DDRs used to predict the embrittlement of the C-Mn steels used in the UK Magnox RPV had been mechanistically based from the 1980s.2 These predict radiation-induced changes in yield stress (hardening) or embrittlement (Charpy impact energy transition temperature or fracture toughness transi­tion temperature) as a function of radiation dose and temperature. The approach adopted has been set out in Jones and Bolton2 and Wootton et al.3,n The advantage of this approach was that the derived rela­tionships could be used with confidence when limited extrapolation was required into regions of neutron dose, dose rate, or irradiation temperatures that were not specifically included in the surveillance database.

It is important to note that it was the advances in understanding that enabled the adoption of a mecha­nistic approach (rather than adopting an empirical approach which had been followed in all other embrit­tlement correlations of this time). More specifically,
the seminal work of Fisher and coworkers in the early 1980s50 assumed that changes in yield stress arose from the combined effects of irradiation damage clusters and copper precipitates. Subsequently,2,3,114 a two-term relationship was finally adopted2 to model both hardening (As) and embrittlement (A T40j) and had the following form:

AT40J

or

As

This relationship follows the model of Fisher and coworkers,50 where Dcopper represents the contribution of nanoscale copper precipitation to the property change and Dmatrix the contribution from matrix hard­ening arising from the production of point defect clus­ters by neutron irradiation. A further simplification was made in developing a DDR that could be applied to operational Magnox reactors. Namely, under the con­ditions of irradiation dose and temperature of interest there was no overaging; that is, the contribution to hardening or embrittlement from Cu cluster formation would reach a peak and then remain constant. Fur­ther, the hardening from Cu clusters could be repre­sented by a constant at all doses of interest, clearly a conservative assumption at doses before which the hardening from Cu clusters had reached a peak. On this basis, mechanistically based DDRs of the form

AT40J 9

or = B V AFtVd [5]

As >

were adopted. In this equation, B represented the material-specific copper precipitate contribution to the property change, with the MD contribution being given by AFtVD. In this term, A is a material specific constant, D is the dpa dose, and Ft is the irradiation temperature dependence factor.2,35 The fact that B is a constant independent of the measured bulk Cu level is consistent with the effect of the low final stress-relief temperature on reducing the variation in the Cumatrix between different materials (see Buswell and Jones70).

DDRs were derived for the different RPV materi­als over the years. They were revised as and when new Charpy impact energy or tensile test data became available or following revisions to the neu­tron doses accrued by the surveillance specimens.114 For example, it was found that SMA welds are much more susceptible to the occurrence of intergranular fracture effects, with manual welds, plates, and forgings
showing minimal effects. DDRs had to be developed that accommodated a nonhardening embrittlement mechanism. In addition, it was established that thermal neutrons could make a significant contribution to the irradiation damage in side-core locations, and that they were not conservatively covered by the DDRs.115,116 This conclusion was reached from an analysis of sur­veillance data from samples irradiated in locations in reactors with different levels of thermal fluxes and also from a well-controlled irradiation in a heavy water moderated reactor in Halden. It was established that to allow for extra displacements from low-energy recoils (^500 eV), a thermal neutron effectiveness fac­tor (k) needed to be introduced to modify the dose term in each material DDR. This meant that the gen­eral form of the two-term DDRs for both embrittle­ment and hardening (eqn [5]) became

AT 9

or = B + AFTJ Df + kDt [6]

As >

In this equation, the definitions of B, A, and FT remained unchanged, but the single dose term, D, was replaced by (Df + kDt), where Df and Dt are the doses of fast dpa (redefined to be from neutrons of energy E > 1 keV) and thermal dpa (from neutrons of energy < 1 keV), respectively; the constant k is the thermal neutron effectiveness factor for the material. The thermal neutron effectiveness factor was found to be material dependent,114 and separate values were estimated for the different RPV materials. It should also be noted that a large-scale sampling and testing program of SMA weld metal removed from a decom­missioned RPV validated the assessment process11 detailed above.

Wotton eta/.114 note that as a result of successfully addressing these and other challenges when the last two steel pressure vessel stations closed in December 2006, they had achieved lifetimes of nearly 40 years. To quote the authors,11 ‘‘This radical approach was subjected to rigorous peer review and its acceptance by the UK Nuclear Installations Inspectorate (NII) regulator was a major achievement.’’ Part of the peer review process may be illustrated by Knott and cow — orkers117,118 which detail the result of independent peer review by the UK Technical Advisory Group on Structural Integrity (TAGSI) of, first, the principles underlying the assessment of mechanical properties of irradiated ferritic RPV steels, and, second, the effects of gamma irradiation dose on the properties of C-Mn steels used in RPVs.

