Intergranular Stress Corrosion Cracking Growth Behavior of Ni-Cr-Fe Alloys in Pressurized Water Reactor Primary Water

Stabilized Type 347 (UNS S34700)⁽¹⁾ stainless steel (347SS) and Type 304 (UNS S30400) stainless steel (304SS) had been used in steam generator tubing (SG tubing) in the first nuclear-powered prototype (S1W) and submarine (S2W) and in the early stages of pressurized water cooled reactors (PWR). The background of the choice of these materials was described in the Corrosion and Wear Handbook for Water-Cooled Reactors,3 which contains a materials database and explains the concept for material selection for the first nuclear-powered submarine, the prototype (S1W), the USS Nautilus (S2W), as well as the first commercial PWR (Shippingport, Pennsylvania).

The main criteria for the selection of materials in those early days were the following:

• —high corrosion resistance in a readily available alloy for the small purification system needed in a submarine

• —good fabrication characteristics

• —extensive experience with alloys in other industries, such as petrochemical and fossil power

Even before 1957, the effect of cold work on stress corrosion cracking (SCC) had been considered for material selection for submarine and first PWR. Surprisingly, they had already found SCC in non-sensitized and cold-worked (>10%) Type 304 that had been exposed to oxygenated pure water at 260°C and 316°C. However, at that point, SCC had not been found in cold-worked non-sensitized Type 304, Type 347, and Type 316 stainless steel (UNS S31600, 316SS) after a test with hydrogenated pure water at 316°C lasting for 38 days. Their initial major concern was secondary-side cracking of steam generator tubing made of stainless steel when exposed to seawater, especially onboard a submarine. To reduce SCC, treatments with phosphate (PO₄³⁻) as an inhibitor for seawater leakage and sulfite (SO₃²⁻) as an oxygen scavenger were recommend for the secondary water chemistry based on the test results3 and the experience in the prototype reactor (S1W). Also, non-sensitized stainless steel (347SS) was recommended as a less-susceptible tubing material. These events are supposed to be the brief history for the material selection for SG tubing material for first USS Nautilus and Shippingport from the historical literature.3 However, in the early stages of PWR, it seems that SCC had been identified4–6 on the secondary side of steam generator tubing made from stainless steel, as a result of seawater leakage and caustic conditions. Research work had started from 1958 to develop the alternative SCC-resistant materials for steam generator tubing considering sea-water leakage. Results showing good SCC resistance of mill-annealed Alloy 600 (UNS N06600, MA 600) against seawater leakage environment have been reported.7–12 On the other hand, Coriou, et al.,1 had reported results in 1959 to show the SCC susceptibility of Alloy 600 in dearated high-temperature pure water simulating PWR primary water. However, no field SCC experiences were recognized in primary environment at that time. Therefore, the results reported by Coriou, et al., seem not to have been considered as a high-priority item for material selection in those days.

Eventually, the United States Navy decided to change the material used in steam generator tubing to Alloy 600 in 1962 as a countermeasure against the SCC failures on the secondary side mainly caused by seawater leakage. However, intergranular SCC (IGSCC) had occurred on SG tubing made with Alloy 600 from the primary side at U bends and mechanically expanded area at tube sheet joint throughout the world. Furthermore, intergranular attack (IGA) had occurred on SG tubing made by Alloy 600 from the secondary side at just above the tube sheet and tube support plate joint with a drilled hole. On those degradation modes of Alloy 600, both in laboratories and operating PWR, were well summarized and discussed in the literature2,13–14 by Staehle and Gorman, including important lessons learned. Staehle13 warned strongly in his writing, “In the past, serious experimental work on obvious corrosion problems was sometimes not undertaken until after previously predictable failures occurred.”

