THE EFFECT OF ENVIRONMENT, CHEMISTRY, AND …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG011.pdfaustenitic stainless steel, and the behavior shown is typical for stainless

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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THE EFFECT OF ENVIRONMENT, CHEMISTRY, AND MICROSTRUCTURE ON THE CORROSION FATIGUE BEHAVIOR OF AUSTENITIC STAINLESS STEELS IN

HIGH TEMPERATURE WATER

L.B. O’Brien1, R.G. Ballinger2, D.J. Paraventi1, L. Yu2, Y. Maruno2, P.W. Stahle2

1Bechtel Marine Propulsion Corporation – Bettis Laboratory, West Mifflin, PA, 15125 2Massachusetts Institute of Technology – H.H. Uhlig Laboratory, Cambridge, MA, 02139

ABSTRACT

Laboratory testing of austenitic stainless steels in Light Water Reactor type environments has shown corrosion fatigue crack growth rates (CFCGR) which are often accelerated over those in air. Recent research that has focused on material chemistry (specifically high sulfur content) and loading conditions has produced rates where the environmental enhancement can be reduced (retarded) relative to the normally enhanced rates in water. Since mechanistic understanding is limited in this area, the effect of sulfur on the CFCGRs of austenitic stainless steel was investigated to understand the mechanisms behind environmental enhancement and retardation. Fatigue tests on low and high sulfur heats of 304L stainless steel were performed in deaerated pressurized water at 288 oC under sawtooth loading with rise times from 5.1 to 5100 seconds with ∆K of 8.6 or 17.1 MPa√m and load ratio, R=Pmin/Pmax, of 0.4 and 0.7. The low sulfur content material experienced increasing crack growth rates with increasing rise time, while the high sulfur material experienced a decrease in rates. The retardation in crack growth rates is thought to occur because of increased corrosion at the crack tip as a result of the presence of sulfur from the material leading to either crack closure at low load ratios or an enhanced creep mechanism, which leads to stress relaxation ahead of the crack at high load ratios. Analytical electron microscopy of the fracture surface was used to evaluate these proposed mechanisms.

Keywords: Corrosion fatigue, 304SS, sulfur

1. INTRODUCTION

The corrosion fatigue crack growth behavior of austenitic stainless steels in high temperature dearated pressurized water (DPW) environments is highly dependent on loading conditions, environment, and material chemistry. Experimental data and lifetime performance in high temperature DPW environments have shown CFCGRs that are significantly accelerated compared to those in air. [1,2,3] Currently, conservative factors (based on experimental data) are applied to the air fatigue crack growth curves from Section XI of the ASME Boiler and Pressure Vessel Code in order to predict component lifetime. [4] Due to a lack of fundamental understanding of environmental effects in DPW on stainless steels, component lifetime predictions are often uncertain, and the conservative factors applied often result in reduced calculated lifetimes and potentially decreased inspection intervals. As a result, there has been recent effort to revise existing fatigue crack growth rate models to account for environmental effects.

In high temperature DPW, environmentally enhanced CFCGRs manifest themselves as an increase in rate with increasing rise time, such as that seen in Figure 1, shown for a range of R from 0.3-0.9 and cyclic stress intensity (∆K) from 4.5-15 MPa√m. [2] Testing was performed on low sulfur (0.001 wt%) austenitic stainless steel, and the behavior shown is typical for stainless steels with low sulfur content. Recent data on fatigue crack growth of high sulfur austenitic stainless steels in DPW has shown a beneficial effect of sulfur on the crack growth rates, such that the environmental enhancement is reduced, or even entirely removed, with increased sulfur content. This phenomenon is referred to as crack growth retardation, and has been observed in both corrosion fatigue (under long cycle times) and stress corrosion cracking (SCC) for high sulfur materials ranging from 0.015-0.030 wt% sulfur, where 0.03 wt% sulfur is the upper limit allowed by ASTM specification A666. [3,5]


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In this paper, the role of sulfur content on the CFCGRs of austenitic stainless steels is further examined in DPW environments to provide additional insight into the likely mechanisms controlling environmental enhancement and retardation. Targeted CFCGR testing and subsequent post-test electron microscopy results are correlated, and future work to examine specific phenomenon is recommended.

