Influence of mechanical and environmental variables on crack growth in PWR pressure vessel steels

Int. J. Pres. Ves. & Piping 24 (1986) 139 173

Influence of Mechanical and Environmental Variables on Crack Growth in P W R Pressure Vessel Steels* D. R. Tice United Kingdom Atomic Energy Authority, Springfields Nuclear Laboratories, Salwick, Preston PR4 0RR, Great Britain (Received: 16 December, 1985)

ABSTRACT Knowledge of the propagation rates of sub-critical cracks exposed to primary circuit coolant is essential for the use of defect assessment procedures for validation of P WR pressure vessel integrity. The results of research programmes conducted by the UKAEA to assess the conditions under which the P WR environment influences crack propagation of low alloy steels under cyclic and steady loading are described. The measured corrosion fatigue crack propagation rates for low and medium sulphur steels in good quality flowing water were well below those predicted by the A S M E Section XI, Appendix A assessment curves, but higher rates were attainable for high sulphur steel or under adverse environmental conditions. No influence of material microstructure due to welding was observed in high flow water. No susceptibility to stress corrosion has been observed in bolt-loaded specimen tests on a range of parent steels and weldments. Cracking in slow strain rate tests was observed only at high potentials, such as were produced by the presence of dissolved oxygen, unless the specimen orientation was such as to maximise access of P W R coolant to sulphide inclusions in the steel. The implications of the data for operating plant and progress with the development of improved methods for assessing crack growth in operating plant are discussed. * Extended version of a presentation made at the 8th International Conference on

Structural Mechanics in Reactor Technology, Brussels, 19-23 August, 1985. 139 ~, UKAEA, 1986.

140

D. R. Tice

1.

INTRODUCTION

Knowledge of the propagation rates of sub-critical cracks exposed to primary circuit coolant is essential for the use of defect assessment procedures for validation of the integrity of the reactor pressure vessel (RPV) of a pressurised water reactor (PWR). Under cyclic loading conditions, the reference crack growth curves for such an analysis, given in Section XI, Appendix A of the ASME Boiler and Pressure Vessel Code, are used. These curves are based on corrosion fatigue data available up to 1979 (Fig. 1), of which the majority were generated by Bamford.1 These data, when individual specimen results are inspected, do not follow a linear Paris relationship between cyclic crack growth rate, da/dN, and stress intensity range, AK, but display a 'plateau' region over which da/d N is nearly independent of AK. Crack growth in the plateau region is 10-/.* o R:071 T 0 0 6 1 (o R=0 2~, TO 0 11 ALL SPECIMENS Icpm SINE WAVE

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Mechanical/environmental factors in crack growth in RPV steels

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dominated by time-dependent processes, similar to those which result in static stress corrosion cracking in appropriate conditions, but the continuous straining under cyclic loading may result in enhanced cracking in environments which will not sustain static stress corrosion. There have been some reported studies of the susceptibility of RPV steels to stress corrosion cracking under steady loading in light water reactor environments. 2 -4 In general, these have shown susceptibility only under oxygenated conditions, but a series of preloaded specimen tests reported by Bamford e t a l . 4 did show cracking in a test environment intended to simulate PWR coolant. It has been suggested that oxygen may also have been a contributory factor in this case. 5 Both corrosion fatigue and stress corrosion are believed to be controlled by the same mechanistic processes, i.e. the rupture o f ' a normally protective oxide film by dynamic straining allowing environmental attack at the crack tip, in competition with chemical repair of the film, i.e. repassivation. The differences in the environmental conditions necessary for steady load stress corrosion compared to corrosion fatigue reflect the different mechanical straining conditions to which the crack is subjected. Work reported in this paper is aimed at an investigation of both corrosion fatigue and stress corrosion, in order to support existing crack growth codes and to aid the development of improved methods of defect tolerance assessment.

2.

EXPERIMENTAL PROGRAMME

Test facilities for both corrosion fatigue and stress corrosion investigations, at the United Kingdom Atomic Energy Authority (UKAEA) laboratories, at Harwell (AERE), Risley (RNL) and Springfields (SNL), were used. Corrosion fatigue testing was performed on 25 or 50 mm compact tension specimens, cyclically loaded by servohydraulic or servoelectric testing machines. Crack length was continuously monitored by either compliance or dc potential drop, with indicated crack lengths being corrected after test by measuring beach marks on the fracture surface corresponding to the start and end of test. The tests were performed in a simulated PWR water environment of high purity and very low oxygen content (Table 1) at elevated temperatures, usually 288 °C. The water environment was continuously refreshed and a water flow rate past the specimen maintained. This varied between

