Effect of high-temperature water and hydrogen on the fracture behavior of a low-alloy reactor pressure vessel steel

Journal of Nuclear Materials 478 (2016) 343e364

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Effect of high-temperature water and hydrogen on the fracture behavior of a low-alloy reactor pressure vessel steel €tig a, Z. Que a S. Roychowdhury a, b, *, H.-P. Seifert a, P. Spa a b

Paul Scherrer Institut, Nuclear Energy and Safety Research Department, Laboratory for Nuclear Materials, 5232 Villigen, PSI, Switzerland Materials Processing & Corrosion Engineering Division, Mod-Lab, D-Block, Bhabha Atomic Research Centre, Mumbai 400085, India

h i g h l i g h t s
Hydrogen content, microstructure of LAS, and strain rate affects tensile properties at 288  C.
Strength affects hydrogen embrittlement susceptibility to a greater extent than grain size.
Hydrogen in LAS leads to strain localization and restricts cross-slip at 288  C.
Possible hydrogen pickup due to exposure to 288  C water alters fracture surface appearance without affecting fracture toughness in bainitic base material.
Simulated weld heat affected zone microstructure shows unstable crack propagation in 288  C water.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 6 December 2015 Received in revised form 14 April 2016 Accepted 22 May 2016 Available online 24 May 2016

Structural integrity of reactor pressure vessels (RPV) is critical for safety and lifetime. Possible degradation of fracture resistance of RPV steel due to exposure to coolant and hydrogen is a concern. In this study tensile and elastic-plastic fracture mechanics (EPFM) tests in air (hydrogen pre-charged) and EFPM tests in hydrogenated/oxygenated high-temperature water (HTW) was done, using a low-alloy RPV steel. 2e5 wppm hydrogen caused embrittlement in air tensile tests at room temperature (25  C) and at 288  C, effects being more significant at 25  C and in simulated weld coarse grain heat affected zone material. Embrittlement at 288  C is strain rate dependent and is due to localized plastic deformation. Hydrogen pre-charging/HTW exposure did not deteriorate the fracture resistance at 288  C in base metal, for investigated loading rate range. Clear change in fracture morphology and deformation structures was observed, similar to that after air tests with hydrogen. © 2016 Elsevier B.V. All rights reserved.

Keywords: Low alloy steel Hydrogen embrittlement Fracture toughness

1. Introduction Structural integrity of the reactor pressure vessels (RPV) of light water reactors (LWR) is of utmost importance with regard to safety of operation and service lifetime. In addition to the known life limiting degradation processes (irradiation embrittlement, environmentally assisted cracking (EAC) and thermo-mechanical fatigue [1e3]), concerns due to detrimental environmental effects remain that can be attributed to hydrogen uptake from primary water and/or corrosion reactions [4]. A corrosion resistant austenitic stainless steel cladding is

* Corresponding author. Materials Processing & Corrosion Engineering Division, Mod-Lab, D-Block, Bhabha Atomic Research Centre, Mumbai 400085, India. E-mail addresses: [email protected], [email protected] (S. Roychowdhury). http://dx.doi.org/10.1016/j.jnucmat.2016.05.033 0022-3115/© 2016 Elsevier B.V. All rights reserved.

present on the RPV internal surface for most of the LWRs, which minimizes corrosion and reduces accumulation of corrosion products in the coolant and crud built-up on the fuel elements. However, direct contact of RPV steel and coolant may occur within cracks in the cladding formed during fabrication or welding (e.g., hot cracks and relaxation cracks in the cladding or cladding/RPV interface) [5], or during reactor operation [6]. RPV feedwater nozzle corner cracking by thermo-mechanical fatigue and corrosion fatigue in boiling water reactors (BWR) or stress corrosion cracking (SCC) in RPV penetration and attachment welds in BWRs and pressurized water reactors (PWR) have been reported earlier [7] which causes direct contact of coolant and the RPV steel. The presence of open, incipient cracks has always to be assumed in safety analysis [8,9]. Hydrogen picked up by the RPV steel during fabrication is reduced to levels below 1 wppm (parts per million by weight) by

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heat treatment [3]. But hydrogen pickup by RPV steel can occur during service due to various reasons: i) contact with hydrogen containing reactor coolant (hydrogen from radiolysis primarily generated in the high neutron flux reactor core region and intentional additions to suppress radiolysis in PWR and to mitigate SCC in BWR with hydrogen water chemistry (HWC) and ii) due to corrosion reactions (uniform or local) which can result in very efficient local hydrogen pickup. Hydrogen uptake increases with increasing dissolved hydrogen concentration and corrosion rate and both are interrelated. Radiation also leads to radiation-induced transmutation and recoil proton injection but this contribution is not significant as compared to the other sources of hydrogen. The effect of irradiation is primarily indirect, by the creation of radiation defects as potential hydrogen traps, thus, increasing residual hydrogen levels in the RPV steel (and increase in yield stress, increasing hydrogen embrittlement susceptibility). The exchange current density of the hydrogen redox reaction in the hydrogenated, high-temperature water (HTW) on low-alloy steels (LAS), stainless steels and nickel base alloys is 1e3 orders of magnitude higher than the uniform corrosion rates. Hence, the corrosion potential of these materials are usually very close to the equilibrium electrode potential of the hydrogen redox reaction. Additionally, low conductivity of LWR coolants further limits cathodic overpotential and possibility for hydrogen supersaturation from local corrosion reactions. Since, the hydrogen diffusion, permeation and release rates are high in RPV steels and since oxide films do not present a significant diffusion barrier, the average bulk content in the steel (and thus the hydrogen enrichment at various traps) close to the surface is usually controlled by dissolved hydrogen content of water as per the Sievert’s law. In case of an intact cladding and under typical steady-state PWR power operation conditions, the bulk (lattice) hydrogen close to the surface from the dissolved hydrogen content in the coolant and corrosion reactions is less than 0.1 wppm, which is significantly lower than the hydrogen levels used in this study. In the RPV, there is a continuous flux of hydrogen from the inner wetted surface through the wall thickness to the “hydrogen-free” outside environment resulting in a corresponding hydrogen concentration gradient. Equilibrium hydrogen segregation with local bulk hydrogen is expected in case of saturable and non-saturable traps (e.g., sub-surface underclad cracks) over the prolonged operation periods. Depending on the nature and concentration of hydrogen traps (such as radiation defects, non-metallic inclusions, oxides etc.), local stresses and plastic strains in the RPV steel, the resulting average bulk hydrogen concentration can be significantly higher. After exposure to primary water in a nuclear reactor under neutron irradiation, measured hydrogen contents in the range of 1e5.8 wppm has been reported earlier [10]. Elevated corrosion rates, e.g., during flow-accelerated corrosion of carbon steels or under certain specific (transient) water or occluded crevice chemistries, inevitably increases hydrogen availability resulting in higher hydrogen pickup. As further discussed in Section 4.2., in oxygenated HTW, aggressive occluded crevice chemistry can be formed in cracks in LAS even in the absence of harmful anion species in the bulk environment due to the dissolution of MnS inclusions and enrichment of the resulting sulphides by the potential-gradient between crack mouth and tip. In hydrogenated HTW, on the other hand, there is no potential gradient and crack crevice chemistry and hydrogen content are similar to the bulk environment. This situation changes in case of stressed and plastically strained crack-tips, e.g., in fracture mechanics or corrosion fatigue tests under suitable strain rate conditions. Here, the local hydrogen pickup is dominated by the evolving complex local, aggressive occluded crack crevice chemistry and the local corrosion processes at plastically strained crack-tip as well as by the crack-tip/

notch mechanics (high hydrostatic triaxial stress state, dislocation transport of hydrogen). The dissolution of the MnS-inclusions that are intersected by the crack flanks and exposed to the crack crevice environment by the growing crack can result in a moderate acidification of the crack crevice even in hydrogenated HTW, if the MnS exposure and dissolution rate is higher than their transport out of the crack by diffusion [11,12]. Very high metal dissolution rates at the continuously plastically strained bare crack-tip may result in moderate cathodic overpotential and a corrosion potential intermediate to the equilibrium potential of metal dissolution and hydrogen reaction. The hydrolysis of metal cations produces additional hydrogen. Owing to the small area of the bare metal exposed at the crack tip compared to the crack flanks, low conductivity and fast diffusion rates of hydrogen in water, shifts in pH, corrosion potential and hydrogen concentration are moderate, but clearly increase the hydrogen availability and uptake. Furthermore, the hydrogen uptake under dynamic plastic straining in carbon and LAS [13] is significantly increased (hydrogen enhanced strain-induced vacancy mechanism) and further enriched in the peak hydrostatic stress region ahead of the crack-tip. The absorbed hydrogen levels in the process zone at crack-tip during fracture mechanics tests can thus be significantly high, eventually reaching the hydrogen levels used in this study achieved by pre-charging. Hydrogen effects on tensile behavior of LAS are strongly influenced by temperature, microstructure, and strain rate [14e19]. Most of the studies on hydrogen effects were done at room temperature, or below, where the effects are most significant, but limited number of test results are available at temperatures in the range of 250e288  C [18,20,21]. The weld HAZ, specifically the high hardness, coarse grain HAZ (CGHAZ), is highly susceptible to hydrogen effects [15,16,22]. Detrimental effects on mechanical properties of base metal and heat affected zone (HAZ) of RPV steels have been reported at hydrogen concentrations (2.5e3.5 wppm) similar to those picked up in the same material, after exposure to primary water in a nuclear reactor (1e5.8 wppm) [10], though calculations indicated hydrogen in RPV to be benign at reactor operating temperatures [23]. Hydrogen effects on EAC behavior of LAS in HTW, and synergistic (or competitive) effects with dynamic strain aging (DSA), are well recognized [2,11,24,25]. These synergistic effects can have a greater significance for weld HAZ material as compared to the base material. Moreover, increase in yield strength due to radiation also increases susceptibility to hydrogen embrittlement. There is now growing experimental evidence that the mechanical properties of most structural materials might be degraded by exposure to reactor coolant (hydrogen effects) in the LWR operating regime [10,26e31]. It is still not fully understood if such hydrogen contents can be achieved in the RPV during LWR service and the limited information available is not fully conclusive in this regard. A few earlier studies have reported detrimental effects on the fracture of LAS due to exposure to 288  C water [30,32,33]. Environmental effects in 288  C water were manifested as a higher number of pop-in events without changing the initiation fracture toughness, attributed to synergy between corrosion generated hydrogen and DSA [33]. In another study, reduction in fracture toughness and crack growth resistance with reducing extension rates was reported for tests in 288  C water and was attributed to the high sulphur content of the steel [32]. Elongation rates used for the fracture toughness tests [32,33] were low and the measured toughness and crack growth resistance reduced with reducing extension rates [32]. This is an indication of a possible apparent reduction in toughness due to a significant role of EAC at slow test extension rates. Moreover, there were no reported fractographic observations after fracture toughness tests using LAS in water at

