Environmentally assisted cracking of solution-annealed Type 316L austenitic stainless steel with two different surface treatments (polished versus ground) was investigated using the constant extension rate tensile test methodology in high-pressure water at 350°C and low-pressure H2-steam vapor at 350°C, 400°C, 440°C, and 480°C, while maintaining electrochemical corrosion potentials in the NiO stability regime. Flat tapered specimens were used to indentify the threshold stress for cracking (∼400 MPa) under these environmental conditions. Intensive oxidation and typical intergranular cracks were observed to initiate from the polished surface at 400°C and 440°C, whereas wavy cracks initiated on the ground surface. The results indicate that the ultrafine-grained layer formed adjacent to the ground surface effectively suppressed intergranular crack initiation under these test conditions. The slow loading under H2-steam vapor at 400°C under the oxidizing condition (NiO) was found to be a suitable high-accelerated test to study early stages of the cracking.

Austenitic stainless steels (SS) are widely used for nuclear power plant components due to their excellent corrosion resistance and good mechanical properties at elevated temperatures1  and their good performance in neutron irradiation environments.2  However, some field cases3-5  have shown that the environmentally assisted cracking (EAC) is an issue for long-term operation because cracks can develop in some unpredictable locations during 20 y to 60 y of operation in pressurized water reactor (PWR) conditions. To understand the long-term damage process, it is important to study EAC by reproducing this degradation phenomenon in the laboratory, and this is only possible by reducing the test time using accelerated environmental conditions.

It is well known that EAC very slowly initiates and then slowly propagates from the surface of SS components during exposure to the high-temperature (HT) water and loaded above the threshold stress, thereby inducing very slow (quasi-static) straining. To accelerate EAC, tests are performed under conditions more severe than those for in-plant operation, i.e., at higher temperature, increased stress and straining, and more aggressive environment. It must be emphasized that cracking behavior under such accelerated conditions must reproduce all of the features of EAC occurring under actual service conditions, especially the crack morphology and the cracking mechanism. During long-term exposure to the PWR cooling water, a thin protective oxide layer builds up on the SS component surface. As a result of local oxidation reaction, the oxide can be locally deeper at some microstructural features, such as grain boundaries. It is widely believed that intergranular cracks, which are typical for EAC that occurs in SS exposed to HT water, initiate at the regions where grain boundaries intersect the surface, with preferential oxidation along the grain boundary.

From an engineering point of view, crack initiation occurs only when the crack can be detected by nondestructive examination (NDE), detection limits of which can be several millimeters and thus sufficiently long to be used in fracture mechanics assessments. On the other hand, with current laboratory techniques it is possible to investigate submicrometer-sized crack formation and propagation. Thus, the EAC initiation can be traced in the early stages, which has been divided into precursor, incubation, and slow growth, all happening in micro- or nanoscale dimensions. This approach had been proven useful in studying effects of industrially-used surface treatments on EAC sensitivity.6-7  The surface treatment effect has repeatedly appeared in attempts to explain the EAC field cases3,8  and to reliably model the degradation process.9  According to current opinion, the early stages of EAC initiation mechanism in the SS and HT water systems include preferential intergranular oxidation, microsized crack formation, and step-like crack growth.9-10 

The long-term EAC degradation can be studied experimentally using accelerated testing methodologies. It has been shown that EAC initiation in Types 304L/316L (UNS S30403/S31603(1)) stainless steel can be accelerated by simulating the PWR primary water environment at 360°C in an autoclave (e.g., Scenini, et al.11 ), and further accelerated by even higher temperatures in steam vapor. The EAC autoclave experiments performed by Economy, et al.,12  showed that the EAC behavior of Alloy 600 (UNS N06600) tested in primary water at 320°C was similar to the behavior of this alloy tested in high-pressure steam at 400°C, thus suggesting that the EAC mechanism was the same for these two environments. Since then, experiments in atmospheric-pressure hydrogenated steam have been used to perform accelerated experiments mostly with Ni-based alloys.10,13-20  This low-pressure system can accelerate oxidation kinetics of Ni-based alloys while maintaining the reducing electrochemical conditions similar to those measured in the PWR primary water environment.13,17  For austenitic SS, the oxidation and EAC experiments conducted in the steam have been limited and their results were inconclusive.3,21 

The aim of this study is to understand the accelerated initiation of EAC cracks in Type 316L austenitic stainless steel. In order to accelerate the crack initiation, EAC tests were performed under mechanical slow straining during environmental exposure in the water at 350°C and in hydrogenated steam at 350°C to 480°C. To further accelerate the cracking, oxidizing conditions were applied in both environments instead of the reduction conditions required in primary cooling systems of PWR plants. To study the effect of surface finish, a special type of flat specimen was used. The flat specimen had a tapered gauge proven to be the most suitable for studying crack initiation conditions, because the tapered gauge allowed the simultaneous examination of a range of applied stresses.22-23  The EAC initiation process resulting from the accelerated testing has been investigated. The observed morphologies of initiated EAC cracks as well as development of oxidation in low-pressure H2-steam have been compared to the previous tests performed in EAC primary water in an autoclave.24-25  The EAC initiation has proven to be very sensitive to the applied accelerating factors (temperature, strain rate, and electrochemical potential), as well as to surface finishing of the flat specimens.

Materials

The Type 316L austenitic stainless steel evaluated in this study was produced by the Industeel of AlcelorMittal Group26  and its chemical composition is shown in Table 1. The material was acquired as a 15 mm thick hot-rolled and heat-treated plate that was solution annealed at 1,050°C to 1,100°C. Although the dominant phase of the material was austenite, pronounced δ-ferrite stringers oriented parallel to rolling direction were present throughout the plate as described elsewhere.25  The material mechanical properties at 350°C were the yield strength (YS) of 168 MPa and the ultimate tensile strength (UTS) of 421 MPa.

Table 1.

