Relatively few experiments have used helium implan­tation to investigate the embrittlement of nickel-based alloys, although this technique has been more widely used for austenitic steels. However, some data for high-Ni alloys have been published by Shiraishi et a/.95 and Boothby.96

Shiraishi et a/.95 compared the effects of helium injection and neutron irradiation on the tensile prop­erties of developmental y0-Ni3(Ti, Al) (alloy 7817) and y"-Ni3Nb (alloy 7818) precipitation-hardened 40Ni-15Cr alloys. A number of alloy conditions, including ST, aged, and cold worked, were tested. Cyclotron injections of helium were made at 650 °C to levels of 5 or 10 appm. Neutron irradiations were made at the same temperature to a fast fluence (E > 1 MeV) of 1.7 x 1024nm-2 and a thermal flu­ence of 5.9 x 1024nm-2. The helium content of the reactor-irradiated specimens was estimated to be ^45 appm, produced mainly from the thermal neutron reaction with 10B. Tensile tests were carried out at the implantation/irradiation temperature at a strain rate of ^5 x 10-4s-1. The results showed sim­ilar trends in helium-implanted and neutron — irradiated specimens, with the total elongation values tending to decrease with increasing tensile strength. Variations in tensile strength for each alloy were largely attributable to variations in the initial heat treatment and working schedules. However, there were some indications of softening and reduced ductility in the neutron-irradiated specimens com­pared to those injected with helium. Overall, the y0- hardened alloy 7817 exhibited relatively high tensile strength (typically >700 MPa) but low ductility fol­lowing helium implantation or neutron irradiation (with total elongation values generally <10% and as low as 1-2% in the highest strength conditions). The y"-hardened alloy 7818 showed lower tensile strength (typically 500-600 MPa) but maintained good ductility, with total elongation values always exceeding 10% and generally being above 20% in STA conditions.

Boothby96 examined the effects of helium and/or lithium injection on the tensile properties of STA (ST 1050 °C, aged 8h at 700 °C) Nimonic PE16. The effects of lithium were examined since this ele­ment is also produced from the 10B(n, a)7Li reaction in neutron-irradiated alloys. Helium and lithium were implanted at ambient temperature, either singly or in combination, to levels of10 appm each. Samples were tested in the as-implanted condition or follow­ing an additional aging treatment of 72 h at 750 °C. Tensile tests at a strain rate of 3 x 10-5 s-1 on as — implanted samples showed no effect of helium or lithium on ductility at 200-550 °C. However, at 650 °C, the as-implanted samples were all embrittled to a similar extent, with the total elongations gener­ally reduced to about half of the unimplanted levels regardless of whether He, Li, or (He + Li) ions were injected. Postimplant aging of samples containing He or (He + Li) resulted in further ductility loss in tests at 650 °C, with significant embrittlement also evident at 550 °C though not at 450 or 200 °C. Postimplant aging of samples containing only lithium, however, resulted in some recovery in ductility compared to the as-implanted condition at 650 °C but some duc­tility loss at 550 °C. Thus, although it was clear that lithium had a detrimental effect on ductility, it did not appear to exacerbate the effects of helium. A mechanism for lithium embrittlement was not identified, though ductility loss in both Li and He implanted samples was associated with an increased propensity for intergranular fracture.

Some additional, previously unpublished, data from a helium injection experiment conducted by Boothby and Cattle are given in Table 2. In this experiment, helium was implanted at ambient tem­perature to levels of 2, 10, and 50 appm into Nimonic PE16 that had been given a two-stage aging treat­ment (ST 1050 °C, aged 4h at 800 °C plus 16h 750 °C). As before, helium-doped samples were either retained in the as-implanted condition or given an additional aging treatment to coarsen the dispersion of gas bubbles. Tensile tests were carried out at 650 °C at strain rates of 3 x 10~5 and 3 x 10~6s~ The results show a significant loss of ductility even at 2 appm helium. In tests carried out at the higher strain rate, the total elongation values decreased progressively with increasing helium con­tent and were further reduced by postimplant aging. The ductility of as-implanted samples was generally lower but less sensitive to helium concentration in tests at a strain rate of 3 x 10~6 than at 3 x 10~5s~ However, there was little effect of the strain rate on the ductility of the postimplant aged samples.