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40 litres rain-1 at AERE, resulting in turbulent flow, to 0.5 litre min(SNL and a few AERE tests), with laminar flow. Although tests were in most cases conducted in water according to the specification in Table 1, in some specific tests the concentrations of the dosing species (Li, B, H2) were intentionally varied or omitted, and in a few cases deliberate additions of oxygen or sulphate anion were made to the water. Two main types of test were used to study stress corrosion cracking (scc). Firstly, precracked specimens were exposed to simulated PWR water to assess susceptibility to crack propagation by stress corrosion under static loading. The specimens were in most cases self-loaded by a bolt, but a few short duration tests under constant applied load were also performed. Secondly, a series of slow strain rate tests on smooth tensile specimens was performed at RNL, in either PWR coolant or oxygenated undosed water. The aim of this work was to investigate the necessary environmental conditions for scc of a range of relevant materials. The effect of flow rate was also studied in these tests. Electrochemical potentials of the specimens were monitored, and some tests were also carried out at imposed potentials by means of a circumferential counter electrode. Further details of experimental facilities and techniques have been described elsewhere. ~' 7 A range of A533B plate and A508-III forging steels were tested, including both modern low sulphur steels and older steels with bulk sulphur contents up to 0.025~o. Some tests were also performed on weldments, attention being devoted to both weld metal and the heat affected zone (HAZ) of the adjacent steel plate or forging. Details of the parent materials and weldments tested are given in Tables 2 and 3 respectively.

3.

CORROSION F A T I G U E TESTS

3.1. Environmental and material variables

In order to investigate the influence of environmental variables for a range of materials, a single frequency of loading, 0-0167 Hz, was used for the majority of tests. This frequency was also used to generate much of the ASME database 1 and so allowed a direct comparison of results between different experimental facilities. Most tests were performed at a stress ratio (R = Kmin/Km~x)of 0-7, but some lower R tests were also carried out.

D. R. Tice

146

The results of several tests on low and medium sulphur steels (Casts A and B) under high flow conditions (Fig. 2) reveal only a small influence of the aqueous environment, and show a negligible effect of the main PWR water additives (LiOH, boric acid and hydrogen).6 A wider range of steels tested in standard PWR coolant again showed little enhancement of crack growth, except for the highest sulphur content steel tested, Cast E (Fig. 3). Moreover, under high flow conditions it proved impossible to sustain crack propagation, even at air rates, below AK-- 15 MPa x/m. The possible influence of the metallurgical structure and composition ofweldments compared with parent steels has been investigated in a series of high flow tests at AERE. Tests were performed on specimens with the precrack either in weld metal or in the heat affected zone (HAZ) approximately 2ram away from and parallel to the fusion line. The 10% FREQUENCY 00167 Hz CASTS A AND B A533B ENVIRONMENT R=0.2R=0.5R=0.7

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results, Figs 4 and 5, indicate that for neither weld metal nor H A Z is there any significant influence of the PWR environment on crack growth. Thus, under high flow rate conditions, no adverse influence of microstructure or composition in weldments is observed. The crack growth rates observed in the AERE high flow rig were much lower than those produced by some other laboratories performing comparable tests, but at lower water flow rates. An international Round Robin test on a medium sulphur steel (at 0.0167 Hz, R = 0.2 in undosed pure water) showed fairly consistent behaviour between laboratories, except for the AERE high flow rate rig where negligible enhancement was observed. In order to confirm the influence of water flow rate, a high flow rig was converted to low flow by removing the circulating pump from the circuit; this resulted in enhanced crack growth rates similar to those

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AERE corrosion fatigue crack growth data for several casts of A533B steel in high purity and high flow rate PWR water.

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obtained at most other laboratories (Fig. 6). The role of steel sulphur content was also demonstrated, since a low sulphur steel (Cast A) did not show enhanced crack growth even at low flow rate. A number of tests at low flow rate have been performed at SNL. No influence of the PWR environment was observed for low sulphur A533B steel (Cast A) at 0.0167 Hz, R = 0.7. Similar behaviour was obtained for this steel and a similar sulphur content A508-III forging (Cast F) at 0.05 Hz (Fig. 7). A test on a higher sulphur steel (Cast D) showed little environmental influence on the macroscopic crack growth rate, However, the a-N and d a / d N A K plots show considerably more scatter at the higher values of AK than is normal for the dc potential drop crack monitoring technique used. Examination of the fracture surface after completion of the test

Fig. 8.

Fracture surface, A533B plate (Cast D), showing Iocalised fan-shaped cracking around inclusions.

Mechanical/environmental factors in crack growth in RPV steels

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(Fig. 8) showed a large number of fan-shaped regions centred on sulphide inclusion sites. These indicated local environmental cracking, but it is not known why these regions did not link up to produce enhanced macroscopic crack growth. The AERE test on the same steel also showed some signs of localised cracking, but there were far fewer fan-shaped regions, and they were much smaller (typically 0 . 1 m m compared to 1-2 mm). One of the A508-III HAZ materials, Cast G, examined at high flow at AERE has also been tested at SNL under low flow conditions. Environmentally assisted cracking was observed above A K = 2 4 M P a x / m (Fig. 7). The enhanced crack growth observed was localised in a D-shaped region, being greatest in the specimen centre and negligible at the edges (Fig. 9). During this test the crack was extended by 2 mm over the range of AK from 21 to 24 MPa x / m in order to reduce the test duration. The period of accelerated cracking occurred immediately after resuming 0.0167 Hz cycling. Towards the end of the test the load was held constant for a period of 200h, at K = K m , x = 1 0 0 M P a x / m . Following this period, resumption of0.0167 Hz cyclic loading led to crack propagation at only the inert environment rate. Crack growth at the

