The simplest way to determine the mechanical strength of pebbles is by performing crush tests with individual pebbles. These tests are very useful for optimization of pebble properties and for quality control during pebble production.

In crush tests, pebbles are arranged between two plane plates. In a pebble bed, the pebbles in contact with the cavity wall experience similar conditions because there is only one contact at the corresponding pebble hemisphere. Pebbles that are in contact only with other pebbles have on an average six contacts86,87 and the contact forces are in general smaller than on the pebble-wall contact.

The crush loads of 0.5-mm diameter Li4SiO4 and 1-mm diameter Li2TiO3 pebbles used for the European HCPB project are in the range of 4-5 and 35-50 N, respectively, as shown by Knitter et at}9 and Roux et at}92 (see Figure 16).90 The crush loads of the Li2TiO3 pebbles were independent of 6Li enrichment.

Assuming coverage of the wall by pebbles of about 0.7, the lower values of 4 and 21 N for the 0.5 and 1.0 mm diameter pebbles correspond to pres­sures on the wall pebbles of about 12 and 15 MPa. These values are significantly above the assessed maximum pressures of about 6 MPa in the HCPB ceramic breeder pebble bed.91

However, there are several effects that decrease this margin:

• In pebble beds, more generally, in granular mate­rials, an inhomogeneous branching of forces occurs (see Jaeger et at.92), with the effect that some
pebbles experience much larger forces than the average value.

• Thermal annealing results in a significant grain size growth in Li2TiO3 pebbles; a slight decrease in crush load during aging was found for the Li4SiO4 and Li2TiO3 pebbles.77

• Neutron irradiation will have the strongest effect to reduce the pebble strength, as mentioned here.

Fortunately, there is one mechanism, as outlined below, that is expected to alleviate the problem con­siderably: thermal creep.

A general remark to be made here is that care is to be taken with respect to statistics: postirradiation tests from the EXOTIC-7 high burnup experiment showed fragmented pebbles as well as intact pebbles maintaining average strength levels.33, This statis­tical aspect is to be addressed more rigorously in optimization of pebble-bed technologies serving pre­dictable and reliable operation of breeding blankets.

The effectiveness of permeation barriers is often assessed by the permeation reduction factor (PRF), which is the ratio of the effective permeability of specimen without a barrier to an equivalent specimen with a barrier; thus, the greater the PRF value, the more effective the permeation barrier.

4.16.3.3.1 Oxides

Most metals form a native oxide layer in the presence of oxygen. Generally, oxides have very low permeabil­ities for hydrogen isotopes, and native oxide layers may reduce permeability by about an order of magni­tude.175 The coefficient of thermal expansion is often very different from that of the underlying metal. This can cause cracks and spalling, which can reduce the effectiveness of the barrier. In environments that are aqueous or are subjected to elevated temperatures in the presence of oxygen, the oxide layer may be replen­ished and may even grow and coarsen. In addition to native oxides, an oxide layer might be deposited onto the base metal or the base metal might be dipped or
otherwise coated with a second metal, which forms a low-permeability oxide. Such coatings may reduce hydrogen permeability by five orders of magnitude or more.176 Chromia, alumina, and rare-earth oxides have been studied extensively.

The low dissociation pressure of Cr2O3 makes it a common native oxide on steels when allowed to form at elevated temperatures and relatively low oxygen partial pressures.177 Chromia is a better barrier (offering a permeation reduction of about an order of magnitude)175,178 than various Cr2MO4 spinels that may also form (M = Ni, Fe, Co).177 Chromia is also present in mixed oxides in chemical densified coatings, which help to give reduction factors of four orders of magnitude.179

Aluminum forms a self-passivating native oxide that has been shown to be resistant to hydrogen iso­tope permeability, because of a very low solubility for hydrogen. Because this layer is very thin (~4 nm) and the hydrogen permeability of the base metal is very low, it remains debatable whether this amorphous native oxide, which can be grown by anodization, has a lower permeability than aluminum or not.180,181 Cleaned stainless steel may be hot-dipped in alumi­num (forming both a relatively pure surface layer and mixed aluminides between the surface and substrate) and then oxidized. This hot-dip aluminizing proces­sing is simple and generally forms coatings that have excellent adhesion properties (although substantially different thermal expansion coefficients),182 which reduce permeation rates by at least one order of
magnitude, and sometimes more than five orders of magnitude.175,183

