Extensive research and studies have been carried out to determine the durability of concrete under various service conditions, and thus information on the pro­gressive changes in the physical and chemical nature of concrete under such conditions is available through technical committees and publications of organizations such as the ACI, International Union of Laboratories and Experts in Construction Materi­als, Systems and Structures (RILEM), and Interna­tional Federation for Structural Concrete (CEB-FIP). Available damage models for reinforced concrete in large measure have addressed corrosion of steel reinforcement in concrete and new construction. Improved damage models and guidelines for their use are desired to predict failure probability of a degraded concrete structure, either at present or at some future point in time. Additional investigations also are desired with respect to synergistic effects involving more than one degradation factor and the interaction of loading and environment.

In the case of zirconium alloys, many authors have studied the postirradiation microstructure by using transmission electron microscopy (TEM). In 1979, an international ‘round robin’ was undertaken consisting of TEM observations of neutron-irradiated recrys­tallized zirconium alloys45 in order to determine the nature of the point-defect clusters. A more recent compilation of observations is given by Griffiths.46 It has been now proved by numerous authors that in zirconium alloys mainly dislocation loops with (a) Burgers vector can be found. Only for high fluence, the (c) component dislocation loops appear. Cavities are observed only in very specific cases.

4.01.1.3.1 (a) Dislocation loops

It is now clearly established by numerous authors45-57 that for commercial neutron-irradiated zirconium alloys (e. g., annealed Zircaloy-2 described in Northwood eta/.45) at temperatures between 250 and 400°C and for irradiation dose lower than 5 x 1025 nm~2, the point-defect clusters that can be observed by TEM (>2 nm) consist of perfect dislocation loops, either of vacancy or interstitial nature, with Burgers vector (a) = 1/3(1120), situated in the prismatic planes with typical diameter from 5 to 20 nm, depending on the irradiation temperature (Figures 3 and 4). These loops are found in very high density, typically between 5 x 1021 and 5 x 1022 m~3 depending on the irradiation temperature (Figure 5). , 1 The three (a) Burgers vectors are equally represented. Thorough studies of neutron damage in zirconium using the high-voltage electron microscope (HVEM) have also been given.53,58,59

The proportion of vacancy loops to interstitial loops depends on the irradiation temperature. Indeed, it is observed that for an irradiation temper­ature of 350 °C approximately 50% of observed loops are vacancy loops, whereas for an irradiation temper­ature of 400 °C, 70% of loops are vacancy loops.45,46 For a low irradiation temperature (below 300 ° C), the majority of loops present in the material are of the interstitial type.

The loop habit plane is close to the prismatic plane, but accurate determination proves that the loops are not pure edge but their habit plane is usually closer to the first-order prismatic plane {10Ї0}. The authors have also observed that for loop diameters lower than 40 nm the loops are circu­lar but for diameters larger than 40 nm the vacancy loops become elliptical with the great axis along the (c) axis, the interstitial loops remaining circular. The (a) loops also appear to be aligned in rows parallel to

the trace of the basal plane.46,50

For an irradiation temperature of 300 °C, no dislocation loop can be observed below a neutron fluence of 3 x 1023 nm~2 in the case of annealed Zy-2 (Zircaloy-2) irradiated at 300 °C.51 However, from this fluence, the loop density increases rapidly with increasing fluence but saturates at a density of 3 x 10 m, from a relatively low fluence of approx­imately 1 x 1024 nm~2 (Figure 5). The loop density saturation has been confirmed by X-ray analysis.60 The loop size exhibits a parabolic increase with fluence but no clear saturation in the evolution of the loop size is seen even after a fluence of 1 x 1026 n m~2.51,67