4.05.5.2 US Mechanistically Guided DDRs

In the late 1990s, mechanistically guided correlations were developed to describe the embrittlement of RPV materials employed in the United States.119,120 In common with other mechanistic DDRs, the form of these correlations is determined by current mecha­nistic understanding, but, in this case, the coefficients employed in the various terms in the correlation are determined by fitting to the extensive mechanical property data on the embrittlement of RPV materials acquired in vessel surveillance programs (Charpy V-notch shift at the 41J transition temperature, AT41 J). These correlations are developed to describe a specific fleet of reactors, namely the embrittlement of RPV materials in US boiling water and pressurized water reactors (BWRs and PWRs).

The initial mechanistically based or guided corre­lation models were presented in NUREG/CR-6551, published in November 1998.119 The models dis­cussed incorporated material chemical composition and various exposure variables to enable predictions of TTS and USE changes. Another embrittlement shift model was developed at about the same time on the same mid-2000 database under the auspices of the Electric Power Research Institute (EPRI) and the American Society for Testing and Materials (ASTM) E10.02 subcommittee (the E900 model120), published as E900-02 in 2002. This was a simplified form but did not have a strong dependence on flux.121

The embrittlement shift model in NUREG/CR — 6551 was updated in July 2000 with additional sur­veillance data collected since the earlier work; this is referred to in this report as the Draft 2000 model.122 Significantly, motivation for a new modeling effort came from the fact that 62 additional low flux BWR shifts became available in 2003 (described below). These data were significantly underpredicted by the previous shift models,119,120 so it was necessary to investigate the cause of the underprediction. Addi­tional pressurized water reactor (PWR) data from surveillance reports (about 140 shifts) were also added to the database in 2003 and 2004. Finally, the reliability of the database was improved when all old and new surveillance data were reviewed for com­pleteness, duplicates, and discrepancies, during the summer and fall of 2004, in cooperation with the ASTM Subcommittee E10.02 on Radiation Effects in Structural Materials.123

The DDR that has been incorporated into the latest USNRC Regulatory Guide on screening limits for pressurized thermal shock124 was produced by

Eason eta/.123 This is the most explicitly mechanistic DDR for MnMoNi steels produced to date, referred to as ‘EONY’ for convenience after the authors. The DDR, which is much more complex than the mild steel DDRs discussed above, is

AT = MF + CRP (in°F) [7]

where

MF = A(1 — 0.001718Ti) (l + 6.13PMn2’47) vfe) [8]

and A = 1.140 x 10-7 for forgings, 1.561 x 10-7 for plates, and 1.417 x 10-7 for welds

Ti = irradiation temperature (° F); P = bulk P (wt%); Mn = bulk Mn (wt%)

ft for f > 4.39 x 1010ncm-2s-1 д 0.259

for f < 4.39 x 1010ncm-2s-1

= effective (flux-corrected) fluence

and	80
t T
CRP = B( 1 + 3.77NiL191) f (Cue, P) g(Cue, Ni; fte) [9]

where B = 102.3 for forgings; 135.2 for plates in vessels manufactured by Combustion Engineering (CE); 102.5 for non-CE plates; 155.0 for welds; 128.2 for plates of the standard reference materials (SRMs)

0 for Cu < 0.072wt%

min[Cuactuai, Cumax] for Cu > 0.072wt%

= effective Cu level [10]

in which Cuactual = bulk Cu level (wt%), Cumax = 0.243 for typical (Ni > 0.5) Linde 80 welds, and 0.301 for all other materials. (Equations [7] and [8] are in °F, reflect­ing units in the original reference. It is to be noted that °F are employed in USNRC regulatory guides, rather than SI units.)

0 for Cu < 0.072

[Cue — 0.072]0′[1] for Cu > 0.072 and P < 0.008

[Cu — 0.072 + 1.359(P — 0.008)]0,668 for Cu > 0.072 and P > 0.008

[11]

g (Cue, Ni, fte)

1 1 unit [ log10 (ft) e + 1.139Cue — 0.448Ni — 18.120′ = 2 + 2tan 0.629

20

о

(a)

dependence of both the matrix and CRP term is broadly consistent with the understanding outlined in the previous section. The concept of f te is partic­ularly important as it both provides a means of allow­ing for flux effects and gives a threshold below which flux effects might be expected.63 These trends are further illustrated in Figure 16.

Overall, for a Cu-containing steel (say 0.2-0.3 wt% Cu), the MD becomes a significant fraction of the damage only at doses beyond the plateau in the shift from CRPs. This is consistent with the hardening from MD inferred from microstructural data. Carter et al. examined the effect of irradiation on microstructure on a high copper Linde 80 flux weld BW2 (0.25 wt% Cu, 0.62 wt%Ni,

0. 017 wt% P),125 and concluded that out of a total hardness of DHvtot 40 ± 6 the hardness from MD was

DHvMatrix 5-10 VPN.

In Section 4.05.2, it was described how irradiation also caused a drop in the Charpy USE. It is to be noted that Eason et a/.22 used the US surveillance power reactor database to investigate the dependence of the USE drop (DUSE) on a number of variables. They demonstrated that there was a strong correla­tion between the DUSE and Charpy TTS at 30 ft-lbs. Eason et al. derived a detailed set of equations that allowed the DUSE to be determined from the TTS for a number of product forms.

