influence of sulfur and ferrite on scc and corrosion...

17th International Conference on Environmental Degradation of Materials in Nuclear Power Systems – Water Reactors August 9-13, 2015, Ottawa, Ontario, Canada

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INFLUENCE OF SULFUR AND FERRITE ON SCC AND CORROSION FATIGUE BEHAVIOR OF MODEL HEATS OF STAINLESS STEEL

E. A. West, C. Tackes, G. Newsome, N. Lewis

Bechtel Marine Propulsion Corporation (BMPC), West Mifflin, PA 15122-0079, USA

ABSTRACT

There is growing evidence that sulfur can have a beneficial effect on the stress corrosion cracking (SCC) and corrosion fatigue (CF) performance of stainless steel in deaerated hydrogenated water despite the well documented detrimental effect it has on the environmentally assisted cracking behavior of low alloy steels in oxygenated water and stainless steels in boiling water reactor normal water chemistry environments. Model heats of stainless steel have been generated with sulfur contents of


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All model heats were generated by making controlled additions of either sulfur powder or iron and chromium to the base 304/304L heat E5174. Heat E5174 was dual certified 304/304L bar stock stainless steel that was produced by Electralloy. The three model heats with elevated sulfur content (E5174-M1, -M2, and –M4) were produced by Bechtel Marine Propulsion Corporation. The E5174-M3 model heat was generated by adding iron and chromium to the E5174 remelt to achieve a high chromium equivalent and low nickel equivalent composition to promote an elevated ferrite content. The model heats underwent vacuum induction melting and were mixed for 10 minutes before casting into cylindrical ingots. The material was homogenized in air at 1093°C for 24 hours and then air cooled to room temperature. Next, the ingots were open-die hot forged to refine the cast microstructure and produce a material form suitable for hot rolling. Hot rolling was performed within the dynamic recrystallization regime to further refine the microstructure and to produce a plate that could be used to fabricate compact tension (CT) specimens. Following hot rolling, the plates were annealed at 1074°C for 2.5 hours and water quenched. Materials intended for deaerated water SCC experiments underwent subsequent single pass cold rolling to impart 19-20% cold work.

Table 1 Composition of 304/304L stainless steel heats included in SCC and CF experiments. Model heats E5174-M0 though –M4 were generated by remelting the base as received heat E5174 and performing additions of sulfur powder or chromium and iron.

Post-processing microstructural analyses of the model heats revealed that they had consistent and acceptable microstructures. All five plates had an average ASTM grain size of 4, three of the plates had an average Knoop microhardness of 157± 2 HK, and TEM analyses on E5174-M2 indicated that the large majority of grain boundaries was boride free and exhibited no chromium depletion. Although heat E5174-M1 had an average ASTM grain size of 4, there were some regions of the material microstructure where a bimodal grain size distribution was observed. It is noteworthy that the hardness of the elevated ferrite heat (E5174-M3) was slightly higher than the other plates, i.e., 167 HK. This higher hardness can be attributed to the material ferrite content which was evenly distributed and determined via feritscope measurements to be 2%. The ferrite contents in the other heats of material were negligible (


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other tests were run with an initial nominal K of 27 MPa√m with periodic partial unloading (PPU) 8.6 times per day with a stress ratio (R) of 0.7 and a trapezoidal waveform with a 9,000 second hold time and 500 second rise and fall times. Crack extension was measured using the in-situ electrical potential drop (EPD) technique and all specimens underwent a post-test destructive examination (DE) to enable correction to the crack growth rates and loading conditions. The DEs were performed by taking eighteen equally spaced measurements of the crack extension across the fracture surfaces on both specimen halves. The SCC growth rate measurements are summarized in Table 2 and plotted in Figure 1. It can be seen that the SCC growth rates of the model heats with sulfur additions are over an order of magnitude slower than the SCC growth rate of the control heat E5174-M0 with