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288  C. In this investigation, two approaches were used to investigate the effect of HTW and hydrogen on tensile and fracture behavior of low-alloy RPV steel at reactor operating temperatures. In the first approach, hydrogen was intentionally introduced by cathodic precharging followed by tensile and fracture mechanics testing in air to identify potential hydrogen effects at reactor operating temperatures and to support the interpretation of the HTW tests. The hydrogen contents introduced by cathodic pre-charging were significantly higher than the average bulk hydrogen contents in the RPV during steady-state power operation. Such high hydrogen concentrations could be reached in the process zone at the crack-tip during a loss of coolant accident or a fracture mechanics test in HTW and may not be relevant to normal reactor operation. In the second approach, fracture mechanics tests were performed in HTW to investigate effects of hydrogen pickup from HTW on fracture behavior, under loading conditions with minimal effect of EAC and similar strain rates as in loss of coolant accidents. The effect of test temperature, strain rate and microstructures was investigated and detailed post-test characterization was done. A mechanistic explanation for hydrogen and HTW effects in RPV steels in the LWR regime and interpretation of the experimental observations is given.

345

followed by air cooling (AC). Tempering martensite reduced the microhardness to 344 HV0.1 and 250 HV0.1 for HT1 and HT2 specimens, respectively. PAG boundaries with an average size of 75 mm for both HT1 and HT2 specimens were partially resolved by etching in 2% nital solution. A similar approach of heat treating steels to simulate weld HAZ has been reported earlier [38,39]. The PAG size, microhardness/yield stress and annealed martensite microstructure after the heat treatment HT1 of the Biblis C base metal are very similar to that of the high hardness CGHAZ (PAG size ~50 mm) of the actual circumferential core girth weld of the PWR RPV that was evaluated by metallographic examination at PSI [35]. The difference in hardness after HT1 is less than 10% of the hardness of the actual weld CGHAZ (typically 320 HV0.1 [35] with locally higher values of up to 350 HV0.1) and no difference in microstructure was observed after metallographic examination in an optical microscope. In contrast to the simulated CGHAZ material (HT1), the actual weld characterized earlier [35] was subjected to a post weld heat treatment after welding, at a ~50  C lower temperature than the Q þ T tempering temperature (545  C/59 h/465  C/600  C/21 h/465  C/595  C/11.3 h/air cooled) to reduce the weld residual stress. The heat treatment HT1 does not exactly match the thermal history of the RPV weld CGHAZ zone and considering the heterogeneous nature of the RPV weld HAZ (in terms of microstructure/strength/hardness), moderate differences may exist between the actual weld CGHAZ and the simulated weld CGHAZ which can only be revealed by detailed TEM studies (dislocation structures, internal stresses, carbide morphologies, etc.). The authors regard HT1 as a fairly reasonable simulation and approximation of the real CGHAZ in a multipass submerged arc weld of thick-walled structures with regard to PAG size, microhardness/yield stress and annealed martensite microstructure (which are believed to be the most important factors with respect to the hydrogen embrittlement behavior) with identical chemical composition and initial microstructure. The heat treatment HT2 aimed at having a microstructure and PAG size similar to HT1 but with strength levels comparable to the AR base metal. This was done to separately evaluate the effects of strength and grain size on measured properties.

2. Materials and experimental procedure 2.1. Materials A low-alloy RPV steel with low to medium sulphur (22 NiMoCr 3 7) was used in this investigation, having a chemical composition similar to pressure vessel steel ASTM grade A 508 Gr.2 Cl.1 [34]. This material is from an actual PWR RPV that was not commissioned. The steel used has been characterized as being moderately susceptible to DSA, and of low EAC susceptibility in previous investigations [11,35]. Tensile properties of this steel are listed in Table 1, along with the chemical composition, as previously reported in Ref. [35]. The steel was austenized and quenched (870e905  C/6.9 h/WQ) and tempered (635e655  C/11.3 h/air cooled) in as-received (AR) condition, having a bainitic microstructure with a prior austenite grain (PAG) size in the range 10e20 mm and hardness in the range 261 ± 4 HV0.1. Specimens with three different microstructures were used: i) AR base metal, ii) a simulated, high-hardness coarse grain heataffected zone (CGHAZ) (HT1) and iii) low hardness simulated CGHAZ (HT2). HT1 and HT2 specimens were produced by a specific 2-stage heat treatment in argon atmosphere to prevent oxidation and the properties are summarized in Table 2. AR specimens were first given a grain coarsening heat treatment (1000 þ 5  C/1 h/ water quenching (WQ)) [36,37], which resulted in high hardness martensite formation (549 HV0.1). Subsequently, the specimens were tempered at 625 ± 5  C for 1 h (HT1) and for 72 h (HT2)

2.2. Specimens Cylindrical tensile specimens of transverse orientation (T), of gauge length 30 mm and gauge diameter 6 mm were used for tensile tests. 1T-C(T) (B ¼ 25 mm) and 0.5T-C(T) (B ¼ 12.5 mm) compact tension specimens of TL orientation were used for elasticplastic fracture toughness (EPFM) tests in air and in HTW. The C(T) specimens were fatigue pre-cracked in air to an initial crack length to ligament ratio a0/W of ~0.5. The load during the final stage of pre-cracking was controlled to keep the maximum stress intensity factor, Kmax, at 20.7 MPa m0.5 and 13.6 MPa m0.5 for the 1T-C(T) and 0.5T-C(T) specimens, respectively. Subsequent to fatigue pre-

Table 1 Chemical composition (weight %) and mechanical properties of 22 NiMoCr 3 7. C

Si

Mn

P

S

Ni

Cr

Mo

V

Cu

Co

Al

N (wppm)

0.22

0.20

0.91

0.008

0.007

0.880

0.420

0.53

0.007

0.04

0.01

0.018

80

Tensile properties Ductility

25  C 288  C

sYS (MPa)

sUTS (MPa)

Total elongation (%)

Reduction in cross-sectional area (%)

467 400

605 578

17 16

72 70

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Table 2 Mechanical properties and microstructural characteristics of AR, HT1 and HT2.

AR HT1 HT2

Grain size (mm)

sYS at 288  C (MPa)

sUTS at 288  C (MPa)

Hardness (HV0.1)

Microstructure

10e20 75 75

367 741 464

525 846 587

261 344 250

Bainite Tempered martensite Tempered martensite

cracking, all C(T) specimens were side-grooved to a thickness reduction of 10% on each side. Fig. 1(a) and (b) shows the dimensions of the specimens used in this investigation. 2.3. Hydrogen charging procedure for tests in air Cathodic hydrogen charging in galvanostatic mode was done at room temperature (25  C) in 1N sulphuric acid þ 30 mg l1 arsenic

trioxide solution. Specimen surface was metallographically polished up to 1000 grit emery paper, followed by ultrasonic cleaning in acetone for 5 min prior to charging. Gauge section of tensile specimens were charged for 1 and 7 h at a current density of 10 mA cm2 and the entire C(T) specimen was charged for 7 and 24 h at a current density of 0.3 mA cm2. Charging was done for two different durations to introduce two different levels of hydrogen in the steel. Specimen surfaces were polished with 1000 grit emery

Fig. 1. a) Schematic showing dimensions of C(T) specimens (0.5T-C(T)/1T-C(T)) along with the configuration for DCPD crack length measurement and b) cylindrical tensile specimen dimensions.

S. Roychowdhury et al. / Journal of Nuclear Materials 478 (2016) 343e364

paper after charging and subsequently copper was electrolytically deposited [18]. Slow diffusion and low solubility of hydrogen in copper result in slow permeability, which is several orders of magnitude lower than in bcc steel. Starting material (AR) is not expected to contain any hydrogen (0 wppm). Hydrogen analysis by hot gas extraction was done after tensile testing at 25  C for two AR and for one HT1 and HT2 samples. The specimens were stored in liquid nitrogen (196  C) immediately after tensile testing to prevent hydrogen effusion. The hydrogen content in segments containing the fracture zone was analyzed to be 2.6 ± 0.4 wppm after 1 h charging and 5.5 ± 0.6 wppm after 7 h charging, for all the three microstructural conditions. During hydrogen charging at 25  C, hydrogen diffuses inside to a distance of approximately 0.3e0.7 mm from the surface (calculated using Fick’s second law of diffusion). During subsequent tensile loading, dislocation induced hydrogen transport occurs which homogenizes the hydrogen across the cross-section. Homogenization is confirmed by fractographic evidences of hydrogen observed across the cross-section after tensile tests (see Section 3.1.1). Tensile specimens tested at 288  C after hydrogen charging were assumed to have the same hydrogen concentration, for the same charging conditions. The measured hydrogen contents represent a realistic level of the concentration present in the fracture zone of the tests at 25  C. Hydrogen content in the high-temperature tensile tests is expected to be lower due to partial effusion of trapped hydrogen during heating, despite copper coating, and effusion from the uncoated fracture surface during cooling down from test temperature. Hence, hydrogen analysis after high-temperature tensile tests was not attempted, although the real concentrations are expected to be significantly lower. Hydrogen content of C(T) specimens (hydrogen precharged specimens tested in air and specimens tested in HTW) was not analyzed, since it is expected to be inhomogeneous during and after tests at 288  C, concentrated primarily at the crack tip under high tri-axial hydrostatic tensile stress. 2.4. Tensile tests in air Specimens were tested immediately after copper deposition and the tests were performed at room temperature (25  C) and 288  C, at strain rates varying from 106 s1 to 101 s1. Tests at 288  C were done using a three zone split type furnace and a thermocouple was attached to the middle of the gauge section of specimens to ensure control of ±3  C for the test duration. Test temperature was reached typically within a period of 80e90 min and the specimens were held at a constant load of 200 N at 288  C for additional 10 min, prior to start of loading. All the tests were repeated 2 (AR, 0 wppm hydrogen) to 3 (hydrogen pre-charged) times under identical test conditions. Subsequently, fractographic examination was done in scanning electron microscope. 2.5. Elasto-plastic fracture mechanics (EPFM) tests 2.5.1. EPFM tests in air Tests using air fatigue pre-cracked 0.5T-C(T) and 1T-C(T) specimens were done at 25, 150 and 288  C and temperature at the notch was controlled within ±3  C by a thermocouple. Hydrogen precharged C(T) specimens (charging duration e 7 and 24 h) were tested after 24 h to allow for through thickness homogenization of hydrogen in the crack-tip region. Hydrogen atoms diffuses inside from the C(T) specimen surface as well as the crack tip region during the homogenization period. Further homogenization occurs as a result of dislocation transport in the crack tip plasticity region during the EPFM test. Load-line displacement (LLD) was recorded using a clip gauge (Sandner EXRC3þ6u) and tests were done at an

