EFFECT OF HIGH-TEMPERATURE WATER AND HYDROGEN …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG034.pdftoughness due to a significant role of EAC at the test extension rates

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

EFFECT OF HIGH-TEMPERATURE WATER AND HYDROGEN ON THE FRACTURE BEHAVIOUR OF A LOW-ALLOY REACTOR PRESSURE VESSEL

STEEL

S. Roychowdhury1, 2, H.P. Seifert1, P. Spätig1, Z. Que1, S. Ritter 1 1Paul Scherrer Institut, Nuclear Energy and Safety Research Department, Laboratory for Nuclear

Materials, 5232, Villigen, PSI, Switzerland 2Materials Science Division, 2nd Floor, Mod-Lab, D-Block, Bhabha Atomic Research Centre,

Mumbai - 400 085, India

ABSTRACT

The effect of high-temperature water and hydrogen on the fracture behaviour of a low-alloy RPV steel with two different microstructures - bainitic base metal (BM) and simulated coarse grain heat-affected zone (CGHAZ) of welds, was evaluated in this investigation. Tensile properties were characterized in air at different strain rates, temperatures and hydrogen levels. Elastic-plastic fracture mechanics (EPFM) tests, followed by fractography and characterization of deformation mechanism, were also performed. The EPFM tests were performed at different temperatures in i) air, with and without hydrogen pre-charging and ii) in hydrogenated and oxygenated high-temperature water in refreshed autoclave systems. The majority of testing was performed with the BM.

Hydrogen in the range of 2 to 5 ppm in the RPV steel resulted in embrittlement in tensile tests, both at 25 and 288 °C. The embrittling effect was most significant at 25 °C and strain rate dependent. Maximum embrittlement was observed at strain rates of 10-5-10-4s-1at 25 °C and at 10-3-10-2s-1 at 288 °C. Embrittlement is related to a strong localization of plastic deformation and not due to a brittle microscopic process (for e.g., cleavage).

In contrast to tensile tests, hydrogen pre-charging did not reduce the initiation toughness and tearing resistance at 288 °C in the bainitic base metal in air in the investigated loading rate range, although a clear change in fracture morphology was observed. At 25 °C, unstable, brittle crack extension was observed after severe hydrogen pre-charging, at a relatively high J value without prior stable crack growth. Under the same loading conditions, neither oxygenated nor hydrogenated high-temperature water resulted in a reduction of initiation toughness and tearing resistance at 150 and 288 °C so far, with respect to corresponding air tests. Although, a clear change in fracture morphology and deformation structures was observed that was very similar to that in EPFM and tensile tests in air after hydrogen pre-charging. The observed change of fracture and deformation modes were not sufficient to significantly modify the fracture behavior at 288 °C, but this effect of hydrogen could change for more susceptible materials, lower temperatures or critical loading/strain rate conditions.

Keywords: Low Alloy Steel, Hydrogen Embrittlement, Fracture Toughness

1. INTRODUCTION

Structural integrity of the reactor pressure vessel (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 [1-3]), concerns due to detrimental environmental effects remain that can be attributed to hydrogen uptake from coolant and/or corrosion reactions [4]. Direct contact of RPV steel and coolant occurs via cracks in the protective austenitic stainless steel cladding formed during fabrication [5], or during reactor operation [6]. Furthermore, the presence of open, incipient, cracks has always to be assumed in safety analysis [7]. The RPV steel picks up hydrogen from the LWR coolant (hydrogen from radiolysis and intentional additions) and corrosion reactions and reaches equilibrium bulk concentrations


of up to a very few ppm within a few weeks or months at 300 °C, which is high enough to affect their mechanical and fracture properties [8-12]. Detriment in mechanical properties of base metal and heat affected zone (HAZ) of RPV steels have been reported at hydrogen concentrations (2.5-3.5 ppm) similar to that picked up in the same material, after exposure to primary water in a nuclear reactor (1-5.8 ppm) [12]. Furthermore, crevices/cracks with aggressive occluded crevice chemistry will result in even higher concentrations of absorbed hydrogen, irrespective of bulk coolant chemistry. 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 [8-11, 13]. Though calculations indicated hydrogen in RPV to be benign at reactor operating temperatures [14], there is experimental evidence indicating the contrary [12, 13].

Hydrogen effects on tensile behaviour of low alloy steel (LAS) are strongly influenced by temperature, microstructure, and strain rate [15-20]. Most of the reported studies were done at room temperature, or below, where the effects are most significant. There are a limited number of test results reported indicating embrittling effects of hydrogen at temperatures in the range of 250-288 °C [19, 21, 22]. The weld HAZ, specifically the high hardness, coarse grain HAZ (CGHAZ), is highly susceptible to hydrogen effects [16, 17, 23]. Hydrogen effects on EAC behaviour of LAS in high temperature water, and synergistic (or competitive) effects with dynamic strain aging (DSA), are well recognised [2, 24-26]. These synergistic effects can have a greater significance for weld HAZ material.

A few earlier studies have reported detrimental effects on the fracture of LAS due to exposure to 288 C water [13, 27, 28]. Environmental effects in water at 288 C were reported to be manifested as occurrence of a greater number of pop-in events attributed to synergy between corrosion generated hydrogen and DSA, but without changing the initiation fracture toughness [28]. In another study, a trend of reduction in fracture toughness and crack growth resistance as extension rates were decreased was noted for tests in 288 C water and this was attributed to the high sulphur content of the steel [27]. Elongation rates used for the fracture toughness tests [27, 28] were low and the measured toughness and crack growth resistance reduced with reducing extension rates [27]. This is an indication of possible apparent reduction in toughness due to a significant role of EAC at the test extension rates. Moreover, there are no reported fractograph after fracture toughness tests using LAS in water at 288 C.

In the present investigation, the effect of hydrogen and high-temperature water on tensile and fracture properties of low-alloy RPV steels at different temperatures, strain rates and microstructures was investigated. Detailed post-test characterisation was done to understand the influence of high-temperature water and hydrogen on tensile and fracture properties at elevated temperatures. The present paper summarises the firstresults from an on-going project which aims to identify critical combinations of microstructure (including simulating weld HAZ microstructures), environmental factors (hydrogen level, temperature) and loading conditions (strain rate), that result in significant high-temperature water and hydrogen effects in the LWR temperature regime. The present paper is an update of a previous publication [11]. It is stressed that effects of high-temperature water on fracture toughness were massively overestimated in that publication due to a systematic error in the evaluation procedure of the EPFM tests and were not confirmed after proper evaluation.