Increasing the irradiation temperature leads to a decrease in the loop density and to an increase of the loop size.45,55,61 Indeed, it was shown by Northwood etal.45 that neutron irradiation performed at 350 °Cof annealed Zy-2 up to a fluence of 1 x 1025 n m~2 leads to a mean loop diameter between 8 and 10 nm and a loop density between 8 x 1021 and 5 x 1022 m~ ; whereas a neutron irradiation of the same alloy per — formedat400 ° Cup to a fluence of 1 x 1025nm~2 leads to a mean loop diameter between 16 to 23 nm and a loop density between 4 x 1021 and 2 x 1022 m~3.45 Above 500 °C, no irradiation damage is formed.52 The (a) loop microstructure is found to be very sensitive to alloying elements such as oxygen. Indeed, for high — purity zirconium with very low oxygen content, the (a) loops are large and in low density, whereas for commercial zirconium alloys (with oxygen content between 1000 and 1500 ppm) the growth speed of loops is considerably reduced yielding smaller loops

in much higher density.45,55

It was also reported from TEM observations that a particular band contrast of alternative black and white was superimposed on the usual radiation damage nor­mally visible on thin foils of irradiated materials. This phenomenon has been connected to the alignment of the loops in the same direction and is believed to be a thin-foil artifact. It has been named ‘corduroy’ contrast by Bell.62 The commonly accepted explanation of this artefact is based on the local elastic relaxation of the internal stresses in TEM thin foils, in areas where pronounced alignment of (a) loops is present.63

Various aspects of behavior resulting in irradiation creep can be illustrated with some examples pre­sented in Figures 68-75.

4.02.9.2 Creep Disappearance

The previous figures demonstrate the swelling-creep correlation at its simplest when swelling is either zero or just beginning, but not yet provoking the next shift in quasi-equilibrium. When looking across a wider

Figure 68 Creep-induced deflections of helical springs constructed from two steels with different composition that were irradiated in DMTR at 100 °C, normalized to the elastic deflection, showing that both the transient and steady-state creep rate B0 are proportional to the stress level. While the transients are different in the two steels, the posttransient creep rate is independent of composition. Reproduced from Lewthwaite, G. W.; Proctor, K. J. J. Nucl. Mater. 1973, 46, 9-22. The maximum dose is ~0.5dpa.

Figure 67 Linear stress dependence of total diametral strain (creep and swelling) for 20% cold-worked PCA (Ti-modified 316 stainless) pressurized tubes irradiated in FFTF at 400°C. Reproduced from Garner, F. A.; Puigh, R. J. J. Nucl. Mater. 1991, 179-181, 577-580. Stress-free swelling is approximately three times the Y-intercept value with the largest swelling at ~8%.

Figure 73 Swelling and creep strains observed in two French steels irradiated as pressurized tubes in PHENIX, showing strong correlation between the two types of strain as the swelling rate increases. Reproduced from Dubuisson, P.; Maillard, A.; Delalande, C.; Gilbon, D. D.; Seran, J. L. In Effects of Radiation on Materials: 15th International Symposium; STP 1125; 1992; pp 995-1014.

range of swelling behavior some unusual behaviors are often observed. An example is shown in Figure 76 where the two-peaked swelling behavior frequently observed in 300 series steels is mirrored in the creep strains, but the relative proportions of the two strains are distorted.172 This is one manifestation of the creep disappearance phenomenon in which the attainment of significant swelling causes irradiation creep to strongly drop in rate or even to disappear under some conditions as seen in Figures 77 and 78.

In early fuel pin studies it was often observed that irradiation creep strains would increase and then abruptly decrease and sometimes stop entirely, even though fission gas pressures continued to increase.173,174 These results were interpreted as evi­dence of fuel swelling very quickly to meet and thereby put stress on the cladding but later the onset of swelling in the clad caused it to out-swell the fuel and break contact. Actually, the driving force

Figure 74 (left) Diametral strains resulting from void swelling at 400 °C in neutron-irradiated stress-free tubes constructed from nine titanium-modified 316 stainless steels, (right) Stress-normalized midwall creep strains observed in three of these steels, showing a strong correlation of swelling and irradiation creep rates in each steel. Reproduced from Toloczko, M. B.; Garner, F. A.; Eiholzer, C. R. J. Nucl. Mater. 1992, 191-194, 803-807.

