CRACK GROWTH RATE AND FRACTURE TOUGHNESS J-R CURVE …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG075.pdf · water reactors. Due to both thermal aging and irradiation embrittlement,

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

CRACK GROWTH RATE AND FRACTURE TOUGHNESS J-R CURVE TESTS ON IRRADIATED CAST AUSTENITIC STAINLESS STEELS

Yiren Chen,1 Bogdan Alexandreanu,1 Wei-Ying Chen,1 Zhangbo Li,2 Yong Yang,2 Ken Natesan,1 and Appajosula S. Rao3

1Argonne National Laboratory, 9700 S. Cass Ave., Lemont, IL 60439, USA 2University of Florida, Gainesville, FL 32611, USA

3US Nuclear Regulatory Commission, 11545 Rockville Pike, Rockville, MD 20852, USA

ABSTRACT Cast austenitic stainless steels are used for components with complex shape in the cooling system of light water reactors. Due to both thermal aging and irradiation embrittlement, the long-term performance of CASS materials is of concern. To assess the extent of embrittlement, crack growth rate and fracture toughness J-R curve tests were performed in this study on thermally aged and unaged CF-3 at ~320°C in simulated LWR coolant environments. The samples were miniature compact tension specimens irradiated to a fast neutron fluence of 5x1019 n/cm2 (E>1 MeV) or 0.08 dpa. While environmentally enhanced cracking was observed under cyclic loading, SCC crack growth rates were low for these CF-3 specimens. Following the crack growth rate tests, fracture toughness J-R curve tests were carried out in the same test environments. The irradiation-induced reduction in fracture toughness was more evident in the unaged than aged specimens. At 0.08 dpa, the fracture toughness values of unaged specimens were comparable to that of aged CF-3. Preliminary microstructural results showed that irradiation-induced microstructures evolved much more rapidly in the unaged than in the aged CASS materials, consistent with that observed in the fracture toughness tests.

Keywords: Cast austenitic stainless steel, thermal aging embrittlement, irradiation embrittlement, environmentally assisted cracking

1.0 INTRODUCTION Cast austenitic stainless steels (CASS) have been used extensively in light water reactors (LWRs) with good field experience [1]. As a near-net-shape manufacturing technique, casting is often used for producing reactor components with complex shapes, such as pump casings, valve bodies, elbows, control rod guide tube spacers, lower support columns, etc. The CF grade CASS alloys are essentially the cast version of 300-series austenitic stainless steels (SSs) and possess excellent mechanical properties and corrosion resistance in reactor coolant environments [2][3]. Cast components are solidified directly from liquid metals and have a characteristic dual-phase microstructure consisting of austenite and delta ferrite. The exact ratio of austenite to ferrite in a CASS alloy is affected by the alloy’s composition and solidification history. While austenite is the main phase, 10-20% delta ferrite is normally present in CASS components. A minimum amount of delta ferrite is required to ensure the soundness of the castings and to prevent “hot cracking” during solidification. Furthermore, the retained ferrite phase also plays a crucial role in strengthening the austenite matrix [4][5] and in improving the corrosion resistance of CASS alloys [6][7].

Despite its positive effects on mechanical and corrosion performance, delta ferrite can have a negative impact on the fracture resistance of CASS alloys. Exposed to elevated temperatures, CASS alloys can experience hardening and embrittlement. It is well accepted that the ferrite phase is vulnerable to thermal aging embrittlement after a long-term aging at 300-500°C [8][9][10]. An increased tensile strength and reduced ductility have been observed for CASS alloys after thermal aging. The mechanisms of thermal aging embrittlement of CASS are attributed to the precipitation and growth of carbides and G-phase, and

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the formation of Cr-rich alpha prime phase through spinodal decomposition [3][8][9][11]. In addition to thermal aging, neutron irradiation can also induce microstructural and microchemical changes in crystalline solids [12][13]. A large population of point defects and point defect complexes can be produced from fast neutron bombardments. The evolution of irradiated microstructure gives rise to deteriorated ductility and fracture resistance in structural alloys. For the CASS components subject to both elevated temperatures and neutron irradiations, a possible interplay between the two degradation processes needs to be evaluated. The combined effects of thermal aging and neutron irradiation must be assessed to ensure the long-term service performance of CASS components in LWRs.