2. EXPERIMENTAL PROCEDURE

2.1 Material and Specimens

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 [29]. This material is from an actual pressurized water reactor (PWR) RPV that was not commissioned. The steel used has been characterized as being moderately susceptible to DSA, and of low EAC susceptibility by previous investigations by the Paul Scherrer Institut [24, 30]. Tensile properties of this steel are listed in Table 1, along with the chemical composition, as previously reported in [30]. The steel was quenched


(870-905°C/6.9h/WQ) and tempered (635-655°C/11.3h/air cooled) in as-received condition, having a bainitic microstructure with prior austenite grain size in the range 10-20 m and hardness in the range 261+4 HV0.1.

The as-received (AR) base metal, a simulated, high-hardness (HT1) coarse grain heat-affected zone (CGHAZ) and low hardness CGHAZ (HT2) were used and produced by a specific 2-stage heat treatment. The heat treatments were done in argon atmosphere to prevent oxidation and Table 2 lists the properties in the heat treated condition. Both the heat treatment stages were followed by microhardness measurement and microstructural evaluation. The specimens were first given a grain coarsening heat treatment at 1000+5 °C (HT) for one hour, followed by water quenching (WQ), which resulted in a prior austenite grain size of ~50 m. Rapid grain growth in pressure vessel steels at 1000 C [31, 32] has been reported earlier. The microhardness increased from 261 to 549 HV0.1 after heat treatment HT (1000+5 °C/1 h/WQ) due to martensitic transformation. Subsequently, the specimens were tempered at 625+5 °C for 1 hour (HT1) and for 72 hours (HT2) followed by air cooling (AC). This resulted in a tempered martensite microstructure, reducing the microhardness to 344 HV0.1 (HT1) and 250 HV0.1 (HT2). Similar approach of heat treatingsteels to simulate weld HAZ has been reported earlier [33, 34].

The microstructure and the microhardness after the heat treatment HT1 are similar to that of the high hardness CGHAZ of the actual circumferential core girth weld of the PWR RPV[30]. The difference in hardness after HT1is less than 10% of the hardness of the actual weld CGHAZ (320 HV0.1 [30]). The heat treatment HT2 aimed at to have similar microstructure and prior austenite grain size as HT1 but with strength levels comparable to the AR base metal. This was done to separately evaluate the effects of strength level and grain size on measured tensile/toughness properties.

Cylindrical tensile specimens of transverse orientation,of gauge length 30 mm and gauge diameter 6 mm were used for tensile tests. 1T-C(T) (B=25 mm) and 0.5T-C(T) (B=12.5 mm) specimens (W/B=2) specimens of TL orientation were used for elastic plastic fracture toughness (EPFM) tests in air and in high temperature water. The C(T) specimens were fatigue pre-cracked in air to an initial crack length a/W of ~0.5. Load during final stage of pre-cracking was controlled to keep Kmax value at 20.7 MPa·m0.5 and 13.6 MPa·m0.5 for 1T-C(T) and 0.5T-C(T) specimens, respectively. Subsequent to fatigue pre-cracking, all C(T) specimens were side-grooved to a thickness reduction of 10% on each side. Data analysis was done according to the standard ASTM E1820 [35].

2.2 Hydrogen Charging Procedure for Tests in Air

Cathodic hydrogen charging in galvanostatic mode was done at room temperature (25 °C) in 1N sulphuric acid + 30 mg·l-1 arsenic trioxide solution. Before charging the specimen surface was metallographically polished up to 1000 grit emery paper, followed by ultrasonic cleaning in acetone for 5 minutes. Gauge section of tensile specimens were cathodically charged for 1 and 7 hours at a current density of 10 mA·cm-2 and the entire C(T) specimen was charged for 7 or 24 hoursat a current density of 0.3 mA·cm-

2. Specimen surfaces were polished with 1000 grit emery paper subsequently and copper was electrolytically deposited [19]. Starting material (AR) is not expected to contain any hydrogen and will be referred to as containing 0 ppm hydrogen. Hydrogen analysis was done after tensile testing at 25 C, for two AR specimens and for one HT1 and HT2 samples.The specimens were stored in liquid nitrogen immediately after tensile testing to prevent hydrogen effusion. The hydrogen content was analysed to be 2.6+0.4 ppm (parts per million by weight) after 1 h charging and 5.5+0.6 ppm after 7 h charging, for all the three microstructural conditions. No effect of microstructure on the hydrogen content was observed due to the limited number of specimens analysed.The measured hydrogen contents represent a realistic lower bound of theconcentration present in the fracture zone during the 25 C tensile testing.The hydrogen content in the high-temperature tensile tests is expected to be lower due to partial effusion of trapped hydrogen during heating. Most of the hydrogen is expected to effuse from the uncoated fracture surface during cooling-down from test temperature. Hence, hydrogen analysis after high temperature tensile tests was not attempted. Hydrogen content of C(T) specimens wasnot analysed since it is expected


to be inhomogeneous during and after tests at 288 C. Hydrogen would concentrate primarily at crack tip under high tri-axial hydrostatic stress.

2.3 Tensile Tests in Air

Specimens were tested immediately after copper deposition and the tests were performed in air at room temperature (25 °C) and 288 °C, at strain rates varying from 10-5 s-1 to 10-1 s-1. 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 80-90 minutes and the specimens were held at a constant load of 200 N at 288 °C for additional10 minutes, prior to start of loading. All the tests were repeated 2 (AR, 0 ppm hydrogen) to 3 (hydrogen pre-charged) times under identical test conditions. Subsequently, detailed fractographic examination was done in scanning electron microscope.

2.4 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 pre-charged C(T) specimens were tested after 24 hours to allow for through thicknesshomogenization of hydrogen in the crack-tip region. Load-line displacement (LLD) was recorded using a clip gauge and tests were done at an extension rate of 0.25 and 0.35 mm·min-1 for 0.5T-C(T) and 1T-C(T) specimens, respectively. EPFM tests and data analysis was done according to ASTM E 1820 [35]. 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 [35]. The unloading compliance technique usually slightly underestimated the final crack length by 5%. A few tests in air without hydrogen pre-charging were done without unloading and crack length measurement by the reversed direct current potential drop (DCPD) method. Table 3 lists the experimental details and results for the tests done in air.