Chemical Composition (wt%) for the Type 316L Austenitic Stainless Steel26 

Chemical Composition (wt%) for the Type 316L Austenitic Stainless Steel26
Chemical Composition (wt%) for the Type 316L Austenitic Stainless Steel26

Specimen

Flat tapered tensile specimens were extracted from the 15 mm plate using electrical discharge machining (EDM) with their gauge parallel to the plate rolling direction (i.e., loading direction during material testing was parallel to the rolling direction). Each specimen was 3 mm thick and its gauge width varied from 4 mm to 6.4 mm along the gauge length (Figure 1[a]). Opposite flat surfaces of the specimens were subjected to different surface finishes: one side was manually ground with the 500 grit SiC papers (wet grinding), so that the grinding marks were inclined by an angle about 30° to 45° to specimen loading axis, whereas the other side was metallographically polished using 1 μm diamond paste. As described elsewhere,25  the following characteristics of the specimen surfaces were measured: (i) the roughness (Ra) of the polished and ground surfaces were 0.005 µm and 0.032 μm, respectively; (ii) the HV0.01 microhardness measured on transverse cross sections was 198 and 209 at cross sections 20 μm below the polished and ground surfaces, respectively (Figure 1[b]); and (iii) a compressive residual stress of −300±60 MPa for the polished surface and from −130±80 MPa (between grooves) to −350±80 MPa (inside grooves) for the ground surface.

FIGURE 1.

(a) Drawing of the tapered specimen used in this study. (b) Microhardness profiles measured from the polished and ground flat sides for the cross-section specimens.

FIGURE 1.

(a) Drawing of the tapered specimen used in this study. (b) Microhardness profiles measured from the polished and ground flat sides for the cross-section specimens.

Close modal

Environments and Equipment

The EAC experiments were performed by testing the tapered specimens in the following two environments and corresponding equipment, representing the HT cooling water simulated in an autoclave and the atmospheric pressure hydrogenated-steam vapor (HSV), respectively.

High-Temperature Water

The Sustainable Energy project (SUSEN) water autoclave system was set up by the UJV Rez (Czechia), consisting of mechanical loading system in a stand-alone frame and a water loop with flow rate of 1.5 L/h to 2 L/h. The high-temperature water (HTW) in the autoclave system contained 1,200 ppbw Li+, 2,700 ppmw BO33–, 484 ppmw B3+, and 1.1 ppb O2 to simulate the PWR primary water without a hydrogen gas addition, i.e., it formed the NiO stability conditions in there, which are more oxidizing than the Ni/NiO transition.27  The low oxygen content in the water was maintained by bubbling N2 gas through the water fill tank. The tapered specimens were tested in the autoclave at 350°C and 17 MPa. More information about the experiments is given elsewhere.24-25 

Hydrogenated-Steam Vapor

The test cell maintaining the low-pressure superheated HSV environment was installed on the Zwick/Roell Kappa SS-CF electromechanical creep testing machine. The HSV system was originally developed by Scenini, et al.,11  and subsequently used in several independent studies.13,17,19,28-29  The oxygen partial pressure (pO2), and hence the hydrogen partial pressure, was controlled and set by adjusting the inlet flow rates of H2O steam and H2 gas. The H2O flow rate was controlled with the HPLC (high-performance liquid chromatography) pump Watrex P102, while the H2 was controlled with the flowmeter Brooks Delta Smart II. A gas mixture of 6.5% hydrogen and 93.5% argon was used. By changing these flow rates, it was possible to directly control the electrochemical corrosion potential. For simplicity, a parameter R is used to describe the environment. The R represents the ratio of the oxygen partial pressure, at the Ni/NiO (pO2 Ni/NiO) for a given temperature, and the pO2 in the system. Further details can be found in Volpe, et al.30  The tapered specimens were tested in the HSV environment at 350°C, 400°C, 440°C, and 480°C and at R = 1/6, i.e., in the environment six times more oxidizing than the Ni/NiO transition. Values of parameters used for setting this environment are shown in Table 2, where PPR represent ratio of steam partial pressure over H2 fugacity.

Table 2.

Parameters of HSV Environment

Parameters of HSV Environment
Parameters of HSV Environment

Specimen Loading

The Type 316L flat tapered specimens exposed to the HTW and HSV environments were loaded with a constant extension rate tensile (CERT) testing. The specimens were exposed to the respective environments for 5 d to 7 d prior to loading. The loading with two extension rates of 2 × 10−6 m/s (S1) and 2 × 10−8 m/s (S2) were used; these rates would correspond to the strain rates of 1 × 10−4 s−1 and 1 × 10−6 s−1 if the specimens had uniform width instead of the tapered one. The first specimen in each environment was loaded up to failure and the further experiments were then interrupted after achieving the maximum load.

Post-Test Specimen Characterization

To evaluate the presence and extent of EAC and specimen surface oxidation after the experiments, the specimens were analyzed using scanning electron microscopy (SEM) and energy dispersive x-ray (EDX) spectroscopy. A TESCAN MIRA3 field emission gun (FEG) SEM and a TESCAN LYRA3 Ga+ focused ion beam (FIB)-FEGSEM equipped with an Oxford Instruments Si drift detector (X-Max 80) for EDX microanalysis were used for the detailed characterization of the tested specimens. The surface oxides were analyzed using SEM-EDX spectrum imaging in the FIB-SEM operated at 10 kV and with 3 nA beam current. EDX elemental maps were extracted from the spectrum image datasets and processed using the AZTec software. The live-time for each dataset acquisition was 120 s. These analyses provided qualitative information concerning the oxides formed on the specimens during the tests.

The morphology of the EAC cracks formed during the tests were analyzed using the FIB-SEM, using the conventional FIB specimen preparation process. To prepare each cross section of a chosen crack or grain boundary, the specific crack or grain boundary locations were first identified and then a protective Pt over-layer was deposited to protect the cross sections from damage during the milling process. The FIB milling was performed using a 30 kV Ga+ ion beam with currents of ∼10 nA, ∼1 nA, and ∼0.2 nA for the final polish of the cross section. The microstructure observation of the near-surface layer, grain boundaries, and cracks was performed using the secondary electron (SE) and the backscattered electron (BSE) imaging modes in the FIB-SEM.