347

LLD rate of 0.25 and 0.35 mm min1 for 0.5T-C(T) and 1T-C(T) specimens, respectively. Crack length was calculated from periodic unloading compliance (LLD interval of 0.1 and 0.2 mm for 0.5T-CT and 1T-C(T) specimen, respectively) and corrected for specimen rotation [40]. EPFM tests and data analysis was done according to ASTM E 1820 [40]. A few tests in air without hydrogen pre-charging were done without partial unloading and with crack length measurement by the reversed direct current potential drop (DCPD) method. This was done to compare the different evaluation procedures and crack length measurement methods. 2.5.2. EPFM tests in high-temperature water EPFM tests were done at 150 and 288  C in simulated BWR and PWR primary water conditions in refreshed autoclaves attached to recirculating water loop maintaining water chemistry, previously described in detail [2,41] and a simplified schematic of the HTW loop is shown in Fig. 2. Reducing BWR hydrogen water chemistry (HWC) was simulated by hydrogenated (dissolved hydrogen (DH) content of 1.4 wppm) and nitrogenated neutral high-purity water  (pH288 C ¼ 5.7, inlet conductivity ¼ 0.055 mS cm1). Oxidizing normal water chemistry (NWC) was simulated by oxygenated (dissolved oxygen (DO) content of 2 wppm) neutral high-purity water. The corresponding electrochemical corrosion potentials (ECP) of the specimens at 288  C were 590 mVSHE (HWC/nitrogenated water) and þ100 mVSHE (NWC). The DO and DH levels used in this study are higher than that usually found in LWRs. PWR primary water was simulated by mildly alkaline borated and lithiated, hydrogenated high-purity water (1000 wppm B as H3BO3,  2.3 wppm Li as LiOH, pH288 C ¼ 6.9, inlet conductivity of 24 mS cm1, DH content of 1.8e1.9 wppm). The corresponding ECP at 288  C was 735 mVSHE. In some experiments in PWR simulated conditions, additional in-situ galvanostatic hydrogen charging was done at a current density of 3 mA cm2 inside the autoclave at 288  C for the entire duration of the test, which shifted the measured corrosion potential to more negative values by ~700 mV (the real polarization and cathodic overpotential might be significantly smaller due to ohmic potential drops because of the low conductivity of the lithiated and borated water of ~200 mS cm1 at 288  C). Table 3 lists the experimental details and results for the tests done in HTW. Tests A1 to A5 were carried out in PWR simulated conditions, B1 to B5 were done in BWR/HWC and test B6 and B7 were done in BWR/NWC simulated conditions and nitrogenated (oxygen- and hydrogen-free) HTW, respectively. Due to load capacity limitations in the PWR loop, only 0.5T-C(T) specimens were tested in PWR simulated conditions. During the experiments, all important mechanical (load, pull rod stroke) and environmental parameters at inlet and outlet (dissolved oxygen and hydrogen, conductivity, pH, temperature, pressure, flow rate) were recorded continuously. The corrosion potential ECP of the specimen and redox potential (platinum probe) were continuously monitored with respect to a Cu/Cu2O/ZrO2-membrane reference electrode. Crack advance was monitored using reversed direct current potential drop (DCPD) with a resolution of approximately 1 mm and crack length was calculated using Johnson’s equation [2] (see also Section 2.6.2). 2.5.2.1. Pre-oxidation in high-temperature water. The air fatigue pre-cracked specimens were initially pre-oxidized for a period of ~12 days to establish a homogenous oxide film and enable hydrogen pick-up from the environment and corrosion reactions. During pre-oxidation, specimens were either cyclically loaded (frequency 0.01 Hz, DK 9.9 MPa m0.5, minimum load/maximum load (R) ¼ 0.7, ~104 cycles, ~12 days) to have an active growing EAC or the specimens were under low constant load (no crack growth,

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Fig. 2. Schematic of HTW loops. Table 3 Experimental details and J values at stable crack initiation (JQ and JDCPD) for tests in HTW with base metal. T ( C) Specimen type LLD rate (mm min1) Environment Test no. Pre-oxidation Duration (days)

Mode of loading Crack growth (mm) In-situ hydrogen JQ

JDCPD C1

C2

A1 A2 A3 B1

12 19 11.5 12

Constant loada 18,177 cyclesb Constant loada 10602 cyclesb

0 0.246 0 0.166

3 mA cm-2 3 mA cm2 no e

346 346 331 393

229 257 214 270

0.471 0.479 0.527 0.531

BWR/HWC

A4 A5 B2

3 mA cm2 no e

290 182 306 218 296 190

393 0.436 400 0.417 402 0.460

BWR/NWC O2& H2 free

39932 cyclesb Constant loada Constant loada þ 10,000 cyclesb 27483 cyclesb Constant loada 10,000 cyclesc 11,842 cyclesb 10,311 cyclesb

0.341 0 0.081

B3 B4 B5 B6 B7

29 12 34 (Constant load)a þ 11 (Cyclic load)b 32 12 20 15 12

0.378 0 0 0.085 0.06

e e e e e

379 385 288 299 299

472 465 388 427 406

150

0.5T-C(T)

1T-C(T)

0.350

BWR/HWC

288

0.5T-C(T)

0.250

PWR

1T-C(T)

0.350

a b c

0.250

J (kN m1) J ¼ C1$DaC2

PWR

256 186 167 184 167

446 447 443 515

0.426 0.526 0.426 0.519 0.465

500 N for 0.5T-C(T), 1000 N for 1T-C(T), (KI < 3 MPa∙m0.5). Cyclic loading, DK ¼ 9.85, 0.01 Hz. Cyclic loading, DK ¼ 2.9, 0.01 Hz.

duration of exposure ~12 days). The EAC growth rates under cyclic load (~1.2$1010e2.2$1010 m s1) were two to three times higher than the corresponding fatigue crack growth rates in air (~8.0$1011 m s1) and correspond to the low sulphur line of the Ford and Andresen model [2,11]. Some experiments were done with higher number of cyclic loading (up to ~4$104 cycles) with higher exposure time (up to 32 days) during pre-oxidation. In one case, cyclic loading was done at a DK of 2.9 MPa m0.5, below the fatigue and EAC crack growth threshold and hence did not induce any active EAC growth but induced only crack tip plasticity. Both pre-oxidation period and cyclic loading/EAC crack growth could influence hydrogen accumulation and initial content in the cracktip process zone affecting the subsequent fracture behavior. Hydrogen pickup from HTW and diffusion inside the C(T) specimens at the test temperature is discussed in Section 4.2.1. 2.5.2.2. EPFM tests in high-temperature water. C(T) specimens were monotonically loaded immediately after pre-oxidation and the pull rod stroke corrected by the elastic compliance of the load train was considered to be LLD of the specimen, which was confirmed by

control experiments with clip gauges. The applied loading rate for 0.5T-C(T) was 0.25 mm min1 (dK/dt-1$102 to 2$102 MPa m0.5 h1) and for 1T-C(T) 0.35 mm min1 (dK/dt e 3$102 to 5$102 MPa m0.5 h1), same as that used for tests in air and similar to loading rates in loss of coolant accident transients. Loading rates induced an average dJ/dt of 4$103 kN m1 h1 for both types of C(T) specimens. Compliance technique for crack length measurement was intentionally not used to prevent any additional EAC effects during loading/unloading [33]. After the tests, specimens were immersed in liquid nitrogen, broken open and were used for fractographic examination. 2.5.3. Calculation of J-integral 2.5.3.1. Tests in air. The load-displacement data was used to calculate the J values by the incremental method of ASTM E1820 [40]. The crack extension was determined by the unloading compliance technique, which usually slightly underestimated the final crack length by ~5%. 2.5.3.2. Tests in high-temperature water. The reversed DCPD

S. Roychowdhury et al. / Journal of Nuclear Materials 478 (2016) 343e364

technique (configuration of wire connections in Fig. 1(a)) was used for crack length monitoring and the crack length was calculated using the Johnson equation (1) where U and Uo are the DCPD signal corresponding to the instantaneous crack length (a) and initial crack length (ao) measured by post-test fractography respectively, 2 y is the distance between the DCPD potential measurement points (Fig. 1(a)) and W is the C(T) specimen width. The crack extension as calculated by DCPD was linearly corrected by the actual crack extension measured by post-test fractography. The initial fatigue or corrosion fatigue pre-crack length a0 from fractography was assigned to initiation of stable crack growth as described in Figs. 3 and 4(a). The crack extension measured by DCPD slightly overestimated the actual crack extension (from fractography) by ~5% and this accuracy was better than for the unloading compliance method. The relative error of the corresponding crack lengths from the load line was much smaller (<1%). The applied procedure for the detection of the onset of ductile crack growth was verified by qualification tests in air and HTW that were interrupted shortly after the onset of ductile crack growth. The difference between the small crack advances of 70e300 mm predicted by DCPD and measured by post-test fractography was less than 10%, confirming the adequacy of the selected approach.

a¼

2W

p

cos1

coshðpy=2WÞ h  .  io py pao cos 2W coshfðU=Uo Þcosh1 cosh 2W

(1)

The methods for the determination of the J-Da curve, blunting line and JDCPD at the onset of ductile crack growth are illustrated in Fig. 3. The blunting line and apparent crack advance Dai due to blunting were calculated with JDCPD and the flow stress sY (sY ¼ (sYS þ sUTS)/2) using equation (2). After onset of ductile crack growth, the J values [40] were calculated with the load, LLD and corresponding crack length, at an LLD interval of 0.1e0.2 mm, by the same method as used for tests in air using unloading compliance. The compliance change of the specimen with crack growth was considered during the calculation [40]. J-Da curves were plotted using the calculated J values, where Da ¼ DaDCPD þ Dai is the increase in crack length during loading. Fig. 4(a) shows the JDa curve with all the exclusion lines, blunting line and the limits on

349

J and Da as specified by the standard [40].