2.5 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, 36]. Reducing BWR hydrogen water chemistry (HWC) was simulated by hydrogenated (dissolved hydrogen content of 1.4 ppm) and nitrogenated neutral high-purity water (pH288°C= 5.7, inlet conductivity = 0.055 S·cm-1) and oxidising normal water chemistry (NWC) were simulated by oxygenated (dissolved oxygen content of 2 ppm)neutral high-purity water. The corresponding electrochemical corrosion potentials (ECP) of the specimens at 288 °C were -590 mVSHE (HWC/nitrogenated water) and +100 mVSHE (NWC). PWR primary water was simulated by mildly alkaline borated and lithiated, hydrogenated high-purity water (1000 ppm B as H3BO3, 2.3 ppm Li as LiOH, pH288 °C= 6.9, inlet conductivity of 24 S·cm-1, dissolved hydrogen content of 1.8 to 1.9 ppm).The corresponding ECP at 288 °C was -735 mVSHE. In some experiments in PWR simulated conditions, additional in-situ hydrogen charging was done at a current density of 3 mA·cm-2 inside the autoclave at 288 °C for entire duration of the test, which shifted the corrosion potential to more negative values.

Table 4 lists the experiment details and results for the tests done in high temperature water. Tests A1 to A5 were carried out in PWR simulated conditions, B1 to B5 were carried out in BWR/HWC and test B6 and B7 were done in BWR/NWC simulated conditions and oxygen- and hydrogen-free high-temperature water, 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 m and crack length was calculated using Johnson’s equation [2] (see also Section 2.5.3).

2.5.1 Pre-oxidation inHigh Temperature Water

Fatigue pre-cracked (in air) specimens were initially pre-oxidised 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, K 9.9 MPa·m0.5, minimum load/maximum load (R) = 0.7, 103 cycles, ~12 days) to have an active growing EAC or the specimens were under low constant load (no crack growth, duration of exposure ~12 days). The EAC growth rates under cyclic load (~1.2·10-10–2.2·10-10 m·s-1) were 2 to 3 times higher than the corresponding fatigue crack growth rates in air (~8.0·10-11m·s-1) and correspond to the low sulphur line of the Ford and Andresen model [2, 24]. 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 very small load amplitude at a K of 2.9 MPa·m0.5, below the fatigue and EAC crack growth threshold and hence did not induce any active EAC growth. Both pre-oxidationperiod and cyclic loading/EAC crack growth could influence hydrogen accumulation and initial content in the crack-tip process zone affecting the subsequent fracture behaviour.

2.5.2 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.Control measurements with clip gauges confirmed the adequacy of this approach. The applied loading rate for 0.5T-C(T) was 0.25 mm·min-1 (dK/dt-1·102 - 2·102 MPa·m0.5·h-1) and for 1T-C(T) 0.35 mm·min-1 (dK/dt - 3·102 - 5·102 MPa·m0.5·h-1), 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·m-1·h-1 for both types of C(T) specimens. The resulting crack-tip strain rates were estimated to be in the range of 10-3 s-1. Compliance technique for crack length measurement was intentionally not used to prevent any additional EAC effects during loading/unloading [28].After the tests, specimens were immersed in liquid nitrogen, broken open and were used for fractographic examination.

The crack length was calculated from DCPD signal using Johnson equation. Initial fatigue or corrosion fatigue pre-crack length (a0) was measured from post-test fractography and was assigned as the point of initiation of stable crack growth, as described in Figure 1. The crack extension was then linearly corrected by post-test fractography total crack extension measurements. The calculated crack extension slightly overestimated the actual crack extension by 5 % and the accuracy of DCPD is thus much better than that of the unloading compliance method. The relative error of the corresponding crack lengths from the load line was much smaller (< 1 %).

The methods for the determination of the J-a curve, blunting line and JDCPD at the onset of ductile crack growth are also illustrated in Figure 1.The blunting line and apparent crack advance ai due to blunting were calculated with JDCPD and the flow stress Y (Y = (YS+UTS)/2). After onset of ductile crack growth, the J values [35] were calculated with the load, LLD and corresponding crack length, at an LLD interval of 0.1-0.2 mm,by the incremental method of ASTM E1820 [35]. The compliance change of the specimen with crack growth was considered during the calculation. J-a curves were plotted using the calculated J values, where a = aDCPD + ai is the increase in crack length during loading.

2.5.3 J at Initiation of Stable Crack Growth

J values at initiation of stable crack growth were calculated by two different techniques. For all the tests J at initiation (JQ) was measured as per the ASTM standard [35], where JQ is the J value corresponding to the point of intersection between the regression line, J = C1·aC

2and the 0.2 mm exclusion line (line parallel to blunting line and offset from origin by 0.2 mm, slope of blunting line = 2·y

288°C for high


temperature tests and = 2·y25°C for 25 C tests), as shown in Figure 2b. For tests using DCPD signal for

crack length measurements, additionally JDCPD at onset of stable crack growth was measured. JDCPD was calculated at the instant when the slope of DCPD potential drop vs. LLD curve changes (Figures 1 & 2) during monotonic loading, in accordance with earlier studies [2]. The JQ values measured from unloading compliance and by DCPD were very similar, the latter usually being slightly smaller. The JDCPD values on the other hand were usually 35 % (25 to 50 %) smaller than the corresponding JQ values. DCPD technique has been successfully used for plotting J-a curves and estimating JQ for a variety of materials with accuracy [43-47] and the crack length measured by both the techniques in this study have also been reported to be comparable [43, 44].

3. RESULTS AND DISCUSSION

3.1 Tensile Test Results

3.1.1 Tensile Behaviour at 25 °C

Tensile tests were performed using specimens having AR microstructure, at strain rates in the range 10-5-10-3 s-1. Figure 3(a) and (b) summarises the tensile properties and clearly indicates that 5.5 ppm hydrogen affects the tensile properties (YS, reduction in area) to a greater extent at the strain rates of 10-5-10-4 s-1 and the effects become less significant at strain rate of 10-3 s-1. There is an increase in YS (Figure 3(a)) and greater reduction in ductility (RA Figure 3(b)) at strain rate 10-5-10-4 s-1. It is reported earlier that hydrogen does not affect the UTS at 25 °C significantly [19].

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 Figure 4(a). In presence of hydrogen, fractography showed regions of ductile fracture due to MVC, quasi-cleavage (QC) and brittle fracture as shown in Figure 4(c-d). Brittle and QC facets were observed to be always associated with inclusions, confirmed by energy dispersive spectroscopy to be predominantly oxides of aluminium. Figure 4(c-d) shows the typical fracture surface after tensile test at strain rate 10-4 s-1 for specimens containing 5.5 ppm hydrogen. In presence of hydrogen, fracture was observed to be dominated by shear (ductile fracture), as indicated by 45° inclination of the plane of fracture to the loading direction (schematic in Figure 4(e)) and by elongated shear dimples (Figure 4(d)). Regions of QC (Figure 4(b)) and brittle fracture (Figure 4(c)) were perpendicular to the loading direction (regions marked “A” and “B” in Figure 4(c)).