In the current study, a CF-3 CASS alloy with high ferrite content was evaluated for environmentally assisted cracking and fracture resistance in low-corrosion-potential environments. The specimens were in as-cast (unaged) and thermally aged conditions and were irradiated to a low dose. The combined effect of thermal aging and neutron irradiation on the cracking susceptibility and fracture toughness of the CASS alloy was evaluated. The test results are discussed in light of microstructural characterization obtained from irradiated and thermal aged CASS specimens.

2.0 EXPERIMENTS 2.1 Material and Specimen A low-carbon CF-3 heat was selected for testing in the current study. The heat was a static cast of a 76-mm-thick slab. The chemical composition of this heat is given in Table 1. Using Hull’s equivalent factors [14], a ferrite content of 21% was estimated for this CASS alloy. A ferroscope measurement of the cast slab gave a slightly higher ferrite content of approximately 24%. Both the as-cast (i.e., unaged) and thermally aged materials were obtained from a previous research program at Argonne National Laboratory (ANL). The thermal aging was conducted at 400°C in an air environment for 10,000 hrs. More details on the casting and thermal aging history of this heat can be found in reference [3]. The samples were polished with SiC papers up to 1200 grit. After a final finish with diamond slurry, the polished surfaces were swab etched with ferric chloride. The unaged and aged microstructures are shown in Fig. 1. Neither the volume fraction nor the morphology of delta ferrite was altered by the applied thermal aging.

A sub-sized compact tension (CT) sample geometry, as shown in Fig. 2, was used for mechanical testing. The sample geometry is a modified 1/4T-CT specimen. The sample was approximately 6.5 mm thick (B dimension) and 14 mm high. The starter notch size (a dimension) was approximately 6 mm. To ensure an in-plane crack growth, side grooves approximately 5% of the sample thickness were added on both sides of the sample.

Neutron irradiation was conducted at the Halden reactor in Norway. Both unaged and aged samples were irradiated in a helium-filled capsule at approximately 320°C to a fast neutron fluence (E >1 MeV) of 5.56 x 1019 n/cm2, or about 0.08 dpa for SS samples. Additional details of the irradiation can be found in reference [15].

2.2 Crack Growth Rate and J-R Curve Tests Crack growth rate (CGR) and fracture toughness J-R curve tests were performed with two servo-hydraulic testing systems in the Irradiated Materials Laboratory (IML) at ANL. The IML is an air-atmosphere beta-gamma hot-cell facility maintained at a negative pressure with respect to the surroundings to maintain a proper radiological barrier. Each hot cell is equipped with its own loading frame, autoclave (1-liter 316 SS), and data acquisition system.

The tests were performed either in simulated pressurized water reactor (PWR) water or in low dissolved-oxygen (DO) high-purity water (corresponding to boiling water reactor (BWR) with hydrogen water chemistry (HWC)) provided by two water recirculating loops. Both environments have low corrosion potentials, which are known to reduce the sensitivity of SSs to stress corrosion cracking (SCC) [16]. The

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simulated PWR water contained approximately 2 ppm lithium, 1000 ppm boron, and 2 ppm hydrogen. The conductivity was approximately 20 µS/cm. The low-DO high-purity water contained less than 10 ppb DO, and was covered with a gas mixture of nitrogen with 4% hydrogen. The conductivity was kept below 0.07 µS/cm. For both systems, water was circulated at a rate of 20-30 mL/min through the autoclaves. The temperature and pressure of the autoclaves were kept at ~320°C and ~1800 psig, respectively, during testing.

All specimens were precracked in the test environment with a triangular waveform at a load ratio of 0.2 and 1 Hz frequency. The maximum stress intensity factors (Kmax) were between 17 and 24 MPa m1/2 for the irradiated samples and approximately 15 MPa m1/2 for the unirradiated samples. Following an initial precracking stage, environmentally enhanced cracking was induced by a series of cyclic loading steps with gradually increased rise times and load ratios. Once environmentally enhanced cracking was stabilized, the cyclic CGR test was transitioned to a SCC test under a static load with or without periodical partial unloading (PPU). SCC CGRs were then measured at several stress intensity factors. As the crack propagated during the test, a near constant-K condition was maintained at the crack tip with load-shedding.