Alloy 600 was then phased out in commercial power plants because of IGSCC failures in both the primary and secondary sides. Eventually, almost all SG tubing material made by Alloy 600 was changed to thermally treated Alloy 690 (UNS N066900, TT690) from 1989. On the other hand, Siemens (now AREVA [Charlotte, North Carolina]) had decided13 to switch to modified or nuclear-grade Alloy 800 (UNS N08800, 800NG) from Alloy 600 in the beginning of 1970 based on results of Coriou's data.1 Likewise, Atomic Energy of Canada Limited (AECL, Chalk River, Ontario, Canada) decided15 in mid-1970 to switch to Alloy 800NG as the SG tubing material for the new CANadian Deuterium Uranium (CANDU) reactors. At present, the materials used in SG tubing are Alloy TT690, Alloy 800NG, and Alloy TT600. Furthermore, Ti-stabilized stainless steel, 08Kh18N10T, equivalent to Type 321 (UNS S32100, 321SS), is used16–18 as SG tubing material in the Russian Vodo-Vodyanoi Energetichesky Reactor (VVER) (Russian-type water-water energetic reactor).

The excellent performance of Alloy TT690 and Alloy 800NG is now widely recognized in primary systems in operating PWR and pressurized heavy water reactors (PHWR) throughout the world. However, Andresen, et al.,19–20 Paraventi and Moshier,21 Toloczko and Bruemmer,22 Alexandreanu, et al.,23 and Arioka, et al.,24–26 recently reported that IGSCC growth occurs in Alloy TT690 in PWR primary water if the materials have been cold-worked more than 17%. On the other hand, enough data on SCC growth of CW Alloy 800 is not available to date in hydrogenated high-temperature water. Furthermore, sufficient knowledge is not available on SCC initiation behavior for all of the SG tubing materials during long-term operation, such as beyond 60 years. Given the history of material degradation in light water reactors (LWR), it is clear that proactive research is crucial to identify problems before they manifest in the field as Staehle and Gorman2,13–14 have pointed out in their recent literature. Sharing results of such proactive work with people who work for LWR is necessary; then, they will be able to decide on the appropriate countermeasures when necessary in the future.

Furthermore, Arioka, et al.,24–26,31 have reported recently that cavity formation, which has occurred from the result of the collapse of vacancies at highly stressed grain boundaries on CW Alloy TT690, CW Alloy MA600, CW Type 316, and CW carbon steel if heavily cold-worked materials are used at high-temperature conditions. Based on those results, Arioka, et al., reported26,31 that it is necessary to take into account the possible change in bonding strength in grain boundaries caused by cavity formation. For this objective, sufficient knowledge should be prepared from a wide range beyond at least 60 years, including possible changes in bonding strength at grain boundaries, local chemical composition at the grain boundaries, and surface.

Based on these incentives, the present authors started this research work as initial steps to avoid the same mistakes that had been sustained in the past; here, countermeasures had been selected only from short-term accelerated testing.

The principal purpose of this research was to compare the SCC resistance of materials for SG tubing in PWR primary water used in the world at the present time. The measured results of 20%CW Alloy 800 are compared here with our already published data of SCC growth on 20%CW Alloy TT690,24–,26 20%CW Alloy MA600,25–,26 and 20%CW austenitic stainless steels27–,30 in PWR primary water. The second purpose is to examine the dependence of nickel and chromium concentrations on SCC in PWR primary water. Such results would provide basic knowledge for developing future SCC-resistant materials for SG tubing. The third objective was to consider whether accelerated testing in high-temperature water or steam is appropriate for developing SCC resistance with respect to initiation and/or propagation.

Alloy 800NG is the main subject of these experiments; test materials were melted in the vacuum condition in the laboratory to satisfy the specification of 800NG. A plate of Alloy 800NG was heat-treated at 980°C for 1 h in air and water-cooled. To examine the effects of nickel and iron concentrations in alloys on the rate of IGSCC growth, a series of laboratory heats of X%Ni-16%Cr-Fe alloys with 16, 20, 25, 32, 40, 50, and 60% Ni concentrations were prepared in the same condition in the laboratory. The sample plate with 60% Ni was solution-treated at 1,075°C for 1 h in air and water-cooled. The other sample plates were solution-treated at 1,030°C for 1 h in air and water-cooled. Cold work in the specimens was produced by rolling at room temperature in one direction to reduce the thickness by 20%. The chemical composition, grain size number, and mechanical properties of the materials are described in Tables 1 through 3. Behavior of carbide precipitation was examined using scanning electron microscopy (SEM) and Auger electron spectroscopy (AES) using compact tension-type SCC test specimens. Measured carbide coverage in TT690 was 52.5% as reported in our literature.24 On the other hand, any trace of carbide precipitation was not identified in Type 316, 16%Ni-16%Cr-Fe, 20%Ni-16%Cr-Fe, and 25%Ni-16%Cr-Fe alloys by 20,000X magnification SEM observations. However, very fine carbide precipitation was identified at grain boundary in 32%Ni-16%Cr-Fe, 40%Ni-16%Cr-Fe, 50%Ni-16%Cr-Fe, 60%Ni-16%Cr-Fe, and Alloy 800 even in the solution-annealed condition. Examples are shown in Figure 1. Observed carbide was assumed to precipitate during water cooling from the solution-annealed temperature due to its low carbon solubility. However, the size of carbide precipitate was very fine, usually less than 100 nm. This is around 10 times smaller than TT690 and MA600.