3.1.2 Tensile Behaviour at 288 °C

Tensile tests at 288 °C were performed using AR and HT1 specimens, at strain rates 10-3-10-1 s-1, and at 10-2 s-1 for HT2 specimens. DSA may occur in LAS at 288 °C and cause strengthening (increase in UTS) and reduce ductility [19, 37]. The negative strain rate sensitivity shown in Figure 5(a) and (b) confirms the occurrence of DSA, in both AR and HT1 specimens, irrespective of hydrogen content. Hydrogen causes softening, as indicated by the reduction in the UTS in AR steel (Figure 5(a)).

In AR specimens containing 5.5 ppm hydrogen, the mean value of UTS drops by 1-4 MPa in the strain rate range of 10-2-10-1 s-1, where as in HT1 specimen a higher drop in UTS was measured (19-21 MPa), for a lower hydrogen content (2.6 ppm hydrogen), at the same strain rate. For HT1 specimens with 5.5 ppm hydrogen, the UTS value measured at a strain rate of 10-2 s-1 was 578 MPa, a drop in UTS of approximately 268 MPa (UTS for HT1 with 0 ppm hydrogen was measured to be 846 MPa).Thus, hydrogen effects become significant at lower concentrations in steel as grain size and strength increases (Figures 5(a) and (b)). In AR and HT1 specimens, the effect of hydrogen on RA is more significant at strain rate of 10-2 s-1 (Figures 5(c)). HT2 does not show significant hydrogen effects as compared to HT1 for the same hydrogen level, as seen in Figures 5(c), indicating that strength affects hydrogen embrittlement susceptibility to a greater extent than the grain size for hydrogen levels used in this study.


Hydrogen resulted in shear dominated fracture at 288 °C, for all the specimens (AR, HT1 and HT2), as indicated by ~45° inclination of the fracture surface to the loading direction (Figure 6 (a, b)) and elongated dimples as shown in Figure 6 (c) (region marked “C” in Figure 6(b)). In addition, other features like macrovoids greater than 100 m in size and secondary cracks parallel to the loading directions can be seen in Figure 6(d). Fracture was predominantly ductile due to MVC though QC facets associated with inclusions were observed in few of the experiments and was not as common as at 25 °C. Regions of “QC” and secondary cracks along with flat featureless regions (region “B” in Figure 6(b)) indicating extensive localized plastic flow of the material.

3.1.3 Effect of Hydrogen and Strain Rate on Tensile Properties

Interstitial hydrogen atoms have mobility due to diffusion, which is several orders of magnitude higher than other interstitial solute atoms such as carbon or nitrogen, due to which it cannot effectively pin down dislocations. But presence of hydrogen atoms does alter dislocation mobility, at both 25 and 288 C, which is manifested as effects on both strength and ductility (Figures3 and 5). This can occur if the mobility of hydrogen atoms due to diffusion (depends on temperature) match the dislocation velocity (depends on strain rate). Hydrogen atoms will then exert a drag effect on dislocations hampering free movement and results in an increase in strength. Additionally, presence of hydrogen is reported to result in shear localization causing strengthening [38]. Thus, both drag effect and shear localisation can cause strengthening. On the other hand transmission electron microscopic studies have shown that hydrogen enhances dislocation mobility [39] primarily due to the shielding effect of hydrogen atoms on interaction between dislocations. The shielding effect will aide in dislocation motion and will have a softening and mitigating effect on DSA (significant at 288 C) resulting in reduction in strength. Both effects (strengthening and softening effects) may thus be simultaneously active having partially opposite effects.

The strain rates at which significant hydrogen effects are observed at 25 (10-5 s-1 - 10-4 s-1) and 288 C (10-3 s-1 – 10-2 s-1) and the reported lower bound hydrogen diffusion coefficient (D) values in bcc steel at those temperatures (D at 25 °C ~ 10-11 m2·s-1, at 300 °C ~ 3·10-9 m2·s-1 [40]), differ by approximately two orders of magnitude. This is an indication that dislocation mobility match hydrogen atom mobility within a certain strain rate range only, depending on the temperature. The effect of hydrogen on dislocation mobility and YS has been widely reported to either increase or decrease YS [38, 39, 41]. Dislocations become mobile at onset of yielding and with increasing strain, dislocation density increases and they start to increasingly interact with each other. At 25 C though hydrogen atoms cannot pin down dislocations, at the onset of yielding, dislocations have to initiate moving through a cloud of interstitial hydrogen atoms, which hinders free movement, increasing the load required for dislocation mobility (increase in YS at 25°C(Figure 1(a))). At onset of yielding, strain rate effects are not very significant (confirmed by weak strain rate dependence of YS in presence of hydrogen) but once the dislocation becomes mobile, strain rate effect becomes significant (Figure 3(b)). At 288 °C, mitigating effect of hydrogen atoms on DSA is evident from Figure 5 (a) and (b) which can be attributed to the shielding effect.

Association of dislocation with hydrogen atoms (depending on strain rate and temperature) also results in dislocation transport of hydrogen atoms and this has been reported earlier to be much higher than lattice diffusion [42]. 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 aids 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.

3.2 Fracture Toughness Test Results

3.2.1 Tests in Air: Effect of Hydrogen

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 a (alimit) [11]. The specimen

*To qualify JQ as JIC, 15 different criteria of ASTM E1820 have to be satisfied. Although the important SSY criteria were always satisfied, other isolated criteria were not fulfilled occasionally in some tests. Therefore JQ is used for all cases in this paper.


dimensional requirements were satisfied, as specified in the standard [35]. Limits on J and a at 288 °C were more relaxed for 1T-C(T) specimen as compared to 0.5T-C(T) specimen. Valid data points meet small scale yielding criteria, thus the measured toughness is independent of specimen dimension.

A very limited number of preliminary fracture toughness tests at 25, 150 and 288 °C were performed in air using 0.5T-C(T) and 1T-C(T) specimens, in as-received condition and after hydrogen pre-charging. These tests provided baseline data indicating the effect of hydrogen on fracture toughness and associated fractographic features. JQ values of 360 60 kN·m-1 were measured for as received material at 25 and 150 °C, respectively, and slightly reduced to 290 20 kN·m-1 at 288 °C. At 25 °C, no reduction in initiation and tearing resistance was observed after “moderate” hydrogen charging for 7 h (Figure 7a). After “severe” hydrogen pre-charging, sudden unstable brittle crack extension occurred at 25 °C at a slightly reduced but high JIC value of 281 kN·m-1 (Figure 7a and b) without prior ductile crack growth, while at 288 °C stable ductile crack extension occurred (JQ = 313 kN·m-1) for the same “severe” hydrogen charging conditions (Figure 7c). No reduction of initiation toughness and tearing resistance was thus observed at 288 °C even after severe hydrogen pre-charging (Figure 7c), although a clear change in fracture morphology was observed in presence of hydrogen (see Section 3.2.3). The change in fracture morphology clearly demonstrates the need for further tests over a broader range of material, hydrogen and loading conditions for firm conclusions on potential hydrogen effects in air.