4.08.4.1 Strength Anisotropy

When JAEA started to develop ODS steels in 1985, the ferritic type of ODS steels was applied.3, These are similar to MA957,34 which is single ferrite phase and does not include the martensite. Based on the results of R&D conducted for several years, three kinds of claddings, 63DSA, 1DK, and 1DS, were manufactured in 1990. Their chemical compositions

are 13Cr-0.02C-3W-0.7Ti-0.46Y2O3 (63DSA), 13Cr — 0.05C-3W-0.5Ti-0.34Y2O3 (1DK), and 11Cr-0.09C- 3W-0.4Ti-0.66Y2O3 (1DS). The manufacturing process is almost the same as the process shown in Figure 13, except for the rolling process and intermediate heat treatment, because cold-rolling processing can be hardly applied to these ODS steels. In the case of the 1DK cladding, six warm drawings at 800-850 °C, followed by four warm rolling passes at 500 °C with intermediate annealing at 1080 ° C, were repeated to manufacture the thin-walled cladding in the dimension of 7.5 mm outer
diameter, 0.4 mm thickness, and 1 m length. In the case of the 63DSA and IDS claddings, only six warm rolling passes at 650-700 °C with intermediate annealing at 1100 ° C were conducted. The temperature of the final heat treatment of 1DK cladding was 1150°C for 60s, and 63DSA and 1DS claddings at 1100 for 3.6 ks.

The uni — and bi-axial creep rupture strengths of the manufactured claddings at 650 °C are shown in Figure 17, where the uni-axial corresponds to the hot working direction and bi-axial belongs to the internal hoop direction.3 It was found that

there is strong strength anisotropy, and the bi-axial creep rupture strength is considerably lower than that of the uni-axial direction. Microstructure obser­vations of these claddings exhibited the elongated grains like a bamboo structure in parallel to the working direction. The strength degradation in the bi-axial/internal hoop direction, which is essential for the fuel elements, should be mainly attributed to the grain boundary sliding and crack propagation due to stress concentration.

Graphite has been used as a nuclear moderator in nuclear fission reactors since the very beginning of the nuclear age.1 Indeed, the Chicago Pile No. 1 (CP-1), constructed by Enrico Fermi under the stands at Stagg Field, University of Chicago, used The National Carbon Company’s AGOT grade graphite. On 2 December 1942, Enrico Fermi and his research team achieved the world’s first nuclear chain reaction in CP-1. Subsequently, the early weapons materials production reactors constructed in the United States, United Kingdom, France, and the former Union of Soviet Socialist Republics (USSR) were all graphite­moderated reactors, as were the first commercial power generating fission reactors.

The core of a graphite-moderated reactor is com­prised of stacks of graphite blocks that are usually keyed to one another to facilitate transmission of mechanical loads throughout the core. Vertical chan­nels penetrate the core into which fuel stringers are placed via the reactor charge face using a refueling machine. The nuclear fuel, which may be natural ura­nium (a mixture of 238U, 235U, and 234U) or enriched uranium, is usually sheathed (clad) in a metallic clad­ding. Typically, the cladding is a light alloy (aluminum or magnesium), but it can also be stainless steel (requir­ing enriched fuel) if a higher fuel temperature is desired (>600 °C). The fuel stringer and cladding material may be one and the same, as in the United Kingdom designed Magnox reactor,1 or the fuel may be in the form of stainless steel clad ‘pins’ arranged in graphite fuel sleeves, which are joined to one another and form the fuel stringer as in the UK Advanced Gas-Cooled Reactor1 (AGR). The metallic fuel clad retains the gaseous fission products that migrate from the fuel during the fission reaction and prevents con­tact between the gaseous coolant and the fuel.

An alternative core layout uses integral fuel/mod — erator elements in which the uranium fuel is placed directly into cavities in the graphite moderator block,

Table 1 Currently available nuclear grade graphites

Grade Manufacturer Coke type Comments

Isostatically molded, candidate for high-dose regions of NGNP concepts

Isostatically molded, candidate for high-dose regions of NGNP concepts

Extruded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite Vibrationally molded, candidate for high-dose regions of NGNP prismatic core concepts Vibrationally molded, candidate for high-dose regions of NGNP pebble bed concepts; PBMR core graphite Extruded, candidate for high-dose regions of NGNP prismatic core concepts

Large blocks for permanent structure in a prismatic core Isostatically molded, candidate for permanent structures in a prismatic core

Isostatically molded, candidate for permanent structures in a prismatic core

and the entire block is discharged from the reactor when the fuel is spent. Fuel elements of this design typically utilize ceramic (UO2 or UC2) rather that metallic fuel so as to be capable of reaching higher fuel temperatures. The ceramic fuel kernel is over coated with layers of SiC and pyrolytic carbon to pro­vide a fission product barrier and to negate the use of a metallic fuel clad (see Chapter 3.07, TRISO — Coated Particle Fuel Performance), allowing the reactor core to operate at very high temperatures (> 1000 “C).1 The coated particle fuel is usually formed into fuel pucks or compacts but may be consoli­dated into fuel balls, or pebbles.1 The US designed modular high-temperature gas-cooled reactor (MHTGR) and Next Generation Nuclear Plant (NGNP), and the Japanese high-temperature test reactor (HTTR) are examples ofgas-cooled reactors with high-temperature ceramic fuel.