Using the SCC crack as a starter crack, a fracture toughness J-R curve test was carried out in the same test environment after the CGR test. The sample was loaded in tension at a constant displacement rate of 4.3x10-7 m/s, and the load and sample extension were recorded continuously. The loading process was interrupted periodically and the sample was held at a constant extension for crack-length measurement. A J-integral value was calculated at each holding point. The obtained crack extensions and J values were fitted to a power-law correlation to obtain a J-resistance curve following ASTM E 1820-8a.

After the J-R tests, the specimens were fractured in air at room temperature. The fracture surfaces of the unirradiated samples and replicas of irradiated samples were examined using a scanning electron microscope (SEM). The CGR and J-R curve results were then corrected with the final crack size measured on SEM images.

2.3 Microstructural Examination Transmission electron microscope (TEM) examination was carried out to characterize the precipitate microstructures of ferrite in unirradiated, thermally aged, and neutron irradiated specimens. The TEM samples were 3-mm diameter disks cut from the same heat where mechanical test samples were obtained. The TEM samples were electrochemically polished with a twin-jet polisher in a 5% perchloric acid and methanol solution. The obtained thin foils were examined with a Hitachi-H9000 TEM at 200 keV.

The aged and irradiated specimens were also examined with atom probe tomography (APT) at the Center for Advanced Energy Studies (CAES) of the Idaho National Laboratory (INL). The APT specimens were fabricated from the ferrite phase of the CASS specimens using focused ion beam (FEI 3D Quanta FEG FIB). An ultraviolet (UV) laser-pulse APT (CAMECA LEAP 400X HR) was used for collecting data, and the 3-D reconstruction and analyses were performed using CAMECA’s Interactive Visualization Software (IVAS 3.6.8).

3.0 RESULTS 3.1 Cyclic and SCC Crack Growth Rates A total of five CGR tests were conducted on the unirradiated and irradiated CF-3 samples in both unaged and aged conditions. The specimens were pre-cracked with cyclic fatigue loading in the test environments. Environmentally enhanced cracking was induced by a series of steps with gradually increased rise times and load ratios. During the tests, crack extensions were monitored continuously with a direct current potential drop (DCPD) method. The measured cyclic CGRs were compared with the

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anticipated fatigue CGRs in air under the same loading conditions. The difference between the CGR in water and the CGR in air, i.e. the acceleration due to environment, is used to assess the extent of environmental enhancement. In general, it was difficult to establish the environmental enhancement in this CF-3 alloy in the low-corrosion-potential environments. The corrosion fatigue responses of the samples were sluggish regardless of the aging or irradiation conditions. Crack growth was often stalled at high load ratios, and repeated attempts were made to re-activate the stalled cracks and to stabilize the elevated CGRs.

Significant scatter in cyclic CGR data can be observed due to the inherent uncertainty of these tests. To assess the effects of thermal aging and neutron irradiation, each set of cyclic CGR data is fitted to a power-law correlation (CGRenv = A*CGRair

m) proposed by Shoji et al. [17]. Following the work by Shack and Kassner [18], the power-law coefficient of SSs and CASS alloys is set to 0.5. Fig. 3 shows the curve fitting results of the cyclic CGR data obtained in this study. For comparison, the corrosion fatigue curve developed for unirradiated SSs in 0.2 ppm DO water is also included [18]. No strong environmental effect can be seen among these tests, and the differences between aged and unaged or irradiated and unirradiated samples are too small to be statistically significant. The effects of thermal aging or neutron irradiation on cyclic CGRs are not clear in the current test environments.

Following the cyclic CGR tests, constant-load CGR tests were conducted with and without PPU at one or two stress intensity factors. The measured SCC CGRs are shown in Fig. 4 as a function of the stress intensity factor. All data points are significantly lower than the NUREG-0313 line, a CGR disposition curve for sensitized SSs in high-purity water [19]. While the constant-load CGRs were very low without PPU, moderate CGRs in the range of 10-11 m/s were observed with PPU in some cases. No elevated SCC susceptibility resulting from neutron irradiation (~0.08 dpa) or thermal aging (400°C for 10,000 hr) can be seen in the current test environments.