3.2.2 Tests in High Temperature Water

J-a plots for 1T-C(T) and 0.5T-C(T) specimens tested at 288 °C in high temperature water and air at various test conditions are shown in Figure 8. Within the investigated parameter range, the high-temperature water did not reduce the initiation toughness and tearing resistance as compared to corresponding air tests (Figures 8 & 9 and Table 3). However, a distinct change in fracture morphology and deformation structures [8] was observed, that was very similar to that in EPFM and tensile tests in air with hydrogen pre-charging (see Section 3.2.3). The observed change of fracture and deformation mode was thus not strong enough to significantly modify the fracture behavior at 150 and 288 °C in the as-received base metal, but this could change for more susceptible materials, lower temperatures or critical strain rates. The toughness was similar for hydrogenated and oxygenated neutral or moderately alkaline high-temperature water and also pre-oxidation period and loading conditions during pre-oxidation did not affect the toughness (Figure 10). Even 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).

So far, only stable ductile crack growth was observed in high-temperature water in as-received base metal. Preliminary tests with CGHAZ (HT1) in high-temperature water increasingly revealed significant unstable rapid, but ductile crack extensions after reaching the peak load and some prior stable ductile crack growth with decreasing loading rates (Figures 11 and 12).The evaluation of these CGHAZ tests is ongoing and it is currently not clear, if there is a significant environmental contribution involved or not, nevertheless, they indicate the need for more systematic studies.

3.2.3 Fractography

In air tests without hydrogen, fracture is ductile at all the test temperatures, with 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. Figure 13 (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 can be seen in Figure 13 (a) in addition to ductile dimples. A macrovoid is marked in Figure 13(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. Macrovoids and localised flat quasi-cleavage 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 (Figures6 and 13) and the role of hydrogen atoms in inducing strengthening (shear localisation, drag effect)/softening (shielding effect) is also similar. Mechanistic studies based on change in fracture surface have been reported earlier specifically for hydrogen effects in LAS [48]. Fractographic features observed in this study in presence of hydrogen are similar to those reported earlier [49].

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. Figure 14(a) and (b) shows a typical fracture surface of a 1T-C(T) specimen tested in BWR/HWC (test B2) with different regions labelled. Figure 14(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, Figure 14(a) and (c)), secondary cracks (Figure 14(b) and inset B in Figure 11(c)) and quasi-cleavage facets (Figure 14(a-c)). Quasi-cleavage facets at inclusions were observed in some experiments, while macrovoids, secondary cracks (some of them with intergranular appearance) and flat featureless regions resembling quasi-cleavage facets, were observed after all tests in high temperature water. The secondary cracks are ductile in nature, as indicated in Figure 14(b), where both quasi-cleavage and ductile dimples can be seen inside the cracks.

Comparing Figures13 and 14, the differences/similarities in fractography are evident for tests done in air with/without hydrogen precharging and tests done in high temperature water. Tests in BWR and PWR simulated conditions resulted in higher roughness of the fracture surface (Figure 14) due to significantly higher macrovoid and secondary crack formation. A cross-section of the fracture surface for test B2 was metallographically prepared and etched and Figure 15 shows a typical macrovoid. The image is from the cross-section (marked by arrows) shown in the schematic in Figure 15 (inset).

4. DISCUSSION ON ROLE OF HYDROGEN ON FRACTURE IN HT WATER

The extension rate applied during EPFM tests was few orders of magnitude higher than usually used to investigate EAC or corrosion fatigue, limiting EAC crack growth (unlike previous reported studies [27, 28]). The air fatigue pre-cracked C(T) specimens were exposed to high temperature water 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). Hydrogen (dissolved in water and corrosion generated) picked up from water during the initial pre-oxidation stage and, in particular, at the crack tip during subsequent EPFM loading, could potentially reduce fracture toughness. Although both initiation toughness and tearing resistance were not reduced in high-temperature water in the investigated narrow parameter range, there is clear evidence for hydrogen effects as indicated by fractography. There is a clear change in fracture morphology after tests in high temperature water which is similar to fractographic features after tensile and fracture toughness tests with hydrogen pre-charged specimens. Reduction of ductility properties (reduction of area, elongation at fracture) at 288 °C in tensile tests in air with moderate hydrogen charging (and additional hydrogen release during the heating phase) also indicates potential embrittling effects of hydrogen at high temperature. Hydrogen effects on tensile properties only occurred in a narrow strain rate range with a peak at around 10-2 s-1, which is a factor of ten higher than the estimated crack-tip strain rate (before initiation) of 10-3 s-1 in C(T) specimen, at a loading rate of 0.25 to 0.35 mm/min.

Since hydrogen diffusion, permeation and release rates are very high in RPV steels and oxide films do not represent a significant diffusion barrier, the bulk hydrogen content in steel is primarily controlled by dissolved hydrogen content of the wateras per the Sieverts law. Hydrogen uptake from uniform corrosion is proportional to the corrosion rate and dependent on time and corrosion rates are usually (with exceptions) very low in high-temperature water. The corrosion rate decreases with time and most corrosion occurs within the first day of exposure. Furthermore, the corrosion rates in very high and very low dissolved oxygen water are quite similar. The bulk hydrogen content close to the surface is thus


finally approximately proportional to square root of the coolant hydrogen fugacity fH2 and thus limited to maximum 0.5 to 2 ppm. Based on Fick’s second law of diffusion and a lower bound value of the diffusion coefficient of hydrogen DH at 300 °C (=3·10-9 m2·s-1), the hydrogen concentration at mid thickness of a 1T-C(T) specimen reaches ~90% of the dissolved hydrogen concentration at the surface in 12 days at 288 °C. Hence, no difference on the EPFM fracture behavior in high-temperature water is expected for pre-oxidation period greater than 12 days and upto 32 days, as used in this study.