Dai ¼

JDCPD 2$sY

(2)

2.5.4. J at initiation of stable crack growth J values at initiation of stable crack growth were calculated by two different techniques, as per the ASTM standard (JQ) [40] and from DCPD signal (JDCPD). JQ was the J value corresponding to the point of intersection between the regression line (J ¼ C1$DaC2) and the 0.2 mm exclusion line (line parallel to blunting line and offset C from origin by 0.2 mm, slope of blunting line ¼ 2$s288 for highy 25 C  temperature tests and ¼ 2$sy for 25 C tests), as shown in Fig. 4(a). To qualify JQ as JIC, fifteen different criteria of ASTM E1820 [40] have to be satisfied. Important small scale yielding criteria were always satisfied but other isolated criteria were not fulfilled occasionally in some tests. Therefore, JQ is reported instead of JIc in this paper. For tests using DCPD signal for crack length measurements, additionally JDCPD at onset of stable crack growth was measured. Fig. 4(b) shows a typical variation in the DCPD signal during tests in HTW. JDCPD was calculated at the instant when the slope of DCPD potential drop vs. LLD curve changes (Fig. 4(b)) during monotonic loading, in accordance with earlier studies [2]. The JQ values measured from unloading compliance and by DCPD (control test in air) were very similar, the latter usually being marginally lower. The JDCPD values on the other hand were usually ~35% (25e50%) smaller than the corresponding JQ values. DCPD technique has been successfully used for plotting J-Da curves and estimating JQ for a variety of materials with accuracy [42e46] and the crack length measured by both the techniques in this study have also been reported to be comparable [42,43]. 2.5.5. Validity of EPFM tests Valid J values were calculated for both specimen dimensions (0.5T-C(T) and 1T-C(T)) at all test temperatures and were within limits for J (Jlimit) and crack growth Da (Dalimit) [29,31] and the

Fig. 3. Applied procedure for plotting J-Da curve by the DCPD technique.

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Fig. 4. (a) J-Da plot in air with validity limits and blunting and exclusion lines and (b) variation in DCPD signal and load during EPFM test in HTW.

limits were more relaxed for 1T-C(T) specimen as compared to 0.5T-C(T) specimen. The specimen dimensional requirements were satisfied, as per ASTM E1820 standard [40]. Valid data points meet small scale yielding criteria; thus the measured toughness should be independent of specimen dimensions. Similar toughness values were measured for both the 0.5T and 1T-C(T) AR specimens. Specimen size criteria are less stringent for the high strength simulated CGHAZ materials HT1. Observation of unstable ductile cracking in air and HTW in initial tests using simulated CGHAZ specimens may be due to increased low-temperature creep of this material condition and indicates a possible size effect. Further testing in heat treated specimens is necessary for a clear understanding. The unloading compliance technique is the standard method for crack extension measurements in ASTM E1820. The alternatively applied DCPD method and procedure for crack growth measurement was verified by qualification tests in air and HTW that were interrupted shortly after the onset of ductile crack growth. The achieved accuracy for crack extension (0e0.05 mm) was much better than the criteria of ASTM E1820, where a difference of ±0.25 mm between the measured and predicted crack extension of 0.5 mm is acceptable [40]. 2.6. Post-test evaluations Fracture surface of all specimens and polished and etched crosssections of selected specimens were analyzed by light and field emission gun scanning electron microscopy (FEG/SEM) with energy dispersive X-ray spectroscopy (EDS). AR specimens with 0 wppm and 5.5 wppm hydrogen, tested at 288  C and strain rate of 102 s1, and having similar uniform strain (~11%) but different total strain to failure and reduction in cross-sectional area (Fig. 5) were selected for transmission electron microscopy (TEM) evaluation. Schematic indicating the location from where the TEM specimens were prepared is shown in Fig. 5 resulting in TEM specimens corresponding to ~11% strain. 3. Results 3.1. Tensile test results 3.1.1. Tensile behavior at 25  C Tensile tests at 25  C were done using AR specimens at strain rates in the range of 106e101 s1. Fig. 6(a) and (b) shows 5.5 wppm of hydrogen resulted in a marginal hardening (increase in yield stress, sYS) and significant reduction in ductility (reduction in cross-sectional area) over the test strain rate range. A positive

Fig. 5. Engineering stress-strain plots for the specimens used for TEM investigations.

strain rate sensitivity of sYS was observed at strain rates greater than 104 s1 both with and without hydrogen. The reduction of area with and without hydrogen as well as the embrittlement does not depend on the strain rate. The apparent recovery of ductility at the slowest strain rate of 106 s1 might be related to partial hydrogen effusion during the long test duration of several days. In absence of hydrogen, fracture is always ductile due to microvoid coalescence (MVC) leading to the formation of equiaxed dimples. Final fracture surface has a “cup-cone” shape and a typical appearance is shown in Fig. 7(a). Fig. 7(bed) shows the typical fracture surface after tensile tests at a strain rate of 104 s1 for specimens containing 5.5 wppm hydrogen indicating regions of ductile fracture due to MVC, quasi-cleavage (QC) and brittle fracture. Brittle and QC facets were observed to be always associated with inclusions, confirmed by energy dispersive spectroscopy to be predominantly oxides of aluminum and had a very similar microfractographic appearance as EAC and QC features at MnSinclusions in HTW tests [2,12,24,25]. Shear dominated (ductile) fracture was always observed in presence of hydrogen, as indicated by ~45 inclination of the plane of fracture to the loading direction and by elongated shear dimples (Fig. 7(d)). Regions of QC (Fig. 7(b)) and brittle fracture (Fig. 7(c)) were perpendicular to the loading direction (regions marked “A” and “B” in Fig. 7(b)). 3.1.2. Tensile behavior at 288  C Tensile tests at 288  C were performed using AR and HT1

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Fig. 6. Variation in a) sYS, and b) reduction of area with applied strain rate at 25  C.

Fig. 7. Fractograph at 25  C a) “Cup-cone” fracture in absence of hydrogen, b) 45 inclination of the fracture surface to the loading direction in presence of 5.5 wppm of hydrogen showing the “QC” and “MVC” region marked, c) magnified view of region “A”, “B” showing brittle fracture (encircled regions) near inclusions and d) elongated dimples in the direction of shear.

specimens at strain rates of 103e101 s1, and at 102 s1 for HT2 specimens. Fig. 8(a) shows the typical engineering stress-strain plots measured for AR and HT1 at 288  C and a strain rate of 103 s1 clearly showing the effect of hydrogen in reducing tensile ductility (strain at failure). The susceptibility to embrittlement in HT1 is higher than in AR, as is evident by the much more pronounced reduction in the total strain at failure for the same hydrogen content of 5.5 wppm (Fig. 8(a)). Fig. 8(b) shows the effect of the hydrogen content on the reduction of area for the AR and HT1 material at the strain rate of maximum embrittlement of 102 s1. The degree of embrittlement is increasing with increasing hydrogen content for both microstructures. The degree of embrittlement is more pronounced for the high strength simulated CGHAZ (HT1) material and the threshold hydrogen content for

embrittlement seems to be lower than in the RPV base metal (AR). The negative strain rate sensitivity of UTS at 288  C (Fig. 9(a) and (b)) and serrations in the stress-stain curves confirm the occurrence of DSA, similar to earlier reported results in LAS at 288  C [18,13]. In contrast to the 25  C tests, hydrogen causes a softening as shown by the reduction in the sUTS in AR (Fig. 9(a)) and HT1 steel ((Fig. 9(b)) and the marginal reduction of sYS in HT1 (Fig. 9(d)). For the same hydrogen content, the softening is more pronounced in HT1 than in AR. In contrast to the room temperature tests, the embrittlement is strain rate dependent in both AR and HT1 specimens with a maximum of embrittlement at a strain rate of 102 s1 (Fig. 9(c)). At 288  C, the embrittlement seems thus to occur in a narrow strain rate range only. The degree of embrittlement is more pronounced

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Fig. 8. a) Typical engineering stress-strain curves for HT1 and AR at 288  C and a strain rate of 103 s1. b) Effect of hydrogen content on reduction of area for HT1 and AR at 288  C and the strain rate of maximum embrittlement of 102 s1.

Fig. 9. Tensile properties at 288  C at different test strain rates. Variation in sUTS for a) AR and b) HT1 and c) reduction in cross-sectional area for AR, HT1 and HT2 microstructure. d) Effect of hydrogen content on sUTS and sYS for HT1 at the strain rate of maximum embrittlement of 102 s1.

at 25  C, where the relative decrease in the reduction of area in the AR material is ~60% (Fig. 6(b)) compared to ~25% at 288  C at the strain rate of maximum embrittlement, for the same hydrogen precharging conditions. A part of this difference might be related to the partial effusion and release of hydrogen during heating up of the HT1 specimens. The HT2 material (simulated CGHAZ with low strength level as

AR) does not show significant hydrogen effects as compared to HT1 for the same hydrogen level of 2.6 wppm, as seen in Fig. 9(c). This clearly indicates that strength of steel affects hydrogen embrittlement susceptibility to a greater extent than the grain size for the hydrogen levels and loading conditions used in this study, and hydrogen effects become significant at lower concentrations in steel as strength increases.