Threshold Stress

The minimum stress enabling the EAC initiation is called threshold stress. Owing to the taper shape of the gauge, the EAC threshold stress of the test specimens could be determined from location of the last EAC crack along the gauge length (as shown in Berger, et al.23 ).

To be considered, the EAC crack had to fulfill several criteria: (i) its orientation nearly perpendicular to applied load, (ii) its surface appearance rather straight, (iii) its length longer than 20 μm, and (iv) a cleavage-like or intergranular fracture mode appearance of the crack walls, if visible in its sufficiently wide opening. The last EAC crack in the direction of the wider gauge end was used to calculate the threshold stress as the maximum load achieved in the given specimen test, divided by the cross-section area at this crack location.

Constant Extension Rate Tensile Data

The CERT tests were performed with two extension rates, S1 and S2, in two environments of the HTW and HSV, and at various temperatures from 350°C to 480°C. Three specimens were loaded at the rate S1 to rupture (LW1, L18, and L21) to investigate fracture development, while the remaining tests were stopped at the maximum load so that it was possible to evaluate cracks in absence of excess strain.

The load–displacement curves associated with the CERT experiments are shown in Figure 2. Specifically, Figure 2(a) shows the curves for S1 test rate and Figure 2(b) for S2 test rate. The test curves appearance was slightly affected by both the test rate and temperature. At 350°C, the curves of S1 and S2 rates showed very similar course in both environments, as well as in the air as discussed elsewhere.25  At higher temperatures in the HSV, the maximum load was affected by test rate. At the higher test rate of S1 (Figure 2[a]), the maximum loads of the 350°C, 400°C, and 440°C curves were similar, while the load of the 480°C curve slightly increased. At the lower test rate of S2 (Figure 2[b]), the maximum load of the 350°C curve was lower in comparison to the loads of the 400°C and 440°C curves and higher in comparison to the load of the 480°C curve. Moreover, serrations appeared in all of the S1 and S2 test curves above 3 kN and 2.6 kN, respectively. Presence of the serrations and the test rate effect on the maximum load indicate dynamic strain aging (DSA) in the steel at those temperature and test rate conditions. It should be noted that DSA had been also detected for the same steel batch loaded in the air at 300°C and 350°C using the S2 test rate,25  implying that the DSA appearance at the experiments described here was not induced by the water and HSV environments.

FIGURE 2.

The load–displacement curves of the CERT tests performed on the Type 316L tapered specimens exposed to HTW and HSV at various temperatures and at two extension rates of (a) S1 and (b) S2.

FIGURE 2.

The load–displacement curves of the CERT tests performed on the Type 316L tapered specimens exposed to HTW and HSV at various temperatures and at two extension rates of (a) S1 and (b) S2.

Close modal

The summary of test results is provided in Table 3 and includes (i) the maximum load, (ii) the corresponding maximum stress, calculated using final dimensions at the minimum cross section, and (iii) the threshold stress evaluated for each of the two surfaces.

Table 3.

Summary of CERT Test Results

Summary of CERT Test Results
Summary of CERT Test Results

It was noted that the entire gauge length was plastically deformed after each test. The level of plastic strain could be only estimated considering the tapered shape of the specimen gauge. The maximum plastic strain, which was achieved at the minimum cross section, was estimated as the percentage of the total elongation divided by the gauge length. The maximum strain was ∼15% to 19% in all of the tests. To estimate the plastic strain achieved at the wider end of the gauge, i.e., the minimum plastic strain in the test, it was assumed that the plastic strain decreased linearly along the gauge from the maximum, achieved at the minimal cross section, to zero plastic strain at the location, where the stress equaled to the tested steel YS. Based on that, the minimum strain was estimated to be about ∼10% in all of the tests.

Characterization of the Specimens After Constant Extension Rate Tensile Autoclave Tests in 350°C Water

After the exposure and loading in the water environment at the test rates S1 and S2, tapered specimen surfaces were covered with a thin oxide slice, formed by a continuous inner layer and outer layer composed of large oxide particles (Figure 3). Short cracks appeared on both the polished (Figure 3[a]) and ground (Figure 3[b]) surfaces. The largest number of these cracks were in the area of maximum stress and strains, corresponding to fracture surface of the LW1 specimen at the S1 test rate and minimum cross section of the LW2 specimen at the S2 test rate, and the number of cracks decreased with increasing gauge width of the tapered gauge. Both ductile tearing cracks and the EAC cracks were observed. On the LW1 specimen, a cleavage-like fracture mode was observed on the faces of several open cracks (Figure 3[b]). On the LW2 specimen, intergranular and cleavage-like cracks appeared around its minimum cross section (Figure 4), while intergranular cracks prevailed toward the specimen wider end (Figure 5). Positions of the EAC cracks along the tapered gauge length were documented to find the last EAC crack initiated at the lowest stress, i.e., the EAC threshold stress. Identification of the last EAC crack for the LW1 and LW2 specimens was not always straightforward due to the high number of cracks and large variance in their morphologies that obscured the observation. The corresponding threshold stresses in LW1 and LW2 specimens were determined for both polished and ground surfaces, as shown in Table 3.

FIGURE 3.

HTW at 350°C by S1 (LW1). SEM-SE micrographs of the (a) polished (macro view including fracture and zoomed view of boxed area) and (b) ground surfaces. The images, showing deformation, open EAC cracks, and small surface cracks, were acquired in 3.6 mm and 1 mm to the minimum cross section, where the specimen failed.

FIGURE 3.

HTW at 350°C by S1 (LW1). SEM-SE micrographs of the (a) polished (macro view including fracture and zoomed view of boxed area) and (b) ground surfaces. The images, showing deformation, open EAC cracks, and small surface cracks, were acquired in 3.6 mm and 1 mm to the minimum cross section, where the specimen failed.

Close modal
FIGURE 4.

HTW at 350°C by S2 (LW2). SEM-SE images of the tapered surfaces acquired at the minimum cross section: (a) polished surface showing deformation and EAC cracks; and (b) intergranular EAC open crack on the ground surface.

FIGURE 4.

HTW at 350°C by S2 (LW2). SEM-SE images of the tapered surfaces acquired at the minimum cross section: (a) polished surface showing deformation and EAC cracks; and (b) intergranular EAC open crack on the ground surface.