Figure 1: Crack growth rates as a function of rise time shown for a range of stress ratio (R) from 0.3-0.9 and cyclic stress intensity factor ∆K from 4-15 MPa√m. Crack growth rates increase with increasing rise time. [2]

2. EXPERIMENTAL

2.1 Material Chemistries and Specimen Preparation

Three heats of dual certified 304/304L stainless steel (certified with the mechanical properties of 304, with the low carbon chemistry of 304L) were used during testing. All materials came from forged bar. 1.0T compact tension (CT) specimens were machined in the L-R orientation per ASTM E399. The material chemistries are shown in Table 1. The E5174 heat also contained a higher than typical boron content of 18 ppm. This resulted in the precipitation of (Cr,Fe)2 B-type borides. While this may have affected the corrosion behavior, there was no observed effect on the CFCGR behavior when compared against the D2739 heat and CFCGRs reported in the literature. Specimens were precracked in room temperature air in accordance with ASTM E647 guidelines before fatigue testing.

Table 1: Material chemistries for 304/304L Stainless Steel (wt%) heats used in testing

Heat of 304/304L C Mn S Cr Ni N Fe

E5174 (Low Sulfur) 0.020 1.60 <0.0025 19.8 10.1 0.085 Bal

A16830 (High Sulfur) 0.021 1.73 0.032 18.5 8.2 0.089 Bal

D2739 (Low Sulfur) 0.019 1.60 <0.0025 18.3 9.4 0.051 Bal

2.2 Fatigue Testing

Two autoclaves were used during this study, with either an MTS servo-hydraulic or an Instron electromechanical fatigue machine. The servo-hydraulic system is composed of stainless steel components, whereas the electromechanical fatigue system is composed of titanium (all wetted surfaces


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with temperature above 150C) and stainless steel components. Prior to each test, the autoclave used for the test was conditioned to produce a uniform film of corrosion product on internal components (due to the autoclave being open to the environment for long periods of time before the test program began, the autoclaves initially required longer conditioning periods). Conditioning was performed by sparging with argon gas until an outlet conductivity of 0.11 µS/cm (desired for test) was reached, followed by sparging with 100% oxygen gas, until a concentration of 3000ppb was reached. Finally, 100% hydrogen was bubbled into the autoclave to establish dearated conditions (the final dissolved hydrogen was 30 cc H2/H2O, with O2 <5ppb). The autoclaves were operated at 288oC and 10.2 MPa for all tests. The room temperature pH for these tests was controlled to between 10 and 10.7.

Fatigue tests were performed under constant K at values of approximately either 8 or 15 MPa√m and R of either 0.4 or 0.7. The applied load was autoclave pressure compensated, due to a 2.7 kN internal pressure-induced force on the specimen. Crack length was measured in situ using a direct current potential drop (DCPD) system, and then destructively evaluated post-test to correct the DCPD indicated crack lengths and, hence, growth rates. Table 2 details the scope of the fatigue testing performed on each specimen, including the loading conditions and crack growth regime that each crack tip was produced to. The production of corresponding crack tips is important for further characterization studies using advanced microstructural and chemical techniques.

Table 2: List of specimens and loading conditions tested

Specimen

Testing Step

K (MPa√m)

Minimum to Maximum Load Ratio

(R)

Rise Time, tR (seconds)

Crack Growth

Rate Regime

5174-LR-11 (Heat E5174)

Step 1 19.0 0.4 5.1

Enhanced Step 2 19.7 0.4 51 Step 3 20.0 0.4 510 Step 4 19.7 0.4 5100

A16-LR-10 (Heat A16830)1

Step 1 20.9 0.4 5.1

Retarded

Step 2 19.1 0.4 51 Step 3 18.8 0.4 510 Step 5 18.7 0.4 5100 Step 6 10.6 0.7 5.1 Step 7 9.8 0.7 51 Step 8 9.6 0.7 5100 Step 9 5.7 0.7 5100

2739-LR-2 (Heat D2739)

Step 1 6.7 0.7 5.1 Enhanced

Step 2 6.9 0.7 51

2739-28 (Heat D2739)

Step 1 9.4 0.7 5.1 Enhanced

Step 2 9.8 0.7 51 A16-32

(Heat A16830) Step 1 7.3 0.7 5.1

Retarded Step 2 7.9 0.7 51

1For specimen A16-LR-10, only pertinent steps are included. For example, beachmark phases are not included. Due to length of experiment, no 510s rise time step was performed at R=0.7. Due to a computer failure, the system pressure compensation was removed, increasing the maximum load during Step 9.


