Thermal creep of copper and copper alloys can be significant at relatively low temperatures, because of copper’s low melting point (0.3 Tm = ^134 °C, Tm is the melting point). Nadkarni51 and Zinkle and Fabritsiev2 compared the 100-h creep rupture strength of copper and several PH and DS copper alloys at elevated temperatures. Copper alloys have significantly higher creep rupture strength than pure copper. Creep rupture strength decreases drasti­cally as temperature increases in PH alloys such as CuCrZr, as well as in pure copper, between 200 and 450 °C. DS alloys such as CuAl25 have superior creep rupture strength even above 400 °C because of their thermal stability at high temperatures.

Li et al?1 summarized steady-state thermal creep data for pure copper and several copper alloys, as shown in Figure 5. Pure copper can suffer significant creep deformation at high temperature even with a very low applied stress. The creep rate of pure cop­per can be as high as ^10-4s-1 at ^100 MPa at 400 °C. The creep resistance of copper alloys is con­siderably higher than that of pure copper. The creep
rates of copper alloys strongly depend on the applied stress and the temperature, and can be described by the Norton power law relation; that is, e = Aan exp(—Q/RT) where e is creep rate, a is the applied stress, n is the stress exponent, Q is the activation energy, R is the gas constant, and T is the temperature. DS copper alloys exhibit unusu­ally high values of the stress exponent, for example, 10-21 in the temperature range of 472-721 °C for GlidCop Al15.52

Because of the time-dependent nature of creep deformation, softening behavior due to overaging and recrystallization must be considered during the creep analysis for PH copper alloys. The creep prop­erties of this class of alloys could be significantly changed during prolonged exposure at elevated temperature.

In the Master Curve brittle fracture model, it is assumed that the steel is macroscopically homoge­neous. In real wrought steels, the macro — and micro­structures are seldom fully homogeneous due to various effects occurring during production and cool­ing, such as segregation of elements and development ofinclusions. The resultant inhomogeneity may man­ifest itself as excessive data scatter. However, in welds or dissimilar metal joints, the material in the asso­ciated heat-affected zone (HAZ) may exhibit strongly distributed ‘inhomogeneity’ so that two or more sub­populations can be distinguished in the test data. Often in the final application of the fracture toughness data, it is only essential to determine a statistically sound conservative estimate for the lower bound frac­ture toughness, which can then be used to construct a conservative reference curve for the material.

In most cases, even those where materials show slight inhomogeneity, the standard Master Curve

estimation is sufficient, without further assessments. For testing the data population for possible inhomoge­neity, several approaches have been developed. A com­mon feature for all of these approaches is that they are based on the standard Master Curve, but are extended to properly take into account the inhomoge­neity. The following extended approaches have been proposed for analyzing inhomogeneous materials:

1. Simplified SINTAP procedure for determining a conservative lower bound estimate of fracture toughness when the inhomogeneity is of randomly distributed type.16,20

2. Bimodal or multimodal models for analyzing data­sets showing distributed inhomogeneity (typically welds and HAZ materials).17,20

3. Model for analyzing randomly inhomogeneous materials, including, for example, macroscopic

17,20

segregations.

Of the three methods, SINTAP is the simplest and the one recommended for initial assessment of the quality of data; it can generally be applied in conservative structural integrity analysis when material inhomoge­neity is suspected. An advantage of the SINTAP procedure is that it can be performed for a small dataset, unlike the bimodal or multimodal analyses, which require 15 or more data points for a valid assessment. It should be noted that the SINTAP analysis does not produce a statistically representa­tive description ofthe whole dataset, but its primary purpose is to determine whether the dataset is homogeneous and secondly, to develop a generally conservative lower bound when possible material inhomogeneity is present. The model for analyzing randomly distributed inhomogeneities gives a statis­tically correct description of the fracture toughness data, but requires a large dataset and is more com­plicated to perform than the SINTAP analysis. When there are a sufficient number of fracture toughness data points, the bimodal/multimodal and the randomly distributed inhomogeneous models should be used to better identify a more statistically correct description for a modified Master Curve.

The simplified SINTAP procedure is described next. The SINTAP analysis is based on the maximum likelihood method like the basic Master Curve pro­cedure and it consists of three steps as follows.

Step 1 is the standard estimation giving the first estimate of T0 and the median fracture toughness according to ASTM E 1921. The resulting T0 is used as an input value for Step 2.