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specimen edges was rather faster, partly overcoming the previous nonuniform crack front. It is not known whether the retardation was due to an environmental effect, e.g. passivation during the steady load period, or a mechanical effect caused by plastic deformation at the high constant load. The effects of dissolved oxygen content in the simulated coolant, and of deliberately added sulphate anions, have been investigated, primarily as a possible explanation for large reported differences in measured corrosion fatigue crack growth rates between laboratories. Oxygen additions to high flow rate water in the range 100-500ppb* resulted in high crack growth rates at 0.0167 Hz, but these rates could not be sustained for more than a few millimetres of crack extension 6 (Fig. 10). The high rates could be restored by a period of high frequency cycling followed by resumption * ppb = Parts per 10 ~.

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of 0.0167Hz loading. This evidence indicates that a high initial dissolution rate due to the presence of oxygen is eventually overcome by crack tip passivation. Reinstating a bare crack tip by high frequency cycling allows the process to be repeated. Concern about the possible effect of sulphate contamination arose due to the discovery of its presence in some grades of boric acid and the possibility of leaching from ion exchange columns or by dissolution of sulphur from the steels being tested or from rig components. Sulphate, as sulphuric acid, was added at 1 ppm levels to high flow PWR water. Considerable enhancement in crack growth was observed for medium sulphur (Cast B) (Fig. 11) but not low sulphur (Cast A) steel. The combined presence of oxygen and sulphate resulted in even higher rates of sustained cracking. Only a preliminary investigation of the effect of the temperature of the coolant has been carried out, since the great majority of tests have been performed at the PWR inlet temperature of 288 °C. One test at SNL on low sulphur A533B (Cast A) steel was started at 288°C, and the temperature was reduced to 200 °C after a non-enhanced crack growth

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rate (at A K = 2 0 M P a x / m ) had been established. The reduction in temperature resulted in an increase in fatigue crack growth rate (Fig. 12) leading to the attainment of a 'plateau' crack growth rate of around 0.5~tmcycle -1 at A K = 2 6 M P a x / m . Examination of the fracture surface (Fig. 13) revealed a dramatic change in its appearance at the lower temperature, compared with anything observed at 288°C. The surface was very rough and non-planar with extensive secondary cracking. The crack had bifurcated for most of its length. The crack growth rates shown in Fig. 12 represent the average rate of advancement of the overall crack front, but the crack bifurcation presumably reduced the cyclic stress intensity factor below the nominal calculated value. 3.2. Mechanical variables

The major mechanical variables under cyclic loading conditions, frequency and R ratio, have been investigated at high flow rate. The results are described in detail elsewhere 8 but indicate that even for the best combination of material and environmental conditions (low sulphur steels in high flow rate, high purity water) environmental cracking can be induced at sutficiently high frequency (1 Hz at R = 0.7), although these high frequencies are not relevant to plant. The clearly defined frequencydependent 'plateau' regions of crack growth in these tests led to the development of mechanistic descriptions of environmentally assisted cracking (see Section 5.4) which have since been supported by data on higher sulphur steels, or in low flow or contaminated water which show EAC at much lower frequency (e.g. Fig. 11). The high flow data also demonstrated a dependence of the threshold AK on stress ratio of the form AKT

(1 -- R)

Only a very limited investigation of the influence of loading frequency has been performed at lower flow rates in the present work. Low flow SNL data on low sulphur steel show plateau behaviour at 0.5 Hz but not at lower frequency. The data, for A533B Cast A, are compared with the AERE data in Fig. 14. Possible interactions between successive loading conditions are of importance in cases where frequency changes are made during a test. Some indications of such effects are suggested by the A508-III Cast G test discussed above (Figs 7 and 9).

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4.

STRESS C O R R O S I O N TESTS

Two main types of test are c o m m o n l y used to investigate stress corrosion. One method is the use of precracked test pieces subjected to steady loading during exposure to the P W R environment. If stress corrosion is observed, this technique yields crack propagation rate data and information on the threshold stress intensity level (Kl~cc) necessary for stress corrosion to occur. Such specimens may either be directly loaded or

Mechanical/environmental factors m crack growth in RPV steels

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self-loaded to a selected K value using a bolt or wedge. The second main method, the slow strain rate test, uses smooth-sided specimens to investigate the range of conditions necessary for scc susceptibility, since individual tests are much more rapid than for other test techniques. However, if it is necessary to perform tests at very low strain rates, test durations may be protracted. Both types of test were used to study the scc susceptibility of A533B and A508-III steels in normal and faulted PWR water chemistry.