3.2 Fracture Toughness Using the crack resulting from the SCC test as a starter crack, fracture toughness J-R curve testing was performed for each sample in the test environment at approximately 320°C. After each test, a J-R curve was constructed by fitting the calculated J values and corresponding crack lengths to a power law relationship, J=C*∆an. The J value at the 0.2-mm offset blunting line (JQ value in ASTM 1820) was obtained from the J-R curve. Note that the JQ value cannot be validated for J1C in the current study since a sufficient constraint could not always be maintained due to the size of the specimens.

Fig. 5 shows the JQ values obtained at different aging and irradiation conditions. As expected, the thermal aging induces a significant reduction in fracture toughness. The JQ of the aged sample decreases by approximately 120 kJ/m2 from the unaged value. Neutron irradiation also has a significant impact on the fracture toughness. The JQ value of the irradiated sample without prior aging is comparable to that of the unirradiated sample aged at 400°C for 10,000 hr. The extent of irradiation-induced embrittlement is much greater in the unaged specimens compared to the aged specimens. With the combination of thermal aging and neutron irradiation, the JQ decreased further from the unirradiated aged value.

3.3 Fracture Morphologies of CGR and J-R Curve Tests After the J-R curve tests, the samples were broke open in air at room temperature under cyclic loading. The tested specimens were examined with SEM, and Fig. 6 shows a typical fracture surface. Fracture morphologies varied as the crack was advanced with different modes of loading. Transgranular (TG) cleavage-like cracking was the dominant failure mode for the CGR test, and the fracture morphology during the J-R curve test was predominantly ductile. At the beginning of the CGR test, coarse deformation ledges resulting from fatigue cyclic loading can be seen next to the machined notch. As the crack advanced deeper into the sample, the fracture surface became gradually smoother, suggesting an increasing trend of environmental contribution to crack propagation. At the later stage of the CGR test, ferrite morphology became visible on the fracture surface. Beyond the CGR test region, ductile dimples

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resulting from microvoid coalescence covered the entire fracture surface. The same fracture morphologies were also observed in irradiated and aged + irradiated specimens, as shown in Fig. 7. In general, no obvious difference was detected among the fracture surfaces of aged or irradiated specimens. Neither thermal aging nor irradiation (at approximately 0.08 dpa) altered the fracture modes of CGR and J-R curve tests.

3.4 G-phase Precipitation and Phase Separation in Delta Ferrite The effect of thermal aging and neutron irradiation on the microstructural evolution of CF-3 was examined with TEM. The initial microstructure of the CF-3 showed a very low dislocation density in both the ferrite and austenite phases. No carbides were observed on the ferrite-austenite boundaries, and no precipitates were present in the delta ferrite. Thermal aging at 400°C for 10,000 hr did not produce any visible change in the austenite. A low density of dislocation loops could be seen in the austenite after neutron irradiation. Microstructural changes resulting from thermal aging or neutron irradiation were mainly observed in the delta ferrite. A high density of precipitates was observed after either thermal aging or neutron irradiation. Fig. 8 shows the centered dark-field (DF) images formed with a G-phase diffraction spot. The precipitate microstructures of the aged and neutron-irradiated CF-3 are very similar. A quantitative analysis was performed using the DF images, and the size distributions of G-phase precipitates are provided in Fig. 9. The average sizes of the G-phase particles are 3.2 nm and 5.1 nm for the aged and irradiated CF-3, respectively. The precipitate number densities are 4.9x1021 m-3 for the aged sample and 2.2x1022 m-3 for the irradiated sample. Apparently, the G-phase precipitates are larger and more numerous in the irradiated specimens compared to the aged specimens.