Figure 15 illustrates grain boundary structures in a tensile-tested PE16 sample which had been aged subsequent to helium injection. Failure in this case appeared to occur by the growth and coalescence of cavities which were nucleated at grain boundary gas bubbles.97 The nucleation of unstable cavities at grain boundary helium bubbles requires the application of a critical tensile stress, which is an inverse function of the bubble radius, normal to the boundary. Cavity growth then occurs via the stress-induced absorption of vacancies. In the as-implanted condition, however,
the helium dispersion was too fine to enable grain boundary cavities to nucleate during tensile testing. Reduced ductility in the as-implanted samples appeared to be associated with grain boundary wedge cracking, where, as discussed by van der Schaaf and Marshall98 in relation to helium embrit­tlement of type 316 steel, the role of helium may be simply to decrease the effective surface energy for fracture.

Although it is evident from simulation experi­ments that helium alone can largely account for irra­diation embrittlement, it is more difficult to assess the significance of other radiation-induced effects such as matrix hardening and grain boundary segregation and/or precipitation. One experiment which examined the effect of the radiation-induced precipitation of the Ni3Si g’ phase on the ductility of a binary Ni-8 at.% Si alloy was described by Packan et al99 In this experi­ment, thin foil tensile specimens were bombarded with either protons or a-particles to damage levels of 0.1­0.3 dpa at 750 K; irradiation with a-particles resulted in the introduction of high helium concentrations of about 750 appm per 0.1 dpa. Proton and a-particle irradiations both resulted in the formation of g layers about 20-30 nm thick at the grain boundaries, but the material remained relatively ductile, exhibiting trans­granular failures, in tensile tests carried out at a strain rate of ^3 x 10~4 s-1 at room temperature and, for the proton-irradiated case only, 720 K. Unfortunately, no tests were carried out at higher temperatures and sam­ples which were irradiated with a-particles only at 750 K were not tested except at room temperature. However, low ductility intergranular failure was

induced in a test carried out at 720 K in a sample which was preimplanted with 1000 appm He at 970 K, then irradiated to 0.3 dpa, introducing an additional 2300 appm He at 750 K. Preimplantation of helium at 970 K produced grain boundary bubbles which were 10-20 nm in diameter, compared to 1.5-2 nm in mate­rial that was only irradiated with a-particles at 750 K. The results of this experiment therefore indicated that the radiation-induced precipitation of y’ at grain boundaries did not give rise to embrittlement unless helium was also implanted into the specimens.

The use of refractory metal alloys in radiation envir­onments can offer high-temperature capabilities not matched in other alloy categories. Refractory metal alloys also offer exceptional compatibility with liquid metal coolants. As described in some detail in this

Figure 26 Comparison of the increase in Vickers Hardness for tungsten and tungsten alloys for similar dose and irradiation temperatures. Reproduced from He, J. C.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2008, 377, 348-351; Kurishita, H.; Kobayashi, S.; Nakai, K.; etal. J. Nucl. Mater. 2008, 377, 34-40.

chapter through mechanical property comparisons, these materials are sensitive to impurity contami­nation during metallurgical processing as well as in-service exposures that can lead to grain bound­ary embrittlement issues. The inherent irradiation response of bcc-structured materials also limits refractory metal use at temperatures >0.3 Tm, with significant degradation in material properties with displacive irradiation doses as low as 0.03 dpa.3

Improvements in the irradiated mechanical prop­erties of refractory metal alloys have been observed in recent experimental work, even at low irradia­tion temperatures. This is in part through improved control over impurity levels and also through ther­momechanical processing techniques that result in microstructures with reduced sensitivity to radi­ation embrittlement. This was discussed with refer­ence to LCAC molybdenum,1 where samples

irradiated in the stress-relieved condition showed improvement over material in the recrystallized condition up to the recrystallization temperature. Further development of HP-LCAC molybdenum has resulted in higher aspect ratio grain morpholo­gies that led to plain strain conditions in the grain lamellae during deformation.82 In addition, reduced
grain sizes or higher aspect ratios decrease dis­tances to defect sinks, further reducing irradia­tion sensitivity. While Mo has traditionally been used to study the behavior of W, the microstruc­tural changes and purity control that have been employed for irradiation studies of Mo have not been incorporated into W.