Step 2 is a lower tail maximum likelihood estima­tion and is performed so that all data exceeding the median fracture toughness are censored by substituting d = 0 and by reducing the corresponding KJc values to the median curve. If the resulting new T0 (T0-2) is lower than the T0 from Step 1 (T0-1), then T)-sintap = T)_i; oth­erwise, Step 2 shall be repeated using the last T0 estimate (T0-2) as a new input value until a con­stant T0 (T0-2adj) is obtained. If the number of valid data is ten or more, T0-SINTAP = T0-2adj, otherwise Step 3 shall be performed.

Step 3 gives the minimum value estimate of T0 and is performed only if the number of valid test data is less than 10. In Step 3, an additional safety factor is incorporated for cases where the number of tests is small. Here, the value of T0 is estimated for each single data point to find the maximum value of T0 (T0 max). If the result­ing T0 max> T0-2 adj + 8 ° Q T0-SINTAP = T0 max, otherwise T0-SINTAP = T0—2 adj.

The steps and the data censoring in Steps 1 and 2 are described schematically in Figure 18. The flow chart of the iterative procedure is shown in Figure 19.

4.15.6.1 Irradiation Damage

Greenwood calculated damage parameters for solid lithium tritium breeder materials irradiated in the fusion breeder reactor (FUBR) and breeder exchange matrix (BEATRIX) series of irradiations in EBR-II and FFTF reactors.186 He showed that most of the displacement damage arises from fast neutron reac­tions such as elastic and inelastic scattering rather than from the 6Li (n, a) T reaction. The displacements per atom is only an estimate of the total energy available for creating displacements in a compound and does not directly relate to observable damage effects due to the high probability of recombination effects.

Leichtle and coworkers149,150 evaluated damage correlation parameters in fusion and fission reactor systems for the ceramic breeder materials Li2TiO3, Li4SiO4, and Li2O. A sophisticated damage calculation taking into account all relevant primary knock-on atoms (PKAs) was performed to obtain displacement damage parameters for these low-mass polyatomic materials. They arrived at a concise fusion-fission correlation of irradiation effects. Such compar­isons have been performed for typical fission neutron flux spectra of a thermal and a fast reactor system. In

preparation for an IEA-framed international program on high fluence breeder irradiation, Fischer et a/.150 performed analyses of lithium burnup and displace­ment damage in Li2O, Li4SiO4, and Li2TO3 for a HCPB-type DEMO blanket design (Figure 63).

There are little data available for irradiated tungsten (see Chapter 1.04, Effect of Radiation on Strength and Ductility of Metals and Alloys). Based on results for other refractory alloys and limited data on tungsten, one would expect neutron irradiation to increase the strength and decrease the ductility of the tungsten armor largely through increases in the DBTT. To minimize the neutron-induced material

degradation, it is reasonable to limit the operational conditions for components in a neutron environment to temperatures above ^900 °C where recovery of tungsten takes place,107 as the ductility loss is more pronounced below about 0.3 Tm. This is possible in the region close to the plasma-facing surface, but it is impossible in the heat sink region as tungsten will be in contact with materials that cannot operate at this temperature and stress concentrations in these ‘cold’ areas have to be avoided.108 In the case of ITER, Cu will be employed in the heat sink while steel is more likely to be used in DEMO, which has a higher operating temperature.2 Hence, a greater under­standing of the irradiation response of tungsten at temperatures between 700 and 1000 °C is

needed.109,110 The effect of embrittlement is alle­viated when operating above 250 °C, although in the presence of He (produced by transmutation reactions) somewhat higher temperatures may be required.83

Although at intermediate temperatures (0.3-0.6 Tn), void swelling and irradiation creep are the dominant effects of irradiation, the amount of volumetric swelling associated with void formation in refractory alloys is generally within engineering design limits (<5%) even for high neutron fluences (^10dpa). Very little experimental data exist on irradiation creep of refractory alloys, but data for other bcc alloys suggest that the irradiation creep will produce negligi­ble deformation for near-term reactor applications.110

A range of particle types, fluxes, and energies strike the PFMs and interact in the near-surface region of the PFM. The most common interactions are with hydrogen fuel ions, ranging in energy from a

.N

0.2 —

о

z

few eV to hundreds of eV. In addition to hydrogen ions, fuel by-product ions such as helium, and impu­rities from the first wall also impact the surface. Severe surface layer damage occurs because of such ion impacts, and significant erosion of surface material additionally occurs. Various mechanisms are responsi­ble for erosion, dependent on the surface temperature of the graphite. The mechanisms can generally be characterized in order of increasing temperature phe­nomenon as physical sputtering, chemical erosion, and radiation enhanced sublimation (Figure 21).43 Above 2000 °C, the vapor pressure of graphite dominates the erosion. In addition to the obvious issue of erosion-degraded component lifetime, removal and redistribution of carbon contributes to a high tritium inventory trapped in redeposited carbon.44