Specimens were machined as 0.5T compact tension-type (0.5T CT) specimens with 12.5 mm thickness. These specimens were not side-grooved. Specimens were prepared in the T-L orientation, i.e., crack growth direction parallel to the rolling direction as shown in our previous literature.24–27 A fatigue pre-crack of about 2 mm was produced using a load ratio (R = K_min/K_max = 0.1) with 8 Hz and at a K_max below the stress intensity for testing. All tests were performed under static constant load conditions. The specimens were tested at a fixed constant load; the test environment and the load were maintained constant during tests. The durations of testing were between about 721 h and 6,903 h. The initial K values of testing were approximately 30 MPa m^1/2. The final K value was less than 40 MPa m^1/2 except for some specimens, such as 16%Ni-16%Cr-Fe tested at 290°C. The fracture surfaces were analyzed after testing using SEM to determine the crack morphology and the depth of SCC. The rates of SCC growth were calculated by Equation (1).

The test apparatus consisted of a recirculating flow loop mainly made by 316SS. The flow rate of test facilities was in the range between 16 L/H and 18 L/h.

Specimens were electrically insulated from loading pins with zirconia sleeves. The SCC growth rates (the 20%CW Alloy 800, a series of X%Ni-16%Cr-Fe alloys, and 25% and 40%Ni-20%Cr-Fe alloys) were measured in simulated PWR primary water, which contained boric acid (H₃BO₃, 500 ppm as B), lithium hydroxide (LiOH, 2 ppm as Li), and dissolved hydrogen (30 cc STP H₂/kg H₂O) in the range of temperature between 290°C and 360°C. The concentrations of hydrogen were adjusted by hydrogen gas bubbling at about 0.1 MPa through the solution in the storage tank at room temperature before the solution was pumped into the autoclaves; the effluent hydrogen and oxygen were measured at ambient temperature using a hydrogen and oxygen monitor. The hydrogen concentration was controlled between 30 cc/kg H₂O and 35 cc/kg H₂O during testing. The oxygen concentration was controlled less than 5 ppb during the test.

The principal objective of these observations was to clarify whether intergranular penetrations occur on IGSCC-susceptible materials from the surface as the IGSCC initiation site resulted from local oxidation. Careful observation by SEM was performed on all the surfaces of the bottom of the CT specimens. Then, some specimens were examined in cross-sectional view to observe the morphologies of intergranular attacks. The samples were prepared by gallium ion sputtering in a focused ion beam (FIB) apparatus. Elemental composition near the intergranular attack (IGA) was measured using AES with the same sample to confirm its oxidation behavior.

Rates of SCC growth were measured in simulated PWR primary water on 20%CW Alloy 800 at 340°C and 320°C for 5,825 h and 6,903 h, respectively, using 0.5T CT specimens. Any trace of IGSCC was not observed after testing at 340°C even after 5,825 h exposure in PWR primary water. Very shallow intergranular cracking (maximum around 200 μm in depth) was observed very locally after testing at 320°C for 6,903 h exposure in PWR primary water. Fracture surfaces of these specimens are shown in Figure 2. Temperature dependence on the rate of SCC growth of 20%CW Alloy 800 were summarized together with other SG tubing materials using already published data26 on 20%CW Alloy TT690 and 20%CW Alloy MA600. Furthermore, the rates of SCC growth of 20%CW 31625–,27 were also compared in the same figure as shown in Figure 3. Then, the observed rates of SCC growth of 20%CW Alloy 800 at 340°C were compared with our already published rates of SCC growth in the same environment of 20%CW 316(TS),25–,27 20%CW Alloy TT690(TL),22–24 and 20%CW Alloy MA 600(TL),23–,24 as shown in Figure 4. From these results, the following important patterns were observed. First, excellent SCC growth resistance was observed on 20%CW Alloy 800 relative to other materials for SG tubing in the range of tested temperature, though the number of the data is very limited to date. The observed excellent SCC resistance of 20%CW Alloy 800 in primary water was supported by the field experiences of more than 40 years in CANDU and Siemens (now AREVA) reactors.