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Hydrogen resulted in shear dominated fracture at 288  C (~45 inclination of the fracture surface to the loading direction similar to room temperature tests), for all the specimen microstructures (Fig. 10(a and b)). Elongated dimples due to shear (at region “C” in Fig. 10(a)) were also observed. In addition, other features like macrovoids greater than 100 mm in size and secondary cracks parallel to the loading directions can be seen in Fig. 10(b). Fracture was predominantly ductile due to MVC though QC facets associated with inclusions were observed in only a few of the experiments. Regions of “QC” and secondary cracks along with flat featureless regions (region “B” in Fig. 10(a)) indicate extensive localized plastic flow of the material. In the room temperature tests, the area fraction of QC and brittle features at inclusions on the fracture surface is much higher and is clearly related to the higher trapping efficiency at lower temperatures at strong trap centres such as the interface between the matrix and inclusions (oxide and MnS). 3.1.3. Deformation behavior and structures TEM investigation after tensile testing at conditions of high hydrogen embrittlement susceptibility (strain rate of 102 s1 at 288  C) clearly indicates that presence of hydrogen causes a greater degree of strain localization (Fig. 11(a)) as compared to specimen without hydrogen (Fig. 11(b)). Fig. 11(c) shows some evidence of strain localization in specimen without hydrogen and locations of possible increase in local stress levels in the carbide-matrix interface as reported earlier [18,47]. Localization of strain results in stress concentration and become the preferred sites for hydrogen accumulation and void nucleation. Dislocation cell formation with an approximate cell size of 200 nm was observed for specimens without hydrogen, with the cell interiors relatively free of dislocations (Fig. 12(a)) while for specimens with 5.5 wppm hydrogen, such cell structure was indistinct (Fig. 12(b)). Ease in dislocation motion and cross-slip is a necessary condition for dislocation cell formation in three dimensions [48]. Efficient association of hydrogen atoms with moving dislocations at a strain rate of 102 s1 hindering free mobility and restricting cross-slip, a reduction in stacking fault energy promoting planar coarse slip, and localization of plastic deformation will restrict dislocation cell formation. 3.2. Fracture toughness (EPFM test) results 3.2.1. Tests in air: effect of hydrogen Fracture toughness tests were performed at 25, 150 and 288  C using specimens in AR condition with 0 wppm hydrogen and after hydrogen pre-charging for 7 and 24 h (Table 4). This provided baseline data indicating the effect of hydrogen on fracture

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toughness and associated fractographic features. For AR material a JQ of 380 ± 60 kN m1 was measured at 25 and 150  C, respectively, and reduced to 290 ± 20 kN m1 at 288  C. At 25  C, no reduction in initiation and tearing resistance was observed after hydrogen charging for 7 h (Fig. 13(a)). After 24 h hydrogen pre-charging, sudden unstable brittle crack extension occurred at 25  C at a J value of 281 kN m1 (Fig. 13(a) and (b)) without any prior ductile crack growth, while at 288  C stable ductile crack extension occurred (JQ of 313 kN m1) for the same hydrogen charging duration (Fig. 13(c)). Though no reduction in JQ and tearing resistance was observed at 288  C after 24 h hydrogen pre-charging (Fig. 13(c)), a clear change in fracture morphology was observed (see Section 3.2.3). 3.2.2. Tests in high-temperature water J-Da plots for 1T-C(T) and 0.5T-C(T) specimens of AR material tested in HTW (288  C) and air at various test conditions (including EPFM tests in air using unloading compliance (ULC) and DCPD) are shown in Fig. 14. For the test conditions used in this study exposure to HTW did not reduce the initiation toughness and tearing resistance significantly, as compared to corresponding air tests (Figs. 14 and 15 and Tables 3 and 4) and only stable ductile crack growth was observed. However, a distinct change in fracture morphology and deformation structure [26] was observed, which was very similar to that after EPFM and tensile tests in air in the hydrogen pre-charged condition (see Section 3.2.3). The JQ value was similar for hydrogenated and oxygenated neutral or moderately alkaline HTW. Duration of exposure or loading conditions during pre-oxidation in HTW did not affect the initiation toughness significantly (Fig. 16). Additional continuous in-situ hydrogen charging under PWR conditions did not reduce the toughness (although this might not be very efficient due to the high electrolyte resistance). Thus, the change in fracture and deformation mode in HTW did not significantly affect the fracture toughness and the tearing resistance at 150 and 288  C in the AR base metal for the test parameters used in this study. In contrast to the base metal AR, initial EPFM tests in HTW and air using simulated CGHAZ (HT1) specimens indicate a potential HTW effect on initiation fracture toughness as shown by the lowering of JQ value from 440 kN m1 (288  C, air, no hydrogen) to 226 kN m1 (simulated PWR conditions, Fig. 17(aec)). So far, all these tests in air and HTW consistently indicated significant unstable rapid crack extensions. Occurrences of sudden unstable crack extensions were instantaneous, after stable crack extension to some extent and always occurred beyond the peak load. Fig. 17(a) shows a representative figure of the unstable crack extension in HT1 specimen and the fracture surface examination showed the

Fig. 10. a) “QC” in regions marked “A” and extensive plastic deformation in regions marked “B”, showing flat featureless regions and b) fracture surface showing large voids and secondary cracks.

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Fig. 11. Microstructure of AR specimen after tensile test at 288  C and 102 s1 applied strain rate (Fig. 5) in presence of a) 5.5 wppm (regions of strain localization marked) and b) 0 wppm hydrogen. c) strain localization at carbide-matrix interface in specimen with 0 wppm hydrogen.

Fig. 12. Dislocation structure after tensile test at 288  C and 102 s1 applied strain rate (see Fig. 9) in presence of a) 0 wppm and b) 5.5 wppm hydrogen.

Table 4 J values at stable crack initiation (JQ and JDCPD) and experimental details for air tests. Test in air with base metal (AR) Tests in air 

J ¼ C1$DaC2

J at initiation 1

Specimen type

LLD rate (mm min

Test condition

JQ (kN m

C1

C2

25 150 288

0.5T-C(T)

0.25

0 wppm H

316 323 267

199 135 168

495 431 385

0.641 0.501 0.501

25 25 25 150 288

1T-C(T)

0.350

0 wppm H

380 332 372 413 313

e e e e e

625 551 601 562 453

0.672 0.773 0.827 0.592 0.571

H precharged

281a (24 h) 440 (7 h) 312 (24 h)

e e e

e 609 441

e 0.573 0.542

a

JIC at sudden unstable fracture.

)

1

T ( C)

25 25 288

)

1

JDCPD (kN m

)

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Fig. 13. J-Da plots in air at (a) 25  C and (c) 288  C in the AR condition (0 wppm) and after hydrogen pre-charging, (b) fracture surface after unstable crack growth at 25  C, after 24 h hydrogen pre-charging.

related to plastic collapse due to low-temperature creep, DSA and strain localization, occurring simultaneously at the test temperature. Further experiments are needed and are currently in progress to evaluate the role of HTW on initiation toughness at 150 and 288  C and the reasons for the occurrences of such rapid unstable crack extensions.

Fig. 14. J-Da plots at 288  C in air and different HTW environments for AR base metal specimens.

crack growth to be predominantly ductile (Fig. 17(b)). Instantaneous crack jumps were not associated with any indication of brittle fracture, although some few isolated local brittle features, usually in the vicinity of inclusions, were observed at the ground of a few macro-grooves. The unstable ductile cracking might be

3.2.3. Fractography after EPFM tests In air tests without hydrogen, fracture is ductile at all the test temperatures, with very limited occurrence of macrovoids and quasi-cleavage facets. The macrovoids are in most cases associated with inclusions as these are preferential sites for void nucleation. Fig. 18(a) shows the fracture surface of a 1T-C(T) specimen after fracture toughness test in air at 288  C without hydrogen. After hydrogen charging and EPFM test in air at 288  C, various fractographic features in addition to ductile dimples can be seen in Fig. 18(b). A macrovoid is marked in Fig. 18(b) and other features such as cracks at inclusions (A), secondary cracks away from the primary crack plane (B) and localized quasi-cleavage facets associated with inclusions (C), are shown in insets, and are similar to those reported earlier [49]. Macrovoids and localized quasicleavage facets were observed to a greater extent in presence of hydrogen. It can be qualitatively concluded that the role of hydrogen in deteriorating mechanical properties in C(T) specimen and tensile specimen is similar, which resulted in similar fractographic features (Figs. 10 and 18). The role of hydrogen atoms in

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Fig. 15. Comparison of JQ and JDCPD values of AR base metal specimens at 288  C in air and different HTW environments.

Fig. 16. Comparison of JQ and JDCPD values of AR base metal specimens at 150/288  C in high-purity, hydrogenated HTW environment for various pre-oxidation conditions for experiments B1eB5 (see Table 3 for test details).

inducing strengthening (shear localization, drag effect)/softening (shielding effect) is also similar. The fractographic appearance after tests in BWR and PWR simulated conditions indicated features similar to those observed after tests in air at 288  C using hydrogen pre-charged specimens. Fig. 19(a) and (b) shows a typical fracture surface of a 1T-C(T) specimen tested in BWR/HWC (test B2) with different regions labeled. Fig. 19(c) shows a typical fracture surface after a test in PWR simulated conditions. Various features in addition to ductile dimples can be seen, such as macrovoids (inset A, Fig. 19(a) and (c)), secondary cracks (Fig. 19(b) and inset B in Fig. 19(c)) and quasicleavage facets (Fig. 19(aec)). Quasi-cleavage facets at inclusions were observed in some experiments, while macrovoids, secondary cracks (some of them with intergranular appearance as shown in Fig. 20(b)) and flat featureless regions resembling quasi-cleavage facets, were observed after all tests in HTW. The secondary cracks are ductile in nature, as indicated in Fig. 19(b), where both quasicleavage and ductile dimples can be seen inside the cracks.