Close modal
FIGURE 5.

HTW at 350°C by S2 (LW2). SEM-SE images of EAC cracks on tapered surface and FIB longitudinal cross-sections: polished surface (a) IG crack (413 MPa); ground surface (b) IG crack (411 MPa), and (c) mixed crack (395 MPa) showing TG initiation across the UFG layer and the PF region. Red arrow marks FIB location at the crack. Loading direction is horizontal.

FIGURE 5.

HTW at 350°C by S2 (LW2). SEM-SE images of EAC cracks on tapered surface and FIB longitudinal cross-sections: polished surface (a) IG crack (413 MPa); ground surface (b) IG crack (411 MPa), and (c) mixed crack (395 MPa) showing TG initiation across the UFG layer and the PF region. Red arrow marks FIB location at the crack. Loading direction is horizontal.

Close modal

Depth of the EAC cracks had been investigated in the specimen centerline longitudinal cross section using SEM, as shown elsewhere.25  That previous observation had shown that maximum depths were about 12 μm and 5 μm for the polished and the ground surfaces, respectively. Analyses of EAC crack morphology of the LW2 specimen were performed using FIB cross-sectioning and the results are shown in Figure 5. The intergranular character of the cracks was confirmed using the site-specific FIB cuts. On the polished surface, a crack with the length of 25 μm on the surface was observed to be 6.5 μm deep (Figure 5[a]); another shorter crack, which was about 13 μm in length on the surface, was deeper at approximately 8.5 μm. At the same time, a thin zone of deformed microstructure was observed under the surface (Figure 5[a]). Considering previous references,7  this zone was marked as a plastic flow (PF) region. On the ground surface, an intergranular crack was discovered with the length of 38 μm and depth of 13.5 μm (Figure 5[b]), using FIB cross-sectioning. Another crack was also observed with unclear character (Figure 5[c]), yet it was showing transgranular (TG) initiation across the ultrafine-grained (UFG) microstructure layer and the PF region, both appearing under the surface.

On both surfaces of the LW2 specimen, thin continuous darkly-imaging features associated with the surface oxide layer and the subsurface deformed region were discovered, visible on SE-FIB images shown in Figures 6(a) and (b). The oxide layer was between 30 nm and 200 nm thick. Total size of the deformed PF regions under the ground and polished specimen surfaces did not differ (Figure 6) and were estimated to be 680 nm to 1,000 nm. However, the UFG subsurface layer was clearly observed only under the ground surface of about 700 nm (Figure 6[b]). Under the polished surface, the grain size did not change in the subsurface deformed region.

FIGURE 6.

HTW at 350°C by S2 (LW2). SEM-SE images of thin oxide surface layers, and UFG and PF subsurface regions in FIB sections formed on the (a) polished and (b) ground surface, close to the last crack.

FIGURE 6.

HTW at 350°C by S2 (LW2). SEM-SE images of thin oxide surface layers, and UFG and PF subsurface regions in FIB sections formed on the (a) polished and (b) ground surface, close to the last crack.

Close modal

Characterization of the Specimens After Constant Extension Rate Tensile Tests in the Hydrogenated-Steam Vapor

Exposure to Hydrogenated-Steam Vapor at 350°C

After the exposure and loading in the HSV, nonuniformly distributed oxide particles were observed on the surface of both ground and polished specimen surfaces (Figures 7[a] and [b]). The extent of plastic deformation was evident from the extensive slip steps emerging from both surfaces; some ductile cracks were identified in these regions (Figure 7[b]). No EAC cracks with their characteristic fracture mode and orientation perpendicular to the loading were observed for these specimens.

FIGURE 7.

HSV at 350°C by S1 (L18). (a) SEM-SE image showing representative portion of the polished surface at the minimum cross section. (b) Higher-magnification SEM-SE image showing the ductile crack initiation sites along the slip steps for the ground surface.

FIGURE 7.

HSV at 350°C by S1 (L18). (a) SEM-SE image showing representative portion of the polished surface at the minimum cross section. (b) Higher-magnification SEM-SE image showing the ductile crack initiation sites along the slip steps for the ground surface.

Close modal

Exposure to Hydrogenated-Steam Vapor at 400°C

In comparison to the lower temperature, the tested specimen exposed to the HSV at 400°C showed more pronounced oxides. Different oxidation phenomena occurred at the two specimen surfaces. On the polished surface, a thin continuous oxide layer appeared (Figures 8[a] through [c]). In addition, distinctive outer oxides appeared, forming crystal particles arranged into rows, likely along slip steps (Figure 8[c]). Additionally, sporadic straight intergranular cracks were observed, as typical for the EAC (Figures 8[a] and [b]). On the ground surface, a thin continuous oxide layer covered most of the surface and aggregates of oxide particles formed around some grinding grooves and cracks (Figure 8[d]). Multiple clusters of short cracks appeared; whereas these cracks were oriented mostly slightly inclined or perpendicular to the loading direction, while crossing or often being inside the grinding grooves (Figure 8[d]), the last of these cracks was narrow (Figure 8[e]), resembling some cracks seen in the water (Figure 5[a]). No other kinds of cracks were observed on this surface.

FIGURE 8.

HSV at 400°C by S2 (L25). SEM-SE images of the polished surface: (a) open intergranular crack showing step-like features inside (485 MPa), (b) the last crack (412 MPa), and (c) oxidation in the center of the wider gauge end (270 MPa); and the ground surface: (d) oxidation and group of cracks (494 MPa) and (e) the last crack (386 MPa). Loading direction was horizontal; the minimum cross section experiencing the maximum stress was in a distance on the left.

FIGURE 8.

HSV at 400°C by S2 (L25). SEM-SE images of the polished surface: (a) open intergranular crack showing step-like features inside (485 MPa), (b) the last crack (412 MPa), and (c) oxidation in the center of the wider gauge end (270 MPa); and the ground surface: (d) oxidation and group of cracks (494 MPa) and (e) the last crack (386 MPa). Loading direction was horizontal; the minimum cross section experiencing the maximum stress was in a distance on the left.

