Low cycle fatigue behavior of alloy 690 in simulated PWR water—Effects of dynamic strain aging and hydrogen

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Low cycle fatigue behavior of alloy 690 in a simulated PWR water-Effects of dynamic strain aging and Hydrogen Jong-Dae Hong, Junho Lee, Changheui Jang, Tae Soon Kim

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Received date: 20 March 2014 Revised date: 25 May 2014 Accepted date: 26 May 2014 Cite this article as: Jong-Dae Hong, Junho Lee, Changheui Jang, Tae Soon Kim, Low cycle fatigue behavior of alloy 690 in a simulated PWR water-Effects of dynamic strain aging and Hydrogen, Materials Science & Engineering A, http://dx. doi.org/10.1016/j.msea.2014.05.069 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

  Low Cycle Fatigue Behavior of Alloy 690 in a Simulated PWR Water Effects of Dynamic Strain Aging and Hydrogen Jong-Dae Hong a, Junho Lee a, Changheui Janga* and Tae Soon Kim b a

Department of Nuclear and Quantum Engineering, Korea Advanced Institute of Science and Technology, 291 Daehak-ro, Yuseong-gu, Daejeon, 305-701, Republic of Korea b

Central Research Institute, Korea Hydro and Nuclear Power Co., Ltd., 1312 Gil 70 Yuseongdaero, Yuseong-gu, Daejeon, 305-343, Republic of Korea

Abstract The low cycle fatigue (LCF) tests of Alloy 690 were performed in air and simulated PWR water under various loading conditions, and the results were analyzed to understand the LCF behavior of Alloy 690 in PWR water. The LCF life of Alloy 690 was shorter in PWR water than in air; however, the degree of reduction was less significant than for austenitic SSs. Through the tests and subsequent analysis, it was observed that dynamic strain aging (DSA) and hydrogen evolution by metal dissolution near the crack tip affect the fatigue crack growth behavior in PWR water. While DSA was observed, its effect on LCF life was not evident. A hydrogen-induced dislocation structure was observed near the crack surface in TEM micrographs, indicating the operation of hydrogen-induced cracking (HIC) in PWR water. However, due to improved corrosion resistance, HIC was less operative for Alloy 690 than for austenitic SSs during the LCF test in PWR water. Consequently, the LCF life of Alloy 690 was longer than that of austenitic SSs in PWR water.

Keywords: Low Cycle Fatigue; PWR Water; Alloy 690; Hydrogen; Dynamic Strain Aging

*

Corresponding author. Tel.: +82 42 350 3824; fax: +82 42 350 3810. E-mail address: [email protected] (C. Jang).

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1. Introduction Austenitic alloys, such as Ni-Cr-Fe alloys (Alloy 600/690 and welds) and stainless steels (SSs), have been widely used in primary piping systems in nuclear power plants (NPPs) due to their good corrosion resistance and mechanical properties. Because primary system pipes and components exposed to high-temperature water are subjected to cyclic stress caused by various transients during operation, corrosion fatigue (also called “environmental fatigue”) is considered one of the key degradation mechanisms. In this context, researchers in the U.S. [1-3] and Japan [4,5] have produced extensive low cycle fatigue (LCF) life test data and developed statistical models to predict the LCF life of metallic materials in primary coolant environments, such as carbon and low-alloy steels, stainless steels (SSs), and Ni-Cr-Fe alloys. Based on these data, U.S. NRC issued Reg. Guide-1.207 [6] requests consideration of the detrimental effects of primary coolant environments in the design and construction of the Class I components of new NPPs. It has been reported that the hydrogen that evolves by the cathodic reaction balancing the anodic metal dissolution at the crack tip affects the environmentally assisted cracking (EAC) observed for austenitic alloys in a high-temperature water environment [7-10]. For austenitic SSs, the environmentally assisted reduction of LCF lives in high-temperature water is most likely caused by hydrogen–induced cracking (HIC) [9,10,11] despite the resistance of these materials to hydrogen embrittlement (HE) due to their high activation enthalpy for hydrogen migration [12]. Evidence of HIC, such as the micro-voids ahead of the crack tip, cleavage facets on the crack surface, and dislocation multiplication was reported for the austenitic SSs, which demonstrate the significant reduction in the LCF life in PWR environments [7-10]. For Ni-Cr-Fe alloys, the effects of key loadings and environmental parameters on the LCF life would be similar to those of austenitic SSs [2-5,13] due to the similarity in the microstructure and corrosion resistance. However, the effects were less significant for 2