The control over precipitate formation in the pre­irradiated condition appears to result in changes to some physical material properties, specifically, swelling and densification in Nb-lZr,25,27 that may lead to variations in mechanical properties. An understanding of the effect of preirradiation thermo­mechanical processing or in-service microstructural changes that occur during irradiation may lead to improved properties or the ability to avoid dangerous embrittlement issues that can occur through precipi­tate development. This may be of particular interest in Nb and Ta-base alloys that incorporate Zr or Hf additions that react with impurity elements and pro­duce precipitates.

Alloying Mo and W with Re results in improved mechanical properties of unirradiated alloys, in­creased radiation hardening, and radiation-induced embrittlement.62,120 However, much of this work is
on recrystallized, high Re concentration material, the purity of which may not be ideal. The effect that RIS has on the degradation of properties of Mo-Re alloys is a matter of concern. Further work is needed on higher purity, lower Re (5-20 wt% Re) concentration material with reduced grain size, or that with a tai­lored aspect ratio similar to that of LCAC-Mo.

Initial results show improvements to the irradiated properties of Mo and W through the incorporation of either rare earth oxide124 or TiC additions.112,143 These additions aid in restricting grain growth, pro­vide sinks for radiation-induced defects, and act as obstacles to or deflection points for crack propagation. Though these results are preliminary, they outline the need for further examination of incorporating stable dispersion strengthening particles to refractory metal alloys.

G. A. Young, M. J. Hackett, J. D. Tucker, and T. E. Capobianco

Knolls Atomic Power Laboratory, Schenectady, NY, USA

© 2012 Elsevier Ltd. All rights reserved.

Crack-like defects can degrade component lifetime by eliminating the initiation stage of phenomena such as fatigue or stress corrosion. Similarly, other flaws (e. g., lack of fusion defects, gas porosity, and inclusions) can act to magnify global stresses, pro­duce locally aggressive environments via their occluded geometry or composition, and initiate cracking. In order to mitigate these flaws, it is critical to differentiate between defect types.

The first step in the prevention of cracking is understanding the temperature range over which the cracking occurs. The primary measure is to determine whether defects are ‘hot cracks’ or ‘cold cracks,’ that is, whether they form above or below the solidus temperature of the alloy. Secondly, the loca­tion of the crack in the weld (composite region, unmixed zone, partially melted zone, heat-affected zone) and in the microstructure (solidification boundary, crystallographic boundary, etc.) must be determined.1 Once these distinctions are made,
strategies to eliminate cracking can be developed via changes to the welding process, weld parameters, filler metal, joint design, fixturing, and/or postweld heat treatment.

The supersolidus/subsolidus distinction, com­bined with the unifying concept of homologous tem­perature, illustrates the commonality of welding defects and degradation mechanisms across alloy sys­tems as shown in Figure 1. Supersolidus ‘hot crack­ing’ defects include solidification cracking, liquation cracking, and hot tearing. Subsolidus ‘cold cracking’ includes precipitation, transformation, and segregation — induced cracking (SIC) mechanisms.

Since the late 1940s, many journal papers, conference papers, and reports have been published on the change in properties in graphite due to fast neutron damage. Many different units have been used to define graphite damage dose (or fluence). It is impor­tant to understand the basis of these units because historic data are still being used to justify models

used in assessments for component behavior in reac­tors. Indeed, some of these historic data, for example, stored energy and strength, will also be used to sup­port decommissioning safety assessments.

Early estimations of ‘graphite damage’ were based on the activation of metallic foils such as cobalt, cadmium, and nickel. Later, to account for damage in different reactors, equivalent units, such as BEPO or DIDO equivalent dose, were used where the dam­age is referred to damage at a standard position in the BEPO, Calder Hall, or DIDO reactors. The designers of plutonium production reactors preferred to use a more practical unit related to fuel burnup (megawatts per adjacent tonne of uranium, MW/Atu). Research­ers also found that the calculation of a flux unit, based on an integral of energies above a certain value, was relatively invariant to the reactor system and used the unit En > 0.18 MeV and other variants of this.

