As mentioned above, hydrogen retention depends on the trapping sites available in the material and their relative energies. Their existence and concentrations are influenced not only by the impinging H-ions, but also by the manufacturing process, thermal pre­treatments, the material’s composition and micro­structure, and the surface quality. Accordingly, the retention increases with the amount of porosity in the material, as it allows a deep penetration of hydrogen and the voids and pores provide the highest

trapping energies264,275 with thermal desorption

occurring at temperatures >700 K.45,262

Another material parameter that increases hydro­gen retention is the number of dislocations,266,286 particularly those introduced during deformation processes used for material densification. However, the recrystallization of the material removes not only dislocations but also vacancies and vacancy clus­ters, which have been introduced by the impinging H-ions286,287 and as grain boundaries. This effectively reduces the trapping sites for hydrogen retention and, consequently, the lowest retention is observed for high-purity SC materials, particularly due to the low diffusion rate compared to polycrystalline tungsten.275,288,289 This low diffusion rate results in a near-surface accumulation of hydrogen, which acts as a diffusion barrier and leads to a saturation of hydrogen retention with increasing fluence.290 Such saturation is not observed for pure polycrystalline tungsten due to the possible hydrogen migration along grain boundaries.291

Finally, the hydrogen retention is influenced by impurities277 and dopants. The addition of La2O3 and TiC particles as well as the formation of pores, for example, by potassium doping, not only introduce traps and increase hydrogen retention,293 but also decrease the diffusion rate.291 In contrast, alloying with up to 10% Re has no measurable effect on the H retention properties of the material,279 as it only creates a slightly deformed crystal lattice structure but no additional hydrogen traps.

In addition to hydrogen retention, material dam­age and particularly blistering is influenced by the material’s microstructure. Blistering occurs preferen­tially when the crystal is oriented with the (111) direction perpendicular to the surface292 and the blisters develop in different shapes from low, large, and spherical to high, small, and dome or cone-

shaped.45,262,267 The blisters in recrystallized materi­als are mainly plateau-shaped, often multilayer structures, which indicate a step-wise build-up, and in few cases also small blisters on top of large ones are formed.263,293 However, it is significant that the blister size is commonly limited by the grain size45,294 indicating that the grain boundaries play an important role in the formation of blisters. Accordingly, SCs and nanostructured materials such as W-TiC provide the strongest resistance against blistering, although the particular reason is different. For SCs, the hydrogen diffusion and accumulation is limited and there is a fast desorption from low — energy traps at elevated temperatures. In contrast, for nanostructured materials the size of individual grains is extremely small and so is the volume for blister formation. Furthermore, the migration of hydrogen is significantly increased by the large num­ber of grain boundaries.292

Further material parameters that reduce blister formation are open porosity and the surface finish, particularly the number of random or artificially introduced scratches that might act similar to grain boundaries.45 In contrast, the introduction of impurities and dopants in commercially available grades of tungsten increases the number of blisters and exfoliation in both their stress relieved and recrystallized states.277,293

4.17.4.4.3 Combined loading conditions

As described above, the damage mechanisms of hydrogen and He-irradiation are rather similar, although they occur in different temperature ranges. Accordingly, their mutual interaction is also strongly influenced by the implantation tempera­ture. Therefore, the testing sequence plays a role in the behavior, as for He preirradiation followed by hydrogen implantation, the implantation tempera­ture of He determines the amount and kind of produced damage and the He-retention, which sub­sequently influences the hydrogen uptake occur­ring as described in Section 4.17.4.4.2. For example, He-implantation at RT either does not change the retention or may increase it due to the formation of additional trapping sites295-297 and the lower diffusion rate of He compared to H. With increasing He implantation temperature up to 800 K, hydrogen retention significantly decreases compared to pure hydrogen irradiation.261,292,296 This may be attributed to the occupation of trap sites by He as a result of its increasing mobility.298 Potential trap sites are the numerous He-induced nanosized bubbles acting as a diffusion barrier.292 A further increase in temperature to 1600 K does create significant material damage by He due to pore and bubble formation or even blistering. This tremendously increases the number of trap sites in the material and leads to He desorption during implantation and accordingly increases the hydro-

299

gen retention.

For simultaneous loading of He and hydrogen, the fraction of He should reach at least 5 at.% to observe significant changes in the material’s response.261,292 Furthermore, for implantation tem­peratures below 900 K, results similar to those described above are observed for sequential ion beam loading.299 However, due to desorption of hydrogen at high temperatures >1000 K, no hydro­gen retention takes place and the damage mechan­isms are dominated by the He-irradiation during such temperature excursions.

Correlated with hydrogen retention, blister for­mation at temperatures <900 K decreases with decreasing hydrogen retention. This is valid until the number of voids and pores, which enhance hydro­gen retention, start to form open porosity and thereby generate small grain structures. These allow a fast hydrogen diffusion through the material and limit the agglomeration of hydrogen necessary to form blis­ters similar to the case of nano-structured materials such as W-TiC described above (Section 4.17.4.4.2).

Investigations of the influence of radiation dam­age (highly energetic hydrogen, neutrons) and impurity irradiation, for example, by carbon atoms, resulted in the depth resolved and particle energy — dependent formation of dislocations, dislocation loops, and even small voids acting as effective trapping sites for hydrogen and influence blister formation.274,276,300-309 Upon annealing, the disloca­tions and dislocation loops were moved and/or

annihilated,310,311 which is positive news as it

would limit the tritium inventory,312 as long as no He is present in the system. In contrast, with the addition of He the dislocations, dislocation loops, and helium bubbles do not vanish at identical annealing conditions, which has a direct impact on the mechanical and thermophysical performance of tungsten. However, He positively influences hydrogen retention in an intermediated tempera­ture range as described above and inhibits the for­mation of a W carbide layer, which is typical for combined hydrogen and carbon loading.311

Finally, the results obtained from the investigation of the mutual influence of ion irradiation and thermal loads are strongly correlated with the choice of the heat source. In the literature, electron beam guns were favored,313,314 which are characterized by heat deposition in a depth range of several micro­meters for W, depending on the acceleration voltage. However, as the thickness of the ion-irradiation — affected surface layer is comparably thin, the majority of the electrons pass through the modified surface layer, which leads to most doubtful conclu­sions. In contrast, lasers are more reliable, as they apply only surface heat loads. The combination of He-irradiation and laser-induced thermal loads (A T = 1400 K, n = 18 000) at high base temperatures (~1700K), resulted in an affected layer thickness (13 pm) about 10 times larger than that without laser irradiation (1-2 pm), which might be attributed to the steep temperature gradient supporting the diffusion of He. This surface modification combined with laser-induced surface roughening, as observed in Section 4.17.4.1.2 for typical thermal shock loads, leads to an enhanced degradation of the thermal diffusivity of W, which further increases the surface roughness and results in local or full melting of the

W surface.199

4.17.5 Conclusion

Along with other favorable properties, tungsten is characterized by the highest melting temperature among all metals, a low energy threshold for sputter­ing, and a low tritium inventory compared to carbon — based materials. These characteristics make tungsten the most promising material for the plasma-facing inner wall of future nuclear fusion devices based on the magnetic confinement principle, and it is also under consideration for inertial fusion applications. Accordingly, it has been selected as the PFM for a large part of the ITER divertor during its start-up phase and will be used for the full divertor as soon as tritium operation starts; in addition, it is the reference material for DEMO.