Figure 75 Creep modulus measured for six austenitic steels irradiated in BOR-60 fast reactor at 420 °C, showing an enhancement of creep versus Ni-equivalent. Reproduced from Neustroev, V. S.; Shamardin, V. K. J. Nucl. Mater. 2002, 307-311, 343-346. This behavior corresponds to the known effect of nickel on void swelling, indicating swelling-enhanced creep.

was primarily increasing levels of fission gas but irradiation creep had disappeared by ~7% burn-up.

Several features of creep disappearance are noteworthy.

1. The combined creep and swelling strain rate in a fuel pin or pressurized tube cannot exceed 0.33%

per dpa or one-third of the eventual steady-state swelling rate.

2. As swelling approaches 1% per dpa the creep rate backs down proportionately to maintain this max­imum rate as shown in Figures 78-80.

3. The limit of 0.33% per dpa is reached before swelling gets to a significant fraction of 1% per dpa, as shown in Figure 80. Some tubes had already reached the maximum strain rate limit, but then lost their gas pressure and continued to swell at less than 1% per dpa.

4. As the creep cessation process gets underway the creep strain loses its responsiveness to the magnitude of the stress. Note in Figures 79 and 80 that doubling the hoop stress did not double the strain rate in the tube.

5. The coupling coefficient D tends to fall to zero rather quickly when swelling-before-creep occurs but falls more slowly in creep-before-swelling scenarios (fuel pins vs. pressurized tubes).175

A consensus explanation ofthe creep disappearance phenomena has not yet been reached. Various models have been proposed involving voids acting to erase the anisotropy of dislocation Burgers vector176,177 and the involvement of precipitate sinks to serve as strong sinks that compete with dislocations.175

4.05.4.4.1 Introduction

The purpose of this subsection is to establish the level of understanding regarding the irradiation-induced formation of CECs, and how this understanding sup­ports the mechanistic framework outlined above. Insight has been developed by characterizing popula­tions of irradiation-induced CECs in Cu-containing steels and, in particular, the dependence of CEC structure, composition, size, and number density on material and irradiation parameters (e. g., steel com­position, fluence, flux, etc.). Measurements have also been made on the level of Cu in the matrix (Cumatrix) which is not associated with any CEC or precipitate. This is an important measure as, clearly, Cumatrix will decrease during irradiation as CECs are formed, and at SOL it may be lower than at the bulk level if Cu is precipitated during the final heat treatment.

Several techniques have been successfully em­ployed to characterize the CECs formed during neu­tron irradiation of Cu-containing pressure vessel
steels (Table 6). TEM was first used to observe CECs, but the cluster sizes were close to the resolu­tion limit for most TEMs, so atom probe (3DAP) and small angle neutron scattering (SANS) became the most commonly used methods to characterize CECs. However, these are not the only techniques; for example, an important development has been the advent of experiments employing a positron annihi­lation technique, coincidence Doppler broadening (CDB) spectra of positrons annihilating in aged or irradiated Cu-containing alloys or steels. CDB pro­vides a means of identifying the elements around the annihilation site.52 It has become the standard practice to characterize the same as-irradiated specimens with a number of techniques.53,54 Part of the logic behind this is that all the techniques have limitations, either in terms of volume analyzed or complexity of interpret­ing experimental data, and that combining techniques provides better information. These microstructural techniques have been discussed in greater detail in English and Hyde.55,56