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2.3 Post Test Sectioning

Upon completion of testing, the specimen was sectioned into three parts, as shown in Figure 2, in order to perform destructive evaluation to correct the electrical potential drop (EPD) indicated in situ crack growth rates (Section B), to evaluate the fatigue surface under scanning electron microscopy (SEM) (Section B), and for further microscopy (Sections A and C), such as atom probe tomography (APT) and transmission electron microscopy (TEM). The method of specimen sectioning is taken from Reference 6.

Figure 2: Sectioning of specimens for further analysis. Cuts were made using electrical discharge machining (EDM) using a copper wire and a water bath.

3. RESULTS

3.1 Low Sulfur Material (Heats E5174, D2739)

3.1.1 Crack Growth Rates

Three tests were conducted on low sulfur materials to produce crack tips for short and long rise times. Figure 3 shows the destructively evaluated (DE) corrected crack growth rates for all low sulfur tests as a function of ∆K. Results are presented along with two types of reference curves: the crack growth rate in air (ASME Section XI) and the crack growth rates in high temperature DPW, as proposed in ASME Draft Code Case N-809 (referred to as “N-809”). [4,7] The experimental crack growth rates for low sulfur material increase with increasing rise time, and match the N-809 curves well, demonstrating full environmental enhancement.

Section

Section

Section


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Figure 3: Experimental crack growth rates as a function of ∆K. (A) Crack growth rates for Heat E5174 (low sulfur), tested at R=0.4. (B) Crack growth rates for Heat D2739 (low sulfur), tested at R=0.7. Crack growth rates are shown along the ASME Section XI air curve (grey) and N-809 high temperature DPW crack growth rate curves for corresponding rise times (shown as black, blue, orange and red for rise times of 5.1, 51, 510, and 5100 seconds, respectively). For both stress ratios, the experimental crack growth rates correspond well to the N-809 reference curves, indicating full environmental enhancement.

3.1.2 Fractography

Fracture surfaces were examined using scanning electron microscopy without removal of corrosion product to evaluate the surface morphology including the effect of sulfur content and loading conditions. In all low sulfur heats, the fracture surface was consistent with full environmental enhancement. Yellow arrows in the bottom left of each image denote macroscopic crack growth direction.

3.1.2.1 Heat E5174

The fracture surface for all rise times is seen in Figure 4, with lines denoting the individual test phases and crack growth direction. Early test phases, such as Step 1 (5.1-second tR, denoted as B), show a heavier corrosion product than later test phases, such as Step 4 (5100-second tR, denoted as E), due to the length of time the specimen was in test in the autoclave; the oxide at the beginning of the test is the oldest on the crack flank. The fracture surface in general had a sparse oxide layer, consistent with fractography reported in References 3 and 7.

A B


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Figure 4: The complete fracture surface for specimen 5174-LR-11. The yellow arrows denote macroscopic crack growth direction. The air precrack is denoted by A, Step 1 with a 5.1-second tR is denoted by B, Step 2 with a 51-second tR is denoted by C, Step 3 with a 510 second tR is denoted by D, Step 4 with a 5100-second tR is denoted by E, and the fatigue apart in air is denoted by F.

Figure 5 shows the fracture surface during the first step of testing, with a 5.1-second rise time. Arrows indicate regions of river patterns (A) and areas where localized cracking occurred perpendicular to the general crack growth direction (B and C). This morphology is a type of attack where cracking occurs adjacent to areas that are crack resistant. [2] The crack growth path in this region will continue along areas of less fatigue resistance, or along a different crystallographic slip plane than the general crack direction.


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Figure 5: Step 1 (5.1-second tR). Arrow A shows an example of river patterns. Arrows B and C show areas where the cracking occurred perpendicular to the macroscopic crack growth direction.

Figure 6 shows the fracture surface during the fourth step of testing, with a 5100-second rise time. Fatigue striations, denoted by arrow B, were within a factor of 2 of the macroscopic crack growth rate reported for this step. Striations are a result of the crack opening in increments as the load is cyclically applied during corrosion fatigue, and occur perpendicular to the crack growth front. Arrow A points most likely to a grain boundary, due to its straight nature, rather than the curvature seen in the fatigue striations.

Figure 6: Step 4 (5100-second tR). Arrow A denotes a possible grain boundary. Arrow B denotes fatigue striations.