Complete oxygen depletion occurs in the crack enclave because the oxygen consumption by corrosion reactions 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 high-temperature water. The crack tip environment thus has an occluded aggressive crack crevice environment in the immediate vicinity of the plastically strained, bare and active crack-tip, and is at a low potential for all tested high-purity, high-temperature water environments. Hydrogen uptake within cracks and its transfer to the process zone at the crack-tip is governed by the local crevice environment, corrosion reactions and crack-tip mechanics. Different environments (BWR HWC, BWR NWC, PWR) and test conditions (pre-oxidation duration,loading during pre-oxidation, in-situ hydrogen charging) showed similar effects on fracture which suggests that hydrogen uptake in the process zone at the crack-tip in the EPFM tests is dominated by the (local) corrosion reactions.

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. The pH of the occluded crevice solution in the moderately alkaline PWR environment is only slightly higher (1.5 to 2 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 SO4

2- impurities or due to dissolution of MnS-inclusions (maximum acidic pH-shifts of 1 to 1.5 units) and potentially very slight alkalization in case of very high-purity water and low-sulphur steels. 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. The higher dissolved hydrogen content in PWR environment is (partially) compensated by the lower corrosion rate in the moderately alkaline PWR water. A similar behaviour in EPFM tests is thus not unexpected from this perspective.

As shown in Figure 16, 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 and hydrogen formation occurs. The oxide film formation on the crack flanks in the direct vicinity of the active crack-tip further generates hydrogen. In addition, hydrogen generated and picked up during oxidation of fresh exposed surfaces at the crack tip will reach susceptible locations (MnS- and oxide inclusions, intersections of slip bands, prior austenite grain 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 [41] in the crack-tip plastic zone. 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. [50].

The hydrogen diffusion rate is two to three orders of magnitude faster than the typical crack growth rates in EPFM tests of 10-7 to 10-6 m·s-1. Exposure (by the growing crack) and dissolution of sulphide inclusion (intersected by the crack front and enclave) to high temperature water at the crack tip is another potential source of hydrogen [51]. However, this may not be significant in the present low-sulphur steel or short-term EPFM tests with high loading rate (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 [51]. Thus, there is a continuous source of hydrogen which can enter the steel during the


fracture toughness tests (as well as during the pre-oxidation stage) and some significant local hydrogen supersaturation can thus be expected to occur under these dynamic crack-tip plasticity conditions. Furthermore, the 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 [51].

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 tap centres may occur by local hydrogen embrittlement mechanisms with rapid propagation back to the main crack-tip (Figure 16). Such hydrogen embrittlement mechanisms that may occur in RPV steels are [51, 52]:

1. Hydrogen-enhanced decohesion embrittlement (HEDE) hydrogen weakens the interatomic bonding brittle or QC micro-cracks at strong traps like MnS or oxide inclusions, intergranular cracking at grain boundaries in combination with grain boundary carbides or metalloid segregation.

2. Hydrogen-enhanced local plasticity (HELP) hydrogen reduces the elastic interactions forces between dislocations and dislocation-obstacles (shielding effect), hydrogen reduces the stacking fault energy and thus cross slip, hydrogen reduces Peierls-stress, but can have a drag effect on mobile dislocations also.

3. Hydrogen-enhanced strain-induced vacancies (HESIV) [53-55] hydrogen reduces formation energy of vacancies and facilitates their agglomeration/clustering and collapse to microvoids very high vacancy and hydrogen concentrations in highly cold-worked regions affects plastic deformation behaviour (climbing of edge dislocations, vacancy clusters and nanovoids as dislocations obstacles) and enhances microvoid coalescence.

4. Adsorption-induced dislocation emission (AIDE)[52] hydrogen facilitates nucleation of dislocations at surface (crack-tip, voids) and their emission into the plastic zone.

Combinations of 1 to 4 are possible in RPV steels under LWR conditions and their relative contributions depend on temperature, strain rate and microstructure, etc. 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 (e.g., due to dislocation pile-ups), e.g., 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 MnS-inclusions observed in EPFM and EAC tests in high-temperature water [51].

Although no high-temperature water effects on the initiation toughness and tearing resistance were observed in these very preliminary tests in a narrow parameter range so far, there is enough experimental and theoretical evidence that hydrogen effects may occur in RPV steels in the LWR temperature range. More pronounced effects can be expected with more susceptible materials (high sulphur steels with high EAC susceptibility, CGHAZ with high yield stress and high plastic welding strains, steels with high DSA susceptibility), at lower temperatures (with more hydrogen-deformation mechanism and stronger effects) 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 from 100 to 200 °C.

5. SUMMARY AND CONCLUSIONS

The effect of hydrogen and high-temperature water on the mechanical and fracture behaviour of a low-alloy RPV steel (bainitic base metal and simulated coarse grain weld heat-affected zone (CGHAZ)) was evaluated by tensile and elastic-plastic fracture mechanics (EPFM) tests in air and EFPM experiments in hydrogenated and oxygenated high-temperature water in refreshed high-temperature water loop autoclave


systems. All the tests were followed by detailed fractographic analysis to establish the mode of fracture. From the results the following conclusions can be made:

i) Hydrogen in the range of 2 to 5 ppm in the RPV steel resulted in embrittlement in tensile tests in air, both at 25 °C (room temperature) and at 288 °C and the embrittling effects was more significant at 25 °C, at higher hydrogen contents in the steel and in the CGHAZ material with higher yield stress.

ii) Maximum embrittling effects were observed at strain rates of 10-5-10-4 s-1 and 10-3-10-2 s-1 at 25 and 288 °C, respectively that were related to the matching of dislocation mobility and hydrogen diffusivity.

iii) In the investigated loading rate range (0.25 to 0.35 mm·min-1), “severe” hydrogen pre-charging, resulted in sudden unstable brittle crack extension without prior stable ductile crack growth at 25°C, while at 288 °C stable ductile crack extension occurred for the same “severe” hydrogen charging conditions. No reduction of initiation toughness and tearing resistance was observed at 288 °C after severe hydrogen pre-charging, although a clear change in fracture morphology was observed.

iv) At the same loading rate, exposure to high temperature water containing dissolved hydrogen (BWR/HWC, PWR), dissolved oxygen (BWR/NWC) or with additional electrochemical in-situ hydrogen charging (PWR) for 12 days (up to 32 days) did not reduce the initiation toughness and tearing resistance with respect to corresponding air tests. A clear change in fracture morphology and deformation structures was observed that was very similar to that in EPFM and tensile tests in air with hydrogen pre-charging. The type of environment and period of pre-oxidation and type of loading during the pre-oxidation did not affect the toughness values.

v) Detailed fractographic analysis revealed that embrittlement in tensile tests in air and the change of fracture morphology are related to (hydrogen-induced) extensive localization of plastic deformation (due to micro-void coalescence) and not to a microscopically brittle process. Besides micro-void coalescence, various amounts of quasi-cleavage and secondary cracking as well as macro-void (> 100m) formation were observed in case of hydrogen pre-charging or in high-temperature water which were manifestations of localization of plasticity.

vi) Although no high-temperature water effects on the initiation toughness and tearing resistance were observed in these very preliminary tests in a narrow parameter range so far, there is enough experimental and theoretical evidence that hydrogen effects may occur in RPV steels in the LWR temperature range. More pronounced effects can be expected, e.g., with more susceptible materials, at lower temperatures or under suitable critical loading/strain rate conditions. Further systematic studies are thus recommended.