4.19.6.1 PFC Design Considerations

Robust engineering solutions are needed for the com­ponents that directly surround the plasma in order to withstand the thermal and mechanical loads during normal and off-normal operation and contribute to protect the outer components from the effect of neu­trons, especially the vacuum vessel and the magnets. Considerations here are limited to the portion of the PFC surface that protects the main chamber (the so-called first wall), which for both JET and ITER is made of beryllium.

The strategy and the criteria adopted to design the PFCs of JET and ITER are substantially different. One of the main differences is the longer plasma pulse duration foreseen in ITER that has required the use of water-cooled structures to handle the power in steady-state conditions. This translates into design solutions that consist of tiles bonded to an actively cooled copper-alloy substrate that must withstand the thermomechanical loads with a good engineering margin and achieve the desired fatigue lifetime. An additional problem in ITER is the deg­radation, albeit limited, of thermal and mechanical properties and the properties of the joints due to neutron irradiation (see Section 4.19.4.4).

In contrast to ITER, technical constraints pre­vented the use of actively cooled first-wall struc­tures in JET and the first-wall protection relies on a series of discrete poloidal limiters. The tiles are inertially cooled, and their power handling perfor­mance is driven by the need to (1) avoid surface melting and (2) reduce thermally induced stresses to give an adequate fatigue lifetime. At 40 mm typi­cal thickness, the tiles are thermally thick for a typical 10 s pulse and will handle up to about 60 MJ m~2 without melting.

Another important driver for designing PFCs is the magnitude of the electromagnetic loads associated with plasma disruptions. In tokamaks, dis­ruptions produce large changes in magnetic field, dB/dt, which induce eddy currents in the conducting materials. The currents interact with the local mag­netic field, B, to produce a torque, which is strongly dependent on the geometry. The design of the PFCs has to manage this torques via a combination of the castellations of the tiles along with cuts, which will interrupt the eddy current loops. The tile assembly must also withstand electromagnetic forces due to halo currents which, during disruptions, pass between plasma and vacuum vessel via the tiles. Enormous attention has been paid at JET and ITER during the design phase to address this problem.

The problems associated with the design of the first wall at JET and ITER are briefly described in Section

4.19.6.1.1 and 4.19.6.1.2, respectively.

AlN is an alternative candidate for the in situ coating because of its high electrical resistivity and high compatibility with Li. In this case, Al and N are doped into Li to enhance AlN formation at the sur­face of vanadium alloy substrates. The coating has shown sufficient resistivity and stability under ther­mal cycling.

In this process, however, Li needs to be saturated with N to prevent the dissolution of the AlN layer. With a high N level in Li, vanadium alloy not covered with the AlN layer will get N from Li and become brittle. This issue was observed in high-temperature
capsule tests with Mo and vanadium alloys.15’16 Thus, unless the vanadium alloy channel walls are fully covered with AlN, the use of this coating tech­nique seems problematic. For the same reason, the use of AlN layers, sheets, or plates in Li as the insulator requires full coverage of the vanadium alloy structures with AlN.

Early developments of ceramic breeder blanket con­cepts were based on fission reactor technologies, and the R&D activities concentrated on pellet or pin type specimens. An example is the European BIT project, using machined annular pellets or pins.2-5

Figure 3 shows a schematic view of the BIT concept, based on stacking annular pellets of

LiAlO2, enriched up to 90% 6Li, operating at tem­peratures between 400 and 600 °C.5 Li2ZrO3 was considered as an alternative material, which has been further developed at the Commissariat a l’Energie Atomique (CEA)29

Other breeding concepts using machined shapes have been developed, for example, in the Russian Federation,30 and recently Sharafat et al31 proposed shaping the breeder ceramic as foams.

In this chapter, only a few R&D results with pel­lets or pins are mentioned.

