. In the fourth step the R-value was kept fixed at 0.7 and the frequency was reduced to 0.001 Hz. The crack was allowed to grow for at least 500–900 lm during fatigue precracking in environment. For SS 304 LN1 the fatigue precracking in the environment was done at a DO level of 260 ppb and subsequently the SCC test was done at 260 ppb and then at 760 ppb. However, as the CGR values measured were extremely low, fatigue precracking in environment was repeated at a DO level of 760 ppb and the SCC test was subsequently started at 760 ppb. For SS 304 LN2 the fatigue precracking in the environment was done at 760 ppb and the SCC test was subsequently started at 760 ppb. 2.4. CGR experiment A schematic of the high temperature high pressure (HTHP) loop which was used for the experiments in simulated BWR environment is shown in Fig. 2. The loop contains a 200 L capacity water storage tank which stores demineralised water of desired water chemistry (specific conductivity of 0.055 lS/cm at 25 °C). There is a 4 L capacity autoclave equipped with a load cell with SSRT facility in the high pressure side with a temperature and pressure balanced external Ag/AgCl (0.01 M KCl) reference electrode, platinum redox electrode and platinum counter electrode for electrochemical measurements. Reversible direct current potential drop (DCPD) technique is used for crack length monitoring at HTHP. There is automatic control of the DO content in the water. All the relevant test parameters and the data are continuously monitored in a computer. Details about the HTHP loop have been reported earlier [45]. The fatigue precracked CT specimen along with a reference specimen for DCPD measurements were loaded in the autoclave. The low pressure side of the HTHP loop is initially operated to ob-

tain the desired water chemistry. The autoclave was then pressurised to 10 MPa and heated to 288 °C. Flow through the autoclave was maintained so that the total autoclave volume was refreshed four times every hour. The recirculation loop was allowed to run in this condition for 7 days prior to applying any load so as to stabilise the water chemistry parameters and achieve an inlet conductivity of <0.07 lS/cm at 25 °C and an outlet conductivity of <0.09 lS/cm at 25 °C. The DO level at the outlet was intermittently monitored. When the inlet DO level was controlled at 760 ppb and 260 ppb, the values at the outlet were found to be 15 ppb lower than the inlet DO values. Trapezoidal loading was applied with 9000 s of hold time and an unloading–loading step in between, of 90 s duration. The crack length was continuously monitored by the reversible DCPD technique. The maximum load was maintained at 5.4 kN to obtain a p starting ‘K’ of 25 MPa m. The DCPD signal was recorded once every 30 s. A current of 10 A was used for the reversible DCPD measurements and the DCPD signal was recorded after 5 s of application of the current to ensure stability of the current value. The method of calculating the crack length from the DCPD data has been described earlier [45]. The SCC experiments were carried out at DO levels of 760 ppb, 260 ppb and in deaerated condition (<10 ppb) after proper fatigue precracking in the environment. The crack was allowed to grow till the DCPD signal variation with time was linear indicating a stable crack growth. Typically crack growth measurements were done at each DO level to achieve stable CGR. The ECP was monitored through the entire test duration in one of the experiments using SS 304 LN1. The DCPD signal was recorded for both the test specimen and the reference specimen. 2.5. Fractographic examination After the completion of the experiment, fatigue post cracking of the CT specimens was done and the fracture surfaces were observed in the scanning electron microscope (SEM) to establish the mode of cracking. The crack length was measured from the fracture surface, for both the materials, by measuring the area of

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Fig. 2. Schematic of the high temperature high pressure loop simulating a BWR environment used for the crack growth experiments.

the fracture surface covered in each region and dividing the total area by the net thickness of the specimen after side grooving. The area of the IG regions was measured by printing a high resolution image of the fracture surface on a standard 1 mm  1 mm grid paper and counting the number of squares lying in the regions with IG facets. The total area covered by the IG facets was obtained by the total number of squares counted. This ‘‘area’’ method of measuring the crack length is expected to be accurate. The crack length corresponding to the end of fatigue precracking in the environment was measured from the fracture surface and was used for calculating the corrected crack length. The CGR was calculated by fitting a straight line to the crack length–time plot for each DO level.