4.1. Precracked specimens under steady load A range of specimens of A533B and A508-III parent materials as well as weld metal and the associated HAZ have been exposed to simulated PWR coolant at 288 °C in autoclaves at AERE and RNL. Exposure times for these specimens ranged from 2000 to nearly 13 000h, nearly threequarters of the specimens having been in the PWR environment for over 7000 h. The 25 mm thick specimens were preloaded by a bolt to an initial stress intensity factor of either 60 or 90 MPa x/m. The specimens were periodically removed from the autoclaves for optical monitoring of crack length. The results are summarised in Table 4. No crack extension has been observed in any specimen to date, although only one specimen has been broken open for examination to confirm that no cracking had occurred away from the edges of the specimen. In the one specimen examined, after 5190 h exposure, approximately 10 ~o relaxation of the applied load during the exposure period was measured. This inherent disadvantage in the use of preloaded specimens can be overcome by the use of an external means of specimen loading, either dead loading using a lever and weight system, or actively by means of a universal testing machine. The latter procedure was adopted in a limited series of short term tests at SNL on 2 5 m m thick specimens manufactured from weldments with a precrack in A533B or A508-III HAZ. Final precracking was performed in the PWR environment at 288°C. Following the achievement of the desired precrack (at A K = Kma x = 20 M P a x / m ) , the load was immediately raised to the required static level ( K = 4 0 or 6 0 M P a x / m ) , thus beginning the period of constant load exposure without returning the specimen to an air environment. Crack length was continuously monitored by dc potential drop throughout the test. No significant crack extension was observed in two of the experiments

D. R. Tiee

158

TABLE 4 Bolt-loaded Specimen Test Results

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Crack growth

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No No

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+

60 90

10 926 I 0 926

No No

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t

60 90

8 808 8 808

No No

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60 60

7 996 8 708

No No

A533B B + W (high S)

D

60 60

12788 12788

No No

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G

60 60

5234 5234

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Marrell Fr6res A533B Kawasaki A508 R R A 4 A533B

A H +

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RRA I H A Z CE HAZ N i l W22 H A Z ENSA H A Z

+ + + G

60 60 60 60

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R R A [ weld C E weld N i l W22 weld ENSA weld

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60 60 60 60

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UKAEA Risley

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AERE Harwell

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Mechanical/enuironmental factors in crack growth in RPV steels

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(A533B, K = 40 M P a x / m , 600 h exposure; A508-III, K = 60 M P a x / m , 210 h exposure). However, the experiment on A533B H A Z at a nominal 60 M P a ~ / m showed 1 m m crack extension during 300h exposure. The rate of growth to achieve this total crack length was transient, the m a x i m u m rate being recorded at the start of the experiment, about 3 x 10 - 9 m s - 1 over the first 50 h, decaying to less than 10- 1o m s - 1 after 250h (Fig. 15). The experiment was not continued long enough to determine whether growth had stopped completely. The fracture surface (Fig. 16) revealed that 1 m m crack growth had occurred over the half of the specimen where the initial precrack was longest (at K = 65 M P a x / m ) but there was no crack growth in the other half of the specimen (K=55MPax/m). Examination by scanning electron microscopy revealed a gradual change in the nature and roughness of the surface, consistent with the observed change in growth rate during the test. A section through the specimen was polished and etched in Kalling's reagent, to reveal plastic deformation. This revealed a narrow plastic zone corresponding to precracking at K = Kmax = 20 M P a x / m , a wider zone corresponding to crack growth at much higher K (1 m m long) and a final region showing no plastic deformation, corresponding to post-test fracture in liquid nitrogen. Hardness measurements were consistent with the variations in width of these plastically deformed regions. Thus the

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observation of crack growth under constant load, following in situ precracking, was confirmed. 4.2. Slow strain rate tests

Slow strain rate testing in simulated LWR environments was performed at RNL in a specially constructed test facility, 7 comprising a recirculating water loop providing close control of oxygen and impurity levels and a small volume test section with a Teflon gland at each end through which the tensile test specimen passed. The specimens had a 3 mm diameter,

Mechanical~environmental factors in crack growth in RPV steels

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15 mm long gauge length and were electrically isolated from the test facility. A range of water flow rates over the specimen from laminar to turbulent was possible. The electrochemical potential of the specimefi under test was monitored by an external Ag/AgC1 electrode with a Luggin probe linking it to the specimen. Some tests were performed under potentiostatic control in PWR water by the use of a platinum counter electrode. An investigation of the role of specimen potential on scc susceptibility was performed under two sets of conditions, i.e. either in deionised water with oxygen added to vary the specimen potential or in PWR water with potentiostatic control. N o susceptibility to stress corrosion in parent steel specimens strained in the steel working (L) direction was observed below a critical potential of ~ - 2 5 0 m V . * Above this potential a reduction in ductility was observed, as measured by time to failure or reduction in area. Results are plotted in Fig. 17 in terms of maximum crack length, determined by subsequent examination of the fracture surface. The * All potentials are quoted with reference to the standard hydrogen electrode.