4.16.3.1 Plasma-Facing Materials

Tritium generated in the fusion-reactor blanket will be fed directly into the plasma in the main vacuum cham­ber. There, the tritium will be partially consumed, but it will also interact with the materials composing the first wall. Materials used to line the first wall will be exposed to energetic tritium and deuterium escaping from the plasma. Particle fluxes in the range of 1021 (D + T) m-2 s-1 will continuously bombard the plasma-facing materials. Materials used for the diver­tor at the top and/or bottom of the torus will be exposed to lower energy particles with a flux of 1023 (D + T) m-2 s-1 or higher. While the neutral gas pressure of tritium will be relatively low at the outer vacuum wall boundary, some minor permeation losses will occur. In reality, the primary concerns in the plasma-facing areas are tritium inventory and perme­ation into the coolant through coolant tubes. While future power reactors are likely to have primarily refractory metals such as tungsten, present-day devices are still using carbon and beryllium. In this section on plasma-facing materials, we examine the interaction of tritium with carbon, tungsten, and beryllium.

4.16.3.1.1 Carbon

In many ways, carbon is ideal for fusion applications. It is a low-Z material with a low vapor pressure and excellent thermal properties. The carbon used in fusion applications comes in two forms, graphite and carbon
composites. Graphite is described in Chapter 2.10, Graphite: Properties and Characteristics; Chapter 4.10, Radiation Effects in Graphite, and Chapter 4.18, Carbon as a Fusion Plasma-Facing Material

and is typically made using the Acheson process.33 Calcined coke is crushed, milled, and then sized. The properties of the graphite are determined by the size and shape of these particles. Coal tar is added to the particles and the batch is heated to 1200 K. This process is repeated several times to increase the density of the compact. The final bake is at temperatures between 2900 and 3300 K and takes ~15 days. The final product is quite porous, with a density of around 1.8—1.9 g cm-3 (compared to a theoretical density of 2.3 g cm-3). Graphite is composed of grains (from the original coke particles) with a size of 5-50 pm, which are in turn composed of graphite subgrains with a typical size of 5 nm. Carbon composites are made by pyrolyzing a composite of carbon fibers in an organic matrix. These fibers have a high strength-to-weight ratio and are composed of almost pure carbon. As with graphite, carbon composites are quite porous with a density of <2 g cm — . Carbon (as either graphite or carbon composite) will not serve as a vacuum or coolant boundary. Therefore, permeation through the carbon will not directly affect tritium release into the environment. Coolant tubes inside the graphite will control the tritium release to the coolant system. The ways in which hydrogen isotopes can interact with carbon are explored in the following sections.

It is difficult to consider hydrogen isotope perme­ation in graphite in the same manner as one would for the metals. As was shown by Kiyoshi eta/.,34 hydrogen passes rather readily through graphite in the molec­ular form. It is when one considers the uptake or retention of atomic hydrogen in graphite or a carbon composite that it becomes more interesting. Most graphites have a Brunauer-Emmett-Teller (BET) or specific surface area of 0.25-1.0 m2 g-1.35 This presents a lot of surface area for the absorption of hydrogen isotopes. Barrer,36 Thomas,37 and Bansal38 are only a few of the many who have reported hydro­gen retention on carbon surfaces. Activation energies varying from 24 to 210 kJ mol-1 have been reported. Once absorbed on a carbon surface, the hydrogen can migrate by jumping from one chemically active site to another. In a fusion reactor, this process occurs in carbon at lower temperatures at which the plasma provides ionic and atomic hydrogen for direct absorp­tion. Molecular hydrogen can dissociate on carbon surfaces, but only at elevated temperatures (>1000 K). As a plasma-facing material, graphite will be exposed to atomic tritium and deuterium, and these hydrogen isotopes will migrate inward along the open porosity. Several research groups have measured the diffusion coefficient for hydrogen on carbon surfaces. Robell eta/.39 inferred an activation energy of 164kJ mol-1 for the diffusion during measurements on the uptake of hydrogen on platinized carbon between 573 and 665 K. Olander and Balooch40 used similar experi­ments to determine the diffusion coefficient for hydro­gen on both the basal and prism plane: 6 X 10-9 exp (-7790/7) m2 s-1 for the basal plane and 6 X 10-11 exp(-4420/7) m2s-1 for the prism plane. Causey eta/.41 used tritium profiles in POCO AXF-5Qgraphite exposed to a tritium plasma to extract a diffusivity for tritium on carbon pores of 1.2 X 10-4 exp(-11 670/7) m2 s-1. An example of the deep penetration of hydro­gen isotopes into the porosity of graphite was reported by Penzhorn et a/.42 Graphite and carbon composite tiles removed from the Joint European Torus (JET) fusion reactor were mechanically sectioned. The sec­tions were then oxidized, and tritiated water was collected for liquid scintillation counting. Relatively high concentrations of tritium were detected tens of millimeters deep into the tiles.