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Figure 1 Results of SCC growth rate tests conducted in 338°C DW (a) with an initial K of 27 MPa√m and PPUs applied 8.6 times per day, and (b) with an initial K of 38 MPa√m and without PPUs. Sulfur additions consistently result in reductions in the SCC growth rate. 3.1 SCC Discussion The beneficial effect of sulfur on SCC performance in DW observed in the current and previous studies by other researchers [1-4] may appear counterintuitive given the well documented understanding of the detrimental effects of elevated sulfur content on the environmentally assisted cracking (EAC) performance of pressure vessel steels and other low alloys steels in oxygenated high temperature water [7-10], the detrimental influence of sulfate impurities on cracking of stainless steel in BWR normal water chemistry, and the move toward intentional sulfur removal from 304 SS over the past few decades. Therefore, an understanding of the mechanism by which sulfur is promoting SCC retardation in stainless steel in high temperature DW is sought here in light of the opposite effect it is has on low alloy steels. MnS inclusions in low alloy steels are known to be soluble in high temperature water and their dissolution generates sulfur anions that can inhibit repassivation of the base metal following oxide film rupture. They can also promote hydrogen assisted cracking through increased hydrogen absorption into the lattice during the delayed repassivation and by acting as hydrogen traps ahead of the stressed crack tips [7, 11]. Therefore, the specimens in the current study were analyzed to search for signs of delayed repassivation along crack paths to determine if the role of sulfur on corrosion of stainless steels has similar characteristics to that of low alloy steels despite the opposite net effect. The SCC specimens from the 338°C DW experiments with PPUs were analyzed to assess the influence of sulfur on crack tip morphology and oxidation. The crack flank oxides consisted of an inner layer of chromium rich spinel and an outer layer of iron rich spinel. Measurements of the chromium rich spinel layer thickness, performed through FIB milling on the fracture surface of the CT specimen along the SCC paths, showed negligible difference between the specimens with different sulfur contents. Furthermore, the grain-to-grain variability in iron rich oxide crystal coverage on the fracture surfaces inhibited an accurate comparison between the material heats. The difference between the high and low sulfur heats was pronounced, however, in the crack tip region as shown in Figure 2. The near crack tip regions of the E5174-M4 and E5174-M2 heats with sulfur additions were generally blunter and filled with more oxide relative to the base E5174-M0 heat without sulfur additions. Transmission electron microscopy (TEM) analyses also revealed that there was a region of sulfur enrichment at the metal/oxide interfaces of the heats with sulfur additions (E5174-M4 and E5174-M2). Energy dispersive X-ray spectroscopy (EDS) maps showing the composition of the dual layer oxide and sulfur enrichment at the metal oxide interface in the heat E5174-M2 SCC specimen are shown in Figure 3. Careful examination of the fracture surface


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of the same CT specimen revealed sites where MnS inclusions had dissolved, as shown in the accompanying scanning electron microscopy (SEM) micrograph, which were identified by the morphology of the crater and EDS detection of Mn and S enrichment. This characterization work on the model heats supports the contention that MnS dissolution also inhibits repassivation of stainless steels. The iron rich oxide crystals shown in the SEM micrographs in Figure 2 are usually associated with a physical depression in the base metal. The thicker oxides in the crack tip region, which penetrate in an irregular manner into the base metal, appear less protective than the thinner more uniform chromium rich oxide on the low sulfur heat E5174-M0. The resulting crack tips are significantly blunter in the high sulfur specimens which reduce the stress concentrations at the crack tip. Despite this apparent inferior corrosion performance caused by the sulfur additions, which is consistent with the corrosion decrement of low alloy steels, the elevated sulfur model heats show a net improvement in the SCC performance. This is likely enabled by the ability of the high sulfur stainless steel to ultimately form a protective chromium rich spinel layer despite the presence of elevated crack tip sulfur levels due to MnS dissolution in the high temperature water. These types of chromium rich spinels are not present in low chromium content low alloys steels (magnetite is the predominant oxide), and so prolonged anodic dissolution of the bare metal interface and corresponding hydrogen embrittlement occurs when sulfur species are present.

Figure 2 Micrographs of SCC tip regions of model heat specimens produced from electrical discharge machining (EDM) slices taken from CT specimens following the experiment conducted in 338°C DW with PPUs. The crack tip/paths of the high sulfur heats E5174-M4 and E5174-M2 are generally wider and contain more oxide than the low sulfur heat E5174-M0. The top images are SEM micrographs and the bottom images (left to right) are bright field TEM, high angle annular dark field TEM, and secondary electron micrographs. Note that the top and bottom images do not necessarily correspond to the same crack tip.


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Figure 3 TEM micrograph and corresponding EDS maps of crack tip region of E5174-M2 (0.012 wt% S) specimen following SCC testing in 338°C DW. Sulfur enrichment is observed at the metal oxide interface, and sites of MnS dissolution are shown in the SEM image on the left indicate that these inclusions are likely contributing to the sulfur enrichment. Nickel enrichment is also observed at the metal oxide interface, which has been shown to be a signature of slow crack growth in DW [12]. Enhanced corrosion of the high sulfur materials may induce crack retardation through environmentally assisted creep. As corrosion reactions occur at the tips of SCC and CF cracks, hydrogen or vacancies can be injected into the material which can facilitate dislocation climb induced creep. Such creep behavior may result in a redistribution of crack tip stresses, effectively lowering the peak stress at the crack tip. This type of crack retardation mechanism in sulfur containing stainless steel was introduced by Mills [2] to explain fatigue retardation behavior, and it’s applicability to SCC was also recognized by the author. Other potential retardation mechanisms include crack tip blunting due to enhanced corrosion and inhibition of crack tip transport due to buildup of corrosion products. 4. CORROSION FATIGUE EXPERIMENTS Two sets of corrosion fatigue experiments were conducted on the model heats specimens; one at a high stress ratio with increasing rise time and another at lower stress ratios where crack closure effects were more pronounced. Both sets of experiments were conducted on annealed 0.6 T CT specimens in DW at 321-338°C with 30 cc/kg hydrogen. Consistent with the SCC experiments, the dissolved oxygen content of the water was