Today, the favored option is to calculate the flu — ence using a reactor physics code to calculate the displacements per atom (dpa). However, in the field of nuclear graphite technology historic units are still widely used in the literature. For example, reactor operators have access to individual channel burnup which, with the aid of axial ‘form factors,’ can be used to give a measure of average damage along the indi­vidual channel length.

Fortunately, most, but not all, of these units can be related by simple conversion factors. However, care must be taken; for example, the unit of megawatt days per tonne of uranium (MWd t- ) is not necessarily equivalent in different reactor systems.

When assessing the analysis of a particular com­ponent in a reactor, one must be aware that a single detailed calculation of a peak rated component in the
center of the core may have been carried out to give spatial, and maybe temporal, distribution of that component’s fluence (and possibly temperature and weight loss). These profiles may have then been extrapolated to all of the other components in the core using the core axial and radial ‘form factors.’ In doing this, some uncertainty will be introduced and clearly, some checks and balances will be required to check the validity of such an approach.

There are three categories of graphite creep experi­ments. The first type was the restrained creep experiments. In these experiments a graphite specimen, usually dumbbell in shape, is restrained from shrinkage by a tube or split collar manufactured from graphite that shrinks less than the specimen of interest. In the case of anisotropic graphite, the tube or split collar could be manufactured with its longitudinal axis aligned with the more dimension­ally stable grain direction; the specimen would be manufactured with its axis perpendicular to this. These types of experiments are relatively easy to deploy but are difficult to assess as the load is not directly measured and a ‘creep law’ has to be assumed in the assessment of the results.88 A variation on these experiments was the graphite spring tests used in Calder Hall to define primary creep.86

A second important type of experiment was the ‘out-of-pile measurements’ technique.89,90 These experiments, importantly, give ‘real-time results.’ However, this type of experiment is difficult to install in a reactor and results are obtained only for one specimen.

The final type of experiment is the in-pile rig loading using a string of samples and usually taking advantage of the MTR flux gradient to obtain data on samples to various levels of fluence. There have been various designs of simple strings of specimens loaded either in tension or in compression. Tensile creep tests are vulnerable, that is, if one specimen fails, results for the whole string of specimens could be lost. However, there have been various rig designs aimed at overcoming this problem.

It is rather surprising that although the most stable loops are the prismatic loops, basal loops are also observed in zirconium alloys. Moreover, these loops are of the vacancy character. According to the usual rate theory, vacancy loops should not grow as a result of the bias of edge dislocation toward SIAs.

Figure 7 High density of c-component loops in the vicinity of the precipitates in a Zy-4 sample irradiated to 6 x 1025 n m~2; at 585 K. The arrow shows the diffracting vector [0002]. Adapted from De Carlan, Y.; Regnard, C.;

Griffiths, M.; Gilbon, D. Influence of iron in the nucleation of (c) component dislocation loops in irradiated zircaloy-4.

In Eleventh International Symposium on Zirconium in the Nuclear Industry, 1996; Bradley, E. R., Sabol, G. P., Eds.; pp 638-653, ASTM STP 1295.

The reason for the nucleation and growth of the (c) component loops in zirconium alloys has been ana­lyzed and discussed in great detail by Griffiths and co-workers.46,56,57,74 The most likely explanation for their appearance46 is that they nucleate in collision cascades, as shown recently by De Diego.66 Their stability is dependent to a large extent on the pres­ence of solute elements, which probably lower the stacking-fault energy of the Zr lattice, making the basal (c) component loops more energetically stable. It is also possible that small impurity clusters, espe­cially iron in the form of small basal platelets, could act as nucleation sites for these loops.74,76 However, according to Griffiths,46 this cannot account for the very large vacancy (c) component loops observed, since the growth of vacancy loops is not favorable considering the EID discussed previously. In order to understand the reason for the important growth of the (c) component loops, another mechanism must occur. As discussed by Woo,44 the growth of (c) component loops is well understood in the frame of the DAD model. Indeed, because ofthe higher mobil­ity of SIAs in the basal plane rather than along the (c) axis (and the isotropic diffusion of vacancies), dislo­cations parallel to the (c) axis will absorb a net flux of SIAs whereas dislocations in the basal plane will absorb a net flux of vacancies. This can therefore explain why the basal vacancy loops can grow. The incubation period before the appearance of (c) component loops can be explained, according to Griffiths eta/.,73 by the fact that the (c) loop formation is dependent on the volume of the matrix containing a critical interstitial solute concentration. This volume increases as the interstitial impurity concentration is gradually supplemented by the radiation-induced dis­solution of elements such as iron from intermetallic precipitates (or р-phase in the case of Zr-Nb alloys).