The diffusion or migration of hydrogen isotopes on carbon surfaces occurs at lower temperatures. The solubility, diffusion, and trapping of hydrogen in the carbon grains are higher temperature processes than adsorption and surface diffusion. At higher tempera­tures (>1000 K), hydrogen molecules can dissociate and be absorbed at chemically active sites on carbon surfaces. Some of these sites are located on the out­side of the grains, but many exist along the edges of the subgrains that make up the larger grains. Hydrogen isotopes dissociating on the outer grain boundary are able to migrate along the subgrain boundaries, entering into the interior of the grain. It is the jumping from one moderate energy site (^240 kJ mol-1) to another that determines the effec­tive diffusion coefficient. Traps on the grain bound­aries pose a binding energy barrier (~175 kJ mol-1) that must be overcome in addition to this normal lattice activation. Atsumi et a/.43 used the pressure change in a constant volume to determine the solu­bility of deuterium in ISO 88 graphite. They deter­mined the solubility to be given by K = 18.9 exp (+2320/7) mol H2 m-3 MPa-1/2 over the tempera­ture range of 1123-1323 K. This solubility is shown in Figure 3 along with two data points by Causey44 at 1273 and 1473 K. A negative heat of solution is seen in both sets of data, suggesting the formation of a bond between hydrogen and carbon.

10

0.65 0.7 0.75 0.8 0.85 0.9

Temperature, 1000/T (K-1)

10-11

10-13

‘co

10-15

>>

10-17 h

10-19 10-21

0.4 0.6 0.8 1 1.2 1.4

Temperature, 1000/T (K-1)

Figure 4 Diffusivity of hydrogen in graphite. The bold line is the best estimate given in Causey44 based on uptake data for tritium in POCO AFX-5Q graphite. Adapted from Atsumi, H.; Tokura, S.; Miyake, M. J. Nucl. Mater. 1988, 155, 241-245; Causey, R. A. J. Nucl. Mater. 1989, 162, 151-161; Rohrig, H. D.; Fischer, P. G.; Hecker, R. J. Am. Ceram. Soc. 1976, 59, 316-320; Causey, R. A.; Elleman, T. S.; Verghese, K. Carbon 1979, 17, 323-328; Malka, V.; Rohrig, H. D.; Hecker, R.

Int. J. Appl. Radiat. Isot. 1980, 31, 469.

The variation in the diffusivities of hydrogen in graphite determined by various researchers is extreme. This variation results primarily from differ­ences in interpretation of the mechanism of diffusion (e. g., bulk diffusion or grain boundary diffusion). Representative values for the diffusion is shown in Figure 4. Rohrig et a/.45 determined their diffusion coefficient using the release rate of tritium from nuclear grade graphite during isothermal anneals. They correctly used the grain size as the real diffusion
distance. Causey et a/46 measured the release rate of tritium recoil injected into pyrolytic carbon to deter­mine the diffusivity. Malka eta/47 used the release rate of lithium-bred tritium in a nuclear graphite to deter­mine the diffusion coefficient. Atsumi et a/43 used the desorption rate of deuterium gas from graphite samples that had been exposed to gas at elevated tem­peratures to determine a diffusivity. Building on the work of others, Causey44 proposed an alternative expres­sion for the diffusivity that he labeled ‘best estimate.’

The result was based on uptake experiments for tri­tium into POCO AFX-5Qgraphite, and the assump­tion that the total uptake is determined by the product of the diffusivity and the solubility. The uptake data were analyzed assuming that the expression given by Atsumi et a/.43 for the solubility was correct. The expression ‘best estimate’ was used because it prop­erly took into consideration that the grain size was the effective diffusion distance, and that diffusivity was more properly determined by uptake than release. Release rates are strongly affected by trapping.