Ni-Cr-Fe alloys, and the cracking mechanism is not well understood, in part due to the limited amount of relevant tests and analyses. In particular, for Alloy 690, which replaced Alloy 600 in advanced reactors including APR1400 to mitigate the problem of primary water stress-corrosion cracking (PWSCC), limited tests and analysis results are available, of which most of them only consist of the LCF life data [2-5]. Additionally, dynamic strain aging (DSA) has been considered to have a significant effect on the LCF life by affecting the cyclic stress response behavior and dislocation structure development [9,10,14-16]. The DSA behavior in a high-temperature water environment would be different from that in an air environment as the behavior is affected by the absorbed hydrogen produced by the corrosion reaction in high-temperature water [9,10,17]. In this regard, the LCF test for Alloy 690 was performed in air and a simulated PWR environment, and the fatigue behavior was investigated to understand the mechanisms of the LCF life reduction in a simulated PWR environment. The cyclic stress response, dislocation structure, and crack morphologies of the fatigue-tested specimens were analyzed and compared with those of austenitic SSs. Finally, the effects of DSA and hydrogen on the LCF behaviors of Alloy 690 in a simulated PWR environment were discussed.

2. Experimental 2.1. Test Material The material used in this study was Alloy 690, heat number 135264 supplied by Goodman Alloys® in the form of forged rod that was solution-annealed at 1060°C for 3 hours followed by quenching in air. The chemical compositions of the test material analyzed by inductively coupled plasma (ICP) method are listed in Table 1, and the tensile properties are summarized in Table 2. The chemical composition and tensile properties were within the specifications of ASME SB-166 [18]. The microstructure is shown in Fig. 1. The average 3

grain size is approximately 50 m, and the grain boundaries are covered with chromium-rich M23C6 carbides. In addition, there are a few carbides and Ti (C,N) inclusions in the matrix, which is typical for Alloy 690 [19]. < Table 1 > < Table 2 > < Figure 1 >

2.2. Test Conditions and Methods LCF tests were performed in accordance with ASTM E-606 [20] using round-bar type specimens with a gauge length of 19.05 mm and a gauge diameter of 9.63 mm. During the LCF test, the strain in the gauge section was indirectly measured using a linear variable differential transformer (LVDT) attached at the shoulders of the specimen. The displacement between the shoulders was converted to the strain in the gauge section using the correlation, which was pre-determined at test temperatures as previously reported [11]. The test set-up for the LCF test in a simulated PWR environment consists of a servo-electric fatigue testing machine, an autoclave, and the recirculation-type water circulation loop. The details of the test set-up are described elsewhere [11]. The details of the LCF test conditions are summarized in Table 3. The LCF tests were performed in a strain-controlled mode with a fully reversed (R = -1) triangular waveform at the various strain rates of 0.4%/sec, 0.04%/sec, 0.033%/sec and 0.004%/sec. The applied strain amplitudes varied from 0.33% to 1.0%. The LCF life, N25, is defined as the number of cycles for the tensile stress to decrease by 25% from the peak value, as used in previous studies [1,4,9]. The test environments were room temperature (RT) air, 310 °C air, and a simulated PWR water (at 310°C, 15.5 MPa and hereinafter denoted as “PWR water”) containing typical concentrations of dissolved boric acid and lithium hydroxide. The levels of DO (dissolved oxygen), DH (dissolved hydrogen), conductivity and pH were monitored and controlled at room 4