4.01.1.3.2 Void formation

Early studies failed to show any cavity in Zr alloys after irradiation.77 From all the obtained data, it is seen that zirconium is extremely resistant to void formation during neutron irradiation (Figure 8).46,52 The effect of very low production of helium by (n, a) reactions during irradiation was mentioned as a possible reason for this absence of voids. But most probably, the fact that in zirconium alloys vacancy type loops are easily formed can be the reason for the absence of void.52 To favor the formation of voids, various studies per­formed, especially on model alloys, have shown that stabilization of voids can occur when impurities are present in the metal. Helium coming from transmu­tation of boron on Zr sponge67 as well as impurities located near Fe-enriched intermetallics are found to favor the stability of voids.54 Irradiations with elec­trons give better conditions to stabilize voids: the main reason is that irradiation doses can be very high — hundreds of displacements per atom can be reached after few hours.19 Moreover, electron irradi­ation on Zr samples preimplanted with He at various concentrations showed the nucleation and growth of voids only for the samples doped with at least 100 ppm of He.78

Although it is well known that it is the deviatoric component of any stress state that drives creep, there were previously very little data to show whether the creep coefficient is identical in both dilational and compressive stress states. Recent papers by Hall,185,186 Neustroev,187 and Garzarolli188 show that creep coefficients are unchanged by the sign of the hydrostatic stress.

As shown in the next section, additional confirma­tion of the independence of creep compliance on the sign of the hydrostatic stress component can be found in some stress relaxation experiments.

4.02.9.6

Stress Relaxation by Irradiation Creep

There are situations where the applied load is initi­ally fixed and then declines during irradiation. There is usually a transient followed by an instantaneous creep rate defined by B0, but the load is constantly falling, leading to an exponentially declining load. Two examples of in-reactor creep relaxation experi­ments are shown in Figure 82, both conducted on a high-nickel alloy Inconel X-750.

Foster and coworkers have very convincingly demonstrated that creep coefficients derived from creep experiments could be used to successfully pre­dict stress relaxation for the same steel in similar neutron spectra.163

Note that the creep coefficient derived for X-750 from the EBR-II experiment is 1.6 x 10-6 (MPa dpa)-1, just slightly larger than B0 and probably enhanced by low levels of voids or bubbles in this high-nickel alloy. In NRU, however, the creep relax­ation proceeded much faster, partially due to a larger transient, but also because the steady-state creep rate is larger. In this experiment the thermal-to-fast ratio was ^10, so there was significant 59Ni enhancement of dpa rate and probably also bubble formation to enhance the creep rate. The greater scatter at very low residual stresses in the EBR-II experiment is mostly due to frictional variations on the compressed
springs and grain-to-grain interactions that come into play at low stress levels.

Stress relaxation experiments can be conducted using a wide variety of specimen types and usually yield similar results, although the transient regimes often vary with specimen geometry, preparation, and texture versus stress field relationship, as shown in Figure 83.

Creep relaxation by irradiation is important in that it can reduce the opportunity for irradiation — assisted stress corrosion cracking. It does so by decreasing internal or surface stresses produced by deliberate or inadvertent damage, as well as by reducing internal stresses arising from welding, abrupt cooling, etc. Figure 84 demonstrates the radiation-induced relaxation that occurs in a weld that proceeds with a creep compliance of B0 that is independent of the sign of the hydrostatic stress.189 Therefore, it appears that the creep compliance B0 can be confidently applied to any stress state.

As a rule of thumb one can anticipate that by 10dpa, >90% of any preload will be relaxed even in the absence of a transient. The fractional unload­ing is not dependent on the magnitude of the preload as long as the bolt or component was not loaded beyond the yield point.