However, tungsten also offers less favorable prop­erties. Related to these, there are some material issues that have to be resolved before operating tungsten in a fusion environment in an economically reasonable way, which means in DEMO and beyond. These are

• recrystallization, which influences the mechanical properties by reducing the ductility and increasing the DBTT

• embrittlement as a result of neutron-induced damages and transmutation

• resistance to crack formation, depending on the mechanical properties, which is particularly important during transient thermal loads

• He-induced sputtering and modification of a thin surface layer, which is influenced by existing mate­rial damage as well as by temperature and temper­ature gradients, for example, those occurring during transient thermal loads

• melting, which is related to crack formation and the degradation of thermophysical properties as a result of He-irradiation-induced surface modifica­tion; melt splashing and droplet ejection will influ­ence the stable operation of the fusion plasma.

As all grades of tungsten investigated so far have their own individual drawbacks, R&D programs world­wide are aiming for a deeper understanding of the parameters that influence the degradation of tung­sten, and the development of new tungsten grades that are capable of dealing with the above-mentioned requirements. Therefore, the materials are character­ized and qualified with regard to their microstructure before and after recrystallization by

• mechanical tests: evaluation of the material’s strength and DBTT

• thermal shock loading: determination of tempera­ture — and power density-dependent damage, cracking and melting thresholds, which are related to the mechanical and physical properties

• thermal fatigue loading: evaluation of the material’s performance as part of an actively cooled component

• neutron irradiation: characterization of the degrada­tion of the material’s strength and the DBTT as well as the thermal shock and thermal fatigue response

• He — and H-irradiation: determination of the damage mechanisms such as blister, void, and bubble forma­tion as a function of ion energy, fluence, and temper­ature as well as addressing hydrogen retention issues.

However, despite all these efforts, a clear answer on the suitability of tungsten for application in a real fusion environment can only be given by ITER, as it is the mutual interaction of all the different types of loading that determine its lifetime relative to the various material degradation mechanisms.

The existence of tungsten beryllide alloys (i. e., Be2W, Be12W, and Be22W) is an excellent example of the importance of mixed-material surface formation in plasma-facing components.114 Figure 9 shows the tungsten-beryllium phase diagram. Each of the ber — yllides shown in the figure exhibits a lower melting temperature than one would expect from a tungsten plasma-facing surface. If plasma containing beryllium impurities interacts with a tungsten surface, there is a possibility of these lower melting temperature beryllide alloys being formed.

In thermodynamic equilibrium, various beryllide alloys of tungsten have been observed to form,115 and their reaction rates have been measured,116 at tempera­tures in excess of 800 °C. However, as was seen with beryllium carbide forming during plasma bombard­ment at lower temperature than expected thermody­namically, the concern exists that tungsten beryllide could form at temperatures below 800 °C as well.

Well controlled laboratory measurements in vac — uum117 and in plasma simulators118 have shown that although thin, nanometer scale, Be2W layers form at the interface between beryllium and tungsten surfaces, their growth below 800 °C is negligible. In addition, above 800 °C, rapid beryllium sublimation from surfaces can act to limit the amount of beryl­lium available for reacting with tungsten and thereby also limit the growth rate of the alloys. In the present low wall temperature confinement devices, modeling shows that the divertor strike point locations are the only areas where significant beryllide growth might be expected and in these regions there does not

appear to be enough beryllium deposition to raise significant concerns.11 One caveat to these predic­tions would be the existence of intermittent events that raise the temperature of surfaces where signifi­cant beryllium deposits are located, thereby possibly allowing the optimized beryllide growth conditions.

Another concern with regard to thin Be2W surface layers on plasma-exposed tungsten is the impact of these layers on tritium retention. While a thin Be2W surface layer is not likely to retain much tri­tium itself, the thin beryllide surface layer could alter the recombination characteristics ofthe bulk material and change the accumulation rate of tritium within the device. To date, there is little or no data available to address this issue.

While it appears likely that the most serious issues of tungsten beryllide formation may be avoided in present confinement devices, the issues associated with these alloys highlight the uncertainties and impor­tance of understanding and predicting mixed-material
formation in plasma environments. Mixed materials often interact with plasma in much different ways than their elemental components. In the case of the beryllium-carbon system (Section 4.19.3.3.1), the mixed material appears to offer the potential for bene­ficial effects, whereas in the case of the beryllium- tungsten system, the mixed material appears likely to be detrimental to the operation of the device. Each mixed-material system must, therefore, be individually evaluated to determine its potential impact on all aspects of operating surfaces in contact with plasma.

Fracture toughness data for PH copper alloys, CuCrZr and CuNiBe, and DS copper alloys, CuAl15 and CuAl25, are summarized in Figure 4.1, -0 CuCrZr has the highest toughness, and CuNiBe the lowest among these alloys. The large scatter in mea­sured fracture toughness values for CuCrZr in differ­ent studies is likely due to different heat treatments, specimen geometry and dimensions, and testing meth­ods. The temperature dependence of the fracture toughness in CuCrZr, while difficult to accurately define, shows an initial decrease with increasing tem­perature, and then a slight recovery at temperatures above 250 °C. The effect of thermal-mechanical treat­ment on fracture toughness of CuCrZr is insignificant in comparison with its effect on tensile properties.14 The minimum value of the JQ for unirradiated CuCrZr is as high as ^100 kJ m~2.

The fracture toughness of DS CuAl15 and CuAl25 is significantly lower than that of CuCrZr, and shows a strong directional dependence. The toughness is higher in the L-T orientation than in the T-L orientation. The fracture toughness decreases

rapidly with increasing temperature. The JQ value for CuAl25 is only 7 kJ m~2 at 250 °C in the T-L

48

orientation.

The effect of ductile tearing on the measured J-integral is addressed in ASTM E 1921 by limiting the amount of ductile crack growth prior to brittle fracture to avoid inaccuracies from excessive plastic deformation. In some standards, specific formulas for correcting the effect of excessive ductile crack growth have also been presented. It is possible to correct in this way for small crack growth relative to the exces­sive plastic deformation that occurs, but the statistical effect associated with the probability of cleavage fracture initiation cannot be corrected.

Ductile tearing preceding brittle fracture affects the measured fracture toughness by increasing the
volume ahead of the crack tip where brittle fracture initiation can occur, which increases the cumulative probability of failure. This statistical crack growth effect is comparable to the statistical size effect, which also is due to the increased volume of material under stress, increasing the probability of brittle fracture initiation as the crack length increases. In addition to increasing the volume of potential cleavage initiators, ductile tearing also tends to change the crack tip stress distribution. On the other hand, a small amount of ductile tearing can be regarded as beneficial since it can increase the stress triaxiality at the crack tip. This triaxiality effect is of minor importance when compared to the statistical effect that is dis­cussed next.