Additional vertical channels in the graphite reac­tor core house the control rods, which regulate the fission reaction by introducing neutron-adsorbing materials to the core, and thus reduce the number of neutrons available to sustain the fission process. When the control rods are withdrawn from the core, the self-sustaining fission reaction commences. Heat is generated by the moderation of the fission frag­ments in the fuel and moderation of fast neutrons in the graphite. The heat is removed from the core by a coolant, typically a gas, that flows freely through the core and over the graphite moderator. The coolant is forced through the core by a gas circulator and passes into a heat exchanger/boiler (frequently referred to as a steam generator).

The primary coolant loop (the reactor coolant) is maintained at elevated pressure to improve the cool­ant’s heat transfer characteristics and thus, the core is surrounded by a pressure vessel. A secondary coolant (water) loop runs through the heat exchanger and cools the primary coolant so that it may be returned to the reactor core at reduced temperature. The secondary coolant temperature is raised to produce steam which is passed through a turbine where it gives up its energy to drive an electric generator. Some reactor designs, such as the MHTGR, are direct cycle systems in which the helium coolant passes directly to a turbine.

The reactor core and primary coolant loop are enclosed in a concrete biological-shield, which pro­tects the reactor staff and public from g radiation and fission neutrons and also prevents the escape of radioactive contamination and fission product gasses that originate in the fuel pins/blocks. The charge face, refueling machine, control rod drives, discharge area, and cooling ponds are housed in a containment structure which similarly prevents the spread of any contamination. Additional necessary features of a fission reactor are (1) the refueling bay, where new fuel stringers or fuel elements are assembled prior to being loaded into the reactor core; (2), a discharge area and cooling ponds where spent fuel is placed while the short-lived isotopes are allowed to decay before the fuel can be reprocessed.

The NGNP, a graphite-moderated, helium cooled reactor, is designed specifically to generate elec­tricity and produce process heat, which could be used for the production of hydrogen, or steam gener­ation for the recovery of tar sands or oil shale.

Two NGNP concepts are currently being consid­ered, a prismatic core design and a pebble bed core design. In the prismatic core concept, the TRISO fuel is compacted into sticks and supported within a graphite fuel block which has helium coolant holes running through its length.1 The graphite fuel blocks are discharged from the reactor at the end of the fuel’s lifetime. In the pebble bed core concept, the TRISO fuel is mixed with other graphite materi­als and a resin binder and formed into 6 cm diameter spheres or pebbles.1 The pebbles are loaded into the core to form a ‘pebble bed’ through which helium coolant flows. The pebble bed is constrained by a graphite moderator and reflector blocks which define the reactor core shape. The fuel pebbles migrate slowly down through the reactor core and are discharged at the bottom of the core where they are either sent to spent fuel storage or returned to the top of the pebble bed.

Not all graphite-moderated reactors are gas — cooled. Several designs have utilized water cooling, with the water carried through the core in zirconium alloy tubes at elevated pressure, before being fed to a steam generator. Moreover, graphite-moderated reactors can also utilize a molten salt coolant, for example, the Molten Salt Reactor Experiment (MSRE)1 at Oak Ridge National Laboratory (ORNL). The fluid fuel in the MSRE consisted of UF4 dissolved in fluorides of beryllium and lithium, which was circu­lated through a reactor core moderated by graphite. The average temperature of the fuel salt was 650 °C (1200 °F) at the normal operating condition of 8 MW, which was the maximum heat removal capacity of the air-cooled secondary heat exchanger. The graphite core was small, being only 137.2 cm (54 in.) in diameter and 162.6 cm (64 in.) in height. The fuel salt entered the reactor vessel at 632 °C (1170 °F) and flowed down around the outside of the graphite core in the annular space between the core and the vessel. The graphite core was made up of graphite bars 5.08 cm (2 in.) square, exposed directly to the fuel which flowed upward in passages machined into the faces of the bars. The fuel flowed out of the top of the vessel at a temperature of 654 °C (1210 °F), through the circulating pump to the primary heat exchanger, where it gave up heat to a coolant salt stream. The core graphite, grade CGB, was specially produced for the MSRE, and had to have a small pore size to prevent penetration of the fuel salt, a long irradiation lifetime, and good dimen­sional stability. Moreover, for molten salt reactor mod­erators, a low permeability (preferably <10~8 cm2 s-1) is desirable in order to prevent the build up of unacceptable inventories of the nuclear poison 135Xe in the graphite. At ORNL, this was achieved by sealing the graphite surface using a gas phase carbon deposition process.1

As previously discussed, graphite crystal structures, in the form ofHOPG, have been observed to swell in
the ‘C axis direction and contract in the ‘a’ axis direction for all measured fluence and irradiation temperatures. Figure 19 shows early data obtained by Kelly eta/.35 It is clear from Figure 19 that the rate of swelling and shrinkage significantly changes between 200 and 250 0C, indicating that the defect population is becoming more stable above this tem­perature range.