Close modal

Exposure to Hydrogenated-Steam Vapor at 440°C

At 400°C, different oxidation morphologies were observed for the two surfaces of the two specimens. On the polished surfaces, the outer oxide was arranged into various morphologies (Figure 9[a]). In addition, different element compositions between the grain interior and boundaries were observed in some areas at the qualitative EDX maps of the oxides prepared for one of the specimens (Figures 9[a] and [b]). Yet, as the EDX maps give volume information up to about ∼2 μm in depth, i.e., significantly deeper than the oxide scale thickness (estimated to 400 nm), the EDX data were used only to highlight the difference in oxidation on the two surfaces. These oxide morphologies were not observed on the ground surfaces, where oxide particles formed small islands over the thin continuous oxide layer (Figure 9[d]). Moreover, long deformation bands occurred in the loading direction over the whole specimen gauges; fine pores appeared in the deformation bands and along some grain boundaries in SEM (Figure 9[c]).

FIGURE 9.

HSV at 440°C by S2 (L24). Surface oxidation in the center of the wider gauge end (285 MPa). Polished surface: (a) SEM-BSE image and the corresponding 150 µm × 150 μm EDX elemental maps (the dashed square); (b) SEM-SE detail of the full line square shown in (a) and the 60 µm × 60 μm EDX map; and (c) detail of the deformation band of (b) showing chains of fine pores. Ground surface: (d) SEM-BSE image and the corresponding 150 µm × 150 μm EDX elemental maps.

FIGURE 9.

HSV at 440°C by S2 (L24). Surface oxidation in the center of the wider gauge end (285 MPa). Polished surface: (a) SEM-BSE image and the corresponding 150 µm × 150 μm EDX elemental maps (the dashed square); (b) SEM-SE detail of the full line square shown in (a) and the 60 µm × 60 μm EDX map; and (c) detail of the deformation band of (b) showing chains of fine pores. Ground surface: (d) SEM-BSE image and the corresponding 150 µm × 150 μm EDX elemental maps.

Close modal

The polished specimen surfaces, with their more intensive oxidation, were also found to be more sensitive to the EAC. On the polished surface of the specimen loaded by the S1 test rate, cracks were observed showing mostly intergranular (IG) morphology (Figures 10[a] through [c]), but on the ground surface only transgranular cracks were observed, arranged in typical clusters, and some ductile cracks (Figures 10[d] through [f]).

FIGURE 10.

HSV at 440°C by S1 (L21). SEM-SE images of the cracks; polished surface: (a) IG cracks opened by slip and many slip lines emerging at the minimum cross section (537 MPa), (b) the last EAC crack (430 MPa), and (c) IG crack surrounded by ordered oxide chains (430 MPa). Ground surface: (d) ductile cracks in a cluster of short surface cracks at the minimum cross section (537 MPa), and (e) the last EAC crack (435 MPa) and its detail (f). Loading direction was horizontal; the minimum cross section, experiencing the maximum stress, was in a distance on the left.

FIGURE 10.

HSV at 440°C by S1 (L21). SEM-SE images of the cracks; polished surface: (a) IG cracks opened by slip and many slip lines emerging at the minimum cross section (537 MPa), (b) the last EAC crack (430 MPa), and (c) IG crack surrounded by ordered oxide chains (430 MPa). Ground surface: (d) ductile cracks in a cluster of short surface cracks at the minimum cross section (537 MPa), and (e) the last EAC crack (435 MPa) and its detail (f). Loading direction was horizontal; the minimum cross section, experiencing the maximum stress, was in a distance on the left.

Close modal

On the polished surface of the other specimen loaded by the S2 test rate, there were many short intergranular cracks observed, more open in the highly strained narrow gauge region and less open toward the wider end of the taper gauge length (Figures 11[a] and [b]); no ductile cracks were observed. Next to one of the intergranular cracks, an oxidized grain boundary was identified by FIB cross-sectioning, indicating initiation phase of this crack (Figure 11[a]). Step-like growth marks were discovered on faces of another open IG crack (Figure 11[b]). On the ground surface, many short cracks were observed, often appearing in clusters. The crack morphology was evaluated close to the minimum cross section, where some of the cracks were opened by the strain (Figures 11[c] and [d]). These cracks were wavy in depth as well as on the surface and so were classified as ductile. Across one of the crack clusters, deeper local oxides were found under the surface cracks in the FIB cross section, showing early stages of the EAC initiation (Figure 11[e]).

FIGURE 11.

HSV at 440°C by S2 (L24). SEM-SE images; polished surface: (a) the images of the surface where dotted line marks the FIB position of a noncracked part and the arrow marks the cracked part of a grain boundary, and the FIB cross section shows oxidation around the grain boundary, i.e., crack nucleus (393 MPa); and (b) more developed open intergranular crack (549 MPa). Ground surface: (c, d) wavy short and open ductile cracks (549 MPa), and (e) the image of the FIB cross-section location and two FIB area details showing localized oxidation around surface cracks and the UFG and PF subsurface regions (429 MPa).

FIGURE 11.

HSV at 440°C by S2 (L24). SEM-SE images; polished surface: (a) the images of the surface where dotted line marks the FIB position of a noncracked part and the arrow marks the cracked part of a grain boundary, and the FIB cross section shows oxidation around the grain boundary, i.e., crack nucleus (393 MPa); and (b) more developed open intergranular crack (549 MPa). Ground surface: (c, d) wavy short and open ductile cracks (549 MPa), and (e) the image of the FIB cross-section location and two FIB area details showing localized oxidation around surface cracks and the UFG and PF subsurface regions (429 MPa).

Close modal

Using FIB cross-sectioning, thin darkly-imaging features associated with surface oxide layers and subsurface deformed regions were made visible (Figures 11[a] and [e]). The oxide layer was up to about 300 nm thick and was continuous on the polished surface and noncontinuous on the ground surface. Total size of the deformed PF regions under the ground and polished surfaces appeared to be about 1,500 nm and 700 nm, respectively. Similarly, as in case of the LW2 specimen tested in the HTW (Figures 5 and 6), the UFG subsurface layer of the size about 600 nm was clearly observed under the ground surface only (Figure 11[b]).