The statistical effect is due to the preceding ductile crack growth and is dependent on the amount of the crack growth.18,21 A simplified expression for the fracture probability (Pf) at stress intensity (Kj), considering a small amount of ductile tearing (Da < 1 mm), is given in the form1 ,22:

2O2Da
KI2(2m + 1)

[28]

where m and O are material-dependent constants, B is specimen thickness, B0 and K0 are scale factors, and Kmin = 20 MPa Vm. The value of O has been esti­mated to be 5500 MPa for medium strength struc­tural steels (ffy = 300—550 MPa). Equation [28] can be simplified by setting (2 m + 1) = 1 when O becomes ^4700 MPa. Parameter m is defined from the ductile — tearing power law function of the form:

f (Da)=J1mmDam [29]

where J1 mm is the value of J-integral at 1 mm crack growth and Da is crack growth.

By substituting the expression for failure pro­bability (eqn [13]) into eqn [28], one can derive a more practical formula giving a corrected value for Kj (Kj eff) due to ductile crack growth as follows:

2O2Da 1/4 r n

Kjeff = Kmin + (Kj — Kmin) 1 + ,7 [30]

Equations [29] and [30] can be used to take into account the increased fracture probability due to small amounts of ductile tearing, in addition to a possible standard correction for plastic deformation. Consideration of the statistical ductile-tearing effect may become relevant for high-strength steels with a
low-strain hardening capacity or for steels exhibiting low ductile-tearing resistance.

An example of analyses corrected for ductile crack growth is shown in Figure 17. The material is a ther­mally embrittled pressure vessel steel (A508 Cl. 3), which therefore has a high T0 (+69 °C). In this case, the correction lowers the fracture toughness in the upper transition region, but has a negligible effect on the value of T0 and the behavior near the lower shelf.

In-pile experiments have the strong advantage that the tritium release characteristics can be studied, as a function of neutron damage and lithium burnup in combination with thermal-mechanical behavior under neutron irradiation. Such in-pile experiments allow steady-state tritium production and release con­ditions to be achieved. In general, such parameters are closer to breeding blanket conditions, as they allow the application of a wide range of temperatures and purge gas conditions and the study of long-term performance issues such as irradiation damage and lithium burnup. At present, such data are limited in terms of fast

neutron damage doses (thermal and mixed spectra MTR only).

Several irradiation programs have been executed around the world, involving thermal, mixed — spectrum, and fast reactors [as in Dido (Germany), Siloe (France), EBR-2 (US), FFTF (US), HFR (The Netherlands), JMTR (Japan), and WWRK (Kazakhstan)]. The European irradiation projects under the acronym EXOTIC, for extraction of tritium in ceramics, commenced during the mid 1980s at the HFR in Petten , , , , , , (see Table 4).

The initial series concerned both closed capsule and vented capsule operation. The materials investigated were mainly Li2O, Li2SiO3, LiAlO2, LiZrO3, Li8ZrO6, and Li4SiO4 in the series EXOTIC-1 to-6.26 The objective of the EXOTIC-7 experiment has been to irradiate candidate ceramic breeder materials in the HFR to a high lithium burnup (target ~ 10%) and to determine the effects on the mechanical integ­rity of pellets and pebble-bed configurations, and those on tritium-inventory and-release characteris — tics.33 The experiment concerned 8 capsules, and during 11 HFR cycles (261 FPD), lithium burnups of 6-18% were achieved. The test matrix comprised pellets of Li2ZrO3, Li8ZrO6, and LiAIO2 and pebbles of Li2ZrO3 and Li4SiO4. Two capsules contained a mixture of Li4SiO4 and beryllium pebbles. To obtain a high lithium burnup within a reasonable irradiation time, the target materials were enriched with 6Li to about 50%.

The tritium residence time of Li2ZrO3 pellets a higher temperatures was not changed during the irradiation. After a lithium burnup of about 5%, the residence time at lower temperatures decreases significantly with increasing burnup.

After selection of the HCPB as the single solid breeder concept in the European Blanket Project, the EXOTIC-8 and-9 series concentrated on Li4SiO4 and Li2ZrO3 pebbles and a range of Li2TiO3 pro­ducts.175 The irradiation test program concentrated on two types of experiments:

1. Tritium release to low or moderate lithium burnups

2. High lithium burnup and mechanical integrity

The typical designs for these experiments are given in Figure 55, with the general layout and an example of a cross-section from postirradiation testing.

Figure 56 shows a sample temperature and tritium quantities in the purge gas for a typical irradiation cycle.

Tritium residence time (t):

I

G

I, tritium inventory (Bq); G, tritium production rate (Bqmin-1).

At steady state:

Tritium release rate (R) = Tritium generation rate (G)

Temperature transients are performed: A T Difference in tritium inventory (area): AI = I2—I1 Difference in residence time: At = AI/G Data set is processed to obtain: t(T)

See Figure 57.

In the EXOTIC-8 program, the tritium release characteristics and mechanical stability of the refer­ence breeder materials for the European HCPB project have been studied, including the effect of long-term neutron irradiation and high lithium burnup.175 The EXOTIC-8 program started in 1997 and ended in 2002. It consisted of ten experi­ments, named EXOTIC-8/1 to EXOTIC-8/10. The irradiations were carried out in the HFR in Petten in peripheral core positions with the typical neutron fluence rate of about 9 x 10 17 m 2s 2 (fast, En > 0.1 MeV) and 5 x 1017m—2s—1 (thermal). The following materials were used: Li2TiO3 peb­bles produced by agglomeration — sintering and extrusion — sintering, provided by CEA; Li2TiO3 pellets produced by cold pressing, provided by CEA; Li2TiO3 pebbles produced by wet processing and sintering, provided by ENEA; Li2ZrO3 pebbles produced by extrusion — sintering process, provided by CEA; and Li4SiO4 pebbles produced by melt spray process, provided by FZK Karlsruhe.

Two types of experiments were targeted:

1. Major focus on tritium release characteristics by determining differential tritium inventories by thermal transients and achieving low to medium lithium burnups of 1-3%.

2. Major focus on high lithium burnup experiments using pebbles with 50% 6Li enrichment and achieving 11% lithium burnup for Li4SiO4 and 17% for Li2TiO3, at relatively constant tempera­tures, and only few data on release, mostly from postirradiation annealing tests.

Tritium release characteristics are measured both in situ by applying the temperature transients and after irradiation in the TPD setup. The tritium release is characterized by the tritium residence time t with the Arrhenius temperature behavior, as depicted in the summary graph; Figure 58.

A correlation has been established between the pebble density and the measured residence time. The obtained results confirm the understanding that open porosity and small grain size are favorable for faster tritium release.