HOPG data is an important input into multiscale models of irradiation damage in graphite.48 For the purpose of understanding irradiation damage in operating reactors, data would be required ideally from 140 to 1400 0C, the maximum fluence being dependent on the irradiation temperature. Unfortu­nately, the dataset is far from complete. The data due to Brockelhurst and Kelly49 is the most complete set of HOPG irradiation data covering the fluence and part of the temperature range appropriate to AGRs (Figure 20). In the same paper, the authors show the effect of final heat treatment, between 2000 and 3000 0C, on the crystal dimensional change rate of HOPG. The data showed that the lower the heat treatment, the faster is the dimensional change rate, indicating that the dimensional change rate ofa poorly graphitized component would be expected to be greater than that of majority of the components.50 Pre­viously, Kelly and Brocklehurst51 had shown that boron doping also significantly increased the dimensional

change rate in HOPG, and this was again reflected in the behavior of doped polycrystalline graphite.50

HOPG high-temperature data was mainly obtained by investigators interested in the behavior of HTR fuel coatings.52 Some of this data is for low-density pyrolytic carbons and it is not always made clear which material the data refers to. Figure 21 shows all the data known to the author, and it is clear that there is some inconsistency.

In the fusion materials development strategy, the candidate structural materials are categorized into reference and advanced materials. As the reference

materials, RAFM steels were selected because they have the most matured industrial infrastructure. Development of the reference materials is crucial for the realization of DEMO (fusion demonstration power plant) in a timely manner. On the other hand, vanadium alloys and SiC/SiC were nominated as

the advanced materials, which will contribute to increasing attractiveness of the fusion system in terms of cost of electricity and environmental benign­ness. It is recognized that the development of the advanced materials must also be enhanced now due to the long lead time necessary for their development. It should also be noted that vanadium alloys are the only nonferromagnetic and ductile materials of the three candidates. If the impact of the ferromagnetism ofthe RAFM on plasma operation should be unavoid­able and the brittleness of SiC/SiC should be deter­mined unaccepted by design studies, vanadium alloys could be the only candidate of low activation structural materials for fusion reactors.

As shown in the summary of critical issues, a 14MeV neutron source is highly necessary for the qualification of vanadium alloys. IFMIF (Interna­tional Fusion Materials Irradiation Test Facility, a 14 MeV neutron source) is under design and is recog­nized to be essential for developing structural materials for fusion reactors. The TBM to be installed in ITER is also considered to be an important milestone for technological integration. Figure 24 shows the design of the V/Li TBM in ITER proposed by Russia.50 The development of vanadium alloys is planned to pro­ceed with IFMIF for qualification of the alloy and ITER-TBM for technology integration, in addition to fundamental studies using fission reactors, etc.

4.12.9 Summary

As to the application in nuclear systems, vanadium alloys were once candidate cladding materials for LMFBR, but, at present, are considered mostly as candidate low activation structural materials for fusion reactors.

Vanadium alloys, with the reference composition of V-4Cr-4Ti, are one of the three candidate low activation structural materials with RAFM and SiC/SiC. They are the only nonferromagnetic and ductile materials of the three candidates and thus are promising for advanced structural materials of fusion reactors. The self-cooled liquid lithium blanket using structural materials of vanadium alloys is an attractive concept because of the high heat transfer capability, high-temperature operation, simple structure, high tri­tium breeding capability, and low tritium leakage.

Recent progress, especially in the fabrication technologies, has successfully reduced the number of critical issues enhancing the feasibility of the alloys for fusion application. Major remaining issues of vanadium alloys are thermal and irradiation creep, transmutant helium effects on mechanical properties, and radiation effects on fracture properties. For con­clusive characterization of the helium effects, the use of IFMIF is essential.

Efforts are also being made to develop advanced vanadium alloys to extend the temperature window and lifetime of vanadium alloys in fusion reactor environments.

Recently, it been discovered that significant levels of hydrogen can be stored in bubbles and voids in both stainless steels and pure nickel when the hydrogen is cogenerated with helium, especially in light water spectra where there are also environmental sources of hydrogen.73-75 It was shown in these studies that this phenomenon is a direct result of the 59Ni nuclear reactions. Previously, it was a long-standing percep­tion that such storage could not occur at reactor­relevant temperatures.

The retained hydrogen levels are in significant excess of the levels predicted by Sievert’s Law and appear to be increasing with both cavity volume and neutron fluence. Since these gases are known to assist in nucleation and stabilization of cavities, it is expected that the nonlinear 59Ni reactions discussed earlier may lead to a rapidly developing, nonlinear, cavity-dominated microstructure in stainless steels irradiated at temperatures characteristic of pressur­ized water reactors.