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3.1.2.2 Heat D2739

Heat D2739 test specimens were only used to produce short rise time (51 seconds) crack tips. Figure 7 shows the fracture morphology for the entire test for specimen 2739-LR-2, with the measurements of the first test step (5.1 seconds) and the entire fatigue test. The two test steps were hard to distinguish from one another based upon visual inspection of the fracture surface; the surface had a uniform corrosion product, and was significantly less decorated than the high sulfur material which will be discussed in the following section. The fracture surface is also crystallographic in nature. The fracture surface of Specimen 2739-28 is identical to Figure 7.

Figure 7: The complete fracture surface for specimen 2739-LR-2. Step 1 (5.1-second tR) is denoted as A, and Step 2 (51-second tR) is denoted as B. Measurements are included to show the location of Step 1 (1.2 mm) and the total fatigue test (2.126 mm).

In addition to fatigue striations and river patterning of the crack growth front, secondary cracking and slip lines were seen on the fracture surface. Examples of a river pattern and slip lines from specimen 2739-LR-2 are seen in Figure 8. Secondary cracking and fatigue striations from specimen 2739-28 are seen in Figure 9. No striations were observed on specimen 2739-LR-2, possibly due to the lower ΔK the test was run at compared to that of specimen 2739-28. At ΔK values below approximately 13 MPa√m, crystallographic faceting dominates, whereas a combination of crystallographic faceting and fatigue striations are more dominant at higher values, likely due to the difference in slip behavior between the large plastic zone that develops at the maximum load and the smaller plastic zone that develops during the reversed cyclic loading. [3]


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Figure 8: Images from Specimen 2739-LR-2. (A) River pattern crack propagation. (B) Example of slip lines. Unlike fatigue striations, these are straight and not parallel to the crack front, as well as the spacing being too large. The measurements of the slip line spacing are 3.093 µm, 4.237 µm, and 2.296 µm, shown in green. Ridges are formed perpendicular to these slip lines as the crack propagates on parallel planes.

Figure 9: Images from Specimen 2739-28. (A) Example of secondary cracking. (B) Example of fatigue striations. Striations are around 250-560 nm (shown in green), which is on the order of magnitude of the macroscopic crack growth rate.


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3.2 High Sulfur Material (Heat A16830)

3.2.1 Crack Growth Rates

Two tests were conducted on the high sulfur material to again produce both short and long rise time crack tips. Figure 10 shows the DE corrected crack growth rates for all high sulfur tests as a function of ∆K, presented alongside the air (ASME Section XI) and high temperature DPW (N-809) crack growth rates. Unlike the low sulfur materials, experimental crack growth rates do not increase with increasing rise time; crack growth rates tend to remain the same or decrease as rise time increases, demonstrating retarded crack growth rates. The N-809 reference curves are for full environmental enhancement, and therefore over predict the experimental crack growth rates for the high sulfur material. The crack growth rates from specimen A16-LR-10 were determined at higher ∆K values than those found for specimen A16-32, though both correspond well to the N-809 reference curves at a rise time of 5.1s.

Figure 10: Experimental crack growth rates as a function of ∆K. (A) Crack growth rates for Specimen A16-LR-10, tested at R=0.4. (B) Crack growth rates for Specimens A16-LR-10 and A16-32, tested at R=0.7. Crack growth rates are shown along the ASME Section XI air curve (grey) and N-809 high temperature DPW crack growth rate curves for corresponding rise times (shown as black, blue, red, orange and red for rise times of 5.1, 51, 510, and 5100 seconds, respectively). CFCGR retardation is seen in both stress ratios.

3.2.2 Fractography

Fractography from the high sulfur material A16830 shows a heavier corrosion product in almost all test steps when compared to the low sulfur materials (Heats E5174 and D2739).

A B


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3.2.2.1 Heat A16830 Short Rise Time Test

Specimen A16-32 was tested at rise times of 5.1 seconds (Step 1) and 51 seconds (Step 2). A distinct change in fracture morphology was seen between the two steps, unlike what was observed in the low sulfur material. Figure 11 shows the transition between the two rise times tested. Step 1(denoted by A) shows well-defined cleavage facets, and a thinner oxide surface, much like the low sulfur material. In Step 2 (denoted by B), the fracture surface becomes featureless in nature, and a heavier layer of oxide is observed. A higher magnification image of the transition between steps is shown in Figure 11 (B).