The corresponding activation energies derived from the temperature transients and from the TPD measurements are in fair agreement. For Li2TiO3 pebbles, Q= 82-93 kJmol—1 and for Li4SiO4 peb­bles, Q= 112-123 kJmol—1. The activation energy for Li2ZrO3 pebbles is derived only from the temperature transients: Q= 84kJmol—1. In TPD experiments, it was shown that the tritium release can involve multiple release processes. Moreover, in some in­stances the release can be limited by recombination at the grain surface, which introduces uncertainties in the measured values of the activation energy.

The results of in-pile tritium behavior during nor­mal operation and during transients in temperature and gas chemistry measured in the latest irradiation experiment from the EXOTIC series, EXOTIC-9/1, are reported in Peeters eta/.180 The Li2TiO3 pebbles produced by extrusion-spheronization sintering at CEA were irradiated in the HFR in Petten (thermal 0.5 x 1018m—2s—1) for 301 FPD and achieved a burnup of 3.8—4.1%. The temperature varied between 613 and 853 K. Based upon the in-pile tritium release measurements and the analysis of the tritium resi­dence time, it was concluded that tritium release in the new batch of the high-density Li2TiO3 pebbles (93.0% TD) is rather slow compared with the ceramics irradiated in the EXOTIC-8 irradiation campaign.174 Thus, the tritium residence time measured at 773 K in the EXOTIC-9/1 experiment was ^30 h, whereas

Table 4 Overview of ceramic breeder irradiation experiments performed in fast or mixed spectrum fission reactors to relatively high levels of thermal and/or fast neutron fluence. Values for the experimental parameters quoted were taken from the mentioned references. Readers are also referred to the text throughout the whole chapter. Experiment ID Materials Shape Li-6 (%) Li-6 burnup (%) Total Li burn up (%) Fluence thermal (1020 cm 2) Fluence E > 1 MeV (1022 cm 2) EFPD (Days) Temp (°C) Reference FUBR-IA Li2O, LiAlO2, pellets 56, 95 1.1-1.2 1.4-1.45 96.7, [186] Li2ZrO3, Li4SiO4 2.0-2.2 3.9-4.1 177.6, 3.1-3.5 2.5-2.7 274.3 FUBR-IB Li2O, LiAlO2, pellets 0.07,7.5,56,95 3.9-4.3 4.7-5.1 341.5 450-1225 [186] Li2ZrO3, Li4SiO4 BEATRIX-I Li2O, LiAlO2, pellets 0.07,7.5,56,95 6.7-7.4 8.2-8.9 599.3 [186] Li2ZrO3, Li4SiO4 10.3-11.3 13-14 940.8 10.3-11.3 13-14 936.9 MOTA-2A Li2O, LiAlO2, rings, 0.07, 0.2, 1.7, 7.6-9.9 0.2, 7.4-8.8 299.7 550-640 [186] (BEATRIX-II) Li4SiO4, Li2ZrO3 pebbles 7.5, 440-1000 56, 95 MOTA-2B Li2O, LiAlO2, rings, 0.2 1.27, 5.3-7.9 0.2, 7.7-7.9 203.3 530-640 [186] (BEATRIX-II) Li4SiO4, Li2ZrO3 pebbles 85 440-1100 95 CRITIC-3 Li2TiO3 pebbles 1.85 0.9 0.57 334 200-900 [208] COMPLIMENT LiAlO2, Li2SiO3, pellets, 7.5 0.25 178 (HFR) 400-450 [209] (ELIMA-2, Li4SiO4, Li2O, pebbles 1 75 (OSIRIS) 650-700 DELICE-3) Li2ZrO3 In-Pile Pebble- Li2TiO3 pebbles 0.01 JMTR 400-610 [140] Bed EXOTIC-1 LiAlO2, Li2SiO3, pellet 0.06 0.004 0.01 0.013-0.025 25 350-700 [26],[210] (BEATRIX) Li2O 0.6 0.035 7.5 0.28 EXOTIC-2 LiAlO2, Li2SiO3, pellet 0.55-0.6 0,1 0.02-04 0.025-0.05 50 350-725 [26],[210] Li2O EXOTIC-3 Li2ZrO3, Li2O, pellet, 0.55-0.6 0.12-0.13 0.03-0.06 0.03-0.08 75 385-650 [26] Li2SiO3 solid EXOTIC-4 Li2ZrO3, Li6Zr2O7, pellet, 0.55-0.6, 7.7 0.13-0.15, 0.04-0.07 0.05-0.1 97 310-680 [26] Li8ZrO6, Li2O, solid 0.8 Li2SiO3 EXOTIC-5 LiAlO2, Li2ZrO3, annular 7.5 1.8-2.1 0.07-0.13 0.08-0.17 166 325-630 [26] Li4SiO4, pellets Li6Zr2O7, and Li8ZrO6 pebbles

EXOTIC-6 LiAlO2, Li2ZrO3, annular 7.5 3.0-3.2 0.08-0.15 0.1-0.2 199 315-520 [26] Li6Zr2O7, pellets Li8ZrO6, Li4SiO4 and pebbles EXOTIC-7 Li2ZrO3, Li8ZrO6, annular 38-50 5.8-18.1 0.27 0.13 261 410-745 [33] Li4SiO4, LiAlO2 pellets and pebbles EXOTIC-8/1 Li2TiO3 pebbles 7.5 1.86 0.046 201 310-570 [175] EXOTIC-8/2 Li2TiO3 annular 7.5 1.34 0.039 173 420-640 [175] pellets EXOTIC-8/3 Li4SiO4 + 2% pebbles 50 2.6-4 0.003-0.013 201 200-700 [175] TeO2, Be EXOTIC-8/4 Li4SiO4 + 2% pebbles 7.5 1.38 0.033 201 280-590 [175] TeO2 EXOTIC-8/5 Li2TiO3 pebbles 7.5 2.1 0.06 200 500-660 [175] EXOTIC-8/6 Li2ZrO3 pebbles 7.5 2.5 0.071 299 400-640 [175] EXOTIC-8/7 Li4SiO4, Be pebbles 51.5 8-17.5 0.025-0.064 648 400-690 [175] EXOTIC-8/8 Li2TiO3, Be pebbles 50 4.1-10.9 0.11-0.38 450 400-700 [175] EXOTIC-8/9 Li2TiO3 pebbles 7.5 3.51 0.1 448 420-560 [175] EXOTIC-8/10 Li4SiO4 pebbles 7.5 0.09 448 520-695 [175] PBA Li4SiO4, Li2TiO3 pebbles 7.5 1.5-2.9 (for 0.04-0.07 0.11-0.17 294 550-800 [142] natural 6Li EXOTIC-9 Li2TiO3 pebbles 7.5, 30.5 3.1 (for 0.15 0.09 300 340-580 [143] Li4SiO4 7.5, 20 natural 6Li) K-578 (WWRK) Li2TiO3 pebbles, 96 18-23 220 500-900 [212], [82] pellets HICU Li4SiO4 pebbles 0.06, 7.5, 20 0.7-13 1.5 0.7 403 600-900 [211] Li2TiO3 0.06,11,30

Figure 56 Examples of tritium release from the EXOTIC-8 experiment series with a number of programmed temperature transients (including reactor power scram at day 19+20).

the same characteristic measured on the Li2TiO3 pebbles obtained from CEA and ENEA in the EXOTIC-8 campaign was 1.3 and 5.2 h, respectively (Figure 59).