temperature. After the DO concentration was reduced below 5 ppb by Ar gas bubbling, the DH concentration in the feedwater was maintained at 2.2 ppm (corresponding to 25 cc-H2/kg-H2O) to simulate the primary water condition of a PWR. < Fig. 2 > < Table 3 > After the LCF tests were completed, the specimens were sectioned, and the fracture surface and sectioned area were examined using a field-emission scanning electron microscope (FE-SEM). To investigate the evolution of the dislocation structures during the crack growth, the specimens for the transmission electron microscope (TEM) analysis were obtained at a location approximately 2/3 of the crack length (~ 2.5 mm) and within 0.1 mm from the fatigue crack surface. Discs with 3.0-mm diameters were machined and mechanically polished to 150 Pm in thickness. Complete thinning of the disc to electron transparency was achieved by jet electro-polishing at -35 ºC and 15 V in a solution of 20 % perchloric acid and 80 % methanol. In addition, a 300-keV field-emission transmission electron microscope (FE-TEM) was used for the dislocation structure analysis.

3. Results and Discussion 3.1. LCF Life Figs. 2 and 3 present the LCF life test results of Alloy 690 in RT air, 310°C air, and PWR water. For comparison, the LCF life trend curve in air for Ni-Cr-Fe alloys of the U.S. [2] and Japanese program [4] are also plotted. The LCF life data of Alloy 690 in RT air are in good agreement with the existing trend curve for Ni-Cr-Fe alloys in air, which mainly consists of data of Alloy 600 and its weld metal and the ANL prediction curve for austenitic SSs [2]. According to the previous data [2-5], no distinct differences would be observed in the LCF life between Alloy 600/weld metal and Alloy 690/weld metal in air. In addition, these 5

existing data imply that there is no temperature effect on the LCF life. In this study, two tests were performed at 310°C air, and the results were more or less consistent with the behavior observed in Fig. 2. < Fig. 2 > < Fig. 3 > As observed in Fig. 3, the LCF life in PWR water is shorter than that in RT air and 310 °C air. While there is some scatter in the test data, as often observed in LCF tests, the LCF life in PWR water decreased with a decrease in strain rate. Accordingly, it is thought that the LCF life of Alloy 690 may be affected by EAC mechanisms in PWR water. However, the extent of the effects of the environment was known to be considerably less significant than austenitic SSs under the same environment and loading condition [3]. In this study, the maximum reduction of the LCF life at a strain rate of 0.004%/sec was approximately a factor of 3, as observed in Fig. 3 (the value would be 6.8 for austenitic SSs according to ref. 3). A possible explanation for the longer LCF life of Alloy 690 compared with austenitic SSs in PWR water will be described in the following section.

3.2. Cyclic Stress Response The cyclic stress response represents the locus of the tensile peak stress of the hysteresis loop as a function of the number of fatigue cycles, which is useful in understanding the hardening or softening behaviors during the LCF test. The cyclic stress responses of Alloy 690 in the conditions of RT air, 310°C air, and PWR water are shown in Fig. 4. In RT air, Alloy 690 exhibited a rapid hardening behavior during the first ~70 cycles followed by a continuous softening behavior to the end of the LCF life, while rather continuous softening was observed for austenitic SSs at similar strain amplitudes [2,10]. The relatively higher value of the UTS/YS of Alloy 690 than austenitic SSs could have caused such different behaviors, as hardening was more pronounced when the UTS/YS ratio was higher [21,22]. Meanwhile, 6