SiC composites are a family of materials of varied constituents and architectures. Up to the point of writing this chapter, nuclear-grade SiC composites (those specifically developed for application in fast neutron environments and exhibiting neutron irradi­ation damage resistance) are more precisely defined as continuous fiber-reinforced ceramic composites. The history of development for these materials has been reviewed in a number of publications.29,77-79 The primary constituents of these nuclear-grade com­posites are the continuous SiC fiber, a fiber/matrix interphase material that can be SiC or pyrolytic graph­ite or a combination of the two, and a matrix of SiC infiltrated into the woven fiber preform. The most common matrix material is derived from chemical vapor infiltration (CVI), and is essentially identical in structure, properties, and irradiation response to the CVD SiC discussed in previous sections. While there has been little direct study on the effects of irradiation on the material properties of the SiC interphase, it can be assumed that it would also behave in a similar manner to the SiC matrix. However, the effect of neutron irradiation on pyrolytic graphite interphase (if used) will be substantially different from that on both matrix and fiber. While the effect of irradiation on the underlying properties of graphite interphase has not been well studied, it can be assumed that the interphase will behave in a similar manner to nuclear graphite (discussed in Chapter 4.05, Radiation Dam­age of Reactor Pressure Vessel Steels).

I™"’,

(b) •

Z

Figure 21 Example of braided nuclear-grade SiC/SiC composite. Fiber: Hi-Nicalon™ Type-S; Interphase: Multilayer SiC with pyrolytic carbon; Matrix: CVI SiC deposited through an isothermal process. Reproduced from Nozawa, T.; Lara-Curzio, E.; Katoh, Y.; Shinavski, R. J. Tensile properties of advanced SiC/SiC composites for nuclear control rod applications. Wiley: 2007; pp 223-234.

An example of an SiC/SiC composite that has been developed for high-temperature gas-cooled reactor control rod applications is shown in Figure 21. The basic textile weaving of the composite is evident on inspection of Figure 21(a). In this case, a ±55° weave is depicted. For the polished section of Figure 21(b), large voids, which are an unavoidable characteristic of chemical vapor infiltrated materials and also the primary reason why it is difficult to produce gas — impermeable SiC/SiC composite, are clearly observed. In Figure 21(c), the complicated structure of the interphase is seen. In this case, alternating layers of SiC and pyrolytic graphite have been applied. The pyrolytic graphite layer between the SiC layers is quite thin (tens of nanometer), with a relatively thick graphite layer in contact with the fiber itself.

From the earliest study of SiC/SiC composites under irradiation, it was clear that the fiber was the key to performance. As with the impure forms of SiC monolithic ceramics (cf. Figure 17), the impure and oxygen-rich early grades of SiC fiber (trade name Nicalon™) were quite unstable under neutron irra­diation.12,80,81 Researchers were able to directly link an irradiation-induced shrinkage of the SiC-based
fibers with debonding of the fiber-matrix interface that severely compromised the ability to load the high-strength fibers.80 Composite mechanical prop­erties such as strength suffered appreciably.

With continued evolution of the fiber systems to increasingly pure, stoichiometric materials, the irradiation stability improved significantly. Presently, there are two commercial fiber systems used in nuclear-grade composites, both of which have rela­tively low impurity contents and are approaching a 1:1 stoichiometry. Specifically, the ~11 micron Hi-Nicalon™ Type-S fiber has the nominal chemis­try of SiC105, 0.2%-O, while the ^7.5 and ~10 pm Tyranno™ SA-3 fibers have the nominal chemistry of SiC1.07, 0.5% Al. Study has revealed that these ‘near stoichiometric’ fibers exhibit irradiation — induced swelling similar to that of CVD,82 thus avoiding the debonding phenomenon mentioned in the previous paragraph. For this reason, composites fabricated from these materials are superior under irradiation to their predecessors. Consistent with the discussion of properties of irradiated monolithic SiC, the following discussion will be limited to the more pure, near stoichiometric fiber materials.