In addition to TEM characterization, the microstructural changes resulting from thermal aging and neutron irradiation were also examined with APT. The main results of the APT study are summarized in Fig. 10 and Fig. 11. APT Cr maps of delta ferrite obtained from the as-cast, aged, irradiated, and aged + irradiated specimens are shown in Fig. 10. The Cr-rich and Cr-depleted regions are clearly visible, implying a separation of alpha and alpha prime phases. Both the spatial distance and the magnitude of Cr fluctuation are comparable between the aged and irradiation specimens, suggesting a similar effect of thermal aging and neutron irradiation on phase separation. The Cr segregation appears to be more extensive when both thermal aging and neutron irradiation are present as shown in Fig. 10. To further illustrate the segregation of Fe and Cr, the statistical distributions of Fe and Cr concentrations are plotted in Fig. 11. Note that the Fe-normal and Cr-normal in the figure are the normal distributions about the expected Fe and Cr concentrations in ferrite. The effect of thermal aging and neutron irradiation on the redistributions of alloying element is evident. After either aging or irradiation, the Fe and Cr profile peaks are shifted in two opposite directions, indicating a non-random distribution of the two elements. The peak concentrations of Fe and Cr are also lowered and the profiles broadened after aging or irradiation. It is interesting to note that these changes in the concentration profiles are nearly identical for the aged and irradiated specimens. Additional changes can be seen in the aged + irradiated specimen, suggesting a more extensive segregation in this sample.

4.0 DISCUSSION The CGR tests performed in this study did not show a significant cracking susceptibility of CF-3. Neither the thermal aging nor the neutron irradiation elevates the cracking susceptibility. This favorable cracking response is anticipated due to the low-corrosion-potential environments used in the current study. In addition to the beneficial effect of the test environments, the presence of delta ferrite in the CASS materials may also play an important role in minimizing crack growth. It has been shown that CASS samples are more resistant to SCC than wrought SSs in high-DO water [6][7]. This superior SCC performance of the duplex microstructure may arise from the deformation behavior of the ferrite phase. Compared with austenite, the ferrite phase exhibits little plasticity during the SCC tests as shown in Fig.

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7. This observation is consistent with nanoindentation measurements where ferrite is found to be much harder than surrounding austenite in CASS alloys [20]. With a higher fraction of the total load being carried by the ferrite phase, plastic flow within austenite can be considerably reduced. As a result, the slip-dissolution process, which is the leading mechanism for SCC in high temperature water, could be hindered to some extent.

While the current study shows a good cracking response of CF-3, a large reduction in fracture toughness is observed after thermal aging at 400°C for 10,000 hr. The mechanisms for thermal aging embrittlement of CASS materials have been studied extensively [8][9][11]. Phase instability of delta ferrite is attributed to the embrittlement resulting from thermal aging. A miscibility gap in the Fe-Cr system provides the driving force for phase separation. The decomposition of ferrite is also accompanied by redistributions of other alloying elements, which can lead to additional nucleation and growth of precipitates (e.g., G-phase). These secondary phases within ferrite or at austenite/ferrite boundaries are believed to be the microstructural origin of thermal aging embrittlement [8][9][10][11]. In the current study, a high density of G-phase precipitates was observed with TEM in the aged specimen. The APT analysis also showed a strong indication of phase separation. These microstructural changes are consistent with those observed in previous studies [1][3], and may explain the decline in JQ after thermal aging.

The J-R curve tests performed in the current study also show a significant impact of neutron irradiation on the fracture resistance. The neutron irradiation was performed over a period of 4,300 hours. A detailed analysis of the temperature monitoring alloys installed inside the irradiation capsule indicated the irradiation temperature was between 318 and 327°C. A temperature gradient resulting from gamma heating was negligible since the volume of the samples was small. Assuming the samples were at 327°C during the entire period of the irradiation, an equivalent aging time at 400°C can be calculated based on the previous work on thermal aging embrittlement [22]. Using the thermal activation energy (65-250 kJ/mole) suggested in reference [22], the anticipated thermally induced embrittlement during the irradiation would be equivalent to that thermally aged at 400°C for ~20 to 1,000 hours. The level of embrittlement (i.e., the loss of fracture toughness) would be insignificant with such short thermal exposures. Thus, neutron irradiation must play an important role in the significant decline in fracture toughness observed in this study. As shown in Fig. 5, the JQ value of the irradiated specimen was reduced to a similar level expected from 10,000 hours thermal aging at 400°C. The extent of irradiation-induced embrittlement was also greater for the unaged specimens compared to the aged ones.