The trapping of hydrogen isotopes at natural and radiation-induced traps has been examined by sev­eral research groups.41,48-50 Causey et a/.41 exposed POCO-AFX-5Q graphite to a deuterium/tritium mixture at elevated temperatures in an examination of the kinetics of hydrogen uptake. For temperatures above 1500 K, it was discovered that increasing the pressure did not increase the retention. It appeared that the solubility hit an upper limit at 17 appm. Anal­ysis of the data revealed that the 17 appm did not represent solubility, but a trap density. The trap was determined to have a binding energy of 175 kJ mol — . Atsumi et a/.48 found that radiation damage increased the apparent solubility of deuterium in graphite by a factor of 20-50 with saturation at a radiation damage level of 0.3 dpa. A significant decrease in the apparent diffusivity was also noted. In a later study, Atsumi et a/.49 reported graphites and composites to vary sig­nificantly in their natural retention values. They saw that the saturation retention was inversely proportional to the lattice constant (which relates to the degree of graphitization, and so grain size). Radiation damage was seen to decrease the apparent lattice constant and increase the saturation retention. 6 MeV C+ ions were used by Wampler eta/.50 to simulate neutron damage to different graphites. The trap density increased with damage levels up to 0.04 dpa at which the saturation retention was 650 appm. In a defining set of experi­ments, Causey eta/.28 examined tritium uptake in unir­radiated and radiated pitch-based carbon composites. Pitch-based carbon composites have a large lattice parameter due to the sheet-like configuration of the fibers. These composites retained significantly less tri­tium before and after irradiation than other carbon materials. From these results, it is apparent that high — energy trapping occurs at the edges of the hexagonal crystals on the prism plane.

Hydrogen isotope permeation in the normal sense does not apply to graphite and carbon composites. There is inward migration of atomic hydrogen isotopes along porosity at lower temperatures. To calculate the potential tritium inventory for this process, one can obtain an upper bound by assuming monolayer coverage of the pore surfaces. Typical nuclear grade graphite has a specific surface area of 1 m2 g-1. Complete loading of that amount of surface area yields 2 x 1025 Tm-3, or about 2 g for a 20 m2 carbon wall that is 10-mm thick. At higher temperatures, molecular hydrogen isotopes, which are moving through the graphite pore system, are able to disso­ciate and enter the multimicron-sized graphite grains. This migration into the grains occurs along the edges of the nanometer scale subgrains. As the hydrogen migrates inward, it decorates high-energy trap sites. The density of these trap sites is higher than the effective solubility derived from the migra­tion rate. Permeation into graphite will not lead to tritium release from a fusion device, but will affect tritium inventory. If one assumes that radiation dam­age from neutrons has increased the concentration of traps with a binding energy of 175 kJ mol-1 up to 1000 appm, and that tritium is occupying 100% of those traps, the same 20 m2 wall 10-mm thick listed above would now contain an additional 100 g of tritium. Occupation of all of the traps is difficult to achieve: at low temperatures, kinetics makes it impos­sible to achieve saturation, while at substantially higher temperatures, the traps do not remain filled.

There is another process called carbon codeposi­tion that can strongly affect tritium inventory in a fusion device. In the codeposition process, carbon eroded from the walls of the tokamak is redeposited in cooler areas along with deuterium and tritium from the plasma. Because carbon codeposition is not a diffusion or permeation process, it will not be covered in this review. The interested reader is referred to a review of this process by Jacob.51

The irradiation of tungsten and tungsten alloys with energetic neutrons (14MeV) resulting from the D-T reaction causes radiological hazards that were already discussed in Section 4.17.3.2.6. In addition, the neu­tron irradiation affects the material composition by transmutation of tungsten to Re and subsequently osmium (transmutation of W isotopes to Ta and Hf are negligible222). The amount of transmutation strongly depends on the applied neutron wall load and neutron spectrum223 and for the W to Re transmu­tation reaction reaches values between 0.3 and 5 at.% per dpa.222 The subsequent transmutation of Re to Os is expected to occur faster than the production of Re from W resulting in a steadily proceeding burnup of Re. The neutron fluence on the first wall varies strongly with location. For the full lifetime of ITER a maximum of ~0.3MWam-2 is achieved22 («1.35 dpa in tungsten225). As the divertor PFCs will be exchanged 3 times and only the last three will operate in a D-T environment, a neutron fluence of ~0.1 MW m~2 is expected during the lifetime of each PFC. For DEMO, an average neutron wall load of 2 MW m~2 is assumed for the main chamber, which would result in ^45 dpa after 5 full power operation years. These conditions yield a transmutation of 100% W into 75% W, 12% Re, and 13% Os.36 For geometrical reasons, that is, larger surface to angular extension ratio, it will be roughly a factor 2 less in the divertor region.

Furthermore, neutron irradiation damages the material properties by the formation of vacancies and interstitials (see Chapter 1.03, Radiation-Induced Effects on Microstructure). Their behavior including analysis of displacement cross-sections,226,227 diffu­sion, mutual recombination, and clustering are being assessed by atomistic modeling.228-231

Both transmutation and defect generation influ­ences the material properties and subsequently the material response to steady state and transient thermal loads.