Exposure to Hydrogenated-Steam Vapor at 480°C

At HSV at 480°C, discontinuous and preferential oxidation was observed. The oxide layer was clearly thicker than the one obtained at lower temperatures. As well as at 400°C and 440°C, the oxidation found on the two surfaces looked different; the longitudinal deformation bands also appeared.

On the polished surface, only intergranular cracks were observed, irrespective of the applied extension rate (Figures 12[a] and [b] and 13[a] and [b]). On the ground surface, the observed crack morphology was affected by the extension rate. On the specimen loaded with the S1 test rate, only the wavy cracks appeared (Figures 12[d] and [e]), resembling the ductile cracks observed at the lower temperatures. But on the specimen loaded with the S2 test rate, straight cracks appeared that could be classified as EAC cracks according the abovementioned criteria (Figures 13[d] and [e]). Moreover, frequent surface cracks were observed, some of them narrow but longer than 20 μm and some of them opened, occurring in clusters within the oxide islands in high strained regions (Figures 13[d] and [e]). The last crack was only slightly opened, resembling a transgranular crack (Figure 13[f]).

FIGURE 12.

HSV at 480°C by S1 (L22). SEM-SE images; polished surface: (a) IG crack (zoomed box area) and deformation at the minimum cross section (567 MPa), and (b) the last crack (320 MPa). Ground surface: (c, d) cracks and deformation at the minimum cross section (567 MPa), and (e) the last crack (405 MPa). Loading direction was horizontal; the minimum cross section was on the left.

FIGURE 12.

HSV at 480°C by S1 (L22). SEM-SE images; polished surface: (a) IG crack (zoomed box area) and deformation at the minimum cross section (567 MPa), and (b) the last crack (320 MPa). Ground surface: (c, d) cracks and deformation at the minimum cross section (567 MPa), and (e) the last crack (405 MPa). Loading direction was horizontal; the minimum cross section was on the left.

Close modal
FIGURE 13.

HSV at 480°C by S2 (L23). SEM images; polished surface: (a) IG crack interacting with slip at the minimum cross section (490 MPa), (b) the last crack (367 MPa), and (c) SEM-BSE image of oxidation in the center of the wider gauge end. Ground surface: (d) straight cracks opened within thicker oxide island (490 MPa), (e) small open crack surrounded by oxide particles and porous continuous oxide scale (490 MPa), (f) the last crack (397 MPa), and (g) SEM-BSE image of oxidation in the center of the wider gauge end. Loading direction was horizontal; the minimum cross section was on the left.

FIGURE 13.

HSV at 480°C by S2 (L23). SEM images; polished surface: (a) IG crack interacting with slip at the minimum cross section (490 MPa), (b) the last crack (367 MPa), and (c) SEM-BSE image of oxidation in the center of the wider gauge end. Ground surface: (d) straight cracks opened within thicker oxide island (490 MPa), (e) small open crack surrounded by oxide particles and porous continuous oxide scale (490 MPa), (f) the last crack (397 MPa), and (g) SEM-BSE image of oxidation in the center of the wider gauge end. Loading direction was horizontal; the minimum cross section was on the left.

Close modal

In the study, the EAC cracking process was accelerated in comparison to cracking in normal operation conditions of components in the PWR primary circuits, utilizing test rate, temperature, and environment electrochemistry. Effects of these accelerating parameters on the cracking process provoked rather complex results and are therefore discussed in detail in the following sections. In addition, surface preparation effects needed to be also included into the discussion, as the cracking mechanism proved to differ for different surface finishes.

Effect of the Test Rate and Temperature

The load-displacement curves given in Figure 2 showed maximum load changing due to the test rates and temperatures, as well as effects of the DSA. These features seem to be interconnected as the maximum load changes the result from the DSA effect being active.

The mechanism underlying the DSA, also known as the Portevin-LeChatelier effect, is the diffusion of specific atoms to gliding dislocations, which affects the deformation behavior of the material by restricting cross-slip and thereby promoting strain localization.31  Typical serrated flow feature of the DSA can occur only for some intervals of accumulated plastic strain, i.e., only for some ratio of mobile to immobile dislocation densities,32  so this feature is observable only in the specific strain rate vs. temperature window. For the Type 316L steel, the DSA effect used to be observed at high temperatures and slow strain rates. For example, in 17% cold-worked (CW) Type 316L, the DSA was found when applying the low-cycle fatigue loading in the air with strain rates around 10−4 s−1 at the temperature range of 250°C to 550°C.33  The peak DSA activity was observed for higher temperatures at the higher strain rates than for lower temperatures. In the steel, the DSA was explained as an effect of diffusion of interstitial solutes, such as C or N, at 250°C to 400°C, and an effect of substitutional Cr atoms at 400°C to 550°C. In another study with a commercial Type 316NG steel (annealed and 20% CW), the DSA effect was demonstrated at 288°C and 400°C while using strain rate of 10−5 s−1, where the role of diffusing N in the DSA was identified by internal friction technique.34 

As mentioned above, the S1 and S2 extension rates correspond to the strain rates of 1 × 10−4 s−1 and 1 × 10−6 s−1, respectively, if applied to a uniform cross-section specimen instead of tapered specimen. At the strain rates and temperatures applied at the tests, the DSA activity is expected in the steel. As it can be seen in Figure 2(b), which shows the test curves obtained for the S2 test rate, the load increase was the highest at 440°C due to the DSA activity at its peak. Material response to applied displacement, i.e., the load, was increased by internal restrictions to dislocation motion. At 480°C, a strain rate higher than the S2 was needed to reach the DSA activity peak; hence, the maximum load achieved by S2 at 480°C was lower than at 440°C. So, use of the test rate S1, two orders of magnitude higher than the S2, determined that the maximum load was higher at 480°C than at 440°C (Figure 2[a]).

Crack Morphology as a Function of the Test Rate and Temperature

A cleavage-like TG as well as IG crack initiation and propagation was observed on the test specimens. At 350°C, the TG was the prevailing cracking mode at the higher test rate S1, while the IG cracks dominated at the slower test rate S2.