The basic properties of hydrogen transport in alu­mina have been characterized and are presented in Figures 18 and 19. Roy and Coble185 hot-isostatically pressed high-purity (>99.99%) alumina powders and charged the dense alumina with hydrogen at elevated temperatures to determine solubility. Fowler et a/.184 obtained diffusion coefficients for single-crystal, poly­crystalline, and powdered alumina, and for alumina that was doped with MgO. They observed faster diffu­sion in powdered specimens, suggesting that the grain boundaries may provide short-circuit diffusion paths. They also noted that the diffusivity of MgO-doped alumina was four to five orders of magnitude greater than that of pure alumina. This suggests that the purity of barrier coatings matters a great deal and transmuta­tion of barriers in a fusion environment may increase the permeability from the ideal case measured in the laboratory.

Yttria and erbia have been deposited on specimens through a number of physical deposition techniques, including plasma spray, arc deposition, and sol-gel deposition.186-188 The advantage of these oxides is not the magnitude of permeation reduction (one to three orders of magnitude), but their high thermal and mechanical stability in a reducing atmosphere.

Independent of the ion energy, blistering (see Figures 11 and 12) due to H occurs only during irradiation at temperatures below 900-950 K and as a function of the ion fluence199,260 at 500 K.261-263 This temperature dependence of blistering is attrib­uted to the formation, movement, and agglomeration
of vacancies containing trapped hydrogen,264 which is dominant at temperatures <500 K, while the detrap­ping of hydrogen from the defects is prevalent at temperatures >500 K.263

The fluence threshold for blister formation increases with decreasing ion energy and signifi­cantly increases to values >1025 D+ m~2 at ion ener­gies <20 eV. This is assumed to be the result of thin (a few monolayers) which act as oxide diffusion bar­riers at the material’s surface.265 Furthermore, with increasing fluence, the size and number of blisters can increase up to a few 100 pm260 until saturation is reached, which is assumed to be related to the rupture of blisters.265-267 The rupture and related dust formation is the result ofhydrogen accumulation and pressure build up, which is effectively released by the failure of the blister cap. However, whether the blister has vented or not, the thermal contact with the substrate has been significantly reduced (see Figure 11). This eventually results in melting or vaporization of the thin blister cap, particularly during transient thermal loading conditions as described in Section 4.17.4.1, which contributes to further materi­al’s erosion and probably plasma contamination.268

(a)

Large blisters

Small blisters

(b)

(b)

Large blisters

r>’

(c)

Crack/void along grain boundary

(c)

Figure 12 Scanning electron micrographs of small blisters appearing at tungsten exposed to a hydrogen fluence of 1026 D m~2 at 480 K (45° tilt). (a) Initial stage; (b) growing; (c) bursting. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386-388, 356-359, with permission from Elsevier.

Figure 11 Scanning electron micrographs of tungsten exposed to a hydrogen fluence of 1026 D m~2 at 480 K (45° tilt). (a) Overall image; (b) cross-sectional image of a large blister; (c) internal image of small blisters. Reproduced from Shu, W. M.; Kawasuso, A.; Yamanishi, T. J. Nucl. Mater. 2009, 386-388, 356-359, with permission from Elsevier.

In contrast, from the point of view of H retention, blistering is favorable because tritium accumulates in blisters, which act as a diffusion barrier for hydrogen, and which can otherwise penetrate deep into the

material even at rt269-274 until it finally ends up in the heat sink structure and the coolant. Accordingly, as tritium retention is, in general, strongly correlated with the generation of blisters,275 it shows a maxi­mum at an irradiation temperature of about 500 K. , , , , However, the retention of

tritium and deuterium is also dependent on the trapping sites existing in the material. These are, in ascending order of their trapping potential, resid­ual impurities, from which slow desorption occurs even at RT,277 grain boundaries and dislocations,
radiation-induced vacancies and vacancy clusters, and pores. Depending on the occurrence and domi­nance of particular sites, the temperature at which the maximum hydrogen retention is observed varies between 450 and 850 K.138,262,264,268,277-281