3. Results and discussion 3.1. Evaluation of sensitization The microstructure of both the SS in the solution annealed condition observed after electrolytic etching in oxalic acid in shown in Fig. 3a and b. The structure indicates that there is no preferential grain boundary attack or ditch formation. Step forms between the grains due to differential dissolution of adjacent grains which leads to a height difference between them. This step formation is characteristic on etching of non-sensitised stainless steel with no grain boundary chromium carbide or chromium depletion. This microstructure is categorized by ASTM A 262 practice A as a ‘‘step’’ structure and is taken as resistant to intergranular corrosion [51]. There is also no grain boundary carbide present. Such type of step structure is expected in an annealed SS. The microstructure observed for both the SS after the sensitization heat treatment (675 °C, 24 h) and after the oxalic acid etching, is given in Fig. 4a and b. Intergranular attack can be seen in Fig. 4a and b indicating that both the SS have sensitised due to the heat treatment. The ‘‘ditch’’ structure is observed in both the SS as at least one grain is completely surrounded by the attack on the chromium carbide/chromium depletion regions [51]. However, it is clear upon comparing both the sensitised microstructures that the extent of intergranular attack is lesser in SS 304 LN2 as compared to SS 304 LN1 indicating that the length (coverage) of regions of chromium depletion is lesser in SS 304 LN2 as compared to SS 304

Fig. 3. Microstructure of: (a) SS 304 LN1, and (b) SS 304 LN2, in the solution annealed condition after etching in 10% oxalic acid as per ASTM standard A262, Practice A.

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LN1. Marginal variation in the width of the depleted regions is also observed. A quantitative measurement of the DOS can be made by the DL-EPR results. The DL-EPR plot is shown in Fig. 5 and the microstructures observed after the DL-EPR test are shown in Fig. 6a and b and support the observation made after the oxalic acid etching. The obtained DL-EPR values are 22.5 and 8.9 for SS 304 LN1 and SS 304 LN2 respectively. The DL-EPR values measured before sensitisation was 0.04 and 0.07 for SS 304 LN1 and SS 304 LN2 respectively. There is a substantial difference in the DOS values for the two SS in sensitised condition and it is also supported by the microstructures observed after the oxalic acid etching (Fig. 4) and after the DL-EPR test (Fig. 6). Sensitisation due to chromium carbide precipitation depends on the chemical composition [7]. Molybdenum affects the sensitisation behaviour significantly only when present in the range of 2– 3% like in type 316/316L stainless steel [23]. The molybdenum content in the SS grades used in this study is not expected to affect the sensitisation behaviour significantly. Kain et al. [7] used the term Creff (where Creff = %Cr + 0.18 (%Ni) – 100(%C), all compositions are in wt.%), which was originally given by Briant and Hall [55], to correlate the chemical composition with sensitisation behaviour for type 304/304L stainless steel. Creff values correlate well with

Fig. 4. Microstructure of: (a) SS 304 LN1, and (b) SS 304 LN2, in the sensitised condition after etching in 10% oxalic acid as per ASTM standard A262, Practice A. Greater coverage of the chromium depleted regions are evident in (a) as compared to (b).

Fig. 5. The DL–EPR plots for both the stainless steels in the sensitised condition.

Fig. 6. Microstructure of: (a) SS 304 LN1, and (b) SS 304 LN2, in the sensitised condition after the DL–EPR test. Greater coverage of the chromium depleted regions are evident in (a) as compared to (b) confirming the oxalic acid etch results shown in Fig. 4.