162

D. R. Tice

critical potential for cracking was similar whether imposed by potentiostatic control or by the addition of small but significant levels of dissolved oxygen. Little effect of flow rate on the critical potential for cracking was observed, but there was evidence that, once initiated, cracks propagated more rapidly in low flow conditions. Fractographic examination showed that the specimens which underwent scc showed transgranular fan-shaped cracking radiating from a number of initiation sites around the specimen circumference. These were frequently associated with pittir~g, and there was evidence that the pits were usually centred on sulphide inclusion sites. Initiation of further cracks from inclusions intersected by the advancing crack front was also observed. An exception to the observation that scc did not occur close to the free corrosion potential in deoxygenated water was provided by a test on a specimen tested in the short transverse (S) direction in which large sulphide inclusion clusters outcropped from the metal surface, allowing ingress of the test environment. A reduction in ductility was observed compared to the ductility in an inert environment. The fracture surface showed a mixture of ductile fracture and fan-shaped environmentally assisted cracking associated with the inclusion clusters. In contrast with results for specimens tested at higher potential, there were no obvious initiation sites around the circumference. Some tests have also been performed under cathodic polarisation, to assess whether any susceptibility to stress corrosion is observed under conditions where iron dissolution is not possible. No evidence for stress corrosion under these conditions has been obtained.

5.

DISCUSSION

5.1. Environmental factors influencing cracking The results reported in this paper indicate a clear influence of several material and environmental factors on the susceptibility of pressure vessel steels to enhanced rates of crack propagation in elevated temperature aqueous environments. There are substantial interactions between the parameters of importance, so that some combination of adverse conditions is usually necessary for significant environmentally assisted cracking (EAC). The major factors identified are as follows.

Mechanical/environmental factors in crack growth in RPV steels

163

Steel sulphur content A large influence of steel sulphur content is apparent in the corrosion fatigue data, little effect of a well-controlled PWR environment being observed for low sulphur steels. The threshold sulphur level at which EAC is observed is dependent on the water flow rate past the specimen, being higher for high flow and lower for low flow. Most of the sulphur in the steel is precipitated as second phase sulphide inclusions. Evidence from slow strain rate tests indicates that dissolution of these inclusions creates a locally aggressive environment, enhancing metal dissolution at the tip of a crack. Under the most adverse conditions, in a slow strain rate test on an S-direction specimen, debonding at the metal-inclusion interface allows access of the water environment to the crevice so created, causing local scc under monotonic loading. However, in general, cyclic loading conditions are necessary to maintain the necessary local environment. The observed influence of sulphur content on EAC susceptibility and its interaction with flow rate indicates that the local chemistry conditions within the crack crevice are important in producing enhanced cracking. This will depend not just on the bulk sulphur content of the steel, but the inclusion type, size and distribution, which will be influenced by the steelmaking process and the degree of sulphur segregation. It is not obvious what causes the transition from localised cracking around inclusions (Fig. 8) to more general enhanced cracking. Water .[tow rate As stated above, a higher sulphur content in the steel is necessary for EAC susceptibility to be observed at high flow rate than under low flow conditions. High flow conditions may reduce the local concentration of sulphur-containing species either by direct flushing (especially in throughthickness cracked specimens) or by modifying the diffusion gradient for transport of dissolved species out of the crack. Slow strain rate tests indicate little influence of flow rate on initiation in smooth specimens but a substantial effect on crack propagation rate, consistent with the corrosion fatigue results. Anionic impurities Two main water chemistry variables have been investigated in the present work, sulphate anion contamination and dissolved oxygen content. Increased EAC susceptibility is produced by l ppm additions of sulphuric acid for medium sulphur steel in high flow water. It is assumed

164

D. R. Tice

that the presence of this sulphate modifies the diffusion rates of sulphurcontaining species between the crevice and bulk environment, and thus exacerbates the effect of sulphur in the steel. In deoxygenated conditions, thermodynamic calculations suggest that any sulphur present may become incorporated in the passive film, depending on the concentration of sulphur-containing anions. Thus sulphur species, either from the bulk environment or by dissolution from the steel, may modify the relative rates of passivation of the crack sides and of crack tip dissolution. The level of sulphur added in these tests was high by plant standards, less than 50 ppb being easily achievable. However, monitoring of sulphate is not normally required in plant operating specifications. The significance of the present observation may be of more direct relevance in explaining interlaboratory variability in corrosion fatigue crack growth data, since inadvertent additions of sulphate with the dosing chemicals have been reported. 9 Dissolved oxygen

A clear influence of dissolved oxygen has been demonstrated in inducing stress corrosion cracking of A533B and A508 steels in slow strain rate tests. However, the controlled potential tests show that this is due to the increase in potential caused by the presence of oxygen rather than to any direct influence of oxygen itself. A clearly defined critical potential for cracking of around - 250 mV is observed. These observations indicate that under these conditions cracking is dissolution controlled, although this does not exclude the possibility of an alternative cracking mechanism, such as hydrogen embrittlement contributing to cracking at lower potential, such as under cyclic loading conditions. Potential-pH diagrams for the F e - H 2 0 system indicate that at the potential required for cracking, haematite is the stable oxide species (Fig. 18). The data of Congleton et al. ~° indicate that the presence of haematite on the fracture surface can be detected at the potentials resulting in cracking, but there is still substantial magnetite present even at higher potentials. 5.2. Role of environmental variables on cracking susceptibility: comparison with data from other laboratories