Stress relaxation in structural components of operating reactors is not always operating in isolation. Frequently, a component experiences time-dependent stresses that develop with time as a result of the growth or movement of adjacent components. In pressurized water reactors there are bolts that join baffle plates to former plates. These bolts are usually cold-worked 316 but the plates they join are annealed 304 stainless, a higher swelling steel. Initially, the bolt will start to relax its preload but if the plates are swelling faster than the bolts, then differential swelling will begin to reload the bolt. Additionally, if a bolt is replaced with a fresh bolt, the reloading can be even stronger due to larger amount of

differential swelling. Figure 85 shows several calcu­lated histories of bolt loading for PWR-relevant tem­peratures and dpa rates.1

While bolts are generally preloaded to a specified level, there is always some range of attained loads. It is difficult to measure the stress level in a bolt while it is still in place, but a rough measure of the remaining load can be made from the torque required to remove the bolt. While this is not an exact measurement with friction, corrosion, irradiation-induced self-welding, and other complications possibly participating to define the torque, Figure 86 shows that the measured torques are in reasonable agreement with predictions of creep equations based on experiments conducted in BOR-60 fast reactor. The fact that most of the data lie above the predictions may indicate that many of the bolts are indeed being reloaded by differential swelling to some degree.

4.05.4.5.1 Introduction

Historically, the hardening observed in low Cu steels has been considered to arise from matrix defects. MD arises from the clustering of irradiation-induced point defects to form either vacancy or interstitial clusters, and/or solute-defect complexes (see, e. g., Odette and Lucas53). In addition, various solutes may diffuse to these clusters giving rise to complex defect-solute configurations. A critical factor in attempting to develop insight is that there is no one technique that allows a direct characterization of matrix defects. Insight has been obtained from indi­rect studies using techniques such as positron annihi­lation, where the nature of matrix defects can be inferred only after data analysis or modeling.

Within these constraints, there are two aspects which should be discussed. First, the nature of MD, in particular whether it is vacancy or interstitial clus­ters that give rise to the observed hardening, and, second, the evolution of MD clusters at different fluxes and increasing fluence. The discussion also needs to include an assessment of whether MD is dependent on the presence of other solutes and whether MD can be treated independently of CEC formation. It should be noted that in generating insight into MD, studies of model alloys (e. g., Fe-Cu alloys) have been particularly important, including ion and electron irradiation studies. This is in contrast to the study of Cu precipitation where the majority of information has arisen from studies of Cu-containing neutron-irradiated steels.

4.08.3.1 Chemical Composition and Microstructure

9Cr-ODS steels are being developed by the JAEA (Japan Atomic Energy Agency) for application to SFR fuel cladding. Their standard chemical compo­sition is 9Cr-0.13C-0.2Ti-2W-0.35Y2O3 (wt%). The chromium concentration was determined to be 9wt% in terms of ductility, fracture toughness, and corrosion resistance based on a series of irradiation data of ferrite steels. The addition of titanium pro­duces the nanoscale dispersion of oxide particles, which leads to a markedly improved high-temperature strength. If titanium is added to excess, however, it creates too much strength, which negatively impacts

cold rolling and cold workability. To achieve a bal­ance between strength and workability, a value of 0.2 wt% was selected. Tungsten of 2 wt% is also added in order to improve high-temperature strength by means of solid solution hardening.

The microstructure of 9Cr-ODS steels13-18 can be easily controlled by a reversible a-g transformation with a remarkably high driving force of a few hundred MJ m~3, as compared with a driving force of irrevers­ible recrystallization with a few MJ m~3 for 12Cr-ODS steels. By inducing reversible a-g transformations, 9Cr-ODS steel cladding for fast-reactor fuel elements is currently being manufactured at the JAEA.