in 310 °C air, Alloy 690 exhibited a continuous hardening (also called secondary hardening) behavior until 90% of the LCF life. In PWR water, secondary hardening was also observed after the rather constant peak stress range, as shown in Fig. 4. For Alloy 690, secondary hardening behavior under cyclic loading was also reported for the specimen tested in 204°C air at low strain amplitude (0.2%) [14]. Chai et al. proposed that various factors such as dislocation multiplication, the interaction between moving dislocations and stacking faults, the dynamic strain aging effect and the formation of nano-twins contributed to the secondary hardening for Alloy 690 under cyclic loading [15]. For austenitic SSs, the regimes of softening or saturation were observed after primary hardening at high temperatures of approximately 300°C, while secondary hardening behavior was observed under low strain amplitude [1,2,10]. These softening behaviors at high temperature could be attributed to the dynamic recovery (or mutual annihilation) of dislocations enabled by cross-slip. The cellularized dislocation structure under cyclic loading was observed in some of studies of austenitic SSs [10,16]. < Fig. 4 >

3.3. Manifestation of DSA for Alloy 690 In several studies, the cyclic hardening behavior of Alloy 690 was induced by DSA. According to DSA theory, the interaction between dislocations and solute atoms limits the freedom of motion of the dislocations; i.e., by pinning portions of the dislocations, the dislocations’ ability to cross-slip is restricted [16]. In the manifestation of DSA, negative strain-rate sensitivity and serrated flows were often simultaneously observed. For Alloy 690, negative strain rate sensitivity was observed with the extent of hardening increasing with decreasing strain rate, as observed in Fig. 5. The serrated flows in the hysteresis curves were also observed in the tested loading condition in 310 °C air and PWR water. DSA in Alloy 690 was often observed in the temperature range of 200-600 °C over a wide range of strain rates 7

[16,23-25]. It was considered that interstitial carbon acted mainly as solute atoms, such that the diffused solute carbon atoms interacted with the mobile dislocations during deformation. The role of carbon was confirmed as the activation energy of DSA for Alloy 600 and 690 (~160 kJ/mol) was in good agreement with the activation energy of carbon diffusion (~170 kJ/mol) in both materials [16,24,25]. Meanwhile, cyclic hardening could be enhanced with an occurrence of DSA, and the ratio of the cyclic hardening could be used to assess the degree of DSA [26,27]. Fig. 6 shows the ratio of cyclic hardening (Vpeak/Vinitial) measured by the ratio of the maximum stress amplitude (Vpeak) to the stress amplitude at the first cycle (Vinitial) for austenitic alloys LCF tested for similar strain amplitudes and strain rates. As the LCF test environment was changed from RT air, 310 °C air and PWR water, the ratio of cyclic hardening increased for both Alloy 690 and 316LN SS. Compared with 316LN SS, the cyclic hardening was much more pronounced for Alloy 690, including in PWR water. DSA could have positive or negative effects on the LCF behaviors by inducing a local concentration of stress at the crack tip and then increasing the dislocation density due to strain localization at the crack tip [9,10,14,15]. Consequently, DSA would accelerate crack initiation/growth and lead to the reduction of LCF life [26]. However, at a high temperature within the regime of DSA, the formation of micro-twins would increase the plasticity and contribute to the localized deformation, which in turn increased the interaction between the moving dislocations and stacking faults, which caused plastic deformation to more difficult and resulted in a longer LCF life [15]. Due to these conflicting effects, the existing data for the LCF life of Alloy 690 show little or no effect of temperature on their LCF lives from RT to the operating temperature of PWRs [2-5]. Therefore, DSA alone would not markedly affect the LCF life of Alloy 690, as illustrated in our tests in Fig. 2. < Fig. 5 > 8