Figure 11: (A) The complete fracture surface for Specimen A16-32. (B) Transition between Steps 1 A and Step 2 B. Step 2 exhibits a heavier corrosion product than Step 1, and the surface is featureless, unlike the crisp faceting seen in Step 1.

During the first approximately 200 µm of the second test step growth (51 s), the surface has a constant layer of thick oxides, shown in Figure 12 (A) as region B. Following this region is an area of alternating ‘bright’ and ‘dark’ bands of oxides, denoted as region C. The length of region B is consistent with the experimental crack length shown in Figure 12 (B). The crack growth rate after this transition stage was around twice that of the crack growth rate during the transition stage (around 4x10-10 m/s rather than 2x10-10 m/s), though still lower than that of the respective test step in the low sulfur materials (around 7x10-10 m/s). This change in crack growth rate was due to the crack growth mechanisms that are operative. The test conditions, such as bulk chemistry, temperature, pH and loading (except for rise times) were not changed during the test. Further investigation is in progress to evaluate potential mechanisms.


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Figure 12: (A) Alternating pattern of oxides seen in Step 2. Step 1 (5.1-second tR) is denoted by A, the transition stage of Step 2 (51-second tR) is denoted by B, and the region of the alternating oxide feature is denoted by C. Measurement shown is length of the transition stage, around 200 µm. (B) Crack growth rate as a function of cycles. The crack growth rate in the transition region is twice as low as that in the alternating oxide feature region.

Figure 13 shows a higher magnification of this feature, along with a corresponding backscatter electron (BSE) image. This alternating feature was observed along most of the second test step.

Figure 13: (A) Alternating oxide feature. (B) BSE image of the region in (A).

Areas of dissolution ‘holes’ were seen throughout the fracture surface, indicative of areas where a MnS stringer in the material was dissolved. An example of a dissolution hole is seen in Figure 14, along with the measured dimensions of the hole. This dissolution hole was located in Step 2.

B


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Figure 14: (A) MnS dissolution hole. (B) Dimensions of the dissolution hole marked in (A). Dimensions shown are 5.741 µm and 2.86 µm.

3.2.2.2 Heat A16830 Long Rise Time Test

Steps in the long rise time test (5100-second crack tip) were less discernible than in the short rise time test, due to the heavy layer of oxides that was present over the entire fracture surface. The first approximately 3.4mm of the test was performed under R=0.4, and crystallographic faceting is still present beneath the oxide, as shown in Figure 15 (A). During the approximately last 2mm of the test, the stress ratio was changed to 0.7. This portion of the test is shown in Figure 15 (B), and lacks the crystallographic faceting seen in the earlier portion of the test. The CFCGRs showed retardation during both stress ratios.

Figure 15: (A) Example of the fracture surface seen in the first 3.4mm of the test performed at R = 0.4. (B) End of test. Step 9 is denoted by an arrow.


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4. DISCUSSION

The crack growth rates reported show the following trends:

In low sulfur 304SS, the CFCGRs in high temperature DPW environments are environmentally enhanced over CFCGRs in air, and rates increase with increasing rise time. The experimental crack growth rates correspond well to reference curves for equivalent experimental reference curves (N-809). [7]

In high sulfur 304SS, the CFCGRs in high temperature DPW environments at short rise times (5.1 seconds) showed environmental enhancement relative to the CFCGRs in air. Rates at long rise times (51, 510, and 5100 seconds) were unchanged or decreased with increasing rise times, indicative of crack growth retardation.

Figure 16 shows the CFCGRs as a function of rise time for all specimens to further exemplify this point.

Figure 16: Normalized CFCGR as a function of rise time for all materials tested. Rates were normalized to the 5.1s rise time step for each test. Both low sulfur heats showed indications of environmental enhancement, while the high sulfur heat shows moderate crack growth retardation with increasing rise times. Circled data represents the final rise time condition for that test, and are the crack tips available for microcharacterization.