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Table 3 Summary of results from CF testing on 304/304L model heats of stainless steel in DW.

(1) The “Model Predicted CFCGR” is calculated using the Mills Model for fatigue crack growth [6], for the indicated loading conditions. (2) Nominal loading conditions are provided. The specimen has not been destructively examined. The results of the R=0.7 corrosion fatigue experiments with increasing rise time are shown in Figure 4, and both the sulfur and rise time dependence of the fatigue crack growth behavior is evident. At short rise times, there is close agreement between the experimentally measured rates and the Mills Model predictions [6]. As rise times increase, however, the rates measured for the elevated sulfur heats begin to drop significantly below the model predictions. For the longest rise times tested, the addition of 0.006 wt% sulfur results in approximately an order of magnitude reduction in the crack growth rate as can be seen through comparison of the blue and green data points. Crack growth in the 0.012 wt% sulfur specimen stalled completely for rise times in excess of 5,000 seconds. It can also be seen that the onset of significant crack retardation behavior occurs at shorter rise times for the 0.012 wt% sulfur specimen, i.e., for the 510 second rise time, both the


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The results of the lower R corrosion fatigue rise time tests performed in the 321°C DW environment are shown in Figure 5, and they also indicate corrosion fatigue sulfur dependence. The rise times were similar for all test phases and so the experimentally measured rates were normalized by dividing by the model predicted rates to account for the significant differences in R and ∆K loading conditions. The crack growth rates consistently decrease with increasing sulfur content for each set of loading conditions, but all rates fall significantly short of model predictions (note that a normalized value of 1 would indicate perfect agreement with model predictions).

Figure 5 Results of corrosion fatigue crack growth rate (CFCGR) testing in DW environments under lower R (0.1-0.5) loading conditions. The measured rates were normalized by dividing by the Mills Model predicted rates for each set of loading conditions evaluated. Elevated sulfur contents are again shown to reduce crack growth rates and it can also be seen that rates fall significantly below model predictions under these lower stress ratio loading conditions. *The crack growth rates plotted for these specimens are based on EPD data, i.e., the data has not been DE corrected.

4.1 Corrosion Fatigue Discussion The mechanism through which sulfur retards corrosion fatigue crack growth may be similar to the sulfur induced SCC retardation mechanism previously discussed. Crack tip and path analyses performed on stainless steel corrosion fatigue specimens have also revealed the presence of sulfur enrichment at the metal oxide interface of crack tips to be associated with retardation behavior at long rise times [3]. Furthermore, SEM analyses performed along the crack path of the 0.006 wt% sulfur specimen which underwent testing in 321°C DW in the current study showed heavier oxide coverage compared to the


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Figure 6 As-polished SEM micrographs of cross sections of the crack paths of heat E5174-M0 (


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were applied, respectively. As expected, it can be seen that the magnitude of the closure effect becomes increasingly prominent with decreasing ∆K and R. Closure effects would not be expected to influence crack growth rates under these conditions unless the Kmin applied to the specimen dropped below Kclosure. Therefore, for the conditions specified in Figure 7, closure effects do not impact fatigue crack growth rates for R=0.95 loading conditions shown.

Figure 7 Results of calculations performed to evaluate potential influence of closure effects on threshold behavior at high and low R values. As expected, closure effects become more prominent with decreasing ∆K and decreasing R. The EPD data were evaluated to determine if there were signs of crack closure influencing crack growth rates under low R loading conditions. Phase 1 of the R=0.14 test phase provided an opportunity to assess the potential development of closure because initially fast crack growth rates were measured at the beginning of the test phase prior to reaching slower steady growth conditions. In order to qualitatively assess the influence of closure on fatigue crack growth rates, the evolution of ∆Keff was compared to the change in crack growth rate. This was performed using the EPD active data as a function of load as shown in Figure 8. Note that although closure measurements are most often made using compliance methods, EPD measurements of closure are not without precedence [14]. When the EPD active signal was plotted as a function of load during the test phase, a clear evolution in the shape of the curve was noted with time. Snapshots of the beginning and end of this evolution are captured in Figure 8. Initially, the active signal increased only slightly more than the noise level as the load was increased to 285-310 lbs and the active signal reached a plateau. During the next 12 days a substantial slope developed in the active signal at low loads, and by day 13 it was necessary to reach loads of 500-510 lbs before the plateau was reached. The shape of these EPD curves is believed to result from a gradual unzipping of the crack surfaces and the beginning of the plateau is expected to mark the complete separation of the crack surfaces. Thus, the knee in the curve represents the crack opening load and can be used to assess ∆Keff. The measured corrosion fatigue crack growth rate (CFCGR) and ∆Keffective are plotted as a function of time in Figure 9. It can be seen that both the crack growth rate and the ∆Keffective decrease rapidly initially and then reach a plateau after approximately 8 days; suggesting that the development of crack closure could be contributing to the reduction in the crack growth rate early in the test.