With increasing temperature (>1000K), such as that occurring at the strike point of the divertor, and with the lack of blisters, continuously decreasing hydrogen retention is observed.2,255,278,282,283 The remaining amount of retained hydrogen might be attributed to the presence of hydrogen as a solute, which depends only slightly on the incident ion energy, but scales with the implantation fluence and which is assumed to be of the same order of mag­nitude as the trapped concentration. It decreases only slightly with increasing temperature and at 1600 K still amounts to about 10% of the initial hydrogen content retained at 300 K.274

In addition to blisters, the high amount of hydrogen out-gassing at temperatures of 873 K and above results

in the formation of bubbles and pores.253,277,284,285

This effect depends on the ion energy and fluence, which determines the amount and penetration depth of trapped hydrogen. Even though a beneficial smoothing effect on the surface quality is observed in comparison to pure annealing without hydrogen impact, at high temperatures up to 2500 K the surface smoothening might be accompanied by detrimental

crack formation.284

A recent review of mixed-material effects in ITER62 provides background information on mixed-material formation mechanisms and plasma-surface interac­tion effects. Here, the focus is on beryllium — containing mixed-material surfaces (i. e., Be/C and Be/W) and the conditions when one might expect these surfaces to dominate the observed plasma — surface interactions. In addition to plasma interac­tions with mixed-material surfaces, which will be discussed here, other aspects such as changes to ther­mal conductivity, material strength, and ductility, the impact of impurities on material joints, etc., must also be carefully evaluated.

4.19.3.3.1 Be-C phenomena

Beryllium and carbon have been observed to begin thermally interdiffusing at a temperature of around 500 °C,56 resulting in the formation of a beryllium carbide layer. However, beryllium carbide has also been observed to form during energetic carbon ion bombardment of beryllium surfaces at room temper­ature.110 As mentioned in Section 4.19.3.1.2, the change in the binding energy ofthe carbide molecule affects the sputtering yield of both the beryllium and carbon atoms. In addition, the formation of beryllium carbide also has a dramatic effect on the chemical erosion properties of a carbon surface bombarded with energetic beryllium ions.67,68,111

The presence ofberyllium carbide on the surface of a carbon sample exposed to deuterium plasma has been shown to correlate with the reduction of chemical erosion of the carbon surface.70 The speculation for the cause of this effect is that the carbide enhances the recombination of deuterium in the surface, thereby lessening the amount of deuterium available to interact with carbon atoms on the surface. This is similar to the impact of small amounts of boron carbide in a graphite surface affecting chemical erosion.112 However, the difference here is that instead of obtaining the carbide through an expensive production technique, the car­bide forms naturally as beryllium ions in the plasma interact with the carbon surface.

A systematic study of the time necessary to sup­press chemical erosion of a graphite surface due to the interaction with beryllium-containing plasma has been carried out.69 Increasing the surface tempera­ture of the graphite was seen to have the biggest impact on reducing the suppression time. Increasing the beryllium content of the plasma also reduced the suppression time in a nonlinear fashion. An increase of the incident particle energy was observed to increase the time necessary to suppress the chemical erosion of the surface, presumably due to an increase in the removal of the carbide-containing surface layer. A subsequent study showed that applying heat pulses to a graphite surface interacting with beryllium-containing plasma, to simulate surface heating due to intermittent events, acted to reduce the time necessary for the carbide surface to form and suppress the chemical erosion of the surface.1