Second, reverse temperature dependence seems to be observed on 20%CW Alloy 800 as shown in Figure 3. Therefore, detailed studies are necessary to examine the exact temperature dependence of CW Alloy 800 in low-potential, high-temperature water. Typical temperature dependence with peak has already been reported27 on 10% and 15%CW 316 in PWR primary water; peak was observed in the temperature range between 320°C and 330°C. Furthermore, peak in the rate of SCC growth seems to be observed also in 20%CW Alloy TT690 and 20%CW 316 between 340°C and 360°C. This trend suggested that more than two major processes affected the rate-limiting process of SCC growth in PWR primary water in the range of the present test temperatures. This result seems to suggest that a peak in SCC growth of 20%CW Alloy 800 may exist around 320°C. In other words, simple 1/T temperature dependence was observed only in 20%CW Alloy MA600. This suggested that accelerated testing at higher temperature, such as at 360°C, seems to be only applicable in the case of heavily cold-worked Alloy MA 600. Therefore, detailed studies are essential to examine the cause of the peak in SCC growth from various kinds of perspectives not only from an engineering point of view, but also from a scientific point of view.

The principal objective of this study was to obtain basic knowledge to consider one of the causes of different rates and temperature dependence of SCC growth of 20% cold-worked materials with different nickel concentrations: Type 316 (10%Ni), Alloy 800 (32%Ni), Alloy TT690 (63%Ni), and Alloy MA600 (76%Ni) as described above. For this objective, dependence of nickel concentration on the rate of IGSCC growth was studied using 20%CW Ni-16%Cr-Fe alloys with different nickel concentrations at 290°C, 320°C, 340°C, and 360°C in hydrogenated PWR primary water. The maximum test duration was 7,980 h. After testing, specimens were fractured by fatigue in room temperature to examine the fracture morphology and SCC depth to obtain the growth rate of SCC. Examples of the fracture surface after testing at 360°C are summarized in Figures 5 and 6. Significant depth of IGSCC was observed on Ni-16%Cr-Fe alloys with a higher nickel concentration more than 60% Ni, and on alloys with lower nickel concentration less than 20% Ni. On the other hand, no IGSCC was observed on 20%CW 25%Ni-16%Cr-Fe alloy, 20%CW 32%Ni-16%Cr-Fe alloy, and 20%CW 40%Ni-16%Cr-Fe alloy at 360°C as shown in Figures 5 and 6. The measured rates of SCC growth at 360°C were summarized as a function of nickel concentration together with the measured rate of 20%CW Alloy 800, and our already published data of 20%CW 316, 20%CW Alloy TT690, and 20%CW Alloy MA600 are shown in Figure 7. The rates of IGSCC growth decrease together with increasing nickel concentrations in the range of nickel between 10% and 25%, and a high SCC-resistant zone was observed on alloys with 25% and 32% Ni, then the rate of IGSCC growth increases with increasing nickel concentration in the range between 40% and 75% Ni.