The synergistic effects of steel sulphur content, water chemistry and flow rate on EAC susceptibility are summarised 11 in Fig. 19. This figure is

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based on UK corrosion fatigue data at 1 cycle min -1 (including those summarised here and data produced by the Central Electricity Generating Board, 12 Babcock Power and Rolls Royce and Associates). For a given water chemistry, a fairly well-defined boundary between susceptible and non-susceptible conditions can be drawn. This boundary moves towards higher water flow rates and lower sulphur steel contents under conditions of adverse water chemistry, such as in the presence of sulphate or dissolved oxygen. Close to the boundary a fairly high threshold AK for EAC is observed; a reduction in this AK threshold is observed as one moves away from the boundary into the EAC susceptible region. The data of Bamford of Westinghouse, Fig. 1, generally show much higher crack growth rates, and lower threshold values of AK, than the data produced in the present work. Most of the former data were

D. R. Tice

166

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obtained on older steels of medium to high sulphur content, under low flow conditions, well into the region of EAC susceptibility in Fig. 19. However, tests on low sulphur steels have also been reported 4 including tests on A533B Casts A and B (Table 2). Reported crack growth rates were still higher than those measured in the present work. In particular, EAC was observed at AK values well below the threshold values seen in any UK tests.11 This suggests a difference in environmental conditions between test facilities. Evidence for this was provided by the identification of haematite on the specimen surface for the Westinghouse test on A533B Cast A, 13'14 suggesting that at least in the early part of the test, the specimen potential was in excess of the free corrosion potential in deoxygenated water. Another major discrepancy between the present work and published data is apparent in the bolt-loaded specimen results reported here and those obtained by Bamford. 4'15 In the latter case considerable crack extensions were observed, after around 2000 h for A533B HAZ, but only after 37000h for plate. No cracking was observed for weld metal or forging. Within the UKAEA programme only the HAZ material has exceeded the exposure time in which Bamford obtained crack

Mechanical/environmental factors in crack growth in RPV steels

167

propagation, the leading HAZ specimens having exceeded 12000h exposure. It is not possible at this stage to decide whether the difference in behaviour is due to differences in material or test environment, but it is important to note that periodic exposure to oxygenated water during rig start-up could be a relevant factor in inducing cracking, especially in view of the observation of haematite on a corrosion fatigue specimen tested in the s a m e rig. 13'14 In contrast to the U K A E A bolt-loaded tests in which no crack propagation was observed, the observation of transient cracking in one constant load test following in situ precracking is of significance. It is possible that during precracking, exposure of sulphide inclusions created a locally aggressive environment, which was gradually depleted during the test. Transient crack growth under low frequency cyclic loading, following higher frequency cycling, has been reported elsewhere 9'12'16 and may be due to the same cause. Such behaviour has implications for plant subjected to a spectrum of loading transients, but the high frequency precracking used in the SNL constant load test would never be experienced by operating plant. 5.3. Effect of water temperature

There have been few reported investigations of cyclic crack growth in simulated PWR coolant at temperatures between ambient and 290°C. Atkinson and Forrest 12 have published results of a test in which the operating temperature was varied in a series of steps. Unlike those reported in the present paper, these tests were in stagnant water, under which conditions enhanced crack growth was obtained at 290 °C even for low sulphur steels. Their variable temperature test, at 0.0167 Hz, R = 0.7 on Cast A steel, was started at 290°C, and plateau crack growth was attained above AK = 30 MPa x/m. The stepwise sequence of temperature changes (210-130-180-290 °C) was then made, with lower crack growth rates at all temperatures below 290 °C. In contrast, in the test reported in the present paper, the temperature reduction was made at lower AK, where no EAC had been observed at 290°C; enhanced crack growth at 200 °C was immediately observed, but the plateau rate achieved was very similar to that reported by Atkinson. There is thus an apparent influence of temperature on both the plateau crack growth rate (higher at 290 °C) and threshold AK for EAC (lower around 200°C). Cullen et al.17 have also studied the effect of water temperature, at low