3.4. EAC Mechanism - Effect of Hydrogen 3.4.1. Characterization of Fracture Surface Fig. 7 shows the fracture surfaces of the fatigue cracks tested in RT air, 310°C air, and in PWR water at strain amplitudes of 0.4-0.6 % with strain rates of 0.033- 0.04 %/sec. To examine the crack surface in the propagation stage, specimens were taken at positions two-thirds of the length from the crack initiation point to the crack tip (2-2.5 mm from the surface). As observed in the figure, similar fracture modes were observed such that the crack grew in an essentially transgranular (TG) mode with no distinct evidence of intergranular (IG) cracking in all of the conditions, which is typical for fatigue cracking. In addition, well-developed striations normal to the crack growth direction during the propagation stage were observed on the fracture surfaces of the specimens tested in all conditions. Even for the specimen tested in PWR water, oxides were hardly visible on the fatigue crack surface. These fracture surface observations would indicate that mechanical factors, not a slip dissolution/oxidation process, may have played an important role on the LCF cracking of Alloy 690 in all of the test conditions. However, despite the similarity, the spacing between striations was greater for the specimen tested in PWR water (average spacing ~ 8.9 m) than those in air (average spacing ~ 5.6 m for RT air and 5.3 m for 310 °C air), which was consistent with the LCF life results observed in Figs 2 and 3. These characteristics on the fracture surface for the specimen tested in PWR water were similar to those observed for 316LN SS, such as the well-developed striations, increased spacing between striations, and TG cracking [2,9,10]. The presence of well-developed striations on the fractured surface could imply that hydrogen-induced cracking (HIC) was operating, as previously reported for austenitic SSs [1,2]. Meanwhile, it was difficult to find additional evidence of HIC, such as micro-voids on the crack tip, cleavage facets or inclusions on the fractured surface of Alloy 690 tested in PWR water. It 9

was thought that the HIC mechanism, which was considered to be a main contributing factor to the reduction of the LCF life of 316LN SS [9,10], could be less active in Alloy 690 tested in PWR water. < Fig. 7 >

3.4.2. Sectioned Area Observation To investigate the effect of PWR environment on fatigue crack growth and the underlying accelerating mechanisms, the sectioned areas were observed for the etched specimen. Fig. 8 shows the sectioned area of the fatigue cracks tested in RT air, 310°C air, and in PWR water at strain amplitudes of 0.33-0.4 % with strain rates of 0.033- 0.04 %/sec. Similar to the fracture surface observation, it is clear that the cracking mode in RT air, 310°C air and PWR water was transgranular. In addition, the fatigue cracks were also fairly straight for all of the test conditions. < Fig. 8 > Some aspects of crack morphology differ between the specimen tested in air and water environments as shown in Fig. 8. For the Alloy 690 specimen tested in PWR water, the crack tip was quite blunt and the crack opening was wider compared with the specimens tested in RT air and 310°C air. Crack tip blunting could be the evidence of the occurrence of metal dissolution during the LCF test, as suggested for 316LN SS [9,10,11]. Therefore, the metal dissolution could be operating during the low cycle fatigue crack growth in PWR water. However, the effect of metal dissolution on LCF in PWR water is not clear as the metal dissolution would promote crack blunting (Fig. 8), which would lower the stress concentration at the crack tip. Meanwhile, evidence of HIC, such as secondary cracks and microvoids, was hardly observed on the crack surface. As shown in Fig. 8, only a few secondary cracks could be observed for the specimens tested in PWR water, while tens of secondary cracks were developed on the fatigue fracture surfaces of austenitic SSs [28]. 10

Therefore, consistent with the crack surface observation described in the previous section, the sectioned area observation also failed to provide evidence of the HIC mechanism during fatigue crack growth in PWR water.