The SEM fractography provides insight into mechanisms that promote enhancement or retardation. The crystallographic-type faceting that appears in the low sulfur material, along with the appearance of secondary cracks, are two features that have been suggested by Mills to be indicative of a hydrogen embrittlement mechanism. [3] The occurrence of fatigue striations with faceting on certain low sulfur materials suggests that a second slip mode is active in this region. Heterogeneous and homogenous slip systems are active during intermediate and higher ΔK loading conditions. [3] The cyclic plastic zone


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ahead of the crack tip experiences heterogeneous slip. In this instance, the planar slip bands are weakened by hydrogen diffusing ahead of the crack tip, generating faceting and river patterns. The striations are created in the smaller fully reversed plastic zone where homogenous slip is promoted ahead of the crack tip under these loading conditions. Furthermore, the appearance of crack growth fronts that diverge from the macroscopic crack growth direction suggest a type of attack where crack resistant regions are attacked at the flanks, as suggested by Mills. During this attack, the resistant regions do not slow the crack growth. Instead, susceptible regions adjacent to the resistant region become paths for cracking, and the crack is allowed to grow around the resistant region. This region is eventually attacked by cracking orthogonal to the primary crack growth direction, as the stresses around the uncracked region increase. [3]

Fatigue surfaces in high sulfur material lack many of the features mentioned above, and additional features promote current proposed theories for crack growth retardation.[3] These theories suggest that the increased corrosion that occurs due to the high sulfur content (generated as the crack front intersects the MnS stringers) is responsible for retardation. Under low stress ratios, the proposed theory for crack growth retardation is the corrosion product buildup on the fracture surface that leads to crack closure, which reduces the effective load the crack tip experiences and therefore reducing the driving force for crack growth. [9] When the mating surfaces are in contact, it is suggested that the corrosion product that has developed on the surface may appear flattened. [3] Figure 17 demonstrates a possible occurrence of corrosion product flattening during comparable R=0.4 conditions. The high sulfur material is shown in Figure 17(A), while the low sulfur material is shown in Figure 17(B). The fracture surface comparison in Figure 17 is similar in appearance to those found by Mills at similar magnification. [3]

Figure 17: Comparison between fracture surface during R=0.4 conditions, with tR>510 seconds. (A) High sulfur material, indicating possible areas of flattened, heavy corrosion product on the surface (yellow arrow). (B) Low sulfur material, with no indications of mating surface contact and a thin, dispersed corrosion product.

One proposed hypothesis for crack growth retardation at high stress ratios is an increase in vacancies being injected into the crack tip, due to the increased corrosion rate. It has been proposed that these vacancies disrupt the planar slip bands ahead of the crack tip that are responsible for promoting crack growth under hydrogen embrittlement, since vacancies are the mechanism by which edge dislocations climb. This disruption in the slip process leads to homogenous slip ahead of the crack and blunting of the crack tip reduces the stresses at the crack tip. [3] The lack of cleavage-like faceting on the fracture surfaces supports this change in crack growth mechanism, as the featureless surface suggests a homogenized slip process. It is also possible that the accelerated dissolution taking place in the high S


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material results in dissolution of the expected faceted features. Work to investigate the dislocation structures under these surface features will be conducted to investigate these possibilities.

The alternating oxide feature in the high sulfur material (specimen A16-32) had three distinct oxide morphologies in the transition region, as shown in Figures 11, 12, and 13. In Step 1 (5.1s tR), the oxide particles were dispersed and larger in size than those in the Step 2 (51-second tR). The oxide layer at the beginning of Step 2 was heavy, homogeneous, and relatively intermediate in size. In the post transition region of Step 2, an alternating area of intermediate/small oxides and fine grained oxides were prevalent. In austenitic stainless steels in high temperature DPW, it has been acknowledged that a two layer oxide film is formed, consisting of an inner layer of finely grained, tightly adherent, and chromium rich oxides, with an outer layer of larger, crystalline, and iron rich oxides. [10] A schematic of this two layer oxide structure is shown in Figure 18. [10] The abrupt change in size and oxide morphology observed between Steps 1 and 2 may correspond with the transition in crack growth rates that were observed, though it is not conclusive whether the crystal structure or composition of the outer layer of oxides changed between the steps, and oxide film analysis must be performed to substantiate this. Initial energy dispersive x-ray spectroscopy (EDX) has been performed on the alternating regions, and preliminary results have suggested that the oxide particles in Step 1 were not rich in Cr or Ni when compared with Step 2, indicating a difference between the oxide layers in the ‘bright’ and ‘dark’ regions. Since the generation volume may have complicated the results, especially in regions where the oxides were thin (such that results may have included both the oxide layer and substrate), further analysis is required to understand the role of this oxide morphology on the crack growth rates.