4.20.4.1 Tensile Properties

The influence of test temperature, strain rate, and thermal-mechanical treatments on the tensile prop­erties of copper and copper alloys has been studied extensively. Figure 3 illustrates the effect of test temperature on the yield strength of pure copper (in the annealed condition), PH CuCrZr and CuNiBe alloys, and DS CuAl25.15-18,28,32-39 The strength of copper alloys decreases with increasing test tempera­ture. The decrease in strength is moderate up to 200 °C. Significant drops in strength occur at higher temperatures, except that the CuNiBe AT alloy shows a relatively small reduction in strength even up to 400 °C. Pure copper has the lowest yield strength. The tensile properties of pure copper strongly depend on the thermal-mechanical treatment and the impurity content.15-1 ,32,33 CuNiBe alloy has the highest strength over the entire temperature range.34 The tensile properties of PH copper alloys are sensi­tive to the thermal-mechanical treatments. CuCrZr in the solution-annealed, cold-worked, and aged con­dition (SA + CW + A) has superior yield strength at low temperatures relative to CuCrZr in the solution- annealed, and aged condition (SAA). However, the strength of CuCrZr SA + CW + A alloy drops more rapidly with increasing temperature.29,34-39 The yield strength of CuNiBe can be quite different, depending on the processing techniques. The tensile ductility of copper alloys also shows strong temperature depen­dence. The uniform elongation of the CuAl25 alloy decreases considerably as the test temperature in­creases, but increases with increasing test temperature above 400 ° C. The CuNiBe AT alloy shows a moder­ate drop of uniform elongation below 200 °C, but a sharp drop in ductility at higher temperature.34 The uniform elongation of the CuCrZr alloy shows the smallest sensitivity to test temperature. Among

the three copper alloys, the CuCrZr alloy has the best ductility over the temperature range, and the ductility of the CuNiBe alloy is the lowest.

Because of the sensitivity of mechanical properties to thermal-mechanical treatments in PH copper

alloys, the strength of large components made of these alloys can be significantly lower. For example, during component manufacturing, CuCrZr often experiences additional thermal cycles, such as braz­ing, welding, or HIPing. While solution annealing
can be conducted during or after a brazing or HIPing process, rapid quenching is not feasible for large com­ponents, and a much slower cooling rate (e. g., furnace cooled or gas cooled) is applied in the manufacturing cycle. Significant reduction in strength due to slow cooling rates has been reported in CuCrZr.30,40-42 A slow cooling rate (50-80 ° Cmin~ ) and overaging at 560 °C/2h significantly reduced the yield stress and the ultimate tensile strength, and tensile elongations of CuCrZr relative to prime-aged CuCrZr.14 Cooling rates >1200 °C min-1 are required to fully quench the Cu-Cr solid solution.43-45

The effect of strain rate on tensile properties for pure copper and PH CuCrZr and CuNiBe alloys as well as DS CuAl25 alloy was studied at temperatures of 20 and 300 °Q14,34,46 All three copper alloys are relatively insensitive to strain rate at room tem­perature. The strain rate sensitivity parameter of m (where ay = Ce’mand C is a constant) is ~0.01 for the CuAl25 alloy at room temperature. The strain rate sensitivity of this alloy increases significantly with increasing temperature as reflected by a strain rate sensitivity parameter of m ~ 0.07 at 300 °C. Stephens et a/.47 reported a strain rate sensitivity parameter of m ~ 0.1 in the temperature range of 400-650 °C for CuAl25. A similar effect of strain rate on ultimate tensile strength was also observed on these materi — als.34,46 Edwards46 investigated the strain rate effect of copper alloys in air and vacuum, and found that
testing in air or vacuum did not appear to change the strain rate dependence of the CuAl25 alloy, but that testing the CuNiBe alloy in air shifted the embrittlement to a lower temperature.

Loading a cracked structure creates a local stress concentration ahead of the crack tip. A situation

where brittle fracture initiation is likely only in this restricted area around the maximum stress corre­sponds to small-scale yielding conditions. The high stress area is localized ahead of the crack tip extending to the border of the crack tip plastic zone or three to five times the distance from the crack tip to the stress maximum.18 In such a situation, the stress distribution ahead of the crack can be described correctly with the У-integral. When loading exceeds small-scale yielding, large-scale yielding is involved and the У-integral can no longer describe the crack tip stress distribution correctly. At that point, the measuring capacity of the specimen has been exceeded.

The limit for measuring capacity is normally given as a function of the material yield strength and specimen ligament in the form of eqn [14]. The value of M has been proposed from a variety of finite element calculations to be from 50 up to 200.18 Based on present understanding and examination of many
experimental studies, a value of M = 30 can be regarded as a realistic estimate for cleavage fracture with most structural steels. Exceeding the load level given by eqn [14] will lead to gradually increasing amounts of ductile tearing before cleavage fracture initiation.