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the sensitisation time at a particular temperature and a higher Creff value indicates more resistance to sensitisation. For the two stainless steels used in this study the Creff for SS 304 LN1 and SS 304 LN2 was 14.1 and 14.6, respectively indicating that the susceptibility to sensitisation at a particular temperature is similar for these two grades of SS. Thus the substantial difference in the DOS values measured in both the SS can be attributed to the higher nitrogen levels in the SS 304 LN2 as compared to SS 304 LN1. Higher nitrogen (0.16 wt.% in SS 304 LN2) leads to lower DOS values as reported earlier [28,29]. 3.2. Fractography The fracture surface as observed in the SEM for both the tested materials is shown in Fig. 7a–d. Fig. 7a and c shows the regions corresponding to fatigue precracking in air, fatigue precracking in environment, and the transgranular (TG) and the intergranular (IG) regions for SS 304 LN1 and SS 304 LN2 respectively. It can be seen that transition of the crack front from TG to IG occurs only in the final stages of the fatigue precracking in the environment. The crack is TG in nature during fatigue precracking in the environment. During the SCC test, cracking in completely IG in nature as is shown in Fig. 7b and d for SS 304 LN1 and SS 304 LN2 respectively. Fig. 8a and b shows the photographic images of the fracture surface of the CT specimen after fatigue post cracking along with the schematic showing the different regions marked. The crack front is uniform for the SS 304 LN1 but for SS 304 LN2 the crack front was not uniform though the crack had grown uniformly during fatigue precracking in the environment. The crack lengths measured by the ‘‘area’’ method and that obtained by DCPD for the IG regions were in good agreement. 3.3. CGR results The first stage of the CGR experiment was carried out at a DO level of 760 ppb for both the materials. Subsequently the DO level was reduced to 260 ppb and after stable CGR was obtained the test water was deaerated in the third stage, for the test done using SS 304 LN1. For the test done using SS 304 LN2 the water was deaerated after sufficient crack growth at 760 ppb DO and the third stage of crack growth was carried out at 260 ppb. The outlet conductivity was <0.09 lS/cm at 25 °C and the water chemistry parameters and the test conditions like temperature and pressure were stable for the entire duration of the test. The CGR was estimated by making a linear fit to the plot of corrected crack length variation with time. The DCPD data recorded during the transition of the DO levels was ignored. A good linear fit was obtained for the crack length vs test time data and a high value of correlation coefficient (0.97–0.99) and a low value of standard deviation (0.006–0.008 mm) was obtained for both the materials at all the DO levels. In the test using SS 304 LN1, the specimen was initially (85– 875 h from the start of the experiment) fatigue precracked in the environment at a DO level of 260 ppb and after 500 h of SCC testing at this DO level it was increased to 760 ppb. The crack length and ECP variation with time during this test duration is shown in Fig. 9. The ECP values measured with respect to the Ag–AgCl reference electrode increased from 110 mV (Vs Ag/AgCl) to 55 mV (Vs Ag/AgCl) with an increase in the DO level from 260 ppb to 760 ppb. It can be seen from Fig. 9 that the increase in the crack length during this period was approximately 27 lm. The CGR value did not vary as the DO level was increased from 260 ppb to 760 ppb and remained stable at a very low value of 3.65  1012 ms1. The very low value of the CGR measured in this case was attributed to the improper transition of the crack front from TG to IG. Subsequently fatigue precracking in environment

Fig. 7. SEM images after fatigue post cracking for: (a) SS 304 LN1, (c) SS 304 LN2 showing fracture surface with various regions marked and (b) SS 304 LN1, (d) SS 304 LN2 showing clear IG facets observed during SCC testing.

was repeated at a DO level of 760 ppb by applying triangular loading and the crack was allowed to grow further by 470 lm to ensure

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proper transition from TG to IG. Subsequently the SCC test was done at a DO level of 760 ppb. The crack length variation with time for the SCC test after fatigue precracking in the environment is shown in Fig. 10a and b for both the materials used in this study. The figures indicate that the CGR values take some time to stabilize. For SS 304 LN1 the CGR at a DO level of 760 ppb (after 950 h from the start of the experiment) was measured to be initially 1.50  1009 ms1 which gradually increased to a value 2.74  1009 ms1 after 100 h and a corresponding growth in crack length of approximately 480 lm was allowed to take place. Similarly for SS 304 LN2 the CGR value stabilised at a value of 4.50  1010 ms1 after 110 h and a corresponding crack growth of approximately 38 lm. The possible reasons for this delay in stabilisation of the CGR are the transition from the TG nature of the crack to IG in the environment during fatigue precracking, the development of stable crack tip solution chemistry or the specimen surface or the crack tip coming to equilibrium with the environment [53,54,47]. The transition of the crack tip from TG to IG is considered to be very important. It has been reported that improper transition of the crack front from TG to IG may lead to no crack growth or very slow CGR as is shown in Fig. 9 [56]. Shoji et al. have reported that depending on the testing history, quite different steady state CGR values are possible for a given material and test conditions [57]. It was shown that the crack growth during the transient period, before a steady state CGR is achieved, correlated

Fig. 8. Photograph of the fracture surface of the specimens after fatigue post cracking along with a schematic with different regions marked for: (a) SS 304 LN2 and (b) SS 304 LN1.

Fig. 9. Variation in the crack with test duration for SS 304 LN1 at DO levels of 760 ppb and 260 ppb indicating very low crack growth rates.