Similar dependencies of nickel concentration on SCC growth were observed also at 340°C as shown in Figure 8. These trends shown in Figures 7 and 8 are quite similar with the famous and classical results reported by Coriou1 and Staehle2 tested by U-bend specimens in dearated pure water at 350°C. However, this trend of dependence of nickel concentration changed almost completely at lower temperatures less than 320°C. Figures 9 and 10 show the fracture surfaces of test materials after testing at 290°C in PWR water. Interestingly, significant depth of IGSCC was observed in all alloys at 290°C as shown in Figure 9. The measured rate of IGSCC growth at 290°C was summarized as a function of nickel concentration in alloys as shown in Figure 11. No clear dependence of nickel concentration in alloys was observed at 290°C. Similar dependencies of nickel concentrations were observed also at 320°C as shown in Figure 12. This is a completely different trend compared with the trend at higher temperatures more than 340°C. Observed remarkable change in dependence of nickel concentrations on SCC growth suggested that the main rate-limited processes on IGSCC growth seems to change in the range of temperatures between 320°C and 340°C. Regarding the effect of chromium addition on IGSCC resistance, no significant beneficial effect was observed at 360°C comparing the measured SCC growth rates of Alloy TT690(30%Cr) and 60%Ni-16%Cr-Fe alloy as shown in Figure 7. However, beneficial effect of chromium addition was observed at 340°C in 60%Ni-16%Cr-Fe alloy and Alloy TT690 (30%Cr) as shown in Figure 8.

Furthermore, a clear beneficial effect was observed at 320°C in 32%Ni-16%Cr-Fe alloy vs. Alloy 800 (20%Cr) and 60%Ni-16%Cr-Fe alloy vs. Alloy TT690 (30%Cr) as shown in Figure 12. This result suggested that the effect of the chromium addition in alloys seems to be affected by the test temperature of PWR water. A similar beneficial effect of chromium additions on IGSCC growth have also been published by Arioka, et al.,26,28 on Fe-X%Cr-14%Ni-2.3%Mo alloys in PWR primary water at 320°C. Sometimes, inter-granular penetrations were observed on the surface of the bottom of the CT specimen on IGSCC susceptible materials. Figure 13 shows examples of SEM images of intergranular attack on the surface after testing at 320°C on 32%Ni-16%Cr-Fe and 60%Ni-16%Cr-Fe alloys in PWR water.

Figure 14 shows one of the examples of the cross-sectional view of shallow intergranular penetration after sampling by FIB on 32%Ni-16%Cr-Fe alloy tested at 320°C. Judging from the result of AES analysis, Cr-rich oxide was formed in the intergranular attack as shown in Figure 15. Figure 16 shows the temperature dependence of solubility of chromium oxide (Cr₂O₃) in PWR water calculated using Criss and Cobble's correspondence principle.32 The result suggested that solubility of Cr₂O₃ in PWR water with hydrogen seems to increase with increasing temperature. Nothing is clear to date on the cause of the temperature dependence of chromium addition on IGSCC resistance. However, observed temperature dependence of the effect of chromium addition seems to be related with its stability of chromium oxide in high-temperature water. Judging from the observed result of nickel and chromium concentration dependence on IGSCC growth, excellent IGSCC resistance of Alloy 800 seems to be affected by not only chromium addition, rather than the combined effect of nickel and chromium concentration in the alloy. More detailed studies are crucial for the development of SCC-resistant materials in the near future. Furthermore, more precise studies are necessary to make clear the cause of the temperature dependence of the effect of chromium addition in alloys on SCC resistance.

The principal objective of this study was to obtain basic knowledge to understand the causes of different temperature dependencies of SCC growth of 20% cold-worked materials with different nickel and iron concentrations: Type 316 (10%Ni), Alloy 800 (32%Ni), Alloy TT690 (63%Ni), and Alloy MA600 (76%Ni) are shown in Figure 4. For this objective, the measured rates of SCC growth were summarized as a function of temperature on three groups of materials with different nickel concentrations, high-nickel alloys, medium-nickel alloys, and low-nickel alloys separately.