Figure 17 presents such a microstructure observed in a PWR flux thimble tube (cold-worked 316 stain­less steel) at ^70 dpa and 330 °C.76 There is a very high density (>1017cm~3) of nanocavities with dia­meters <3 nm in both the alloy matrix and especially on grain boundaries. The measured concentrations of 600 appm He and 2500 appm H in this specimen are thought to reside primarily within the cavities. Most importantly, these cavities are essentially invisible
under well-focused imaging conditions and can only be imaged using very large levels of under-focus. This implies that previous studies on similar mate­rials may have overlooked such cavity-dominated structures.

When this specimen and near-identical specimens were subjected to slow strain rate testing after irradi­ation, the fracture surface was indicative of ^100% intergranular stress corrosion cracking (IGSCC), with lower doses and gas levels producing propor­tionally less IGSCC.77 As hydrogen is known to be a contributor to grain boundary cracking, it appears plausible that hydrogen storage may accelerate the cracking process and that higher exposure will lead to an increasing susceptibility to cracking. This issue may therefore become increasingly important as PWRs previously licensed for 40 years are being considered for life extension to 60 and possibly 80 years.

4.04.4.1 Dislocation Structures

Dislocation structures in irradiated pressurized tube samples were examined by Gelles et al.72 The materi­als which were examined included stressed and unstressed samples of ST PE16, and stressed samples of ST and STA Inconel 706. A subsequent paper by Gelles73 extended these investigations to the stressed samples of PE16 in STA and OA conditions. Further details of this work were also provided by Garner and Gelles74, and by Gelles.70

Examination of ST PE16, which was irradiated at 550 °C to 2 x 1026nm~2 (E> 0.1 MeV) at hoop stres­ses of 0 and 167 MPa, revealed that the distribution of Frank dislocation loops was similar on all the four {111} planes in the unstressed sample but was aniso­tropic in the stressed material. In the stressed sample, the loop density on any particular {111} plane increased with increasing magnitude of the normal stress component on that plane. A near-linear rela­tionship between the loop density and the normal component of the deviatoric stress tensor, OdN (= sN — Oh, where sN is the normal component of the applied stress on a particular plane and sH is the hydrostatic stress), was found for PE16. This result is in line with the SIPA loop growth model described by Garner eta/.75 No such correlation was found in the similarly irradiated and stressed Inconel 706 samples, however, possibly because the low creep rate of this material at 550 °C did not allow the relaxation of internal stresses.

Unfaulting of Frank dislocation loops with a/3 {111} Burgers vectors proceeds via interaction with a/6{112} Shockley partials to produce perfect a/2 {110} line dislocations. Gelles70 described how this occurs via a two-step process, with the necessary partial dislocations (two per interstitial loop) first being nucleated by an interaction of the faulted loop with a suitable perfect dislocation and then sweeping across the loop to reestablish the perfect dislocation. Gelles73 examined the distribution of Burgers vectors among the six possible a/2{110} perfect dislocation types in irradiated pressur­ized tube samples of PE16. The samples examined included the stressed ST condition irradiated at 550 °C, and STA and OA conditions which were both irradiated at 480 °C to a fluence of 8 x 1026nm—2 at a hoop stress of 331 MPa. The results showed highly anisotropic distributions in the Burgers vectors of perfect dislocations in all the three heat-treated conditions, with dislocation den­sities of the various types differing by factors of up to 10-40 in each sample. The level of anisotropy pro­duced in the population of perfect dislocations was significantly greater than in the dispersion of Frank loops. This is a feasible outcome since, in principle, all loops may be unfaulted by just two variants of the six a/2{110} perfect dislocation types. In effect, the development of anisotropic dislocation structures is a response of the material to produce the strain which is required to accommodate the applied stress. Furthermore, it was found that the perfect disloca­tions in the irradiation creep samples of PE16 were primarily of edge type lying on {100} planes rather than {111} slip planes, indicating that they could only contribute to the creep strain via climb (i. e., by the SIPA mechanism) and not by processes involving dislocation glide.

Two earlier reviews of the irradiation-induced prop­erties of Mo and TZM have been presented as part of the UWMAK-III fusion reactor study86 and for the SP-100 space nuclear power program.19 Much of the known swelling data on irradiated Mo is contained in these reviews, with the majority of data for irradia­tions <10dpa and temperatures below 1073 K. The swelling data available are considerably scattered, with little coherence to examinations on the swelling as a function of temperature or dose.