Figure 18: Illustration of two layered oxide film on austenitic stainless steels. [10]

5. CONCLUSIONS AND FUTURE WORK

5.1 Conclusions

1. Low sulfur materials (Heats D2739 and E5174) show environmental enhancement of CFGRs in high temperature DPW. Experimental data corresponded well to reference data for environmental enhancement in DPW environments.

2. High sulfur material ( Heat A16830) shows environmental retardation of CFCGRs in high temperature DPW. Experimental CFCGRs were either constant or decreasing with rise time. Crack growth retardation was seen during rise times as low as 51 seconds.

3. Fracture surfaces for the low sulfur materials show cleavage-like faceting with a sparse oxide layer, indicating a hydrogen embrittlement mechanism. Environmental enhancement in CFCGRs corresponded with surface features such as river patterns, cleavage-like facets, and a thin oxide layer. Fatigue striations were observed in 2 of the 3 low sulfur specimens; striations were not observed during testing at a ∆K = 6.7 or 6.9.

4. Fracture surfaces for the high sulfur materials demonstrate a featureless surface with a heavy oxide layer. Environmentally retarded CFCGRs were associated with fatigue surfaces that lacked the crystallographic cleavage-like faceting that was seen in enhanced regions. Fatigue surfaces in the high sulfur material had a heavy corrosion product on the surface, due to the dissolution of the MnS stringers.


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5. A transition region from crack enhancement during a short rise time (5.1s) step to crack retardation during a long rise time (51 seconds) step exhibited an abrupt change in oxide morphology. The change in oxide morphology in the high sulfur specimen A16-32 suggests a dependence of the CFCGRs on the formation of oxide layers. The observation of alternating ‘bright’ and ‘dark’ bands of oxides observed in Step 2 is not fully understood, and more analysis on the underlying surface is needed.

REFERENCES

[1] Chopra, O.K. and Shack, W.J. (2007). Effect of LWR Coolant Environments on the Fatigue Life of Reactor Materials - Final Report. NUREG/CR-6909, February 2007.

[2] Tice, D., Platts, N., Rigby, K., Stairmand, J., & Swan, D. (2005). Influence of PWR Primary Coolant Environment on Corrosion Fatigue Crack Growth of Austenitic Stainless Steels. Pressure Vessel and Piping Division Conference, 2005 (pp. 193-205). Denver, CO. ASME.

[3] Mills, W.J. (2013). Accelerated and Retarded Corrosion Fatigue Crack Growth Rates for 304 Stainless Steel in an Elevated Temperature Aqueous Environment. 16th International Conference on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors.

[4] American Society of Mechanical Engineers (2010). Rules for Inservice Inspection of Nuclear Power Plant Components. Boiler and Pressure Vessel Codes, Section XI.

[5] Platts, N., Tice, D., Mottershead, K., McIntyre, L., & Scenini, F. (2012). Effects of Material Composition on Corrosion Fatigue Crack Growth of Austenitic Stainless Steels in High Temperature Water. 15th International Conference on Environmental Degradation of Materials in Nuclear Power Systems-Water Reactors (pp.561-579). Colorado Springs, CO.

[6] Gibbs, J. (2011). Stress Corrosion Cracking and Crack Tip Characterization of Alloy X-750 in Light Water Reactor Environments. PhD thesis, MIT

[7] R. Cipolla. “Case N-809 – Reference Fatigue Crack Growth Rate Curves for Austenitic Stainless Steels in Pressurized Water Environments”, draft Revision 5, 14 May 2014.

[8] Wire, G. L., & Mills, W. J. (2004). Fatigue Crack Propagation Rates for Notched 304 Stainless Steel Specimens in Elevated Temperature Water. Journal of Pressure Vessels, ASME (pp. 318-326)

[9] Lias, P.K. (1988). Overview of Crack Closure at Near-Threshold Fatigue Crack Growth Levels. Mechanics of Fatigue Crack Closure, ASTM STP 982 (pp. 62-92). ASTM, Philadelpia, PA.

[10] O. K. Chopra, “Mechanism and Estimation of Fatigue Crack Initiation in Austenitic Stainless Steels in LWR Environments”, NUREG/CR-6787 ANL-01/25, 2002.

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THE EFFECT OF ENVIRONMENT, CHEMISTRY, AND …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG011.pdfaustenitic stainless steel, and the behavior shown is typical for stainless