The basic analysis procedure described in ASTM E 1921 is intended for specimens and structures for which at least a moderate level of constraint (triaxial stress state) is achieved. With deep-cracked speci­mens or thick-wall structures including internal cracks, the constraint is typically high enough for the standard T0 analysis. The situation is different with low-constraint geometries like surface cracks and thin-sheet structures where the small-scale yielding condition may not prevail (note that the limit for this condition also depends on material strength properties). In principle, the basic Master Curve approach can also be applied for such low — constraint conditions, but the estimation may be overly conservative due to high plastic deformation which is outside the applied fracture model assump­tions. In ASTM E1921 test conditions, sufficient constraint is assured by defining valid tests only as those that exhibit brittle fracture initiation below or at the capacity limit value (eqn [14]) and by limiting the ductile crack extension prior to brittle fracture initiation.

The probability of cleavage initiation is controlled by the narrow zone ahead ofthe crack tip where small — scale yielding condition prevails. Several approaches like the Tstress, the Q-parameter, and small-scale yield­ing corrections have been developed to account for the effect of plastic deformation due to low constraint.19 The Tstress analysis is relatively straightforward since a simple elastic stress analysis can be used instead of a numerical large-scale yielding model. Another advan­tage of using the Tstress is that it may be performed assuming no change in temperature dependence which allows the constraint effect to be described only as a shift in T0. Consequently, the Tstress yields conservative estimates compared to the more complex Q-parameter approach. On the other hand, the Q-parameter is more accurate for very low-constraint situations. On the basis of a simplified linear-elastic analysis, the correc­tion to T0 due to low-constraint SE(B) geometry (a0/W> 0.1) with a negative Tstress has been expressed in the form19:

Tt 10MPa °C~

Tstress < 0 (for SE(B) specimens only)

where T0 is the corrected value and T0deep is the value measured for a deep crack case.

Equation [26] is an empirical result from data consisting of only SE(B) specimens (Figure 16). Additional work has been conducted to refine the expression by including test data from C(T) speci­mens and by a comparison of solutions based on Tstress and the Q-parameter. Thus, a more accurate formula for estimating the T0 from Tstress and the T0 of a deep crack case has been proposed in the form20:

T

T0 « T0 deep + — 1 for Tstress <300 MPa

0 0 deep!2MPa °C_1

(for SE(B) and C(T) specimens) [27]

The main difference between eqns [26] and [27] is that the latter also covers positive values of Tstress up to 300 MPa, whereas in eqn [26], it is assumed that only a negative T«ess has a marked effect on the value of T0.

In the majority of cases, lithium ceramic samples are irradiated in research reactors with thermal or mixed neutron spectra. Extensive studies on tritium retention in neutron irradiated lithium ceramics using the TPD
method have been reported in the literature.152-178’201 The chemical form of released tritium from Li4SiO4 (from FZK), LiAlO2 (from JAERI), Li2TiO3 (from CEA), and Li2ZrO3 (from MAPI) was studied in the out-of-pile tritium release experiment under various purge gas conditions. The pebbles were irradiated for a few minutes in the fluxes 4 x 1017m-2s-1 at Japan Research Reactor 4 (JRR-4) or 1.65-2.75 x 1017m-2 s-1 at the Kyoto University Reactor. The tritium was

released in the purge gas of dry nitrogen, nitrogen with 0.1% of helium, and nitrogen with 0.1% water vapor. Even if hydrogen was added to the purge gas, a considerable fraction of tritium was released in the molecular form of water, HTO. Addition of water vapor to the purge gas greatly enhanced the release rate of tritium. It was concluded that the isotope exchange of tritium with water at the exposed surfaces of the grains is much faster than the isotope exchange of tritium with hydrogen. The water is adsorbed at the grain surfaces from the water vapor present in the purge gas. Even small traces of water of ^30 ppm in dry purge gas canbe enough to promote tritium release in the HTO molecular form. Another source of water at the surfaces of the grains is the reduction reaction that takes place in the H2 reducing atmosphere. Among the studied materials, Li2TiO3 showed the largest water formation capacity.161 Figure 49 shows out — of-pile annealing tests for Li4SiO4 with 0.1% H2/N2 sweep gas and 0.1% H2O/N2 sweep gas, and tritium doped for 2 min at a neutron flux of 2.75 x 10 cm s.