Swelling in Mo is expected to begin around 573­673 K and continue to temperatures near 1573 K.19 Maximum swelling in pure Mo remains below 4% for fluences up to 1 x 1023 ncm~2 (E > 0.1 MeV), ^50dpa, with peak swelling at irradiation tempera­tures near 900 K. Attempts at consolidating the reported swelling data as a function of irradiation temperature through normalizing the fluences proved
to be inaccurate in determining the upper bound limit for maximum swelling. 9 The swelling data col­lected from numerous sources,53,87-89 including those contained in the review work of Brimhall et a/.90 for irradiated Mo as a function of dose and irradia­tion temperature, are provided in Figure 12. Void swelling was found to be < 1% in Mo irradiated to 8 x 10 n cm, E > 0.1 MeV at temperatures between 673 and 1173 K by Evans.53 Void swelling studied by Stubbins et a/.88 in 3.1 MeV 51V+ ion-irradiated Mo between 1173 and 1393 K up to 50dpa remained below 4%, while irradiations between 1523 and 1780 K were near 10%.

Void ordering has been observed in both neutron-

irradiated28,89,91 and ion-irradiated Mo88 at tempera­tures between 700 and 1373 K. Garner and Stubbins89 examined the irradiation and material conditions that contribute to void ordering. Irradiation tempera­tures near 700 K delineate the lower boundary tem­perature for void lattice formation at irradiations above 20 dpa. At lower doses, void lattice formation was not observed. The void superlattice constant, mea­sured as the distance between void centers along the <100> direction in the material, is found to increase with temperature from ^2.4 nm at 700 K to 4.5 nm
at 1176 K.91 Swelling is expected to reach a maximum of 3-4% on the development of the void lattice struc­ture, based on an attainment of an equilibrium ratio of void diameter to void superlattice parameter.92 At temperatures >1423 K, void lattice formation is no longer observed, leading to the high values of swelling observed in the material ion irradiated to high doses.88

The onset of void growth in neutron-irradiated material appears to be accelerated in cold-worked materials compared to annealed materials, reaching a maximum in swelling at doses near 40 dpa for tem­peratures below 873 K and 20 dpa at higher tempera — tures.89 At higher doses, swelling decreases through void shrinkage, with swelling values approaching those of annealed materials. Void shrinkage has also been reported by Bentley eta/.93 and Evans53 to occur because of changes in the void sink bias89 presum­ably due to the segregation of transmuted species at the void surfaces, making them more attractive for interstitials.

Irradiation-induced swelling in TZM has been reported53,94-97 and generally shows similar temper­ature dependence as the pure metal. The fluence and temperature dependence of swelling of TZM was examined by Powell et a/.95 and Gelles et a/.,94 with

results from the latter shown in Figure 13. Peak swelling in TZM following irradiation to 1.78 x 1023 ncm~2 and 873 K remained below 4%, though the data are limited to irradiation temperatures below 923 K. Only limited data are available on direct com­parisons between TZM and pure Mo, with Bentley and Wiffen96 reporting 1% swelling in Mo-0.5%Ti and TZM alloys and 0.6% swelling in pure Mo under the same irradiation conditions. Similarly, 4% swelling was observed in TZM and 3% in pure Mo following irradiation to 5.4 x 1022ncm~2 at 923 K.97 In examining Mo and TZM of different preirra­diated material conditions, Evans53 observed equal or greater swelling in TZM compared to Mo follow­ing irradiation to 3.5 x 1022ncm~2 (E> 0.1 MeV) at 823 and 873 K. However, in the materials irradiated at 723 K for the same fluence, the TZM alloy showed lower swelling, except in the carburized condition. The Ti and Zr atoms not tied up as carbides are assumed to have played a role in reducing void size in the material at the lower temperature.

There is little information on the swelling behav­ior of Mo-Re alloys. Measured swelling of 0.44% in Mo-50Re irradiated to 5.3 x 1022ncm~2 (E > 0.1 MeV) at temperatures which rose during irradiation from 1128 to 1329 K was reported.26 For irradiated Mo-Re alloys, radiation-induced segregation (RIS) and transmutation can lead to precipitation of
equilibrium or nonequilibrium phases, which can be detrimental to mechanical properties. This is examined in the next section.

Electrical resistivity changes to 5.4 dpa irradiated Mo at 733 K were examined by Zakharova et a/.98 using single crystal samples. Increases in resistivity of 10-14% and 92-110% were measured at postirra­diation test temperatures of 298 and 77 K, respec­tively. The largest resistivity changes were measured in the [100] direction. A residual 10% increase in resistivity was measured following annealing above 0.6 Tm associated with the accumulation of trans­muted radionuclides.