3.4.3. Dislocation structure To investigate the environmental effects on the movement of a bulk dislocation structure near the crack during the propagation stage, TEM observations were performed. Fig. 9 shows the dislocation substructures near the crack surface in the specimen tested at a strain amplitude of 0.33% in RT air and at a strain amplitude of 0.4% at 310°C (Air and PWR water). Figs. 9 (a) and (b) show the dislocations structure of the specimen tested in RT air. The well-developed cell structures as well as the planar arrays of dislocations were observed. As described in the previous section, this cellularization results from the mutual annihilation of dislocations and the reorganization of remaining dislocations into a lower energy structure as cyclic stresses were applied [16]. However, as observed in Figs. 9 (c) and (d), the dislocation structure of the specimen tested at 310°C revealed a tendency toward dislocation planarization, with predominantly linear dislocation structures with partly tangled structures. In the specimen tested in PWR water (Fig. 9 (d)), pinning of dislocations by solute atoms was extensively observed. As the TEM specimen was obtained during the secondary hardening regime, the planar arrays of dislocation and dislocation pinning could be induced by DSA. Similarly, Hänninen et al. reported that DSA restricts dislocation cross-slip and climb, which favors planar slip, which again increases the flow stress [25]. Thus, the average dislocation density for the specimen tested at 310°C both in air and PWR water would be much higher than that tested in RT air. To quantitatively verify this tendency, the dislocation density was measured on the TEM micrographs using a line-intersection method suggested by Roy et al. [29], and the results are summarized in Table 4. As shown in the table, the increase in dislocation density is in good agreement with the increase in DSA observed in Fig. 6. 11

3.4.4 Effect of Hydrogen Produced by Metal Dissolution In a corrosive environment, hydrogen is evolved by the cathodic reaction of the local corrosion cell. The amount of hydrogen evolution is proportional to the amount of electron produced by metal dissolution (anodic reaction, M-> Mn+ + ne-). The hydrogen evolution occurs in the region near the crack tip, where the bare metal is exposed to water due to oxide rupture. The evolved hydrogen would be absorbed into the metal during cyclic loading and concentrated at strong trapping sites such as stress-concentrated localized regions [9,10,30], which would affect the fatigue crack growth behaviors in a PWR environment. Previous studies reported that the absorbed hydrogen may serve to reduce the friction and repulsive force between the mobile dislocations and the other dislocations [30], thereby causing an increase of the dislocation density. Likewise, the increased dislocation density of Alloy 690 tested in PWR water by hydrogen-induced strain localization was observed, as illustrated in Fig. 8. In addition, planar arrays of dislocations on the TEM observations essentially resulted from the absorbed hydrogen as well as from DSA. The absorbed hydrogen leads to a linear dislocation structure by promoting planar slip and inhibiting cross-slip [20,31]. Meanwhile, it is reported that the hydrogen absorbed into the metal may induce weakening of atomic bonds as well as enhancing the dislocation mobility by reducing the repulsive force between the dislocations [30]. In short, the dislocation structure tested in PWR water contains planar arrays with a higher density of dislocations mainly induced by the absorbed hydrogen, which was produced by metal dissolution at the crack tip. Therefore, the acceleration of fatigue crack growth and the reduction of the LCF life in PWR water are attributed to the HIC mechanism, similar to austenitic SSs. For austenitic SSs, the reduction in the LCF life in high-temperature water is primarily due to the HIC mechanism [1,3,9,10]. Some evidence for HIC was observed in the fatigue crack tested in high-temperature deoxygenated water such as the micro-voids ahead of the main crack and their tendency to 12

coalesce into the main crack [9]. The absorbed hydrogen was also reported to reduce the dislocation spacing [9,10]. However, the hydrogen evolution rates of Ni-Cr-Fe alloys would be less than those of austenitic SSs because nickel is nobler than iron and chromium. For further quantification, it was reported that in PWR water, the corrosion current density (icorr) of Ni-Cr-Fe alloys (icorr of Alloy 690 ~ 5 x 10-6 A/cm2, Alloy 600 ~ 5.5 x 10-6 A/cm2) [32,33] are two to three times smaller than that of austenitic SSs (icorr of Alloy 316 ~ 1.318 x 10-5 A/cm2) [34]. In addition, under similar loading conditions in PWR water, as observed in Fig. 10, Alloy 690 exhibited a relatively narrower fatigue crack tip than that of 316LN SS, which confirmed the reduced metal dissolution in Alloy 690. Therefore, the contribution of HIC to the reduction of LCF life was not significant for Alloy 690 compared with austenitic SSs due to the reduced hydrogen production from metal dissolution. Consequently, a possible explanation for the longer LCF life of Ni-Cr-Fe alloys including Alloy 690 than austenitic SSs could involve reduced hydrogen evolution caused by the improved corrosion resistance of Ni-Cr-Fe alloys. Table 5 summarizes the comparative results of these behaviors between Ni-Cr-Fe alloys and austenitic SSs in PWR water. < Fig. 10 >