Changes in the tritium inventory resulting from the variation of the H2 concentration in the purge gas (from 0.1% to 1.0%) appeared to be much smaller
than those resulting from the temperature transients. From this observation, it was concluded that the tritium inventory was determined by the thermally activated processes taking place in the bulk of the material (dissociation from traps, diffusion) rather than by recombination and isotope exchange with hydrogen at the surface.

1000/T (K)

Tritium recovery characteristics of a binary bed containing 0.3 and 2 mm diameter Li2TiO3 pebbles were studied under continuous (20 h) and pulsed (200, 400, and 800 s) neutron irradiation in the JMTR.1 5 From the temperature transients from 573 to 623 K, the tritium residence time was estimated as 3h (63.2% of the steady-state value). The complete recovery of the steady-state conditions was achieved after 20 h. The tritium recovery behavior under the
pulsed operation was almost the same as under contin­uous operation, except for the modulations introduced by the pulse operation, which did not exceed 20% of the total signal variation (Figures 60-62).

Effects of irradiation temperature, purge gas flow rate, and hydrogen content in the purge gas on the tritium release characteristics of the Li2TiO3 pebbles were studied in the in-pile irradiation experiment in the JMTR.185 The Li2TiO3 pebbles were fabricated

by the rotating granulation method. The irradiation continued during one cycle of 25 days with the tritium generation rates of 6 x 101°Bqd~1 or 11 x 101°Bqd~1 depending on the position in the core. The following findings were reported: an increase in the purge gas flow rate accompanied by a temporal increase in the tritium release, which was followed by a swift recovery (<10h at 773 K). A decrease in the flow rate produced

an opposite effect. The hydrogen concentration in the purge gas varied from 10 to 10 000 ppm. It was found that at hydrogen partial pressures <100 Pa, the tritium desorption is controlled by the surface reactions and at higher partial pressures by bulk diffusion.

Chikhray et a/.82 performed irradiation tests of Japanese Li2TiO3 ceramics with 96% enrichment of isotope 6Li in the WWRK reactor. Three types
of ceramic samples were examined simultaneously using a system for in-pile tritium monitoring: one (pebbles) — under constant temperature of 650 °C, and two (pebbles and pellets) — within temperature change ranges from 500 to 900 °C. Lithium burnup reached 23% for the active ampoule pebbles, 20% for the passive ampoule pebbles, and 18% for the pellets. The tritium measurement system permitted the

tritium yield rate to be determined under long-term irradiation of lithium ceramic Li2TiO3. Postirradia­tion testing also included mechanical testing.

Mechanical properties of W strongly depend on vari­ables such as production history, alloying elements, impurity level, thermomechanical treatment, and form of material. Depending on the production his­tory and heat treatment, W and W-alloys could have anisotropic mechanical properties. This is expressed by showing significantly better properties in the direction of elongated grains (by rolling, forging, or due to deposition processes for coatings) but poorer properties in other directions.70 While reported data on single crystals (SCs) (e. g., Gumbsch62) and for isotropic materials (e. g., Kurishita et a/.71) give a clear indication of the material’s performance, typi­cally the reported data refer to the best orientation of the material as shown for fusion relevant tungsten grades in numerous publications.44,51,57,72-81 The prop­erties in other directions, particularly the DBTT, could significantly differ.76 This will affect the operational performance, which is reflected by the orientation — dependent thermal shock response.82

Tungsten is a body-centered cubic (bcc) refrac­tory metal, with a comparatively low fracture toughness,61,83 high DBTT, and poor machinability, which is directly correlated to the material’s low ductility and low grain boundary strength.67 How­ever, DBTT is an ill-defined property and depends strongly on purity, alloying elements, thermo­mechanical treatment, and, most essentially, the testing/loading conditions due to its deformation rate dependence.62,63 The obtained values vary over a broad temperature range from room temper­ature (RT) to several hundreds of degrees Celsius. The exact value depends on the stress state, for example, a three-dimensional state of stress in the sample leads to a lower DBTT.

Although many other parameters influence the fracture of bcc metals, the DBTT is usually asso­ciated with the thermal activation of dislocation kink pairs. Below this characteristic temperature the separation of a screw dislocation into three partial dislocations (which cannot easily recombine and are therefore more or less immobile) is responsible for the brittle behavior. Increasing temperature leads to thermal activation of the kink mechanism and increased ductility due to shielding of the crack tip.84 There is an empirical correlation between tem­perature and activation energy for brittle-to-ductile transitions in single-phase materials suggesting that the ratio between the activation energy and the DBTT gives approximately a value of 25.63

Another factor is the occurrence of interstitial solute elements, such as oxygen, carbon, and nitro­gen, which even in very small amounts tend to segregate at grain boundaries, promoting intergran­ular brittleness and increasing the DBTT. Two ways can be used to get rid of or mitigate the negative effects of interstitial impurities: either a reduction of the grain size,84 to dilute their effect on a larger grain boundary surface, or the complete elimination of grain boundaries, as in SCs. The development of W-alloys essentially follows the first route, as the SC technique, although effective, is too costly. The conventional method to decrease the grain size of tungsten or tungsten alloys is to deform the material at an intermediate temperature, above the DBTT and below the recrystallization tempera — ture.81,84-86 The formation of oxides and carbides of the alloy constituents helps to stabilize the grain boundaries and to dispersion strengthen the matrix at high temperature. Recently, mechanical alloying followed by powder densification has been applied to refractory alloys. Materials with a stabilized fine-grained structure and with the grain boundary strengthened by even finer dispersoids of TiC improve the low-temperature impact toughness of refractory alloys, leading to increased ductility even down to RT and create superplasticity at high

71,87-89

temperatures.