Later work on palladium deposited as catalyst on orthosilicate indicated that almost all tritium was released as tritiated water vapor from lithium ortho­silicate pebbles and tritium at higher temperatures remains slow.160 In contrast, it was also found that a considerably larger amount of tritium was released as the molecular form (HT) from the lithium

orthosilicate pebbles already deposited with palla­dium at lower temperatures (see Figure 50).

Alvani and coworkers172,173 correlated TPD after short irradiations in the Casaccia reactor (~2 x 1021m~2) with those from long-term irradia­tion in the HFR; Petten (thermal ^0.5 x 1025m~2) revealed two peaks, at 770 and 941 K (b = 5 Kmin-1) (see Figure 51). It was proposed that the second peak is related to tritium trapping at the oxygen vacan­cies located along the grain boundary interface. The concentration of these trapping sites is signifi­cantly increased by the reduction effect of the R-gas, which results in a shift in the release peak to higher temperatures for the pretreated pebbles. The observed effect is more pronounced for the pebbles with finer

grains. It was also found that a thermal pretreatment at 473 K for 2 h removes only the environmental H2O and CO2 contamination from the surface of the pebbles without affecting the H2O desorption at higher tem­peratures. The observed H2O release above 1173 K is ascribed to the reduction of Li2TiO3 to Li2TiO3_x with x reaching a steady-state value of xeq = 0.01.

Later work by Casadio eta/.174 concerned a batch of Li2TiO3 pebbles (ENEA code FN5) prepared following the ‘citrate’ route. Analysis of TPD spectra gave the correct order of magnitude of the time constants characterizing the main desorption sites, in rough agreement with the residence times obtained by the in-pile step-perturbation methods performed during EXOTIC-8/9 experiment (see Figure 52).175 Pure helium purge increases the tritium inventory; during the last cycle of this irradiation experiment, variations of the H2 concentration in the He purge showed an increase in tritium release rate from Li2TiO3 pebbles that was found to be proportional to PH34 at 473 °C, Figure 53.

The effect of open and closed porosity on tritium release behavior in Li2O single crystal and sintered pellets was studied by Tanifuji eta/.176-179 The pellets had densities in the range of 70-98% and grain sizes from 10 to 60 pm. Irradiations were performed by ther­mal neutrons in JRR-4, JAERI, up to 2 x 1023 m~2. The porosity dependence of tritium release behavior from the Li2O sintered pellets has been investigated through isothermal heating tests, and the results are shown in Figure 54.

For 88% TD specimens irradiated up to neutron fluences of 2 1 022 and 2 x 1023nm~2, no

irradiation effects on the tritium residence time have been observed.

As the high atomic number is an intrinsic material property that cannot be changed, the only possibility to avoid plasma contamination by tungsten is to adapt to the loading realities, that is, thermal loads and plasma wall interaction conditions, and the energy of the incident plasma particles. In particular, surface crack formation, loosening of particles, and particle ejection or melting are addressed (see Section 4.17.4). Concerning the latter, the addition of suitable alloying elements or dispersoids (see Section 4.17.3.3) reduces the material’s thermal conductivity causing a reduction of allowed applied heat fluxes. From this point of view only low-alloyed grades should be considered and the best grade is tungsten of high purity.