Fig. 10. Variation in the crack length with time at different DO levels for sensitised (a) SS304LN1, (b) SS304LN2.

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well with the grain size indicating the importance of the TG to IG transition [57]. At different test temperatures, the transient crack growth behaviour was correlated to the changes in the crack tip reaction kinetics [57]. The CT specimens used in this study did not satisfy the conditions for linear elastic fracture mechanics for the applied ‘K’. The 1T CT specimens follow the conditions for linear elastic fracture mechanics when the material is in the sensitised condition. However, the two materials used in this study were in the sensitised condition and the specimen dimensions used for the CGR studies were also the same; hence comparison of the CGR values measured in both the materials is pertinent.

3.3.1. Effect of dissolved oxygen The stable CGR measured (during 1070–1110 h from the start of the experiment) for SS 304 LN1 at a DO level of 760 ppb was 2.74  1009 ms1 and the crack was allowed to grow by approximately 510 lm at this DO level as shown in Fig. 10a. Subsequently when the DO level was reduced to 260 ppb the crack grew by another 900 lm and the stable value of CGR measured during this period was 2.81  1009 ms1. On subsequent deaeration (DO level < 10 ppb) the CGR reduced to a stable value of 2.09  1010 ms1 and the total crack extension measured in the deaerated condition was approximately 200 lm. The ECP values measured in the totally deaerated condition was 739 mV (Vs Ag/AgCl). The total crack extension, after fatigue precracking in the environment, measured for the entire test duration (950– 1340 h from the start of the experiment) was approximately 2.34 mm. After the initial delay a stable value of the CGR measured for SS 304 LN2 at 760 ppb was 4.5  1010 ms1 as shown in Fig. 10b. The total crack extension at 760 ppb was measured to be 560 lm. On subsequent deaeration (DO < 10 ppb), the CGR value reduced to a value 6.22  1011 ms1 and the total crack extension during this period was 150 lm. When the DO level was increased to 260 ppb subsequently the CGR value (7.76  1011 ms1) was initially low and was of the same order as that measured in the deaerated condition. After the crack grew by 35 lm the CGR value increased and stabilised at a value of 4.5  1010 ms1. This delay in the stabilisation of the CGR value was observed only when the DO level was increased from a totally deaerated condition for SS 304 LN2. When the DO level was changed from 760 ppb to 260 ppb no such delay in the stabilisation of the CGR value was observed. This is related to the time taken for stabilization of the crack tip water chemistry. The total crack extension measured for SS 304 LN2 was approximately 900 lm. The CGR values showed a marginal variation for the tests done at DO levels of 260 ppb and 760 ppb, for both the materials and reduced by an order of magnitude on deaeration (DO < 10 ppb). The ECP values (in mV Vs Ag/AgCl and corresponding values Vs standard hydrogen electrode (SHE)) measured in the experiment for SS 304 LN1 are given in Table 2. Similar ECP values are expected for SS 304 LN2 at the same DO levels. The ECP is the mixed potential associated with the redox reactions occurring on the surface of the SS and has a strong influence on the IGSCC susceptibility of austenitic SS, both in the sensitised and the non-sensitised condition, in high-temperature water. The ECP value for SS shows

a sigmoidal dependence on the DO level on the water [16]. The CGR values markedly decrease when the water is deaerated which corresponds to a low value of ECP (typically below 230 mVSHE) and increases non-linearly with an increase in the DO levels (due to the corresponding increase in the ECP) [16]. An increase in the CGR values with increase in the DO level is always expected however the increase in the DO level should be sufficient enough to appreciably increase the ECP value so as to have any effect on the CGR. A small increase in the DO level from the deaerated condition will cause a large increase in the ECP value and correspondingly an appreciable increase in the CGR values. But at higher levels of DO the extent of increase in the ECP with increase in DO is much lower. In this present study, the CGR did not vary much when the DO level was increased from 260 ppb to 760 ppb which corresponds to an increase in the ECP by 55 mV. A change in the DO level from the totally deaerated condition (<10 ppb) to 260 ppb or 760 ppb results in the change in ECP by 629 mV. Due to this appreciable change in the ECP the CGR value changes by an order of magnitude on deaeration and also it takes more time for the CGR values to stabilize when DO levels are changed drastically.