Another reliable method to increase the ductility at low temperatures and therefore reduce the DBTT is to alloy tungsten with the rather expensive element rhenium, which is a substitutional solute in the W lattice.67,83

As mentioned before, both material deformation and heat treatment influence the DBTT. A heat treatment slightly below the recrystallization tem­perature is able to significantly reduce the DBTT. In contrast, annealing above the recrystallization temperature reduces strength and hardness and increases the DBTT.67

4.17.3.2.2 Component fabrication:

CTE mismatch with heat sink

A mismatch between the coefficients of thermal expansion (CTEs) can lead to thermal stresses at the interface, which are detrimental to the compo­nent lifetime. This can occur with either Cu-based alloys or steels (steel is more likely to be used in case of coatings) such as that used for water-cooled designs, or to W and W-alloys in the He-cooled design. In particular W and W alloys, in the cold — worked and stress-relieved condition, tend to delam­inate in the direction parallel to the direction of deformation. Such delamination can occur during machining or during operation. To avoid failure due to delamination, the orientation of the texture has to be perpendicular to the surface of the joints,90 raising the question of the suitability of plasma-sprayed W coatings. Two possible options are recommended to mitigate the thermal stresses, that is, reducing the joint interface by introducing castellations or using smaller tiles,9 — 3 or introducing soft and chemically

stable interlayers94,95 or graded layers.96-101

Despite the fact that surface finish has no direct effect on the performance of ITER-related compo — nents,94 it is recommended to avoid possible crack initiators in the armor design, such as castellations ending in the tile and to ensure accurate surface finishing.102-104 Designs that have been proven to reduce the tile and interface thermal stresses and to extend the component lifetime beyond the design limits are the macrobrush or the monoblock. The latter is the reference design for ITER105 because it provides the most reliable attachment and therefore a reduced risk of catastrophic cascade failure.106

Finally, the thermal treatment of W during joining manufacturing cycles might have an influence on the material’s properties. While the process tempera­tures during joining of W and Cu do not lead to any significant change of the W properties, in the case of high-temperature brazing of W to W alloys for the He-cooled divertor design,1 2 the recrystallization temperature of W has to be taken into account.

CFCs for fusion applications are specifically designed to maximize thermal conductivity and for this reason irradiation-induced thermal conductivity degradation
is of primary importance. As with ceramics, graphite thermal conductivity is dominated by phonon trans­port and is therefore greatly affected by lattice defects, such as those caused by neutron irradiation. The extent of the thermal conductivity reduction is therefore directly related to the efficiency of creating and annealing lattice defects, and is therefore related to the irradiation temperature.

The effect of neutron irradiation on the thermal conductivity of graphite has been widely studied. The majority of the literature10, , — 1 in this area has been in support of the gas-cooled, graphite-moderated, fission reactor program in the United States and United Kingdom and has focused on ‘nuclear’ gra­phites as well as more fundamental work on pyrolytic graphite.8,27,32,33 In recent years, the emphasis of graphite radiation effects research has switched to its use in PFCs of graphite fusion reactors.10,11,34

Because of the significant advances in carbon — carbon composite (CFC) processing and fiber develop­ment, very high thermal conductivity materials have been recently demonstrated and they have become attractive for high heat flux applications. The highest thermal conductivity CFCs are made from highly crys­talline graphite fibers having intrinsic conductivities approaching those of pyrolytic graphite. For example, vapor-grown carbon fibers35 have a thermal conductiv­ity of 1950 W m-1 K- . Along with advances in fiber properties, improvements have occurred in both mono­lithic graphite and the CFC matrix-processing areas, which also have enhanced thermal conductivities.

The physical processes governing the thermal conductivity of graphites, as well as the mechanisms responsible for the radiation-induced degradation in conductivity, are well established.8 For all but the poorest grades of carbon, thermal conductivity is dominated by phonon transport along the graphite basal planes and is reduced by scattering ‘obstacles’ such as grain boundaries and lattice defects. For graphites with the largest crystallites, that is, pyro­lytic graphite or natural flake, the in-plane room temperature thermal conductivity is approximately 2000 W m-1 K-1.36

The thermal conductivity of graphite-based mate­rials can be written as a summation of the thermal resistance due to scattering obstacles:

K (x) = b(x)

where b(x) is a coefficient that includes terms due to orientation (with respect to the basal plane), porosity, and some other minor contributors. This coefficient
is assumed in most cases to be constant with tem­perature, with a value of around 0.6. The first two terms inside the parentheses are the contributions to the thermal conductivity due to umklapp scattering (Ku) and grain boundary scattering (Kgb). The grain boundary phonon scattering dominates the thermal resistance (1/Kgb) at low temperatures and is insig­nificant above a few hundred degrees Celsius, depending on the perfection of the graphite. The umklapp scattering, which defines the phonon — phonon scattering effect on the thermal conductivity, dominates at higher temperatures and scales nearly as T2.8 The umklapp scattering therefore defines the upper limit to the thermal conductivity for a ‘perfect’ graphite. Following Taylor’s analysis,37 the umklapp — limited thermal conductivity of the graphite crystal would be ^2200Wm~1K~1 at room temperature, in close agreement with the best pyrolytic graphites or the vapor grown carbon fibers mentioned earlier.

The third term in eqn [7], Kx, is the contribution to the basal plane thermal resistance due to defect scattering. Neutron irradiation causes various types of defects to be produced depending on the irradia­tion temperature. These defects are very effective in scattering phonons, even at flux levels that would be considered modest for most nuclear applications, and would quickly dominate the other terms in eqn [7]. Several types of irradiation-induced defects have been identified in graphite. For irradiation tem­peratures lower than 650 °C, simple point defects in the form of vacancies or interstitials, along with small interstitial clusters, are the predominant defects. Moreover, at an irradiation temperature near 150 °C,27 the defect that dominates thermal resistance is the lattice vacancy.

Due to its sensitivity to the presence of defects, the temperature at which graphite is irradiated has a profound influence on the thermal conductivity degradation. As an example, Figure 15 shows one of the most complete sets of irradiation data on Pile Grade A nuclear graphite.38 This graphite is a medium­grained, extruded, anisotropic material with a room temperature thermal conductivity of 172 W m—1 K— in the extrusion direction. Figure 15 presents the normalized room temperature thermal conductivity of this graphite of various irradiation temperatures. It is seen that as the irradiation temperature is decreased, the degradation in thermal conductivity becomes more pronounced. For example, following irradiation at 150 °C, the thermal conductivity of this graphite appears to approach an asymptotic thermal conductivity of ~1% of the original. The reason for

this is that as the irradiation temperature is decreased, the fraction of vacancies surviving a cascade event increases, and thus the number of vacancies available to scatter phonons is much higher for the lower temperature irradiation.

Data have been published for CFCs whose thermal conductivities are similar to those of nuclear graphites, showing degradation similar to that expected from the graphite literature. For example, Burchell34 has shown that the saturation thermal conductivity for a 3-directional composite (FMI-222, Kunirr = 200 W m—1 K—1 at RT) is reduced to ^40% of the origi­nal room temperature conductivity following fast neutron irradiation at 600 °C. Published data for the degradation of thermal conductivity in highly con­ductive CFCs have led to the conclusion that a higher initial conductivity composite results in higher abso­lute conductivity after irradiation.39,40 Figure 16 demonstrates this point. At the extremely damaging irradiation temperature of 150 °C, it is observed that the absolute reduction (lunirr — Kirr) is substan­tially greater for the high thermal conductivity mate­rials compared to the lower conductivity CFCs and graphite, as seen in Figure 16, although the normal­ized fraction (Kirr/Kunirr) is approximately the same for all the carbon materials in the figure. Moreover, saturation in thermal conductivity degradation occurs at a neutron dose of ~1 dpa. Data for higher irradia­tion temperatures11, 1 show that the higher thermal conductivity materials have a slightly larger fractional
change in thermal conductivity (Kirr/Kmirr) compared to lower conductivity materials, although the absolute value of the irradiated thermal conductivity is still greater for the higher conductivity materials. A com­parison of thermal conductivity degradation for a nuclear graphite (CH-45) with the composites FMI-222 and MFC-1 is given in Figure 17.31

For the low-dose regime relevant to machines such as ITER (less than 1 dpa or about 500 h in this figure), the conductivity is seen to decrease by a factor of two for the highest conductivity material (MFC-1) and by about 30% for the nuclear graphite.