Comparing Figs. 18 and 19, the differences/similarities in fractographic features are evident for tests done in air with/without hydrogen pre-charging and tests done in HTW. Tests in BWR and PWR simulated conditions resulted in higher roughness of the fracture surface (Fig. 19) due to significantly higher macrovoid and secondary crack formation. A cross-section of the fracture surface (inset in Fig. 20(a)) for test B2 was metallographically prepared and etched and Fig. 20(a) shows a typical macrovoid. Fig. 21 shows the typical appearance of flat features surrounded by shear lips (marked by arrows), after tests performed in PWR simulated conditions with in-situ hydrogen charging. Such flat features occur due to extensive localized plastic deformation during crack propagation and are not a microscopic evidence of embrittlement as reported earlier [50]. Elongated flat regions, as seen in Fig. 21, may be attributed to hydrogen accumulation at elongated inclusions (either sulphide stringer or band of oxide inclusions) and resultant localized plasticity manifested as tear ridges (marked in Fig. 21). Occurrence of “pop-ins” in LAS during fracture toughness tests in HTW was reported in Ref. [33] and in some of our tests as short bangs, which corresponded to the formation of quasicleavage facets. Macrovoids (Fig. 20(a)) are evidence of enhanced localized plasticity, that occur in regions of preferential hydrogen accumulation, such as inclusion matrix interface, lath boundaries and regions of high localized strain. Accumulation of hydrogen at such locations results in extensive localized plasticity, which leads to formation of deep voids of diameters greater than 100 mm (Fig. 20). Features indicating enhanced localized plasticity were a common observation in all the tests in HTW and were significantly higher in number as compared to air tests in hydrogen pre-charged condition (Figs. 18(b) and 19). This can be attributed to a continuous supply of hydrogen in water, whereas a fixed quantity of hydrogen is available at the beginning of the tests in air for the hydrogen precharged specimens and may be subsequently partially lost from fresh exposed crack surfaces at the test temperature. Macrovoids of similar dimensions has been reported earlier in pressure vessel steel in presence of hydrogen and has been attributed to preferential accumulation of hydrogen at bands of carbide segregation [49].

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Fig. 17. a) Representative plot showing unstable crack extension in HT1 (simulated CGHAZ) specimens during EPFM test in PWR environment at loading rate of 0.025 mm min1, b) ductile crack growth in HT1 showing unstable crack growth and c) comparison of J-Da curve of HT1 in air and HTW under otherwise identical conditions indicating a potential environmental effect on initiation toughness.

Fig. 18. Fractograph after fracture toughness test (288  C, air); a) specimen without hydrogen charging and b) hydrogen charged specimen showing ductile dimples, macrovoid (black arrow), cracks at voids (A), secondary cracks (B) and localized quasi-cleavage facet at an inclusion (C).

4. Discussion 4.1. Effect of hydrogen and strain rate on tensile properties Figs. 6 and 9 clearly indicate that hydrogen affects the plastic deformation behavior in tensile tests at both 25 and 288  C, which is possible only if interstitial hydrogen atom affected dislocation mobility. Furthermore, hydrogen reduces the formation energy of vacancies and facilitates their agglomeration/clustering and collapse to microvoids. This can result in softening due to the

increased climbing of edge dislocations or hardening by vacancy clusters and nanovoids. At 25 and 288  C, a marginal hardening and softening is observed, respectively. The corresponding effect of hydrogen on the ductility properties (strain at failure and reduction of area) is much more significant at 25  C and is independent of strain rate in the investigated strain rate range. At 288  C embrittlement only occurred in a narrow strain rate range with a maximum at 102 s1. In presence of hydrogen, a more shear dominated failure mode (45 inclination of fracture surface) is observed at both the test temperatures. Any softening will

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Fig. 19. (a) Typical fracture surface of a 1T-C(T) specimen after a test in BWR simulated conditions. Macrovoids (inset A), quasi cleavage fracture regions associated with inclusions (inset B) and localized flat quasi-cleavage facets (inset C) can be seen, b) ductile secondary cracks and c) typical fracture surface of a 0.5T-C(T) specimen tested in PWR simulated conditions showing macrovoids (inset A), secondary cracks (inset B) and quasi-cleavage (inset C).

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Fig. 20. a) Cross-sectional optical microscopic image after test B2 showing a macrovoid. Regions ‘A’ - fatigue precrack, ‘B’ - fracture surface during fracture toughness test and ‘C’ post-test fracture surface and a) secondary cracks indicating intergranular cracking.

dislocations in pile-ups and thus the local strain accumulation that facilitates micro-crack and microvoid formation. The shielding and drag effects are dependent on temperature, strain rate and strain. Shielding and drag effects are significant if the hydrogen diffusion rate (dependent on temperature) and dislocation velocity (dependent on strain rate) are comparable [26] which occurs for critical temperature-strain rate combinations only. At high temperatures or very low strain rates, there are no effects; since no hydrogen atmosphere at dislocation cores and strong trap centres are formed (hydrogen atom mobility exceeds dislocation mobility significantly). On the other hand, there are no such effects at very low temperatures or very high strain rates (e.g., in impact loading), since hydrogen cannot follow the dislocations (dislocation mobility exceeds hydrogen atom mobility significantly).

Fig. 21. Typical fracture surface after test in PWR simulated condition showing flat regions of quasi-cleavage surrounded by shear lip (black arrow) and elongated flat region showing tear/shear ridges (black arrow). Intermediate region shows ductile fracture due to micro-void coalescence. Crack growth direction is from top to bottom.

inevitably result in localization of plastic deformation. The microscopic localization of plastic deformation without hydrogen usually results in an increase of macroscopic yield stress [47]. In presence of hydrogen, the microscopic localization may result in an increase or decrease of the macroscopic yield stress depending on the degree of strain localization and magnitude of hydrogen enhanced dislocation velocity [47]. Because of its large molar partial volume and strong electronic interaction with the metal matrix, hydrogen preferentially accumulates in high hydrostatic stress and dilatation strain fields and at trap centres like interfaces of second phase particles or dislocation cores and may reduce the cohesive strength and increase the mobility of dislocations [26]. Due to its very high mobility, hydrogen atmospheres at the dislocation core can move with the dislocations and affect their interactions with other dislocations and obstacles [6,47]. In absence of mobile hydrogen atmosphere, the hydrogen behaves as any other weakly interacting solute. In steels various hydrogen-dislocation interactions were discussed in literature [47,48,51e54]: Hydrogen reduces the elastic interactions forces between dislocations and dislocation-obstacles and thus increases dislocation mobility (shielding effect, softening). Hydrogen atmospheres at (edge and mixed) dislocation cores can reduce the dislocation mobility similar as carbon and nitrogen in DSA (drag effect, hardening & softening). Hydrogen reduces the stacking fault energy and tendency for cross slip and thus increases planar, coarse slip and strain hardening in case of multiple active slip systems. It also reduces the distance between

4.1.1. Effects at 25  C After pre-charging, hydrogen is present as freely diffusible interstitial atoms and at various reversible and irreversible trap sites (immobile) and its distribution in the specimens is (macro- & microscopically) inhomogeneous [26]. At 25  C, hydrogen trapping at strong irreversible and deep trap centres like oxide and MnS inclusion-matrix interface, is very efficient and hydrogen atoms are predominantly trapped at these second phase particles and other traps like grain boundaries, thus resulting in a relatively low lattice bulk hydrogen content available for hydrogen mobile dislocation interactions. A part of the freely diffusible hydrogen atoms may also be partially lost despite the copper coating, due to fast diffusivity in steel. The reduction of ductility is mainly related to the formation of brittle QC features around oxide inclusions, which covered a significant proportion of the fracture surface. Contribution of hydrogen atom-dislocation interaction in reducing ductility was minimum because of which hydrogen effects were observed to be strain rate independent. The degree of embrittlement correlated fairly well with the amount of inclusions and QC features on the fracture surface [26]. The marginal hardening and increase in yield stress is probably related to the moderate hydrogen-induced localization of plastic shear deformation as confirmed by the presence elongated shear dimples. The slight increase of yield stress might also be caused by hydrogen that is trapped at the dislocation cores after pre-charging similar to static strain ageing that occurs in these steels at 25  C (upper and lower yield points or yield plateaus). 4.1.2. Effects at 288  C At high temperature the initial freely diffusible hydrogen present at interstitial bulk lattice positions after pre-charging is desorbed during heating to a greater extent than at 25  C. At 288  C, hydrogen trapping is not efficient and trapped hydrogen is

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increasingly released from the strong traps and diffusing out of specimen during the heating and tensile test phase. The mean diffusion distance x ¼ (2$D$t)1/2 of hydrogen is about 4e8 mm at 288  C in 1 h [26,52]. The time for heating and tensile test varied between 1.3 (102 s1) to 2 h (104 s1), a sufficient amount of lattice hydrogen is thus still available for interactions with the dislocations in the high-temperature tensile tests. Embrittlement is observed in a narrow strain rate range at 288  C only with a maximum at 102 s1. Maximum in embrittlement may be related to the matching of hydrogen atom and dislocation mobility at this strain rate. Further tests at lower temperatures (e.g., 200 and 250  C) would be necessary for verification. Since trapping is inefficient at 288  C, only very few isolated brittle QC features were observed at 288  C and the moderate embrittlement was mainly related to the hydrogen-induced localization of plastic deformation that was also confirmed by preliminary TEM studies (Section 3.1.3) and fractography (e.g., larger number of macrovoids, Section 3.1.2). Since more freely diffusible lattice hydrogen might be available with respect to room temperature tests, probability of dislocationhydrogen interaction is higher, as confirmed by the strain rate dependence, resulting in more significant shielding effect and localization of plastic deformation. The predominance of enhanced dislocation mobility in the strain localized regions and added shielding effect on DSA as compared to strengthening due to drag effects and strain localization (Fig. 9(a) and (b)) result in a slight net softening. A marginal softening and mitigating effect on DSA is observed. Besides shielding, the latter might be partially caused by a competition between hydrogen and carbon and nitrogen for the lattice places in the dilatation zone of the dislocation cores also. At the onset of yielding, shielding effects are moderate (Fig. 9(d)). With increasing strain, dislocation density increases and dislocations start to increasingly interact with each other and shielding and strain rate effects become more significant, e.g. on the reduction of area (Fig. 9(c)). Regions of (hydrogen-induced) high plastic deformation and strain localization are preferred locations for hydrogen accumulation which further leads to extensive localized plastic deformation. This results in earlier saturation of the capacity for uniform elongation (lower uniform strain) and earlier initiation of necking. Subsequently, microvoid formation occurs and hydrogen will result in enhanced localized plasticity in the intermediate regions of the voids. These regions become preferred sites for further hydrogen accumulation and enhanced hydrogen effects. Hydrogen effects will thus be more significant after necking, when plastic deformation is confined to a small volume of the material. This thus explains the more pronounced effect of hydrogen on the reduction of area than on yield stress or uniform elongation strain. Additionally, hydrogen reduces the formation energy of vacancies and thus facilitates their agglomeration/clustering and collapse to microvoids [55,56] and this may be a further contributing factor towards reducing ductility. Very high vacancy and hydrogen concentrations are observed in highly strain-hardened/ cold-worked regions that may affect the plastic deformation behavior (climbing of edge dislocations, vacancy clusters and nanovoids as dislocations obstacles) and enhance microvoid coalesce. Autocatalytic, positive feedback effects occur for both plastic strain localization and local hydrogen accumulation and for hydrogen enhanced strain-induced vacancies and local hydrogen accumulation. Furthermore, association of dislocations with hydrogen atoms results in dislocation transport of hydrogen atoms and it is reported to be much higher than lattice diffusion [54]. Hydrogen atoms associated with dislocations will easily reach susceptible locations such as inclusion (oxides, carbides, sulphides)/matrix interface,