3.3.2. Effect of nitrogen in stainless steel Typically, the worst microstructure developed in the heat affected zone (HAZ) of a weldment of austenitic SS is simulated by a heat treatment at 675 °C for 1 h. However, such a heat treatment does not result in significant DOS for the SS containing nitrogen. Therefore, to clearly reflect the influence of sensitization on IGSCC, a more severe sensitization at 675 °C for 24 h was chosen for this study. The results therefore, reflect comparative CGR values for the two SS containing different levels of nitrogen and do not in any way suggest the actual CGR at the HAZ of the weldment in BWRs. The variation in the crack length with the test time for both the materials tested along with the CGR values are shown in Fig. 10a and b. The CGR values clearly indicate that in sensitised SS 304 LN1 the values are about one order of magnitude higher than that measured in SS 304 LN2. Both the SS have almost identical chemical composition, the only difference being in the level of nitrogen. The heat treatment given to both is the same. However, it resulted in an appreciably higher level of the DOS value in the SS 304 LN1 as compared to SS 304 LN2. This difference in the IGSCC CGR values in both the SS, which is affected by the DOS values, can be attributed to the difference in the nitrogen levels. The CGR values for both the SS at different DO level are plotted in Fig. 11 illustrating the effect of nitrogen in the SS on the CGR values. The extent of the sensitisation of the austenitic SS depends upon the coverage of the chromium depleted zones (proportion of the grain boundary length covered by the chromium depleted

Table 2 Variation in the ECP with DO level for SS 304 LN1 at the test temperature. DO (ppb)

ECP (in mV vs. Ag/AgCl)

ECP (in mV vs. SHE)

760 260 <10

55 110 739

+283 +228 401

Fig. 11. IGSCC CGR for SS304LN1 and SS304LN2 as a function of DO in high temperature (288 °C) high pressure (10 MPa) pure water.

S. Roychowdhury et al. / Journal of Nuclear Materials 410 (2011) 59–68

region), width and the depth (minimum level of chromium in the depleted zones) [6]. One or more of these three parameters defining the level of sensitisation influences the susceptibility to IGSCC and these parameters have been characterised and correlated to the susceptibility to IGSCC and intergranular corrosion [6] for an austenitic SS. Kain et al. have shown that the width of the attacked regions after the EPR test is much more than the width of the chromium depletion regions and is not a measure of the DOS in stainless steel [6] or in alloy 600 [58]. They also showed that in a sensitised type 304 stainless steel, the width of the chromium depleted regions had no effect on the susceptibility to IGSCC [6]. The coverage and the minimum chromium level in the chromium depleted regions were the dominant factors determining the susceptibility to IGSCC [6]. Similar results were also reported by Bruemmer and Was [33] for sensitised austenitic SS in oxygenated, high-temperature water. It was reported that a minimum of 23% of the grain boundaries need to be depleted below the threshold chromium concentration for IGSCC to occur since isolated depleted regions will not lead to continuous path of susceptible material [6,59,60]. Hence, factors which govern the coverage of chromium depleted region or the minimum chromium levels will have an influence on the IGSCC susceptibility and also the IGSCC CGR. Presence of nitrogen in austenitic SS has been reported to affect the IGSCC susceptibility to a large extent. Mozhi et al. [29,61,62] have reported that presence of nitrogen to the extent of 0.16 wt.% in austenitic type 304 stainless steels increases the SCC resistance and reduces it when nitrogen level increases further. Nitrogen addition in austenitic type 304 stainless steels retards the sensitisation kinetics as for a constant carbon level, the grain boundary chromium concentration in equilibrium with the carbide increases with increasing nitrogen content. At nitrogen levels in excess of 0.16 wt.% discontinuous chromium carbide precipitation at the grain boundaries occur resulting in grain boundary migration leading to wider (hence shallower) chromium depleted regions and enhanced sensitisation [61]. The width of the chromium depleted regions in both the SS used in this study are marginally different as is evident from the width of the attacked regions after the EPR test as shown in Fig. 6a and b however, the coverage of the chromium depleted regions is appreciably higher in SS 304 LN1 as compared to SS 304 LN2. Higher nitrogen level in SS 304 LN2 also results in higher chromium levels in the depleted regions [61]. In the SS used in this study the same heat treatment thus resulted in a higher DOS value for SS 304 LN1 as compared to that for SS 304 LN2. A higher CGR in SS 304 LN1 can thus be attributed to the greater coverage and lower levels of chromium in the depleted regions. The earlier reported studies [63] using SSRT on the similarly heat treated SS also showed a similar trend of nitrogen leading to a lower susceptibility to IGSCC in sensitised type 304LN stainless steels. The present study establishes the beneficial effect of nitrogen on CGR in a quantitative manner under controlled experimental conditions. However, it is to be noted that a study on the same nitrogen added type 304L stainless steels, in warm rolled non-sensitised condition, had shown [45] a threefold increase in CGR in simulated BWR environment (at the three levels of DO as used in the present study). Therefore, while nitrogen addition of 0.16 wt.% to type 304L stainless steel improves its resistance to sensitization induced IGSCC CGR in simulated BWR environment, the susceptibility to IGSCC in strain hardened (non-sensitised) condition increases. Nuclear grade stainless steels with a carbon content lower than 0.02 wt.% will have a better resistance to sensitisation for the same nitrogen levels used in this study. However, more data on CGR in BWR simulated environment is required on type 304LN stainless steels, especially on heats that contain different levels of carbon and nitrogen to conclusively establish the chemical composition to be specified for use in BWRs.