4. Conclusions Low cycle fatigue tests for Alloy 690 were performed in air and PWR water, and the crack-accelerating mechanisms in PWR water were investigated. The LCF life of Alloy 690 was shorter in PWR water than in air; however, the degree of reduction was less significant than that in austenitic SSs. DSA was observed in all tests to various degrees, and evidence of metal dissolution was observed near the crack tip in PWR water. However, DSA would not have affected the LCF life of Alloy 690 considering the conflicting effects on the LCF life, such as the enhanced planar slip, which would slow the crack growth rate, and the strain 13

localization near the crack tip, which would accelerate the crack growth. Meanwhile, despite the lack of evidence of the HIC mechanism on the crack surface and crack morphology, a hydrogen-induced dislocation structure, such as a high density of planar arrays of dislocations, was observed near the crack surface in TEM micrographs, which indicated the operation of HIC in PWR water. However, as less hydrogen was produced at the crack tip due to the improved corrosion resistance of Alloy 690, HIC was less operative for Alloy 690 than for austenitic SSs during fatigue crack growth in PWR water. Consequently, the LCF life of Alloy 690 was longer than that of austenitic SSs in PWR water. .

Acknowledgments This study was mainly supported by the Korea Hydro and Nuclear Power Co., Ltd. as an Environmental Fatigue Project. Part of the funding was provided by the Nuclear Research & Development Program (No. 20121610100040), the KETEP-International Collaborative Energy Technology R&D Program (No. 20128540010010) of MOTIE of the Republic of Korea. Financial support for one of the authors was provided by the BK-Plus Program of the MSIP of the Republic of Korea.

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Table captions Table 1. Chemical Compositions of Alloy 690 (wt.%) Table 2. Tensile Properties of Alloy 690 Table 3. Low Cycle Fatigue Test conditions Table 4. Quantified Dislocation Density by Line-intersection Method in RT Air, 310 °C Air and PWR Water Table 5. Comparison of Low Cycle Fatigue Properties of Austenitic Alloys in PWR Water

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Table 1. Chemical Compositions of Alloy 690 (wt.%) C

Ni

Cr

Fe

Si

Mn

P

S

Al

Cu

Ti

0.01

59.31

29.76

9.30

0.39

0.21

0.003

0.001

0.36

-

0.26

Table 2. Tensile Properties of Alloy 690 Y.S. (MPa)

U.T.S. (MPa)

Elong. (%)

RT

290

664

55.1

310°C

226

582

55.8

Table 3. Low Cycle Fatigue Test Conditions Test materials

Alloy 690

Condition

RT Air, 310°C Air, PWR water (310°C)

Wave

Fully reversed triangle waveform (R=-1)

Control mode

Strain control

Strain rate (%/sec)

0.4 - 0.004%/sec

Strain amplitude (%)

0.33 - 1.0 %

PWR Water Dissolved

< 5 ppb

-chemistry

Oxygen (DO) Dissolved

2.2ppm (~25 cc H2/kg•H2O)

Hydrogen (DH) Conductivity

20-25 PS/cm

at RT

(1200 ppm H3BO3 + 2.2 ppm LiOH)

pH at RT

6-7 19

Table 4. Quantified Dislocation Density by Line-intersection Method in RT Air, 310 °C Air and PWR Water Condition

Dislocation density (x 1014 m-2)

RT Air

1.8

310 °C Air

3.74

PWR Water

5.96

Table 5. Comparison of Low Cycle Fatigue Properties of Austenitic Alloys in PWR Water Austenitic SSs

Ni-Cr-Fe alloys

Wider crack opening/Blunted crack tip

Less blunted crack tip

-

Large amount of metal dissolution

-

High hydrogen evolution rate

-

Longer LCF life than austenitic SSs

Clear evidence of HIC mechanism -

Smaller amount of metal dissolution

Dislocation multiplication,

-

HIC mechanism

-

Planar array with high density of dislocations

Micro-void at crack tip and cleavage -

crack

20

Low hydrogen evolution rate

Figure captions Fig. 1. Microstructure of Alloy 690. Fig. 2.