Figure 17 shows the results on Alloy 600 (76% Ni), 60%Ni-16%Cr-Fe, and 50%Ni-16%Cr-Fe alloys. Peaks of the growth rates of IGSCC were observed in 60% Ni and 50% Ni alloys in the range of temperature between 320°C and 340°C. Figure 18 shows the result on 40%Ni-16%Cr-Fe and 32%Ni-16%Cr-Fe alloys. Peaks of the growth rate of IGSCC were recognized around 320°C. Figure 19 shows the result of 25%Ni-16%Cr-Fe, 20%Ni-16%Cr-Fe, and 16%Ni-16%Cr-Fe alloys. In those alloys with low-nickel concentrations, clear peaks were not observed in the range of temperature between 290°C and 360°C. The growth rate of IGSCC seems to increase with decreasing temperature; this is opposite from Alloy 600. Then, results of some alloys with different nickel concentrations were summarized in the same figure to consider the effect of nickel concentration in alloys on the peak temperature as shown in Figure 20. From the result shown in Figure 20, the peak seems to shift to a temperature together with decreasing nickel concentration in alloys. Arioka, et al., had reported29 the similar temperature dependence with peaks in 10%, 15%, and 20%CW non-sensitized 316SS in PWR water. And, they reported that the peak temperature seems to be influenced also by the degree of cold work based on the observed results in non-sensitized CW316 in PWR water.

One of the important points of these findings, as shown in Figures 17 and 20, are that the order of the rate of SCC growth of the materials at 360°C is different from the rate measured at 290°C. Therefore, we have to give careful attention to predict SCC susceptibility at operating temperatures from the results of accelerated testing at higher temperatures such as 360°C or testing in a steam environment more than 400°C. Another point of these results on temperature dependence with a peak suggested that more than two rate-limited processes affect the process of IGSCC growth including the oxidation process. Better understanding of the previous results is crucial. Detailed and comprehensive studies on the mechanism are essential not only from the engineering but also from the scientific points of view.

• ❖ Excellent SCC growth resistance was observed on 20%CW 800NG in the range of temperature between 320°C and 340°C. Excellent IGSCC resistance of CW Alloy 800 seems to be affected by the chromium concentration rather than the effect of nickel concentration. The observed excellent SCC resistance of CW 800NG in primary water has been supported by the field experiences of about 40 years separately in CANDU and Siemens reactors.

• ❖ Significant dependence of the rate of IGSCC on nickel concentration was observed at temperatures of 340°C and 360°C. The rate of IGSCC growth decreases with increasing nickel concentration in the range of Ni concentration between 10% to 25% Ni; furthermore, the rate increase with increasing nickel concentration in the range of nickel concentration between 50% to 76%. This trend is considerably similar to the results reported previously by Coriou, et al.,1 and Staehle and Gorman2 testing in dearated pure water at 350°C.

• ❖ Completely different dependencies on nickel concentration were observed at 320°C and 290°C where no dependency was observed. On the other hand, significant dependency was observed between 320°C and 340°C.

• ❖ Significant beneficial effects of chromium additions in alloys on IGSCC resistance was observed at 320°C, comparing the result of 32%Ni-16%Cr-Fe alloy with Alloy 800 (32%Ni-20%Cr-Fe) and 60%Ni-16%Cr-Fe alloy with Alloy TT690 (63%Ni-30%Cr-Fe). However, no beneficial effect of chromium additions in alloys was observed at 360°C between 60%Ni-16%Cr-Fe alloy and Alloy TT690 (63%Ni-30%Cr-Fe).

• ❖ Shallow intergranular perforation was observed on the free surface on some of the IGSCC susceptible materials even without stress. This result suggests that the then observed temperature dependence of chromium additions in alloys seems to be affected by the equilibria and kinetics of oxide film formation.

• ❖ The peak temperature in the growth rate of IGSCC was observed in almost all test materials except for 20%CW Alloy 600. The peak temperature seems to decrease to lower temperature with nickel concentration decreasing. This result suggested that the process in IGSCC growth was controlled by more than two rate-limiting processes.

• ❖ The important engineering meaning of the complicated temperature dependence with the peak is that the rate-limited process of IGSCC growth at higher temperature is different from that at the operating temperature. Furthermore, the order of SCC resistance at higher temperature is not the same at operating temperature. This means that careful attention should be given to assessing SCC susceptibility and the rate of SCC growth at operating temperatures.

The author acknowledges the financial support of this work by Kansai Electric Power Co. Inc. The author is most grateful to R.W. Staehle for helpful suggestions and detailed discussions. The authors appreciated R. Tapping and M. Wright in AECL for their helpful discussion regarding this work. Careful and reliable experimental support by the staff at INSS, K. Murakami and M. Hirao, is also gratefully acknowledged.