An algorithm has been developed to predict the thermal conductivity degradation in a high thermal conductivity composite (^555Wm~1K~1 at room temperature) as a function of radiation dose and temperature.41 The absence of irradiation data on CFCs of this type required the use of data from intermediate thermal conductivity materials and pyrolytic graphite to derive an empirical radiation damage term.24,2728,39,42

An analysis of the effects of temperature and neu­tron dose on the thermal conductivity is shown in Figure 18. Specifically, the algorithm assumed the nonirradiated properties of the unidirectional fiber composite MFC-1 material compiled with an empir­ical radiation damage term. As with the experimental data of Figures 15 and 16, it is seen in Figure 18 that an enormous loss in thermal conductivity occurs at low irradiation temperatures. Presently, only a few data points exist that are relevant to the validation of this algorithm, and these are also plotted on the figure.39 The data do agree within the errors of irradi­ation temperature and thermal conductivity measure­ment, with the algorithm predictions. However, they

Time in reactor (h) Time in reactor (h)

1-3dpa irradiation

Figure 19 Effect of neutron irradiation on thermal conductivity-driven temperature evolution in a monoblock and flat-plate divertor design.

are insufficient to validate the algorithm and the need clearly exists for additional data for this purpose.

To illustrate the usefulness of such an algorithm, and the significance of the issue of thermal

conductivity degradation to the design and operation of PFCs, this algorithm has been used to construct Figure 19, which shows the isotherms for a mono­block divertor element in the nonirradiated and

irradiated state and the ‘flat plate’ divertor element in the irradiated state. In constructing Figure 19, the thermal conductivity saturation level of the 1 dpa given in Figure 18 is assumed, and the flat plate and monoblock divertor shown are receiving a steady state flux of 15 MW m~ . Both composite materials have been assumed to be in perfect contact with a copper coolant tube or plate. Figure 19 clearly shows two points. First, a very high conductivity composite is required to handle the extreme heat fluxes expected if the temperature is to be limited to <1200 °C (Section 4.18.4). Second, the effect of neutron irra­diation on the temperature is significant. In the case of the flat plate divertor, the temperature rise (AT) changes from ^200 to ^500 °C following irradiation, while for the monoblock, it increases from ~-350 to ^900 °C. It should be noted that the larger tempera­ture increase for the monoblock design is due not to the larger path length of graphite in that configura­tion, but rather to the larger amount of graphite material that is irradiated in the highly damaging low temperature regime (see Figures 15 or 18). The larger temperature increase for the monoblock design could be unacceptable from an erosion stand­point, as will be discussed in Section 4.18.4.

Because of the serious thermal conductivity deg­radation in graphite, scenarios to limit the issue (such as baking the PFM) have been considered. Upon annealing above the irradiation temperature, some interstitial atoms become mobile and can recombine with vacancies, restoring the thermal conductivity of the lattice. It is therefore conceivable that intermit­tent annealing of the PFC could regain some of the irradiation-induced thermal conductivity degradation. Bake-outs are typically conducted between operating cycles of a fusion system for plasma impurity (usually oxygen) control. However, the wall-conditioning temperatures are typically limited to less than 300 °C and for various reasons cannot be significantly increased. Inspection of data such as those given in Figure 2041 indicates that little recovery in thermal conductivity is possible unless bake-out temperatures approach 1000 ° C, and thus in situ annealing can be of only marginal benefit.

Thermal fatigue durability is a necessary but not sufficient prerequisite for the Be/CuCrZr alloy joint. Neutron irradiation is also expected to affect the high heat flux durability of these joints. This is particularly true for the first-wall components, where the joint will experience rather high neutron fluence.

To investigate the effects of neutron irradiation on the joint properties there are two possible ways. The first method is postirradiation testing, that is, irradia­tion of small scale mock-ups in fission reactors and postirradiation testing in heat flux test facilities. The second method is to try to reproduce more faithfully the situation in ITER and to apply a cyclic heat flux during irradiation. This consists of in-pile heat flux testing of small mock-ups in fission reactors. Given the technical difficulty of achieving simulta­neously representative values of heat flux and neu­tron irradiation and to determine exactly what the heat fluxes are within a reactor, it is suggested that the effect of neutron irradiation on the mock-ups be determined by pre — and postirradiation heat flux tests on mock-ups. This method is probably conser­vative as the postirradiation tests are performed on materials and joints having accumulated damage corresponding to the total neutron irradiation dose. However, this should be supported by analysis and material testing.

Most of the tests conducted in the past were done for DS-Cu as a heat sink. Studies of neutron irradia­tion effects on the durability ofthe Be/Cu-alloy joints have been performed in at least two of the ITER parties: Europe,150,162 and the Russian Federation.173

For the Russian experiment, the irradiation con­ditions were 0.3 dpa at 350 °C. The irradiated Be/Cu mock-ups were then tested in the JUDITH facility. The results of the postirradiation high heat flux test­ing of the different Be/Cu mock-ups are presented elsewhere.174 On the basis of the results of these tests, it was concluded that the effect of the neutron irradi­ation is not critical for the joints. Metallographic inspections did not show any significant changes in the braze joint after neutron irradiation. For the CuMnSnCe, a small intermetallic phase was observed in the middle of the braze layer for unirra­diated and irradiated samples. No crack formation was found in the intermetallic. In addition to thermal fatigue tests, shear tests were conducted and it was found that for CuMnSnCe the shear strength decreased after neutron irradiation from 200 to 155 MPa, whereas for the InCuSil braze no irradia­tion influence was observed (^300 MPa). Neverthe­less, the thermal performance of the joints during high heat flux tests was very similar to the perfor­mance of unirradiated mock-ups.