An extensive irradiation defect microstructure (i.e., dislocation loops or defect clusters) is not anticipated at such a low dose exposure (approximately 0.08 dpa), therefore, the observed embrittlement must be attributed to irradiation-induced precipitates or secondary phases. Both the TEM and APT studies have showed that the neutron irradiation introduced a large number of G-phase precipitates similar to those induced by the 400°C thermal aging. Redistribution of alloying elements is also observed in the APT analysis in neutron-irradiated specimens, suggesting the occurrence of phase separation. These irradiation-induced microstructural changes (with respect to the unaged microstructure) are almost identical to that occurring during thermal aging, but, at much higher temperatures. It is clear that the neutron irradiation conducted in this study enables or facilitates the precipitation or phase separation that would otherwise be absent or would occur too slowly at this irradiation temperature. Since the same precipitation and phase separation processes are responsible for the embrittlement induced by thermal aging and neutron irradiation, the two degradation mechanisms are not independent, but related at the microstructural level for CASS alloys.

Furthermore, the difference in the irradiation-induced embrittlement between the unaged and aged samples can also be explained by the microstructural evolution of CASS alloys. Competing with ballistic mixing, irradiation-enhanced diffusion plays a crucial role in the precipitation and phase separation processes. Excess point defects resulting from displacement damage contribute to the diffusion and accelerate reactions. The effect of the enhanced diffusion is not only dependent on irradiation temperature and displacement damage rate, but also affected by initial microstructure. As shown in the

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TEM and APT results, a large number of G-phase precipitates are present and the phase separation starts with thermal aging prior to irradiation. Since the pre-existing precipitation microstructure can affect the defect migration by providing additional sinks for irradiation defects, the microstructure evolves slower in the aged specimens compared to the unaged specimens. In another ion-irradiation study conducted on CF-8 [23], the accumulation rate of G-phase precipitates and dislocation loops were found to be much lower in the aged specimens compared to the unaged specimens. This difference is consistent with the observation that the decline in fracture toughness is evident more in the unaged specimens.

CONCLUSIONS Crack growth rate tests have been carried out on CF-3 specimens with 24% ferrite. The samples were irradiated to 0.08 dpa and the CGR tests were performed at approximately 320°C in low-corrosion-potential environments. Environmentally enhanced cracking was difficult to obtain regardless of thermal aging or irradiation history and only moderate SCC CGRs in the range of 10-11 m/s were recorded in some cases. Overall, the CF-3 alloy showed good resistances to both corrosion fatigue and SCC in the low-corrosion-potential environments.

Neutron irradiation had a significant impact on the fracture toughness of CF-3. With irradiation, the fracture toughness of the unaged CF-3 was reduced considerably, to a level comparable to that of the specimen aged at 400°C for 10,000 hr. An additional reduction in fracture toughness was also observed for the aged CF-3 after irradiation. The extent of irradiation-induced embrittlement was greater in the unaged specimens compared to the aged specimens.

TEM and APT examinations revealed similar precipitation microstructures in both aged and irradiated specimens. A high density of G-phase precipitates was observed with TEM and alloying element segregations, possibly due to phase separation, were detected by APT. The neutron irradiation accelerated the development of precipitate microstructure which is detrimental for retaining fracture resistance. The two degradation mechanisms, i.e., thermal aging and neutron irradiation, are connected and may interact with each other at a microstructural level. Thus, the combined effect of thermal aging and irradiation needs to be evaluated carefully for CASS alloys subjected to both thermal aging and neutron irradiation.

ACKNOWLEDGEMENTS The authors would like to thank Drs. O. K. Chopra and W. J. Shack for their contribution to this project. Our special thanks go out to Ms. T. M. Karlsen, OECD Halden Reactor Project, Halden, for assistance with irradiations and specimen transfer. Dr. Xuan Zhang and Mr. Chi Xu are acknowledged for their contributions to the experimental effort. The TEM work was performed at the IVEM-Tandem Facility funded by the US Department of Energy Office of Nuclear Energy. The APT study was conducted at the CAES facility through the ATR-National Scientific User Facility Rapid Turnaround Program. This work is sponsored by the U.S. Nuclear Regulatory Commission, under Job Code V6454, and by the U.S. Department of Energy, under contract # DE-AC02-06CH11357.