The effect of neutron irradiation on the Weibull mean strength of individual ‘near stoichiometric’ fibers is given in Figure 22.83’84 Within inherent sta­tistical scatter, no change in strength is observed for either the Hi-Nicalon™ Type-S or the Tyranno™ SA-3 bare fibers. The numbers inset to the figure indicate the irradiation temperature of the SiC fibers, with no apparent function of irradiation temperature on strength observed. From the same study, the effect of irradiation on composite properties is also observed. Figure 2 3 67 gives the proportional limit stress for which the load departs from elastic behavior and the ultimate tensile strength. As with the fiber data, and the data for monolithic CVD SiC (Figure 18), the composite strength does not exhibit any statistically meaningful change. Supporting studies14,82,83,85-87 on the strength in tension or bending of neutron-irradiated stoichiometric fiber composites support the fact that at least up to ^40 dpa, composite strength is not significantly affected by irradiation. A recent study88 on the fracture toughness of irradiated and unirradiated Hi-Nicalon™ Type-S composites also reports no appreciable change. However, a minor dif­ference in the fracture surface (length of fiber pull out) and a trend in the fiber-matrix interphase properties are reported,89 suggesting that mechanical property evolution may occur at higher doses.

In the unirradiated state, the thermal conductivity of SiC composites is dependent on variables includ­ing the fibers and matrix constituents, processing, and the level of porosity. For the nuclear composite considered here, there is considerable thermal conduc­tivity anisotropy and temperature dependence typical of all ceramics. This is demonstrated in Figure 24, which gives the measured and calculated thermal conductivity for the two nuclear-grade SiC compo­sites.90 Presented are Hi-Nicalon™ Type-S fiber and Tyranno™ SA fiber composites, each matrix infiltrated through CVI.58 Architectures included balanced (1:1:1 for x:y:z) and unbalanced (1:1:4) 3D forms and 2D laminates (SW: satin weave, PW: plain weave.) In each case, a pyrolytic graphite interphase was applied. The conductivity for all materials is presented in the through thickness direc­tion (perpendicular to the plate and the fabric for the 2D composite.) This typically represents the low — conductivity direction.

As evident from Figure 24 and the supporting analysis by Katoh,90 the fiber makes a significant con­tribution to the thermal conductivity of these highly stoichiometric fiber composites, and this conductivity is a fairly strong function of temperature. However, the absolute conductivity is only a fraction of that for the highest thermal conductivity CVD SiC (cf. Figure 7.)

Figure 23 Effect of neutron dose on tensile proportional limit and ultimate tensile stresses for composites. Data labels indicate the nominal irradiation temperature in °C for Hi-Nicalon™ Type-S (upright) and Tyranno™ SA-3 (oblique) composites. Reproduced from Katoh, Y.; Snead, L. L.; Nozawa, T.; Kondo, S.; Busby, J. T. J. Nucl. Mater. 2010, 403, 48-61.

80

70

60

7

E

§, 50

>.

40

ra

TO

о 30

га

E

20

10 0

0 200 400 600 800 1000 1200

Temperature (°C)

Figure 24 Thermal conductivity of representative nuclear-grade SiC/SiC composite in unirradiated condition. Reproduced from Katoh, Y.; Nozawa, T.; Snead, L. L.; Hinoki, T.; Kohyama, A. Fus. Eng. Des. 2006, 81, 937-944.

As with the CVD SiC discussed in section 4.07.3, silicon carbide composite also undergoes significant degradation in thermal conductivity because of neutron irradiation. The data is somewhat limited;

however, Figure 25 gives the ambient through­thickness thermal conductivity for a plain weave Hi-Nicalon™ Type-S, multilayer SiC interphase, and CVI SiC matrix composite. It is noted that, in

Effect of neutron irradiation on the through-thickness thermal conductivity of Hi-Nicalon™ Type-S, CVI matrix

Figure 26 Comparison of the thermal defect resistance for neutron irradiated CVD SiC and Hi-Nicalon™ Type-S, CVI matrix composite.