At 400°C and 440°C, the IG crack morphology was very similar to that observed at 350°C, regardless the extension rate, but the cleavage-like open TG cracks disappeared. Instead, narrow short TG cracks occurred in typical clusters. With decreasing test rate, character of the TG cracks changed from straight to slightly wavy. With temperature increased up to 480°C, IG cracks and wavy ductile TG cracks (Figures 12[d] and [e]) were frequently initiated at the S1 test rate. Additionally, at the test rate S2, the IG initiation was very likely assisted by grain boundary creep because shapes of some cracks resembled coalescence of several grain boundary cavities (e.g., Figure 13[e]).

Effect of Electrochemical Corrosion Potential

It should be noted that to stimulate the cracking process, the CERT testing was performed at the environment with NiO stability regime, while the PWR primary water, typically of 25 cc/kg of H2 at 320°C, is maintained at the Ni stability regime environment. If the content of H2 is much lower in the water at 350°C, electrochemical conditions shift to the NiO stability regime and about 20 cc/kg of H2 would be needed to achieve the Ni/NiO transition, as discussed by Attanasio and Morton.27  Using water without H2 addition, the CERT testing was performed in region of the NiO stability. To maintain the same electrochemical conditions in water and HSV, the oxygen partial pressure was used. In the HSV, the required electrochemical conditions were set so that value of the oxygen partial pressure was six times higher than the partial pressure of the Ni/NiO transition.

When the temperature at the HSV increased from 400°C up to 480°C, the required electrochemical conditions were kept but corrosion potential increased. As a result, oxidation of the specimens was intensified. On the polished surfaces, the changes observed in oxide morphologies and indicated by the EDX chemical composition on some grain boundaries were likely related to intensive oxidation. At the same time, the oxidation was also intensified owing to diffusion stimulated by stress and strain. Even if EDX is only illustrative, it indicated Cr, Ni, and Fe enrichment in oxides; this enrichment was the most distinctive in the direction parallel to the loading and at the deformation bands, where even Mo enrichment appeared. In addition, fine pores observed on the grain boundaries and in the longitudinal bands indicated that the creep were likely under way in the specimens at these conditions (Figure 9). The Cr, Ni, and Fe belong to main constituents of double-layer oxides built in the steam/water systems, where the inner layer consists of a chromium-rich spinel and the outer layer consists of magnetite or iron-nickel spinels.21,35  The specific electrochemical conditions set for the experiments stimulated the oxide development. As these changes of the oxide structures were found on the polished surfaces but did not appear on the ground surfaces, it indicates a large influence of the surface finish on the oxidation process. However, discussions of effects of deformation on oxide formation exceeds the scope of this paper.

Effect of Surface Preparation

It had been previously discussed that the surface preparation influenced the EAC crack initiation of the austenitic stainless steels in HTW.6-7,10  In SSRT tests of Type 304L steel in simulated PWR primary water environment, Chang, et al.,37  reported that more uniform oxidation and enhanced EAC initiation resistance was associated with a thick UFG subsurface layer formed in the steel during its surface preparation, in contrast to the EAC performance for a highly-polished nearly deformation-free surface. Based on their EAC test results of a solution-annealed Type 316L steel with three different surface finishes at 288°C water, Abe, et al.,7  showed that the UFG layer and PF region presence should not be detrimental to EAC resistance; the amount of local strain accumulation near the grain boundary just beneath the UFG layer was found to be the most significant for IG EAC initiation.

It should be noted that the importance of an accumulated strain under the UFG layer cannot be verified using the CERT testing. Clearly, the strain history of the material under the UFG layer was overwritten by the applied loading. This effect was likely the reason why in the current experiments it was found only that the UFG layer was important, while nothing of this kind was discovered regarding the PF region presence. In these experiments, a very thin work-hardening layer, i.e., the PF region, developed under the surface of the flat tapered specimens during their EDM. An additional subsurface deformation resulted from the specimen surface polishing or grinding. The FIB cross-sectioning revealed that the UFG layer formation was likely induced by the surface grinding, as it was not observed under the polished surfaces. It proved that the UFG layer played the most important role in the crack initiation at the oxidizing HSV conditions. The observation of the experimental specimens’ surfaces clearly showed that large EAC IG cracks initiated on the polished surface, but not on the ground surface. An about 600 nm to 700 nm thin UFG layer was able to prevent the original microstructure of the material, including its grain boundaries, from a direct contact with the water/vapor environment. Instead, oxides formed homogeneously over the whole UFG layer and thus only ultrafine IG cracks could initiate.

The Threshold Stress as a Function of Temperature and Environment

The type of specimen with the tapered gauge was chosen for the study in order to be able to evaluate the stress and strain conditions of crack initiation. Knowing that, the threshold stress needed for the EAC initiation could be determined.

It should be noted that no single type of cracks was observed at the specimens. Hence, the two following crack types have been selected as a base for evaluation of the threshold stresses in the specimens (Table 3): (i) the EAC crack defined by the criteria mentioned in the Threshold Stress section, and (ii) the NC cracks observed on the ground surface at 440°C and 480°C. Polished and ground surfaces of each specimen were observed by SEM to find the last crack location, which was then used to determine the threshold stress. In the graph in Figure 14, the threshold stress data are plotted versus temperature together with the maximum stresses of the tests; in addition, the graph includes the YS, UTS, and flow stresses, which were acquired from standard tensile tests performed with the same material in the air and while using higher strain rate (1.6 × 10−4 s−1).36  It shows that the maximum stresses achieved in the CERT testing are always higher than the UTS, due to slower test rates inducing the DSA effect. Therefore, there is no straightforward comparison of mechanical behavior at different strain rates inside and outside the DSA effect area.

FIGURE 14.

Meaningful stresses gained from the tests, including the threshold stress of crack initiation on the ground and polished specimen surfaces, in relation to temperature. The YS and UTS data are taken from Van den Bosch.36 

FIGURE 14.