bainite lath boundaries and locations of high localized stress/strain. This aid in build-up of local hydrogen concentration at susceptible locations and when local hydrogen concentration exceeds a critical value, the steel locally loses load bearing capacity resulting in embrittlement. 4.2. Role of hydrogen on fracture in high-temperature water The extension rate applied in this study during EPFM tests in HTW was few orders of magnitude higher than usually used to investigate EAC or corrosion fatigue, limiting EAC crack growth (unlike previous reported studies [32,33]). The air fatigue precracked C(T) specimens were exposed to HTW containing dissolved hydrogen (BWR/HWC, PWR), dissolved oxygen (BWR/NWC) or with additional electrochemical in-situ hydrogen charging (PWR) for usually 12 days (and up to 32 days). The C(T) specimens, when tested in HTW, pickup hydrogen (dissolved in water and corrosion generated) during the initial pre-oxidation stage and, at the plastically strained crack tip during subsequent EPFM loading. 4.2.1. Hydrogen pick-up and transportation to susceptible locations Hydrogen (dissolved in water and from corrosion reactions) is picked up on the outer surface of C(T) specimens exposed to the bulk environment and also inside the crack. However, there can be a significant difference in the water chemistry at these two locations (see Section 4.2.3), which significantly affects the amount of hydrogen picked up and its transport to crack-tip susceptible locations. The various sources of hydrogen on the specimen outer surface during the pre-oxidation stage are described below: 4.2.1.1. Hydrogen gas dissolved in water (from radiolysis or intentional additions). Hydrogen diffusion, permeation and release rates are high in RPV steels and oxide films do not represent a significant diffusion barrier. Hence, the bulk hydrogen content in steel is controlled by dissolved hydrogen content of water as per the Sievert’s law (hydrogen in steel ∝√fH2 (hydrogen fugacity)). Hydrogen oxidation reaction on oxidized steel surfaces is another potential source [57]. 4.2.1.2. Hydrogen from corrosion reactions. Hydrogen generation during corrosion/oxidation occurs as a result of the reduction reaction and oxide film formation. In deaerated conditions (BWR/ HWC, PWR, nitrogenated high purity water) water acts as the oxidizing agent and the reduction reaction occurs according to equation (3). In presence of dissolved oxygen, the reduction reaction occurs predominantly according to equation (4) [57,58].

2H2 O þ 2e / H2 þ 2OH

(3)

O2 þ 2H2 O þ 4e / 4OH

(4)

In HTW hydrogen uptake from uniform corrosion is proportional to the corrosion rate and exposure time and corrosion rates are usually (with exceptions) very low in HTW. The corrosion rate decreases with time and most corrosion/oxidation of steel occurs within the initial days of exposure. Furthermore, the corrosion rates in very high and very low dissolved oxygen water are quite similar. Due to very low corrosion and very high diffusion, permeation and release rates, the bulk hydrogen content close to the surface may become controlled by the hydrogen fugacity of the HTW only. Based on Fick’s second law of diffusion and a conservative lower bound value of the diffusion coefficient of hydrogen (DH) at 300  C (¼3$109 m2 s1), the hydrogen concentration at mid thickness of a 1T-C(T) specimen reaches ~90% of the hydrogen concentration at the surface in 12 days at 288  C, assuming no

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significant loss of hydrogen from the surface. Thus, significant effects of pre-oxidation duration is not expected in HTW for preoxidation periods greater than 12 days and up to 32 days, as observed in this study. Hydrogen generation and pick-up also occurs at the crack tip during pre-oxidation and, in particular, subsequently during EPFM test. Complete oxygen depletion occurs in the crack enclave because the oxygen consumption by corrosion reactions on the crack flanks is much faster than its transport by diffusion in the crack crevice, and as a result, the crack-tip is at low corrosion potential in oxygenated and hydrogenated HTW [58]. The crack tip environment thus has an aggressive occluded crack crevice environment (see Section 4.2.3) in the immediate vicinity of the plastically strained, bare and active crack-tip, and is at a low potential for all tested high-purity, HTW environments (both oxygenated and hydrogenated). Hydrogen uptake within cracks and its transfer to the process zone at the crack-tip is governed by the local occluded crevice environment, corrosion reactions and crack-tip mechanics and is favourable for hydrogen pick-up. As shown in Fig. 22, during the EPFM fracture mechanics test with standard loading rates, the protecting oxide film is continuously ruptured at the plastically strained crack-tip, exposing bare metal to the crack crevice environment. Fresh exposed bare metal undergoes anodic dissolution and subsequent hydrolysis of metal cations leads to hydrogen formation [58,59]. The oxide film formation on the crack flanks in the direct vicinity of the active cracktip further generates hydrogen. Additionally, weakly trapped hydrogen in the vicinity of a crack (e.g., in the plastic wake) can very quickly move by fast diffusion or dislocation transport to the moving crack-tip region with high hydrostatic stress, plastic strains and high concentration of deformation-induced vacancies as suggested in Ref. [60]. The hydrogen diffusion rate is two to three orders of magnitude faster than the typical crack growth rates in

361

EPFM tests of 107e106 m s1. Hydrogen thus can easily “follow” the advancing crack-tip. Exposure (by the growing crack) and dissolution of MnS inclusions (intersected by the crack front and enclave) in HTW at the crack tip is another potential source of hydrogen [12]. However, this may not be significant in the present low-sulphur steel or short-term EPFM tests with high loading rate (due to slow dissolution rate of MnS inclusions). The absorption of sulphur anion species from the dissolution of MnS-inclusions or as bulk environment impurities on the bare metal surface may act as recombination poison and also retard the reformation of the protecting oxide films aiding hydrogen pick-up [12]. Hydrogen thus picked up at the crack tip will reach susceptible locations (MnS and oxide inclusions, intersections of slip bands, PAG boundaries) within the high stress triaxiality region ahead of the crack-tip easily. This is because of significantly reduced diffusion distance and stress gradient driven hydrogen diffusion or accelerated transport of hydrogen by mobile dislocations [50] in the crack-tip plastic zone. MnS-inclusions in the region of maximum hydrostatic stress ahead of the crack-tip may act as strong hydrogen traps and location for brittle crack initiation [12,61]. Thus, there is a continuous source of hydrogen which can enter the steel during the EPFM tests in HTW, and significant local hydrogen accumulation can be expected to occur under these dynamic crack-tip plasticity conditions. As soon as the local hydrogen concentration in the process zone reaches a critical value over a critical volume, brittle or ductile/shear crack initiation in the process zone at strong trap centres may occur by local hydrogen embrittlement mechanisms (as discussed in Section 4.2.5 and Fig. 22). 4.2.2. Effect of strain rate Accumulation of hydrogen in the crack tip process zone could potentially reduce fracture toughness depending upon the EPFM

Fig. 22. Localized hydrogen uptake, hydrogen-deformation interactions and hydrogen embrittlement in the crack-tip system during an EPFM test.

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testing parameters. Although both JQ/JDCPD and tearing resistance were not reduced in HTW in the investigated range of test parameters, there is a clear change in fracture morphology which is similar to fractographic features after tensile and fracture toughness tests with hydrogen pre-charged specimens. Hydrogen effects on tensile properties at 288  C only occurred in a narrow strain rate range with a peak at around 102 s1 (Fig. 9). The crack-tip strain rate in C(T) specimen (before initiation) was calculated as per earlier reported studies [62] to be ~7$104e103 s1, for the extension rates applied during the EPFM tests (0.25e0.35 mm min1). The calculated crack-tip strain rate is thus an order of magnitude lower than the strain rate of maximum hydrogen effects at 288  C, thus potentially reducing the embrittling effects of crack tip hydrogen accumulation. Further testing at higher loading rates is thus needed. 4.2.3. Effect of water chemistry (PWR/BWR) on fracture toughness The dissolved hydrogen contents in crevice environment are similar to the bulk hydrogen contents; since relatively little dissolved hydrogen is consumed in the crack crevice and due to its fast diffusion rate. The pH of the occluded crevice solution in the moderately alkaline PWR environment is only slightly higher (1.5e2 pH unit) than in the neutral, high-purity bulk BWR environment. Moderate acidification (and thus higher hydrogen uptake) may occur in oxygenated water in case of presence of Cl or SO2 4 impurities or due to dissolution of MnS-inclusions (maximum acidic pH-shifts of 1e1.5 units) and potentially very slight alkalization in case of high purity water (PWR primary water) and very low-sulphur steels [58]. The crack-tip environment conditions in high-purity water, which govern the corrosion reactions and hydrogen uptake at the crack-tip in EPFM tests, are thus not fundamentally different for all environments. Different environments (BWR/HWC, BWR/NWC, PWR, nitrogenated HTW) and test conditions (pre-oxidation duration, loading during pre-oxidation, in-situ hydrogen charging) showed similar effects on fracture in this study, which suggests that hydrogen uptake at the crack-tip in the EPFM tests is probably dominated by the (local) corrosion reactions and crevice environment during the EPFM test. 4.2.4. Effect of microstructure EPFM (Fig. 17) and tensile test results (Fig. 9) clearly indicate that hydrogen or environmental effects are more significant in the high strength simulated CGHAZ material HT1 than in the RPV base metal AR. Additionally, tensile test results indicate that the strength of the steel is a more important criterion affecting susceptibility to hydrogen embrittlement than the PAG size, for the experimental conditions used in this study. This quite common observation is explained as follows: Due to its large molar partial volume, hydrogen tends to accumulate in high hydrostatic stress & dilatation strain fields and at trap centres. The higher yield stress results in a higher peak and hydrostatic stress in the process zone ahead of crack-tip (~3 sYS under plane strain conditions) and thus facilitates enhanced hydrogen accumulation in this region. Furthermore, due to the smaller plastic zone size with increasing yield stress, the available hydrogen is concentrated at smaller volume and the diffusion distances for hydrogen become significantly shorter. 4.2.5. Mechanism of hydrogen effects in high temperature water As soon as the local hydrogen concentration in the process zone reaches a critical value over a critical volume, ductile/shear or brittle crack initiation in the process zone at inclusions or strong trap centres may occur by various local hydrogen embrittlement mechanisms (Fig. 22). Provided that there is sufficient dissolved hydrogen, it enables the embrittling mechanisms that the microstructure and stress/strain levels “permit” [55]. The magnitude of