comparison to the conductivity shown in Figure 24 (second from lowest curve), the ambient through­thickness thermal conductivity for the material of Figure 25 is relatively low (10.2 ± 2.2Wm~1K~1). This is mostly ascribed to the large porosity for that composite. Nevertheless, the figure clearly shows a significant, irradiation temperature-dependent reduc­tion in thermal conductivity as a function ofirradiation dose. The fact that this is temperature dependent suggests that the degradation is due to the produc­tion of stable point defects and clusters, as discussed in Section 4.07.3, although this may not be the sole factor determining the degradation. Figure 26 provides the accumulated thermal defect resistance at the lowest and highest irradiation temperature for the composite materials of Figure 25, compared with high-conductivity CVD SiC. It is interesting to note that the thermal defect resistance for the composite, while accumulating in the same manner as that of the CVD SiC, is about an order of magnitude greater than that of CVD SiC at a given dose (at least prior to saturation.) This greater accumulation of thermal defect resistance has been recently observed by Katoh67 The reason for this is unclear, although it is plausible that, in addition to defect production, propa­gation of internal interfaces (e.g., cracks) in the com­posite is occurring under irradiation. It is also possible that the defects population responsible for phonon scattering for the composite material is stabilized at a higher level than that of the highly pure CVD SiC.90

Fusion welding leads to residual stresses and strains (distortion) via thermally induced stresses, solidifica­tion shrinkage, and phase transformations. Thermal stresses arise from the large temperature gradients inherent to fusion welding and from differences in the coefficients of thermal expansion (CTE) between materials that make up the weldment (Table 1). Thermal contraction generates stress and distortion during on-cooling, with the maximum residual stress often being the flow stress at which the lowest tem­perature distortion occurs.5,45

Dissimilar metal welds are regions of special con­cern for nuclear power systems as the residual stres­ses are often higher than for similar metal welds.18 For example, in ‘safe end’-type welds, the CTE dif­ference between a pressure vessel low-alloy steel and a corrosion-resistant austenitic alloy leads to higher stresses in these welds and in fact, these locations are known to be at increased risk of stress corrosion cracking for this reason.46—48 Solidification shrinkage is a lesser effect, as the liquid cannot support appre­ciable stress but can affect the local bead contour, the deformation at the weld face, and lead to stress raisers (e. g., concave beads or cracks).5

Phase transformations can also have appreciable effects on the sign, magnitude, and distribution of residual stresses in welds. Becker et al. have shown that for accurate prediction of the residual stresses in pressure vessel-type steels, it is critical to account for the on-cooling phase transformations that occur from welding.49 Furthermore, phase transformations dur­ing postweld heat treatments and from service expo­sure must also be considered. For example, nickel alloys can be susceptible to the development of

=1500 mm; BC +	5tep’= 5Wi;Grid8’00 x 100
=1500 um; E123; Step = 5 um; Grid800 x 100
= 1500 mm. BC + GB; Step = 5 mm. Grid800 — 100
=1500 um; E123; Step = 5 um; Grid800 x 100
GB; Step = 5 um; Grid800 x 100
0.5 cm
Figure 13 Example of a hydrogen crack produced in EN82H from the use of 95%Ar-5%H2 shielding gas and abusive welding practice (refuse welding). The top figures show the crack in cross-section and the corresponding electron backscatter diffraction strain maps. The bottom fractographs show the intergranular/interdendritic nature of the hydrogen crack.

50 цт

50 цт

Figure 14 Illustration of some welding defects in commercially available Ni-30 wt% Cr alloys. (a) Lack of fusion defects in an SMAW (left) and a GTAW (right), (b) unmelted NbFe2 alloying addition in an SMAW (left) and a slag inclusion from an SMAW (right) and (c) surface (left) and internal (right) oxides from a laser weld.

short — and long-range order. The ordered structure typically has a smaller lattice parameter than the bulk alloy and can lead to increased residual stresses with service exposure.10’50-54 Another point to note is that the welds typically have considerable texture that can be a significant factor in both the macroscopic and microscopic (intergrain) stresses and strains in welds.