LCF life of Alloy 690 in RT air and 310°C air

Fig. 3. LCF life of Alloy 690 in PWR water Fig. 4. Cyclic stress behavior of Alloy 690 in the conditions of RT air, 310 °C air, and PWR water Fig. 5.

Magnitude of hardening of Alloy 690 in PWR water with varying strain rate

Fig. 6. The ratio of cyclic stress amplitude (Vpeak/Vinitial) of Alloy 690 and 316LN SS in various environments: strain amplitude of 0.4 % and strain rate of 0.04 %/sec for Alloy 690, 0.46 ~ 0.65 % and 0.024 %/sec for 316LN SS Fig. 7. Fracture surfaces of Alloy 690 in the fatigue propagation stage tested in (a) RT Air (0.033%/sec, 0.5%), (b) 310 °C Air (0.04%/sec, 0.4%), and (c) PWR water (0.04%/sec, 0.6%) Fig. 8. Fatigue crack of Alloy 690 tested in RT Air (0.033%/sec, 0.33%), 310 °C Air (0.04%/sec, 0.4%), and PWR water (0.04%/sec, 0.4%) Fig. 9. Bulk dislocation structure near the crack in (a), (b) RT Air (0.033%/sec, 0.33%), (c) 310 °C Air (0.04%/sec, 0.4%) and (d) PWR water (0.04%/sec, 0.4%) Fig. 10.

Fatigue crack of (a) Alloy 690 (0.04%/sec, 0.4%) tested in PWR water and (b) 316LN SS (0.024%/sec, 0.46%) tested in 310 °C deoxygenated water

21

Figure(s)

  Fig. 1.

Fig. 2.

Microstructure of Alloy 690

LCF life of Alloy 690 in RT air and 310°C air

22

Fig. 3.

Fig. 4.

LCF life of Alloy 690 in PWR water

Cyclic stress behavior of Alloy 690 in the conditions of RT air, 310 °C air, and PWR water 23

Fig. 5.

Magnitude of hardening of Alloy 690 in PWR water with varying strain rate

Fig. 6. The ratio of cyclic stress amplitude (ΔVpeak/ΔVinitial) of Alloy 690 and 316LN SS in various environments: strain amplitude of 0.4 % and strain rate of 0.04 %/sec for Alloy 690, 0.46 ~ 0.65 % and 0.024 %/sec for 316LN SS 24

(a)

(b)

25

(c) Fig. 7.

Fracture surfaces of Alloy 690 in the fatigue propagation stage tested in (a) RT Air

(0.033%/sec, 0.5%), (b) 310 °C Air (0.04%/sec, 0.4%), and (c) PWR water (0.04%/sec, 0.6%)

26

Fig. 8.

Fatigue crack of Alloy 690 tested in RT Air (0.033%/sec, 0.33%), 310 °C Air (0.04%/sec, 0.4%), and PWR water (0.04%/sec, 0.4%)

27

(a)

(b)

28

(c)

(d) Fig. 9.

Bulk dislocation structure near the crack in (a), (b) RT Air (0.033%/sec, 0.33%), 310 °C Air (0.04%/sec, 0.4%), and PWR water (0.04%/sec, 0.4%) 29

(a)

(b) Fig. 10.

Fatigue crack of (a) Alloy 690 (0.04%/sec, 0.4%) tested in PWR water and (b) 316LN SS (0.024%/sec, 0.46%) tested in 310 °C deoxygenated water

30