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4. Conclusions SCC experimental measurements of CGR in simulated BWR environment (high purity water at 288 °C and 10 MPa pressure) were done using type 304LN stainless steels with two different levels of nitrogen (0.08 wt.% and 0.16 wt.%). The SS were sensitised at 675 °C for 24 h. Experiments were done at different DO levels. The following conclusions can be made based on the results obtained: (1) Both the grades of type 304 LN stainless steels (after a severe sensitization heat treatment) are susceptible to IGSCC in high purity water at 288 °C at DO levels of 760, 260 and <10 ppb. Improper fatigue precracking in environment resulted in very low CGR values. Proper fatigue precracking in environment was shown to be essential for obtaining meaningful CGR values. (2) The CGR values changed with time at the start of the experiment at 760 ppb DO and reached a stable value after the crack grew to some extent. Similar behaviour was observed for SS 304 LN2 stainless steel when the DO level was increased from <10 ppb to 260 ppb. (3) CGR values did not vary much with reduction in the DO from 760 ppb to 260 ppb. This was attributed to the measured nominal change in the ECP value by 55 mV on reducing the DO level. Substantial reduction in the CGR values was observed on reducing the DO levels to <10 ppb due to a substantial reduction in the ECP value by 629 mV. (4) The CGR values were higher by an order of magnitude in the SS 304 LN1 as compared to SS 304 LN2 at the three levels of DO used in the test. The greater coverage of the chromium depleted regions and lower chromium levels in the chromium depleted regions (due to higher nitrogen content) are the main reasons for enhanced CGR in SS 304 LN1 as compared to SS 304 LN2.

Acknowledgements The author’s acknowledge with thanks the suggestions and the invaluable discussions with Prof. Tetsuo Shoji, Prof. Zhanpeng Lu, and Prof. Yoichi Takeda on various occasions. The author’s are also thankful to Dr. P.L. Andresen for his valuable comments and suggestions. References [1] [2] [3] [4]

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Supratik Roychowdhury is working as a Scientific Officer in the Materials Science Division, Bhabha Atomic research Centre, Mumbai, India with more than 12 years of experience in the area of corrosion science and engineering. The major areas of his interest include environmentally assisted cracking of stainless steels, failure analysis. He has completed his Masters in Technology from the department of metallurgical engineering and materials science, Indian Institute of Technology, Powai, Mumbai in 2005 and is currently registered for Phd in the same department. He has a number of publications in international journals to his credit.

Vivekanand Kain is the Head, Corrosion Science Section in the Materials Science Division, Bhabha Atomic research Centre, Mumbai, India with more than 24 years of experience in the area of corrosion science and engineering. His major areas of interest include stress corrosion cracking of stainless steels, corrosion behaviour of various alloys used in nuclear reactors, material selection for various applications and failure analysis. He has a large number of publications to his credit in his area of interest. He is also a recipient of a number of awards in recognition to his contributions.

R.C. Prasad is a professor in the department of metallurgical engineering and materials science, Indian Institute of Technology, Powai, Mumbai. His major areas of interest include mechanical behaviour of materials, development and characterisation of composite materials, environmental assisted cracking of various materials, integrity assessment if line pipe steels and failure analysis. He has supervised a large number of masters and Phd theses. He has also conducted large number of workshops and courses in his area of interests. He has a number of publications in national and international journals to his credit.