6. ACKNOWLEDGMENT

Funding from the European Community's Seventh Framework Programme (FP7/2007-2013) 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 L. Nue, B. Baumgartner, D. Stammbach, R. Schwenold and H. Wiese from PSI.

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8. 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 (V) DSA Dynamic strain aging EAC Environmental assisted cracking HAZ Heat affected zone HEDE Hydrogen-enhanced decohesion embrittlement HELP Hydrogen-enhanced local plasticity HESIV Hydrogen-enhanced strain-induced vacancies HWC Hydrogen water chemistry IR Irradiation embrittlement LAS Low alloy steel LLD Load line displacement at specimen (= crack opening displacement) LWR Light water reactor MVC Micro void coalescence NWC Normal water chemistry 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) UTS Ultimate tensile stress YS Yield stress y

25°C Flow stress at 25 °C (MPa)

y288°C Flow stress at 288 °C (MPa)

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 = (W - ao) (mm) C Concentration of hydrogen in steel CS Surface concentration of hydrogen D Diffusion coefficient (m2·s-1) J Fracture toughness parameter defined in ASTM E 1820 (kN·m-1) JDCPD J at initiation of ductile crack growth measured from DCPD signal (kN·m-1) Jel Elastic component of J integral (kN·m-1) Jlimit Maximum allowable J integral of C(T) specimen (kN·m-1) Jpl Plastic component of J integral (kN·m-1) JQ Fracture toughness at initiation of stable crack growth (kN·m-1) at a = 0.2 mm


= 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) ppm Parts per million by weight pH -log([H+]), where H+ is hydrogen ion concentration in water pH288 °C pH of water at 288 °C R Ratio Kmin/ Kmax VH1 Vickers microhardness at load 1kg W C(T) specimenwidth (mm) dK/dt Rate of change in stress intensity factor during fracture toughness test (MPa·m0.5·h-1) dJ/dt Rate of change in J integral during fracture toughnesstest (kN·m-1·h-1) 1T-C(T) Compact tension specimen of 25 mm thickness (B) 0.5T-C(T) Compact tension specimen of 12.5 mm thickness (B) K Kmax - Kmin(MPa·m0.5 ) Poisson’s ratio a Increase in crack length (m) alimit Maximum allowable increase in crack length (m)


9. TABLES AND FIGURES

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 (ppm)

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

YS (MPa) UTS (MPa) Ductility Total elongation

(%) Reduction in cross-sectional area (%)

25 °C 467 605 17 72 288 °C 400 578 16 70

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

Grain size (m)

YS at 288 °C (MPa)

UTS at 288 °C (MPa)

Hardness (HV0.1)

Microstructure

AR 10-20 367 525 261 Bainite HT1 50 741 846 344 Tempered martensite HT2 50 464 587 250 Tempered martensite

Table 3: Experimental details and J values at stable crack initiation (JQ and JDCPD) for air tests.

Test in air with base metal (AR) Tests in air J at initiation J = C1·aC2

T (°C)

Specimen type

LLD rate (mm·min-1)

Test condition JQ(kN·m-

1) JDCPD(kN·

m-1) C1 C2

25 334 199 523 0.612 150 0.5T-C(T) 0.25 0 ppm H 306 135 439 0.598 288 255 168 384 0.533 25

1T-C(T) 0.350

0 ppm H

380 625 0.672 25 332 551 0.773 25 372 601 0.827

150 413 562 0.592 288 313 453 0.571 25

H pre-charged 281# -- --

25 440 609 0.573 288 312 441 0.542

#JIC at sudden unstable fracture


Table 4: Experimental details and J values at stable crack initiation (JQ and JDCPD) for tests in high temperature water with base metal.

Tests in HTW

T (°C)

Specimen type

LLD rate (mm·min-1)

Environ-ment

Test no.

Pre-oxidation J (kN·m-1) J = C1·aC2

Duration (days) Mode of loading Crack growth (mm)

In-situ hydrogen

JQ JDCPD C1 C2

150 0.5T-C(T) 0.250 PWR

A1 12 Constant loada 0

3 mA·cm-2

321 229 438 0.535

A2 19 18177 cyclesb 0.246 3 mA·cm2

347 257 448 0.487

A3 11.5 Constant loada 0 no 326 214 444 0.548 1T-C(T) 0.350 BWR/HWC B1 12 10602 cyclesb 0.166 --- 353 270 491 0.628

288

0.5T-C(T) 0.250 PWR A4 29 39932 cyclesb 0.341 3 mA·cm2

273 182 399 0.492

A5 12 Constant loada 0 no 299 218 401 0.439

1T-C(T) 0.350 BWR/HWC

B2

34 (Constant load)a

+ 11 (Cyclic load)b

Constant loada +

10000 cyclesb 0.081 --- 290 218 411 0.508

B3 32 27483 cyclesb 0.378 --- 326 256 447 0.471 B4 12 Constant loada 0 --- 331 186 464 0.556 B5 20 10000 cyclesc 0 --- 251 167 389 0.582

BWR/NWC B6 15 11842 cyclesb 0.085 --- 299 184 407 0.639 O2& H2 free B7 12 10311 cycleb 0.06 --- 293 167 400 0.440

a500 N for 0.5T-C(T), 1000 N for 1T-C(T), (KI< 3 MPa·m0.5), bCyclic loading, K=9.85, 0.01 Hz, cCyclic loading, K=2.9, 0.01 Hz


Fig. 1.Applied procedure for development of J-a curve by the DCPD technique.

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

325

350

375

400

Onset of stable ductile crack growth

Stable ductile crack growth

IIIII

DC

PD

U

[

V]

Load line displacement LLD [mm]

I

Blunting

22 NiMoCr 3 7288 °C, BWR/HWC1T C(T), a

0/W = 0.5

dLLD/dt = 0.25 mm/min

0

10

20

30

40

50

60

70

Load

[k

N]

0.0 0.5 1.0 1.5 2.0 2.50

200

400

600

800

1000

JQ=267 kN/m

alimit

Jlimit

J

[kN

/m]

Crack extension a [mm]

JDCPD

=168 kN/m

J = 384 a0.501

R2 = 0.997

Air, 25 °C, DCPD0.5T-C(T), a/W0.5

0.25 mm/min

2y

0.2 mm

(a) (b)

Fig. 2. Examples of a DCPD curve during EPFM test in high-temperature (a) and of a J-a plot in air with validity limits and different blunting and exclusion lines (b).