embrittling effect is finally governed by the nature, source and availability of hydrogen, its transport to susceptible locations and the nature and concentration of trap centres and defects in the susceptible location. The hydrogen embrittlement mechanisms that may occur in RPV steels in HTW are [12,63,64]: 1. Hydrogen-enhanced decohesion embrittlement (HEDE): According to this mechanism hydrogen atoms weaken the interatomic bonding leading to brittle or QC micro-cracks at strong traps like MnS/oxide inclusions or intergranular cracking at grain boundaries in combination with grain boundary carbides or metalloid segregation. 2. Hydrogen-enhanced local plasticity (HELP): According to this mechanism hydrogen reduces the elastic interactions forces between dislocations and dislocation-obstacles (shielding effect) and the stacking fault energy restricting cross slip and favouring shear localization. 3. Hydrogen-enhanced strain-induced vacancies (HESIV) [54,55]: Hydrogen atoms reduce the formation energy of vacancies and facilitate their agglomeration/clustering and collapse to microvoids. Very high vacancy and hydrogen concentrations in highly cold-worked regions, e.g. at the crack-tip, affect the plastic deformation behavior (climbing of edge dislocations, vacancy clusters and nanovoids as dislocations obstacles) and enhance microvoid coalescence. 4. Adsorption-induced dislocation emission (AIDE) [64]: Hydrogen atoms facilitate nucleation of dislocations at surface (crack-tip, voids) and their emission into the plastic zone. Combinations of 1e4 are possible in RPV steels under LWR conditions and their relative contributions depend on variables such as temperature, strain rate and microstructure. Mechanism 2 to 4 are not only enhancing ductile failure by microvoid coalescence (MVC), but also facilitating brittle cleavage or quasi-cleavage cracking by various mechanism (such as due to dislocation pileups), at interfaces such as second phase particles or grain boundaries or along preferred highly active slip planes and their intersections. Mechanism 1 (HEDE) may explain the quasi-brittle features around oxide and MnS-inclusions observed in our EPFM and EAC tests in HTW [12]. 4.2.6. Concluding remarks In summary, EPFM tests in HTW did not show a significant effect on initiation toughness JQ and tearing resistance for RPV base metal AR and showed limited effects in HT1 microstructure in the investigated parameter range. A change in fracture morphology clearly demonstrates the need for further tests over a broader range of material, hydrogen level and loading conditions for definite conclusions on potential hydrogen effects in air and HTW. Thus, there is enough experimental (e.g., fractography) and theoretical evidence that hydrogen effects may occur in RPV steels in the LWR temperature range. More pronounced effects are expected with more susceptible materials (high sulphur steels having high EAC susceptibility, CGHAZ with high yield stress and high plastic welding strains, steels having high DSA susceptibility), at lower temperatures (with stronger effects of hydrogen) or under suitable critical loading/ strain rate conditions. Synergies (or competition) with other embrittlement effects in RPV steels like dynamic strain ageing (DSA), temper embrittlement (TE) or irradiation embrittlement (IE), and in particular, the potential for shifts in the ductile to brittle transition temperature should be further evaluated. Tests at lower temperatures are indispensable, since the peak stress intensity factors in loss of coolant accidents for different (postulated) incipient cracks in the RPV are usually reached in the temperature range

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from 100 to 200  C.

363

Baumgartner, D. Stammbach, R. Schwenold, J. Chen and H. Wiese from PSI.

5. Summary and conclusions The effect of hydrogen and high-temperature water (HTW) on the mechanical and fracture behavior of a low-alloy RPV steel was evaluated by tensile and EPFM tests in air with electrochemical precharging and EFPM experiments in hydrogenated and oxygenated HTW in refreshed HTW loop autoclave systems, respectively. The tests were complemented by detailed post-test evaluations on the fracture and deformation mechanism by optical, scanning and transmission electron microscopy. Based on the results, the following conclusions can be made: i) 2e5 wppm of hydrogen in the RPV steel used in this study caused a marginal hardening and softening (yield stress) in tensile tests in air at 25 and 288  C, respectively, and a significantly more pronounced reduction in ductility (reduction of area). The resulting embrittlement was more significant at 25  C, at higher hydrogen contents in the steel and in the simulated coarse grain weld heat-affected zone material with higher yield stress. ii) At 288  C, a strain rate dependent embrittlement was observed with a maximum at a strain rate of ~102 s1 that was attributed to the matching of dislocation and hydrogen atom mobility. iii) During EPFM tests (LLD rate of 0.25e0.35 mm/min) in air at 25  C after 24 h hydrogen pre-charging, unstable brittle cracking and a moderate reduction in initiation toughness was observed for the base metal (AR). Under the same conditions at 288  C, stable ductile crack growth with no reduction in initiation toughness and tearing resistance occurred. iv) Exposure to HTW (BWR/HWC, PWR, BWR/NWC, 288 or 150  C) did not reduce the initiation toughness and tearing resistance of the base metal (AR). Additional electrochemical in-situ hydrogen charging (PWR) for 12 days (up to 32 days) did not affect the toughness properties probably due to inefficient cathodic polarization as a result of high electrolyte resistance. v) A clear change in fracture morphology and deformation structures was observed after tests in HTW that was very similar to that in EPFM and tensile tests in air with hydrogen pre-charging. vi) Hydrogen and HTW effects on fracture morphology are related to extensive (hydrogen-induced) localization of plastic deformation and not to a microscopically brittle process. vii) Experimental and theoretical evidence clearly indicate that hydrogen effects may occur in RPV steels in the LWR temperature range. More pronounced hydrogen and HTW effects on the fracture behavior can be expected with more susceptible materials (high strength, coarse grain size), at lower temperatures or under suitable critical loading/strain rate conditions. Further systematic studies are thus recommended. Acknowledgment Funding from the European Community’s Seventh Framework Programme (FP7/2007e2013) under grant agreement no. 290605 (PSI-FELLOW), and from the SAFE-I & -II project from the Swiss Federal Nuclear Safety Inspectorate (ENSI) is gratefully acknowledged. The authors would like to express their gratitude for the experimental contributions of S. Ritter, S.G. Rao, L. Nue, B.

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Abbreviations and symbols AC: Air cooled AIDE: Adsorption-induced dislocation emission AR: As received ASTM: American Society for Testing and Materials Standards BWR: Boiling water reactor CGHAZ: Coarse grain heat affected zone DCPD: Direct current potential drop (mV)

DSA: Dynamic strain aging EAC: Environmental assisted cracking EDS: Energy dispersive X-ray spectroscopy HAZ: Heat affected zone HEDE: Hydrogen-enhanced decohesion embrittlement HELP: Hydrogen-enhanced local plasticity HESIV: Hydrogen-enhanced strain-induced vacancies HTW: High-temperature water HWC: Hydrogen water chemistry IR: Irradiation embrittlement LAS: Low-alloy steel LLD: Load line displacement at specimen (¼crack opening displacement) LWR: Light water reactor MVC: Microvoid coalescence NWC: Normal water chemistry PAG: Prior austenite grain PWR: Pressurized water reactors QC: Quasi cleavage RA: Reduction in cross-sectional area (%) RPV: Reactor pressure vessel SEM: Scanning electron microscope TE: Temper embrittlement TEM: Transmission electron microscope E: Young’s modulus (GPa) sUTS: Ultimate tensile strength sYS: Yield stress 

C s25 : Flow stress at 25  C (MPa) y C s288 : Flow stress at 288  C (MPa) y

a: Crack length (mm) ao: Initial crack length (mm) a/W: Ratio of crack length (a) and C(T) specimen width (W) B: C(T) specimen thickness (mm) BN: C(T) specimen net section thickness (mm) bo: Initial uncracked ligament, bo ¼ (W  ao) (mm) C: Concentration of hydrogen in steel CS: Surface concentration of hydrogen D: Diffusion coefficient (m2 s1) J: Fracture toughness parameter defined in ASTM E 1820 (kN m1) JDCPD: J at initiation of ductile crack growth measured from DCPD signal (kN m1) Jel: Elastic component of J integral (kN m1) Jlimit: Maximum allowable J integral of C(T) specimen (kN m1) Jpl: Plastic component of J integral (kN m1) JQ: Fracture toughness at initiation of stable crack growth (kN m1) at Da ¼ 0.2 mm ¼ JIC, if all 15 requirements of ASTM E1820 are fulfilled K: Stress intensity factor (MPa m0.5) Kmax: Stress intensity factor (K) at maximum load during cyclic loading (MPa m0.5) Kmin: Stress intensity factor (K) at minimum load during cyclic loading (MPa m0.5) wppm: Parts per million by weight pH: log([Hþ]), where Hþ is hydrogen ion concentration in water pH288

C

: pH of water at 288  C

R: Ratio Kmin/Kmax VH1: Vickers microhardness at load 1 kg W: C(T) specimen width (mm) dK/dt: Rate of change in stress intensity factor during fracture toughness test (MPa m0.5 h1) dJ/dt: Rate of change in J integral during fracture toughness test (kN m1 h1) 1T-C(T): Compact tension specimen of 25 mm thickness (B) 0.5T-C(T): Compact tension specimen of 12.5 mm thickness (B) DK: Kmax - Kmin(MPa m0.5) n: Poisson’s ratio εf: Total strain to failure Da: Increase in crack length (m) Dalimit: Maximum allowable increase in crack length (m)