The microstructure of 9Cr-ODS steel cladding is basically tempered martensite. However, it has been recognized that 9Cr-ODS steel cladding manufactured in an engineering process possesses a dual-phase struc­ture that comprises both tempered martensite and ferrite phases. An example of their microstructure is shown in Figure 4. The ferrite phase appears white, and the elongated phase is indicated by arrows. Their size is about 30-60 pm in length x 3-10 pm in width. The formation of a ferrite phase in 9Cr-ODS steel is somewhat unusual, because only the full martensite phase can be expected in 9Cr-ferritic steel without yttria under normalizing and air-cooling conditions. Moreover, the high-temperature strength of manu­factured 9Cr-ODS steel is significantly improved by the presence of the ferrite phase.1 — 1 This is obvious from the creep rupture data shown in Figure 5.20

Therefore, the control of ferrite phase formation is a key to the realization of high-temperature strength in 9Cr-ODS steel cladding.

In addition to solidification segregation, the local com­position of welds can change in-service via thermal exposure and via radiation-induced transmutation and radiation-induced segregation (RIS). Thermally induced embrittlement is most notable in low-alloy steels that are high in tramp elements, in locations where welding produces local compositional enrich­ment, and in higher nickel grades (e. g., A508 Gr4N), which are intrinsically more susceptible because ofthe cosegregation of nickel and phosphorous.79-82

RJS is a phenomenon in which irradiation-created defects cause spatial redistribution of alloying ele­ments as they diffuse to, and get trapped in, sinks

 

(AEcPNi/NiO ~ x0) 2

 

1 + b • exp

 

0.5

 

c

 

Effective R — T

 

[1]

 

• exp

ШШ ‘ r І g. fj’

A600heat affected zone

(a)

Far from A600 HAZ weld

EN82H GTA weld

Hardness indents

HAZ

Base

01 1378

6 018000

Figure 19 Comparison of Alloy 600 heat-affected zone and base metal structure, chemistry, and strain: (a) cross-section of stress corrosion sample showing the location of the cracking, (b) grain boundary chromium profiles for base metal (blue) and the HAZ (red), (c) qualitative strain map for the base metal, HAZ/weld interface, and (d) typical grain boundary microstructure for the HAZ (sparse M23C6 and M7C3) and base metal (extensive M7C3).

(e. g., voids and grain boundaries) (Chapter 1.18, Radiation-Induced Segregation). RIS can occur when a point defect flux interacts preferentially with a certain elemental species in the alloy, causing
that element to be enriched or depleted near the defect sinks. This preference can be driven kineti — cally (migration barriers) and/or thermodynamically (binding/ordering). RIS is segregation that occurs

1. 0E-10

л*.

о

‘ "O

(D

О

1.0E-11

CL

Figure 20 Illustration of the predicted crack growth rates for Alloy 600 HAZ material as a function of potential, stress intensity factor, and temperature.

in addition to thermal segregation. Like thermal seg­regation, this local change in composition can result in detrimental changes to mechanical and corrosion

83

properties.

RIS has been a concern in the nuclear industry for over 30 years and is considered one of the many factors that lead to irradiation-assisted stress corrosion cracking.6-9 This phenomenon was first predicted by Anthony84 and has been observed in a number of dif­ferent alloys and steels used in nuclear reactors.85-87 In both the iron — and nickel-based face-centered cubic Fe-Ni-Cr alloys, experimental RIS trends at grain boundaries are generally chromium depletion, nickel enrichment, and possible compensation through iron enrichment or depletion.85,86 In body-centered cubic ferritic-martensitic steels, both chromium enrichment and depletion have been reported.88,89 RIS in welds has not been extensively researched but the possibility ofgrain boundary depletion of chromium in corrosion-resistant
alloys that would act in addition to the chromium depletion that occurs during solidification segregation (e. g., Ni-Cr-Nb and Ni-Cr-Mo alloys) is of significant concern.

4.09.3.3 Microstructural Changes

In addition to the relatively short-time microstructural changes that can occur on-cooling or with postweld heat treatment (typically <10h) and induce PIC as discussed earlier, long-time microstructural changes can also occur. While most nuclear alloys have had sufficient vetting to preclude these concerns, recently developed high-alloy nickel-based filler metals raise the concern that topologically close-packed (TCP) phases (e. g., sigma) or long-range ordering could occur in Ni-Cr-Mo welds.10,50 The degradation of toughness by the formation of TCP phases is well established in the superalloy literature and the

formation of long-range order can lead to increased residual stresses and decreased resistance to EAC, most notably to hydrogen embrittlement.50,52,53,90