Meaningful stresses gained from the tests, including the threshold stress of crack initiation on the ground and polished specimen surfaces, in relation to temperature. The YS and UTS data are taken from Van den Bosch.36 

Close modal

As shown in Figure 14, the level of the threshold stress seems to be quite high in comparison to yielding, and on average not temperature dependent. The EAC threshold stresses for both surfaces exposed to the water at 350°C were determined to be about 400 MPa, while using the last IG and TG EAC crack location. At the HSV, this threshold stress level of 400 MPa can be also seen at 400°C for both surfaces of the single specimen tested by the S2 test rate; about the same stress levels have been again seen at 440°C and 480°C on the specimen ground surfaces, even though they were evaluated using different crack morphology, the wavy ductile or NC cracks. At the same time, the threshold stress values seem to be subjected to a scatter of about ±30 MPa, increasing with temperature. This scatter can be caused by several reasons: (i) determination of the last crack using only the surface observation by SEM; (ii) last crack size differences; and (iii) uncertainty of finding a narrow crack of 20 μm in length on an oxidized surface, which can be overlooked easily, especially when cracks density is very low. Likely, this scatter makes it impossible to assess an effect of the surface preparation.

Despite all of this, the threshold stress value of 400±30 MPa is reasonable, considering very rare occurrence of the EAC in PWR primary cooling systems. This stress value, and the plastic strain ∼10% associated with this threshold, is much higher than the allowed operation stresses and strains in plant designs. It implies that the EAC initiation in the field conditions has to be governed by local stresses and strains on the microstructure level.

Similarity of Environmentally Assisted Cracking Initiation in the Field and High Accelerated Conditions

Finding a way to set similar electrochemical conditions in water and vapor environments opens a promising way of how to accelerate the EAC initiation process, which is normally occurring over a long period of time in the plant cooling circuits at temperatures around 300°C. At the same time, the principle of maintaining the unchanging cracking mechanism during its acceleration in laboratory has to be always observed. Of all of the characteristics, which need to be retained to achieve adequate resemblance of the cracking observed in the field conditions and the accelerated EAC initiation process, the grain boundary preferential oxidation, microsized crack formation, and step-like growth were found to be the best traceable.

As discussed above in the Crack Morphology as a Function of the Test Rate and Temperature section, all of the crack morphologies observed in these experiments, especially the morphology of the EAC IG cracks, as they appeared on the polished specimen surfaces in conditions of accelerated testing at all temperatures up to 440°C, resembled the crack morphology known from field conditions.3  However, in the accelerated testing at 480°C, the IG cracks initiation was likely assisted by grain boundary creeping and the resulting morphology was thus very different from the field conditions.

The first stages of the EAC initiation mechanism consist of the formation and rupture of surface oxides and/or metal/oxide interface, as discussed elsewhere, e.g., for the Type 304 SS.6  The intergranular cracking precedes a preferential oxidation of grain boundaries.15-17,19  In this study, the first step of the EAC development was observed on the surface of the specimen tested in the HSV at 440°C. In Figure 10(a), the FIB cross section reveals deep oxidation of a grain boundary, which will likely develop to an IG crack. This future development to the crack can be seen in the other part of this grain boundary, which has already cracked.

Microcracks arising after the surface oxides rupture are developing further in nano- or microsteps,38-41  until a mature crack of a critical size is able to propagate steadily. These steps can be found on the faces of surface cracks if they are opened enough to allow SEM imaging (Figures 4[b] and 8[a]). The fine steps can be also found on the FIB cross sections, e.g., of the specimens exposed to loading in the water (Figures 5[a] and [b]). In Figure 5(a), the IG crack faces close to surface are covered by darker oxides into a depth of 1 μm to 2 μm and the surface of this crack shows fine steps, each of them less than 1 μm in size, decorated by single oxide crystals. Each of these steps likely presents one point in the EAC step-like process.

It can be concluded that the discussion given above documents that the EAC mechanisms developed within the test conditions—at least the IG crack process—provide all the characteristics typical for the EAC development in the field.

The EAC response of the Type 316L austenitic stainless steel was investigated in the environment of high-temperature water at 350°C and H2-steam vapor at 350°C, 400°C, 440°C, and 480°C. There were experiments performed on flat tapered specimens with two different surface finishing, while maintaining electrochemical corrosion potentials in the NiO stability regime and utilizing constant extension rate tensile methodology. From this study, the following conclusions can be drawn:

  • In the autoclave water environment at 350°C using the test rate of 2 × 10−8 m/s, typical EAC intergranular initiation and propagation was successfully reproduced on the ground and polished specimen surfaces. The EAC threshold stress was evaluated to be 400±30 MPa.

  • In the H2-steam environment at 350°C, tests did not induce EAC initiation.

  • In the H2-steam environment at 400°C and 440°C, oxidation increase and typical EAC intergranular cracks were observed, initiating from the polished surface, whereas wavy EAC mixed ductile cracks initiated on the ground surface. The results indicate that the ultrafine-grained layer formed adjacent to the ground surface effectively suppressed intergranular crack initiation under these test conditions.

  • In the H2-steam environment at 480°C, observed oxidation and cracking morphologies exhibited high-temperature features, not corresponding to the EAC mechanism at PWR operation conditions.

  • The high accelerated testing suitable for studying early stages of the EAC, for use in PWR long-term operation, was found to be testing in the H2-steam environment at 400°C and in the oxidized condition, using the strain rate of 1 × 10−6 s−1 or less.

(1)

UNS numbers are listed in Metals & Alloys in the Unified Numbering System, published by the Society of Automotive Engineers (SAE International) and cosponsored by ASTM International.

Trade name.

This work was supported by the EC Horizon 2020 MEACTOS GA No. 755151. This work has been realized within the SUSEN Project (established in the framework of the European Regional Development Fund (ERDF) in project CZ.1.05/2.1.00/03.0108 and of the European Strategy Forum on Research Infrastructures (ESFRI) in the project CZ.02.1.01/0.0/0.0/15_008/0000293, which is financially supported by the Ministry of Education, Youth and Sports—project LM2015093 Infrastructure SUSEN. Helpful discussions with Professor M. Grace Burke, Dr. Fabio Scenini, and Dr. Liberato Volpe from the Materials Performance Centre of the University of Manchester are also acknowledged.

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