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Figure 8 Active EPD signal as a function of load during the first and thirteenth day of Phase 1 of the R=0.14 test phase on the E5174-M0 (


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Although there is qualitative agreement between the reduction in the corrosion fatigue crack growth rate and ∆KEff, substitution of the ∆KEff values into the Mills Model only predicts a 2X reduction in the CF crack growth rate compared to an approximately 30X measured reduction during the first 12 days of the test phase. This difference could be due to an underestimate of the EPD indicated extent of crack closure, the Mills Model development being based on steady fatigue crack growth rate data (where some influence of closure at low stress is likely included in the data), or it could indicate that there are other mechanisms in addition to closure that are causing the reduction in the CF crack growth rate such as a large contribution of crack tip blunting. The EPD indicated closure measurements do not show large differences in the crack opening load between the


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2. WJ Mills, "Accelerated and Retarded Corrosion Fatigue Crack Growth Behavior of 304 Stainless Steel in an Elevated Temperature Aqueous Environment," in Proceedings of the 16th International Conference on Environmental Degradation of materials in Nuclear Power Systems - Water Reactors, Asheville, NC, 2013.

3. EA West, N Lewis, and R Morris, "Stainless Steel Corrosion Fatigue Crack Retardation Behavior at Long Rise Times," in TMS Annual Meeting, San Antonio, TX, 2013.

4. S Nouraei, et al., "Effects of Thermo-Mechanical Treatments on Deformation Behavior and IGSCC Suceptibility of Stainless Steels in PWR Primary Water Chemistry," in 15th International Conference on Environmental Degradation of Materials in Nuclear Power Systems - Water Reactors, Colorado Springs, CO, 2011.

5. EA West, et al., Influence of Ferrite, Cold Work, and Sensitization on SCC of Stainless Steel, in Annual Meeting of the International Cooperative Group on Environmentally Assisted Cracking (ICG-EAC). 2012: Quebec City.

6. WJ Mills, "Critical Review of Fatigue Crack Growth Rates for Stainless Steel in Deaerated Water - Part 2," in EPRI MRP Expert Panel Meeting, Tampa, FL, 2010.

7. HP Seifert, J Hickling, and D Lister, "Corrosion and Environmentally-Assisted Cracking of Carbon and Low-Alloy Steel," Comprehensive Nuclear Materials Volume 5: Nuclear Materials and Corrosion/Waste Materials, 2012, p. 105-142.

8. T Aria and M Mayuzumi, "Effects of Temperature, Dissolved Oxygen and Material Factor on SCC Behavior of Low Alloy Steels in High Temperature Water," Corrosion Engineering, Vol. 49, 2000, p. 243-248.

9. J Kuniya, H Anzai, and I Masaoka, "Effect of MnS Inclusions on Stress Corrosion Cracking in Low-Alloy Steels," Corrosion, Vol. 48, 1992, p. 419-425.

10. GL Wire, "Cessation of Environmentally-Assisted Cracking in a Low-Alloy Steel: Theoretical Analysis," Nuclear Engineering and Design, Vol. 197, 2000, p. 25-44.

11. H Hanninen, Corros. Sci., Vol. 23, 1983, p. 663-679. 12. EA West, DS Morton, and N Lewis, "Oxide Film Characterization Along Crack Paths in Stainless

Steel in Aerated and Deaerated Water Environments," in Proceedings of the 15th International Conference on Environmental Degradation of Materials in Nuclear Power Systems - Water Reactors, Colorado Springs, CO, 2011.

13. TL Anderson, Fracture Mechanics Fundamentals and Applications, Third Edition ed, Boca Raton, CRC Press, 2005.

14. CP M Andersson, S. Melin, "Experimental and numerical investigation of crack closure measurements with electrical potential drop technique," International Journal of Fatigue, Vol. 28, 2006, p. 1059-1068.

influence of sulfur and ferrite on scc and corrosion...

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