10-5 10-4 10-3405

410

415

420

425

430

435

440

AR, 0 ppm

Strain rate [s-1]

YS [

MP

a]

AR, 5.5 ppm

T = 25 °C

(a)

10-5 10-4 10-30

10

20

30

40

50

Strain rate [s-1]

AR, 0 ppm

Red

uct

ion

in a

rea

[%

]

T = 25 °C

AR, 5.5 ppm

(b)

Fig. 3.Variation in a) YS, and b) reduction in cross-sectional area, at varying test strain rates at 25 °C.


(a) (b)

(c) (d)

(e)

Fig. 4.Fractograph at 25 C a)”Cup-cone” fracture, b) 45 ° inclination of the fracture surface to the loading direction showing the “QC” and “MVC” region marked, c) magnified view of region “A”, “B” showing brittle fracture (encircled regions) near inclusions, d) elongated dimples in the direction of shear and e) schematic showing the formation of inclined fracture surface in presence of hydrogen.


10-3 10-2 10-1510

520

530

540

550

AR, 0 ppm

UT

S

[M

Pa]

Strain rate [s-1]

T = 288 °C

AR, 5.5 ppm

10-3 10-2 10-1800

810

820

830

840

850

860

870

UT

S

[M

Pa

]

Strain rate [s-1]

HT1, 0 ppm

HT1, 2.6 ppm

T = 288 °C

(a) (b)

10-3 10-2 10-110

15

20

25

30

35

40

45

50

AR, 0 ppm

Red

uct

ion

in a

rea

[%

]

Strain rate [s-1]

AR, 5.5 ppm

HT1, 0 ppm HT1, 2.6 ppm

HT2, 0 ppm HT2, 2.6 ppm

(c)

Fig. 5.Tensile properties at 288 °C at different test strain rates. Variation in UTSfor a) AR and b) HT1 and c) reduction in cross-sectional area for AR, HT1 and HT2 microstructure.


(a) (b)

(c) (d)

Fig. 6.a) Predominant shear fracture in presence of 5.5 ppm hydrogen, b) evidence of “QC” in regions marked “A” and extensive plastic deformation in regions marked “B”, showing flat featureless regions, c) elongated dimples in the region marked “C” in (b) and d) fracture surface showing large voids and secondary cracks.


0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.50

200

400

600

800

1000

1200

1400

Air Air, H pre-charged for 7 h Air, H pre-charged for 24 h

380 kN/m

J

[kN

/m]

Crack extension [mm]

440 kN/m

25 °C, air, unloading compliance0.5 C(T), 1T C(T), a/W 0.5

0.25 mm/min or 0.35 mm/min

281 kN/m, unstable brittle crack extension

JQ

(a)

(b)

0.0 0.5 1.0 1.5 2.0 2.50

100

200

300

400

500

600

700

Air, 288 °C, H2, ULC

Air, 288 °C, corr. ULC

J

[kN

/m]


Air, 288 °CUnloading compliance1T-C(T), a/W 0.50.35 mm/min

312 kN/m313 kN/m

JQ

(c)

Fig.7. J-a plots in air at 25 °C (a) and 288 °C (c) in the as-receivedcondition (0 ppm) and after hydrogen pre-charging. Unstable brittle fracture surface after “severe” hydrogen charging at 25 °C (b).


0.0 0.5 1.0 1.5 2.0 2.50

100

200

300

400

500

600

700

N2, O

2- & H

2-free, B7

HWC, H2, B5

PWR, H2, A5

NWC, O2, B6

Air, 288 °C, DCPD Air, 288 °C, H

2, ULC

Air, 288 °C, corr. ULC

J

[kN

/m]


T = 288 °C0.5T- & 1T-C(T) 0.25 or 0.35 mm/min

Fig. 8. J-a plotsat 288 °C in air and different high-temperature water environments for as-received base metal specimens.

Air

BW

R/N

WC

BW

R/H

WC

BW

R/H

WC

BW

R/H

WC

HP

W, n

o D

H &

DO

PW

R, B

/Li

PW

R, B

/Li

0

100

200

300

400

500

cycl

ic lo

adin

gn

o cr

ack

gro

wth

cycl

ic lo

ad

ing

cons

tan

t lo

ad

cons

tant

load

cycl

ic lo

ad

ing

cycl

ic lo

adin

g

cycl

ic lo

ad

ing

pH288°C

= 6.9, alkaline

nitr

ogen

ated

0 pp

m D

H &

DO

1.4

ppm

DH

1

.9 p

pm D

H90

0 pp

m B

, 2.3

ppm

Li

J Q o

r J D

CP

D

[kN

/m]

2 p

pm D

O

in-s

itu H

cha

rg.

pH288°C

= 5.7, neutral

T = 288 °C

Fig. 9. Comparison of JQ and JDCPD values of AR base metal specimens at 288 °C in air and different high-temperature water environments.


Fig. 10.Comparison of JQ and JDCPD values of AR base metal specimens at 288 °C in high-purity, hydrogenated high-temperature water environment for various pre-oxidation conditions.

0.0 0.5 1.0 1.5 2.0 2.5 3.00

5

10

15

20

25

234.50 234.75 235.00 235.25 235.50200

210

220

230

240

DC

PD

U

[

V]

Time [s]

Load

[k

N]

Pull rod stroke [mm]

PWR, 288 °C0.5T C(T), 0.025 mm/min

22 NiMoCr 3 7, CG HAZ (HT1)

unstable, ductile crack growth 180

200

220

240

260

280

DC

PD

U

[

V]

Fig. 11. Example of unstable crack extension in CG HAZ (HT1) in EPFM test in PWR environment at a slower loading rate of 0.025 mm·min-1.


Fig. 12.Unstable rapid, but ductile crack growth in CGHAZ (HT1) in PWR environment.

(a) (b)

Fig. 13.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).


(a)

(b)

(c)

Fig. 14.(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 localised 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).


Fig. 15.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.

Fig. 16. Localized hydrogen uptake, hydrogen-deformation interactions and hydrogen embrittlement in the crack-tip system during an EPFM test.

Documents

EFFECT OF HIGH-TEMPERATURE WATER AND HYDROGEN …envdeg2015.org/final-proceedings/ENVDEG/papers/ENVDEG034.pdftoughness due to a significant role of EAC at the test extension rates