4.17.3.2.1 Recrystallization

Recrystallization is a thermally activated process. Therefore, it is expected that the activation energy of nucleation is dominated by small angle grain boundaries. The activation energy of grain growth is dominated by large angle grain boundaries.64 The temperature of recrystallization depends mainly on the deformation history, that is, the higher the degree of deformation, the lower the recrystallization tem — perature,65,66 and the chemical purity. When heated above the recrystallization temperature, the structure of tungsten is altered due to grain growth causing an increase in DBTT and reducing other mechanical properties, that is, strength and hardness.67

There are several possibilities for increasing the recrystallization temperature. Particle reinforcement and controlled formation of porosity are the best and most investigated options.68 For example, the higher recrystallization temperature of dispersion strength­ened alloys results from the interaction between dispersoids and dislocations during hot-working; the higher the amount of hot-work, the finer are the dispersoid particles and the higher is the recrystalli­zation temperature. During recrystallization, these particles prevent secondary grain growth and conse­quently, the recrystallization temperature of disper­sion strengthened alloys may increase compared to pure W.67 Another example is highly creep-resistant doped/nonsag materials with aligned porosity acting as obstacles for dislocation movement as they are used in the lighting industry.69

Experience shows that incomplete recrystalliza­tion often helps to achieve the desired balance in material properties. If the operating temperature is well known, controlled recrystallization during application might be feasible as well.67 However, for operational conditions in nuclear fusion devices, it is expected that the very high thermal strain rates experienced in the thin layer heated by plasma dis­ruption or any other transient thermal event will significantly affect the material’s microstructure and properties.

Significant increases in both strength and elastic modulus occur in graphite at dose levels as low as 0.01 dpa.8 This increase continues to high displace­ment levels until volumetric expansion and extensive micro cracking occurs. For graphite, the reversal to property degradation typically occurs at tens of dpa depending on the graphite type and irradiation tem­perature. The increase in modulus is a result of dislo­cation pinning by lattice defects produced by neutron irradiation. The magnitude of the increase is depen­dent on the perfection of the graphite. For most graph­ite types, a maximum modulus increase of 2-2.5 times the nonirradiated value is typical for irradiation tem­peratures less than 300 °C, with the change becoming less pronounced at higher irradiation temperatures. Irradiation-induced increase in strength occurs in a similar fashion as in the elastic modulus.

Several authors17-23 report the effect of neutron irradiation on the elastic modulus of CFC. For example, Sato18 reports an increase of 42% and 30% in modulus following neutron irradiation to 1.2 x 1025nm~2,

E > 0.18 MeV in the temperature range of 750-810 °C for a 2D pitch fiber and PAN fiber composite, respec­tively. Similar to the irradiation-induced increase in strength, the absolute increase and percent increase in elastic modulus is highly dependent on starting material and irradiation condition. The irradiated and nonirradiated mechanical properties of some can­didate ITER PFC materials are shown in Table 4 for ITER relevant temperatures and doses somewhat higher than ITER neutron doses. Specifically, these materials were irradiated at approximately 1000 °C to a dose of about 2 dpa.20 The change in properties is relatively small because of the high irradiation tem­perature and the relatively low dose.

As with elastic modulus, reported data on the effect of irradiation on the strength of CFCs are somewhat sparse.16-22 Snead23 has reported the strength and elastic modulus of the 3D pitch fiber composite FMI-222 for doses higher than expected for ITER, or more consistent with a fusion power reactor. Figure 13 gives the modulus as a function of dose to 32 dpa at 800 ° C, exhibiting a marked increase to at least 10 dpa followed by a degradation by the 32 dpa

Dose (dpa)

value. The same samples, as seen in Figure 14, exhibit more than a 50% increase in strength, which is retained even at the 32 dpa value. This is particularly remarkable given that the composite had undergone significant dimensional change in this dose range.