168

D. R. Tice

R (0.2). Again there is evidence of a much lower plateau rate, but a lower threshold AK at 204°C compared to 288°C. In all three laboratories (Atkinson and Forrest, la Cullen et al. 17 and the present work) a much rougher fracture surface is observed around 200°C. Congleton e t al. 1° have demonstrated an influence of temperature in slow strain rate tests. Two different mechanisms by which temperature may influence cracking susceptibility are discussed. First, the oxygen content necessary to produce a potential above the critical value for scc ( ~ - 2 5 0 mV) is temperature dependent, more oxygen being required at 288 °C than at 200-250 °C. This is not relevant to corrosion fatigue testing in deoxygenated conditions. Secondly, there is some evidence for hydrogen-assisted cracking at around 150 °C in the observation of some scc in low oxygen water, well below - 2 5 0 m V . The change in fracture surface appearance and increase in crack branching seen in corrosion fatigue tests is consistent with a contribution from a different mechanism of crack advance, such as hydrogen cracking, below 200°C. 5.4. Influence of mechanical loading variables The identification of the material and environmental conditions producing EAC makes it possible, in theory, to specify conditions necessary to avoid EAC altogether. However, in practice it is not possible to guarantee this will be true for all regions of plant, so it is necessary to ensure that the resulting enhanced crack growth rates are acceptable under plant transient loading conditions. The observed frequency dependence of cyclic crack growth indicates that use of cyclic crack growth laws, such as ASME XI, is not soundly based, despite the existing code being a conservative description of all available data for low and medium sulphur steels in good quality PWR water. This requirement has led to the development of a model 6'8 relating environmentally assisted crack growth rates, converted to a time base, to the crack tip strain rate, defined for cyclic loading as ,i~ = - T i n [1 - ½(1 - R) 2] where T is the rise time of the loading cycle. The essence of the model is that, above a threshold strain (characterised by the threshold AK for EAC), environmental crack propagation occurs by dissolution of bare metal created by oxide rupture. The time

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dependence of crack growth is then predicted to be determined by the processes controlling the dissolution rate. This leads to a relationship between the time-based crack growth rate (a) and the crack tip strain rate of the form

This is borne out by experimental data (Fig. 20), although different values of n are observed depending on the environmental conditions (water chemistry and flow rate). It has also been shown 1°']8 that, providing allowance is made for multiple cracking in slow strain rate tests, the crack growth rates from such tests fit extrapolations of lines on the d-~ c plot through corrosion fatigue data (Fig. 20). Thus, slow strain rate data obtained above the critical cracking potential (either in oxygenated water or potentiostatically imposed) lie on an extrapolation of the gc~/2 line through corrosion fatigue data in oxygenated water. The absence of cracking at low potential is consistent with an extension of the high flow ~c line, allowing for the limit of detection of cracking in short duration slow strain rate tests.

170

D . R . Tice

The above model is based on the description of the environmental effect on crack growth as an anodic dissolution process, although an a-e; relationship would still be expected for other cracking mechanisms providing oxide rupture is the rate determining process. There is some uncertainty about the quantification of crack tip strain rate; alternative derivations have been discussed by Lidbury. 19 An alternative approach is to utilise the inert environment crack velocity as an empirical measure of the strain rate. Despite similarities, the various models differ in their predictions of the influence of mechanical loading parameters such as R ratio and AK, leading to different extrapolations of experimental data to plant conditions. Progress in model development has recently been reviewed. 2° 6.

IMPLICATIONS FOR OPERATING PLANT

In order to compute the extent of propagation of defects in a PWR pressure vessel during operational service, a means of quantifying crack propagation rates is required. For buried defects, and those protected from the aqueous environment by the austenitic stainless steel clad, inert environment crack propagation data are used; the ASME XI, Section A code includes a fatigue crack growth law for this case (ASME XI dry, Fig. 1), although alternatives have been suggested. 5'2° In those instances where defects can be exposed to the operating PWR environment, the ASME XI wet code is currently used; this is conservative compared to the present data generated for low to medium sulphur steels in good quality flowing water. However, for high sulphur steels, or in the presence of dissolved oxygen or sulphur anion contamination, higher crack propagation rates are observed. These observations emphasise the benefits to be gained by the use of steels of low sulphur content for plant construction, and the need to avoid segregation of second phase sulphide inclusions. Modern steels meeting these requirements are being specified for current generation PWRs. The increase in crack growth rates produced by the presence of oxygen and sulphur anion contamination confirms the need for close control and monitoring of operating water chemistry, and for the maintenance of hydrogen overpressure to remove oxygen, during plant operation. The necessary water quality is readily achievable in modern plant with existing chemistry control procedures although care must be taken to prevent intrusions of ion exchange resinfl 1 Minimum crack propagation is

Mechanical/em, ironmental factors in crack growth in R P V steels

171

expected in those regions of plant subjected to turbulent flow, although further work is required to quantify the effects of flow rate for realistic defect shapes. No evidence for sustained stress corrosion cracking in deoxygenated PWR operating conditions has been obtained, suggesting that the test environment in reported work where scc was observed 4 was not representative of normal PWR operating conditions. The observed cracking in slow strain rate tests when environmental exposure of sulphide inclusions was maximised required a specimen orientation not relevant to the usual fabrication route of reactor pressure vessels. While the use of the ASME wet code should allow a conservative evaluation of crack propagation rates, a sounder procedure would make use of the mechanistically based models now being developed. Thus crack propagation would be assessed as a combination of inert fatigue and timedependent environmental crack growth. The latter would only contribute when the necessary environmental and mechanical conditions exist for EAC. Work is continuing to define these conditions.

7.