The exact mechanism of radiolytic oxidation in a car­bon dioxide-cooled reactor is complex and has been a matter of debate for some time; the most satisfactory explanation has been given by Best et a/.26 However, in its most simplistic form the mechanism can be described as follows:

In the gas phase,

ionizing radiation

CO2 —————— ^ CO + O* [I]

CO + O* ! CO2 [II]

where O* is an activated state-oxidizing species.

Thus, after ionization the carbon monoxide and oxidizing species rapidly recombine back into carbon dioxide and to an uninformed observer, carbon

dioxide would appear to be stable in an irradiation field. However, in the presence of graphite which typically contains ~10% porosity, 10% of which is initially accessible to the carbon dioxide gas, at the graphite pore surface (mainly internal) carbon atoms are oxidized. This can be simplistically described as

O* + C! CO [III]

The principal oxidizing species is still under debate, but the most favored candidate is the negatively charged ion, CO^.

4.11.7.3 Inhibition

The rate of oxidation can be reduced by the addition of carbon monoxide (CO) and moisture (H2O) and can be greatly reduced by the addition of methane (CH4), as illustrated in Figures 9 and 10. As described above the radiolytic oxidation process pro­duces CO and if CH4 is added, moisture will be one of the by-products of the reaction.

Various tritium breeding fusion blanket concepts have been studied with different combinations of structural materials, tritium breeding materials, and cooling materials. Vanadium alloys have been used in most cases with liquid lithium as the breeding and cooling materials (self-cooled V/Li blankets) for advanced concepts of DEMO (fusion demonstra­tion power plant) and commercial fusion reactors.7,8 Because of high atomic density of Li atoms in liquid Li relative to Li-ceramics, Li-Pb, and molten-salt

Flibe, V/Li systems can obtain high tritium breeding ratio (TBR) without using the neutron multiplier Be. A neutronics calculation showed that ‘tritium self sufficiency’ can be satisfied without Be both in Tokamak and Helical reactor systems.9 Without the necessity of using beryllium as a neutron multiplier, the replacement frequency of the blanket will be reduced because the blanket system is free from the periodic replacement due to the lifetime of Be, which can lead to enhanced plant efficiency.

V/Li blankets can be designed with a simple structure as schematically shown in Figure 1. The blanket is composed of Li cooling channels made of vanadium alloys, reflectors, and a shielding area, which is in contrast to more complex solid breeder blankets that need a solid breeder zone, a neutron multiplier beryllium zone, cooling channels using gas or water, and tritium recovery gas flow channels in addition to reflectors and shielding.

A self-cooled Li blanket using neutron multiplier beryllium was also designed in the Russian pro­gram.10 This concept can downsize the blanket area because of efficient tritium generation per zone. However, the blanket structure must be more

Flowing liquid lithium

Blanket

Figure 1 Illustration of self-cooled Li blanket with V-4Cr-4Ti structural material.

complex than V/Li and new issues need to be solved such as Li/Be compatibility.

General requirements for structural materials of fusion blankets include dimensional stability, compat­ibility with breeder and coolants, high-temperature strength and low-temperature ductility during irradia­tion. For vanadium alloys, issues concerning industrial maturity such as developing large-scale manufacturing technology need to be resolved.

Vanadium alloys could be a candidate structural material for molten-salt Flibe (LiF-BeF2) blankets. For this application, a concept was proposed to dis­solve WF6 or MoF6 into Flibe for corrosion protec­tion of the wall surfaces by precipitation of W or Mo and reduction of the tritium inventory in vanadium alloys by enhancing reaction from T2 to TF, which is more highly soluble in Flibe than T2.11 The TBR of Flibe/V blankets may be marginal, but the neutron shielding capability for the superconductor magnet systems may be superior relative to V/Li according to neutronics investigation.12 In this system, precipitates of W or Mo formed as a result of reaction from T2 to TF needs to be recovered from the flowing Flibe.

Table 1 summarizes the blanket concepts using vanadium alloys with the advantages and critical issues.