The changes in electrical resistivity of LCAC-Mo over a 353-1373K irradiation temperature range up to 3.3 dpa were examined by Li eta/.99 and Cockeram et a/, with the latter examining the recovery of

resistivity following isochronal anneals. The room temperature resistivity for 353 K irradiated LCAC — Mo rapidly increases between 0.01 and 0.1 dpa saturating near 0.2 dpa for an ^42% increase over the unirradiated value.99 Increases in room tempera­ture resistivity of 10-12% were reported following 0.5-1.2 dpa irradiation at 543 K, and 3.3-5.3% after 1.4-2.4dpa at 878 K. At irradiation temperatures >1208 K, little (<3%) to no net increase in resistivity was observed for irradiations up to 3.3 dpa. This is reflected in the higher mobility of vacancies and

interstitials formed during irradiation to diffuse to sinks where annihilation occurs, reducing the electri­cal scattering effects that these defects have at lower irradiation temperatures. The small increases mea­sured at the higher irradiation temperatures were pri­marily due to transmutation products. As is shown in the next section, the changes in electrical resistivity with increasing irradiation temperatures also correlate with changes in measured hardness, though at a greater level of sensitivity. This is controlled by microstruc­tural changes, as the small dislocation loops and voids of high distribution density appearing at the lower irradiation temperatures coarsen into larger and fewer defects that have less interaction with deforma­tion dislocations.

4.08.8.1 Simulated Irradiation

Testing that involves the simulated irradiation of 9Cr-ODS steel was conducted by Allen et al. at the

Environmental and Molecular Science Laboratory at Pacific Northwest National Laboratory, using 5 MeV Ni ions at 500, 600, and 700 °C with a damage rate of 1.4 x 10~3dpas~ The results regarding measured particle size distribution as a function of dose are plotted in Figure 38 for irradiation at 500, 600, and 700 °C.60 Due to TEM’s limited resolution of the images, particles smaller than 2 nm were not detected. At all temperatures, the size of the oxide particles decreases as the dose increases. At higher tempera­tures (600-700 °C), the average size appears to reach a value of ^5 nm. At all three temperatures, the density increases as the radiation dose increases. The decrease in size takes place faster at 600 and 700 °C than at 500 °C, indicating that the reduction in size is not strictly a ballistic effect and that a diffusion-based mechanism is also involved in the dissolution.

Allen extensively reviewed previous papers that presented different approaches to the irradiation of ODS ferritic-martensitic steels that employed various ion beams, electrons, and neutrons; the results are summarized in Table 3.61 A great many findings asserted that oxide particles are stable under radiation. However, as shown in Table 4, the dissolution of oxide particles at higher temperatures and doses has been reported in other studies. Dubuisson62 and Monnet63 reported that small oxides dissolved under radiation at higher temperatures and doses, but did not dissolve at a lower irradiation dose. Their data will be dis­cussed in detail in the following section. In material irradiated in the JOYO fast reactor at temperatures 450-561 °C to doses of 21 dpa, Yamashita found that small particles disappear and average particles increase slightly in size with increasing temperature or dose.64 Monnet supplemented neutron radiation studies with the electron irradiation of yttrium oxides and magnesium oxides in the EM10 alloy at tempera­tures between 300 and 550 °C, and to doses of 100 dpa. In these studies, the yttrium oxides were stable at 400 °C when irradiated with 1.0 MeV electrons, but dissolved under 1.2 MeV electron irradiation.

Allen59 pointed out that the displacement energy for Y and O in yttrium oxide is 57 eV65,66 while that for iron is 40 eV. Assuming similar displacement energies in the Y-Ti-O oxide, the radiation-induced vacancy concentration should be larger in the metal matrix, providing a driving force for a net vacancy flux to the precipitate. This could drive the precipi­tate mass loss if vacancy absorption frees a precipitate atom. From a comparison between electron irradia­tion (Frenkel pairs) and ion irradiation (displacement cascades), Monnet63 also concluded that the ballistic ejection of atoms alone cannot be responsible for the loss of diameter in oxide particles. Free point defects and their diffusion-based mechanism are therefore of major importance and play a dominant role in the dissolution of oxide particles.

The coefficient of thermal expansion (CTE) as measured for natural graphite and HOPG is temper­ature dependent (Figure 2) and the data from a number of authors has been collated by Kelly.6 The room temperature values of CTE are about 27.5 x 10—6K—1 and —1.5 x 10—6K—1 in the ‘c’ and V directions, respectively.

4.11.2.2 Modulus

The crystal elastic moduli6 are Cn (parallel to the basal planes) = 1060.0 x 109Nm—2, C12 = 180.0 x 109Nm—2, C13 = 15.0 x 109Nm—2, C33 (perpendicular to the basal planes) = 34.6 x 109Nm—2, and C44 (shear of the basal planes) = 4.5 x 109Nm—2 as defined by the orthogonal co-ordinates given below:

[1]

The strength of the crystallite is also directly related to the modulus, that is, the strength along the basal planes is higher than the strength perpendicular to the planes, and the shear strength between the basal panes is relatively weak.

4.11.2.3 Thermal Conductivity

The thermal conductivity of graphite along the basal plane ‘a’ direction is much greater than the thermal conductivity in the direction perpendicular to the basal plane ‘c.’ At the temperature of interest to the nuclear reactor engineer, graphite thermal con­duction is due to phonon transport. Increasing the temperature leads to phonon-phonon or Umklapp scattering (German for turn over/down). Imperfections in the lattice will lead to scattering at the boundaries.