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TABLE Table 1. Chemical composition of CF-3 tested in this study.

Cast Grade

Composition (wt. %) Ferrite Content

Mn Si P S Mo Cr Ni N C Measured 1 Calculated 2

CF-3 0.63 1.13 0.015 0.005 0.34 20.18 8.59 0.028 0.023 24% 21% 1 Measured with ferroscope. 2 Calculated with Hull’s equations.

FIGURES

Fig. 1. Optical microscope image of unirradiated CF-3 with 24% ferrite.

Fig. 2. Modified 1/4T-CT specimen used in this study (units in mm).

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Fig. 3. Best-fit curves of cyclic CGR data of unirradiated and irradiated CF-3 with 24% ferrite.

Fig. 4. SCC CGRs of unirradiated and irradiated CF-3 with 24% ferrite.

10-11

10-10

10-9

10-8

10-7

10-11 10-10 10-9 10-8 10-7

CG

Ren

v (m

/s)

CGRair (m/s)

CF-3, 24% δ, tested in PWR or low-DO high-purity water320oC

Red: unaged + irradiated, in PWR water.

Blue: unaged + irradiated, in Low-DO water.

Black: aged + irradiated, in PWR water.

Purple: aged + unirradiated,in PWR water.

Brick: unaged + unirradiated, in Low-DO water.

Corrosion-fatigue of SSs in 0.2PPM DO water [19]

10-13

10-12

10-11

10-10

10-9

10-8

10 15 20 25 30 35

Spec. A-N1, unirr., PPU 2 hr, Low-DO water Spec. A-1, 0.08 dpa, PPU 2 hr, PWR waterSpec. A-2, 0.08 dpa, PPU 2 hr, Low-DO waterSpec. A-2, 0.08 dpa, PPU 1 hr, Low-DO waterSpec. A-2, 0.08 dpa, w/o PPU, Low-DO water Spec. B-N1, unirr., PPU 2 hr, PWR waterSpec. B-N1, unirr., w/o PPU, PWR waterSpec. B-1, 0.08 dpa, PPU 2 hr, PWR water

CG

R (m

/s)

K (MPa m1/2)

NUREG-0313Curve

CASS CF-3 with 24% ferrite low-DO high-purity or PWR water318 - 320oC

Open = UnagedClosed = Aged

Blue = UnirradiatedRed = 0.08 dpa

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Fig. 5. Fracture toughness of unirradiated and irradiated CF-3 with 24% ferrite.

11


Fig. 6. Fracture surface of the unirradiated aged CF-3 tested in simulated PWR water.

CG

R te

st

Fatigue fracture

TG

Dimple fracture

Cra

ck a

dvan

ce

J-R

cur

ve te

st

Post-JR fatigue

Machined notch

12


(a) Unaged + Irradiated, CGR test region (b) Unaged + Irradiated, J-R test region

(c) Aged + Irradiated, CGR test region (d) Aged + Irradiated, J-R test region

Fig. 7. Fracture surface replicas of the unaged + irradiated sample in (a) CGR test region and (b) J-R test region, and the aged + irradiated sample in (c) CGR test region and (d) J-R test region. (cracks propagate

from the bottom to top)

γ

δ

δ

γ

13


(a) Unaged (bright-field)

(b) Aged at 400°C for 10,000 hr (c) Irradiated at 320°C to 0.08 dpa

Fig. 8. Precipitate microstructure in the ferrite of CF-3 under (a) unaged, (b) aged, and (c) irradiated conditions. (a- bright field image; b and c - centered dark field with G-phase reflection)

(a) Aged at 400°C for 10,000 hr (b) Irradiated at 320°C to 0.08 dpa

Fig. 9. The size distribution of G-phase precipitates in (a) aged and (b) irradiated specimens.

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Fig. 10. APT Cr maps in the ferrite.

Fig. 11. Shifts of Fe and Cr concentration peaks resulting from thermal aging or neutron irradiation in ferrite (Cr-normal and Fe-normal are normal distributions of the expected Cr and Fe concentrations in

ferrite).

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