CONCLUSIONS

1. Measured corrosion fatigue crack propagation rates for low and medium sulphur steels, in good quality flowing PWR water, are well below those predicted by the ASME Section XI, Appendix A assessment curves. 2. Increased rates of crack propagation can be observed in high sulphur steels, or under adverse environmental conditions, i.e. the presence of sulphur-containing anions or dissolved oxygen, especially at low flow rate. 3. Corrosion fatigue crack growth rates in weld metal and heat affected zones, in high flow PWR water, are no higher than for parent steels. 4. No crack propagation in bolt-loaded specimens of A533B and A508-III, nor in the associated weld metal or HAZ, has been observed for exposure periods up to 13 000 h. 5. Limited non-sustained cracking has been observed at constant load, following in situ 'high' frequency precracking. 6. No susceptibility to stress corrosion cracking has been observed in A508-III or A533B steels, when stressed normal to the principal working direction, at a potential characteristic of low (less than 5 ppb) oxygen water at 290°C.

172

D. R. Tice

7. For cracking to occur in such specimens the potential must be displaced in the positive direction by at least 400 mV. The role of oxygen is seen to be mainly to control the potential and a threshold level for stress corrosion of ,-~200ppb has been determined under high flow rate conditions. 8. The influence of sulphide inclusions on environmentally assisted cracking has been demonstrated. For example, steel tested in the short transverse direction under slow strain rate loading shows some cracking susceptibility, due to access of water to the inclusions. 9. Data obtained under both cyclic and steady loading conditions can be combined using currently available mechanistically based models. Development of these models should provide a sound basis for predicting crack growth rates when environmental interaction is present. 10. The results presented in this paper indicate that environmental enhancement will be minimal for plant designed and operated to the best modern standards, i.e. constructed from low sulphur steels and operated with closely controlled water chemistry. While use of the ASME XI code should be conservative, a more realistic assessment of crack propagation under these conditions will be provided by the mechanistically based descriptions of environmental cracking currently being developed.

ACKN O W L E D G E M E N T S Dr P. M. Scott (AERE Harwell) and Dr P. Hurst (Risley Nuclear Laboratories) provided data and useful discussions. The work was partly funded by the Central Electricity Generating Board. REFERENCES 1. Bamford, W. H., J. Pres. Ves. Tech., 102 (1980), 433. 2. Choi, H., Beck, F. H., Szklarska-Smialowski, Z. and MacDonald, D. D., Corrosion, 38 (1982), 76, 136. 3. Lenz, E., Stellwag, B. and Wieling, N., Proc. IAEA ,S~vnq~. on Corrosion am/ Stress Corrosion o['Steel Pressure Boundary Components, Espoo, Finland, 1983, p. 243. 4. Bamford, W. H., Moon, D. M. and Ceschini, L. J., Corrosion 83, Anaheim, CA, 1983, paper 12. 5. Marshall, W., Second Study Group report on P W R pressure t'essel integrity, UKAEA, 1982.

Mechanical/environmental factors in crack growth in RPV steels

173

6. Scott, P. M. and Truswell, A. E., Proc. I A E A Specialistd Meeting on SubCritical Crack Growth, Freiburg, 1981, NUREG CP0044, vol. 2, p. 91. 7. Hurst, P., Appleton, D. A., Banks, P. and Raffel, A. S., paper presented at European Federation of Corrosion Conference, Munich, 1984: published in Corrosion Science, 25 (1985), p. 651. 8. Scott, P. M. and Truswell, A. E., J. Pres. Ves. Tech., 105 (1983), 245. 9. Tice, D. R., paper presented at European Federation of Corrosion Conference, Munich, 1984: published in Corrosion Science, 25 (1985), p. 705. 10. Congleton, J., Shoji, T. and Parkins, R. N., paper presented at European Federation of Corrosion Conference, Munich, 1984: published in Corrosion Science, 25 (1985), p. 633. 11. Tice, D. R., Atkinson, J. D. and Scott, P. M., Proc. IAEA Specialists' Meeting on Sub-Critical Crack Growth, Sendai, Japan, May 1985. 12. Atkinson, J. D. and Forrest, J. E., paper presented at European Federation of Corrosion Conference, Munich, 1984: published in Corrosion Science, 25 (1985), p. 607. 13. Congleton, J., cited by B. Tomkins, UKAEA Report ND-R-848(S), HMSO, London, 1982. 14. Atkinson, J. D., Private communication. 15. Lloyd, G. J., Proc. I A E A Specialists' Meeting on Sub-Critical Crack Growth, Sendai, Japan, May 1985. 16. Emanuelson, R. H. and van der Sluys, W. A., Corrosion 84, New Orleans, LA, 1984, paper 171. 17. Cullen, W. H., Torronen, K. and Kemppainen, M., NUREG/CR 3230, 1983. 18. Congleton, J. and Hurst, P., Proc. I A E A Specialists' Meeting on SubCritical Crack Growth, Sendai, Japan, May 1985. 19. Lidbury, D. P. G., Symposium on Localised Crack Chemistry and Mechanics on Em'ironment Assisted Cracking, Fall Meeting of TMS AIME and MSD ASM, Philadelphia, PA, 1983. 20. Scott, P. M., Proc. 1AEA Specialists" Meeting on Sub-Critical Crack Growth, Sendai, Japan, May 1985. 21. Ljungberg, L. G. and Cubicciotti, D., Corrosion 84, New Orleans, LA, April 1984.