For the European experiments, the first irradia­tion campaign (named Paride 1 and Paride 2) of small scale Be/Cu mock-ups and Be/Cu joints took place in 1996-1999. Mock-ups were fabricated from a sin­gle 10-mm thick Be tile (grade S-65C) of dimensions 22 mm x 60 mm, HIPped on a 20-mm thick CuAl25 substrate (grade IG1) with a drilled 10-mm diameter cooling channel (Figure 14). HIPping was done at 830 °C for 2 h with the use of a 50 pm Ti interlayer. One mock-up was neutron irradiated in the test reactor high flux reactor (HFR) and then high-heat flux tested in JUDITH. The neutron irradiation was done at about 200 °C up to a neutron dose of about 0.6 dpa in the Be material. The neutron dose expected at the end of life in the Be armor of the first-wall panels is about 1 dpa (Be). The irradiated mock-up was tested for 1000 cycles at 1.6 MW m-2 plus 100 cycles at 1.9 MW m-2 plus 100 cycles at 2.4MW m-2 plus 1000 cycles at 2.8 MWm-2 plus 100 cycles at 3.3 MW m-2 without any sign of failure. It failed dur­ing the first cycle at 4.25 MW m-2 with a partial detachment of the Be tiles on one end (Figure 14). An unirradiated mock-up was tested for 1000 cycles at

1.5 MW m-2 plus 1000 additional cycles at 3 MW m-2 without any sign of failure but a detachment of the Be tile occurred during the first cycle at 4.5 MW m-2. It was therefore concluded that no significant degrada­tion of the Be/CuAl25 joint was observed up to a neutron dose of about 0.6 dpa. Neutron irradiation test experiments are ongoing or in preparation with Be coated first-wall mock-ups made from CuCrZr alloy to confirm the above result with this Cu alloy.

The first experiment was a joint European/Rus — sian irradiation test campaign. It was prepared by the Efremov Institute of St. Petersburg. The original objective was to perform thermal fatigue testing of two first-wall mock-ups at about 0.5 MW m~2 simul­taneously to neutron irradiation. A failure of the surface heating system made from graphite after about 5000 cycles resulted in the discontinuation of the thermal fatigue testing and a continuation of the campaign with only neutron irradiation. The irradia­tion campaign has been stopped with the achievement of a neutron dose of 0.75 dpa. The first wall mock-ups, one made with Be tiles HIPped at 580 °C and a Cu interlayer (Figure 15) and another with brazed Be tiles with STEMET 1108 braze alloy, will then be high heat flux tested together with unirradiated reference first wall mock-ups. Two other test cam­paigns are in preparation at the NRI of Rez (Czech Republic) and at Petten (The Netherlands) with the objective of thermal fatigue testing three first wall mock-ups in parallel to neutron irradiation.

The question as to whether the correlation between fusion and fission neutron spectra assumed in many of the above measurements is valid or not needs to be discussed. Comparison of changes in the mechanical properties, especially at low temperature, needs to be made with the same He to dpa ratio to ensure that the results will be valid for ITER.

Be/Cu alloy mock-ups have been tested in an electron beam for 1.5 s under a deposited energy density of 60 MJ m~2(132) to simulate Be damage dur­ing a VDE. For a 6 mm thick Be tile, the melt layer was ~1.5mm, while that calculated for the same

condition is 1.25 mm. A cross section of the Be tile after the simulated VDE event is shown in Figure 16. The embrittlement of the Be due to neutron irradi­ation increases the loss of material particles, espe­cially at low irradiation temperature. A clear pore formation (which is expected to be He filled) has been observed in the melt layer of all neutron irra­diated specimens after thermal shock loading (see Figure 16). An area of possible concern is the small surface cracks that form when molten metals reso­lidify. These resolidification cracks could serve as thermal fatigue crack initiation sites and accelerate this type of damage. While this effect has not been extensively studied because of the difficulty of simulating disruptions in the laboratory, it may not be a critical issue as thermal fatigue cracks form after a few hundred cycles in most materials and they grow only to depths where the thermal stress level is above the yield stress.175

High heat flux tests of neutron irradiated mock — ups conducted in the past did not reveal any damage in Be and in Be/Cu joints,176 although the irradiation conditions were not fully ITER relevant (damage dose of ^0.3 dpa instead 1 dpa for the end of life and also lower He per dpa). An increase in crack formation and erosion rate has been observed in the surface of irradiated Be at 350 and 700 °C.177 The S-65C grade presented the lowest damage after

irradiation. High heat flux tests (at more severe con­ditions than needed for the first wall) of the cracked unirradiated Be did not reveal any detrimental behav­ior and loss of material due to cracking.178 From an engineering point of view, to avoid possible crack formation and delamination of the brittle Be, it is recommended to use Be tiles without any stress concentrations.

This technology aims at using a corrosion product that forms at the interface between Li and structural materials as an insulating layer. By careful control of the corrosion reaction, the insulation layer can be formed uniformly on the inner surfaces of complex components. The corrosion layer may also be formed on the cracked area of the coating, thereby repairing insulation defects. The in situ coating with CaO and AlN have been studied in the United States11 and the Russian Federation,12 respectively.

Figure 3 Schematic illustration of the mass transport for in situ oxide coating in Li. Ov: oxygen in vanadium substrate; MLi: metal doped in Li for producing oxide coating.

A CaO insulator layer forms during the immersion of vanadium alloys in Ca-doped Li. For enhancing this reaction and stabilizing the layer, the O level in the alloy was increased by prior doping. The pro­cess is schematically shown in Figure 3. Careful control of the Ca level in Li and the O charging condition of the vanadium alloys made in situ for­mation and healing of the CaO coating possible at <500 °C. Because of the successful demonstration, in situ CaO coating was adopted as the means to mitigate the MHD pressure drop in the ARIES-RS power plant study.1

This technology was, however, shown to be inap­plicable at temperatures exceeding 600 ° C because of the unacceptably short lifetime of the coating induced by increased dissolution.2

When ceramic breeder research was initiated in the 1970s, relevant data for lithium-based ceramics were scarce or nonexistent. Initial screening of candidates was mainly based on examination of the physical and chemical characteristics and neutronic behavior. An extensive R&D effort focused on determining the properties of unirradiated materials and on designing irradiation experiments to understand and quantify the effect ofneutron irradiation on material properties and on recovery ofgenerated tritium. With the publication of these data, the relative merits of the candidates became known and interest changed accordingly. With the evolution of the INternational TOkamak Reactor (INTOR) and International Thermonuclear Experimental Reactor (ITER) projects, much of the international R&D effort was focused on the opportu­nities for implementing breeding units or driver blan­kets in such a device and its role in the technological development ofpower reactor blanket systems.

Several breeder blanket design options had been developed such as high-temperature water-cooled and helium-cooled concepts for DEMO and power reactors, and low-temperature water-cooled concepts for ITER.2-21 The ceramics under consideration exhibit different characteristics, which can make one ceramic more adaptable to a specific blanket concept. In this chapter, the issues being addressed in R&D in support ofcurrent blanket design studies are high­lighted. In this chapter no reference is made to R&D for other fusion systems like inertial confinement.

Two major types ofceramic breeder material con­figurations have been developed based on pressed and sintered pins or pellets or as a collection of packed spheres or pebble beds. The actual arrange­ment of pebbles though may be in tube, which may lead to some confusion when ‘breeder-in-tube’ (BIT) is mentioned. The paper by Ihli et a/.21 provides an overview of blanket design developments and refer­ences ongoing work by various parties.

Until 2010, pebble-bed concepts were the preferred options for all parties involved in ceramic breeder test blanket module (TBM) programs for ITER as reported by Giancarli and coworkers.23,24 Typically, inert gas with hydrogen addition, whose characteristics are dis­cussed later in this chapter, is used for the extraction of the tritium produced from lithium ceramics.25-2