Actively cooled first-wall mock-ups have been tested under relevant heat fluxes in electron beam facilities (accelerated fatigue tests) or in low-heat flux facil­ities, which are capable of delivering moderate heat flux (< 1 MW m~2) via electrical heaters for a very large number of heat pulses (>10 000) of larger dura­tion to investigate the thermomechanical performance (including integrity and possible creep phenomena) of the bonded interfaces and fatigue lifetime. Because of intrinsic limitations, cost, or the limited availability of testing facilities, electron beam facilities are used with the shortest duration of the heating cycle, com­patibly with the conditions of near thermal steady — state. In many cases, to reduce the testing time, the tests are performed at a heat flux well above the nomi­nal one. These tests are often used to select the best joining conditions and then confirmation and final selection are done including tests at representative moderate heat fluxes but for the specified large number of cycles. Because of the health hazards linked to the use of beryllium material, only a limited number of test facilities were available or built for testing beryllium components. As far as Be-compatible electron beam facilities are concerned, two are available in Europe, the Juelich divertor test equipments in hot cells (JUDITH 1 and JUDITH 2) at the Forschungszen — trum Juelich (FZJ) (Germany), one in the Russian Federation at the Efremov Institute in St. Petersburg and one in the United States at Sandia National Labo­ratory in Albuquerque (New Mexico). For the perfor­mance of thermal fatigue tests at moderate heat fluxes and large number of cycles, three facilities were built in Europe, one at the Ispra Joint Research Center in Italy, one at the ENEA Research Center of Brasimone in Italy, and one at the Nuclear Research Institute (NRI) of Rez close to Prague in the Czech Republic. From these three facilities, only the last one is still under operation.

The best high heat flux test results from represen­tative first-wall mock-ups, namely, 10-mm beryllium tiles joined onto 10 mm CuCrZr heat sink layer with embedded 10/12 mm diameter stainless steel pipes and the assembly joined onto a stainless steel backing plate (see Figure 11), were achieved with Ti and Cu interlayers and HIP temperatures of about 580 °C as described in Section 4.19.5.1.1.2. The per­formance limit of this assembly is presently at about 3 MW m~2.162

As far as the thermal fatigue tests are concerned, representative first-wall mock-ups with Be/CuCrZr

joints made as above have been successfully tested up to 30 000 cycles at about 0.6 MW m-2 (see Figure 12). Ultrasonic testing of the Be/CuCrZr joints after testing did not show any indication of defects. These mock-ups will be tested at higher heat fluxes representative of transient and off-normal events to check the available performance margins.

Originally, six ITER Domestic Agencies were candidates for the procurement of the ITER first wall: China, European Union (EU), Japan, South Korea, Russian Federation, and United States. (The seven members of the international ITER project have all created Domestic Agencies to act as the liaison between national governments and the ITER Organization. The Domestic Agencies’ role is to han­dle the procurement of each member’s in-kind con­tributions to ITER.) However, as stated in the Final Report of Negotiations on ITER Joint Implementa­tion of 1 April 2006, a prequalification ‘‘… will be needed for the critical procurement packages shared by multiParties…’’, such as the blanket first wall. Well in advance of the assumed start of the procure­ment, each ITER Domestic Agency shall first dem­onstrate its technical capability to carry out the procurement with the required quality, and in an efficient and timely manner. For the first wall system, this is achieved via a two-stage qualification process: a mock-up qualification stage and a semiprototype qualification stage. Each stage is also split into two phases: a manufacturing acceptance phase and a heat flux testing acceptance phase. The successful manufacturing and testing of two first wall mock — ups for stage I (see Figure 13) demonstrating in particular the know-how to assemble beryllium (Be)

tiles on a CuCrZr alloy and stainless steel bimetallic structure is the prerequisite to be eligible for stage II.

The qualification tests for stage I have been split between the United States and EU in three test facilities: at the SNL for the US and at the NRI in the Czech Republic and the FZJ in Germany for the EU. At the SNL facility, the qualification test programme consists of the performance of 12 000 cycles at 0.88 MW m-2 for 1.6 min followed by 1000 cycles at 1.4 MW m-2, while in the EU test facilities it consists of the performance of 12 000 cycles at 0.62 MW m-2 for 5 min (at NRI) followed by 1000 cycles at 1.75 MW m-2 (at FZJ). To be qual­ified, a Domestic Agency shall fabricate two mock — ups which pass both tests.

The first wall mock-ups fabricated by the EU Domestic Agency have successfully achieved the above test programme conditions without any indi­cation offailure. Additional tests were also performed on these mock-ups to assess the limit and tests were performed for 200 cycles at 1.7 MW m-2 at the SNL and for 100 cycles at 2.25 MW m-2 plus 100 cycles at 2.75 MW m-2 at FZJ without any indication of fail­ure. Tests were stopped so as not to exceed the maximum acceptable Be temperature in the test facilities. The progress on the fabrication and thermal tests is described elsewhere.163-171