Ceramic breeder materials have been tested under various conditions to determine their irradiation

response in terms of tritium production and release and their microstructural, thermal, chemical, and mechanical stability. The nuclear loading of ceramic breeder affects their performance in various ways, of which the most important are the following:

evaluation of the probability distribution of the maxi­mum contact forces for different loading conditions (see Figure 34) 205 A comparison of results obtained by DEM and tomography is shown in Figure 35133: the packing factor distributions agree quite well, and the radial and vertical positions of particles show the same structure as shown in Figure 17(b). Other approaches are found like those of Aquaro and co-

123,199,202,203

workers.

• Lithium burnup and other neutron transmutation reactions gradually change the material composi­tion and affect the chemical and physicochemical interactions.

• Transmutation reactions of major and minor con­stituents gradually change the radioactivity levels relevant for in-tokamak operation, hot-cell opera­tions, and management of waste, including recy­cling options.

•
Atomic displacements are induced by neutron impact, most effectively by fast neutrons, causing significant lattice defects, damage, swelling, and so on.

• Generated tritium affects the physicochemical and mechanical behavior, and the inert, 3He decay product stays within the material if it is not desorbed earlier into the purge gas flow.

• Generated helium resides in the material, and through formation of clusters of bubbles will generate stresses leading to macroscopic effects such as swelling, altered thermal transport, and/or fracture.

Numerous irradiation experiments have been per­formed on ceramic breeder pebbles using material test reactors with thermal, mixed, or irradiation phe­nomena fast neutron spectrum. As the 6Li cross­section, in particular, is much higher for thermal neutrons, many of the irradiation phenomena can be effectively studied in thermal or mixed-spectrum reac­tors, that is, without using 14MeV neutrons. This
section concentrates on the thermal-mechanical behavior, while tritium production and transport are dealt with in the following section.

High fluence and high lithium burnup were achieved in fast reactor irradiations at experimental breeder reactor II (EBR-II) and fast flux test facility (FFTF) facilities, as reported by Hollenberg and cow­orkers.28,134 The bulk of these experiments concerned Li2O, LiAlO2, and LiZrO2, with only a few data on Li4SiO4 shaped as annular pellets and LiZrO3 peb­bles. Good tritium release behavior of Li2O and LiZrO3 has been reported, even for temperatures higher than 1000 °C; pellet thermal conductivity of Li2O and LiAlO2 was decreased at lower irradiation temperatures but appeared fairly unaffected when operated over 400 °C. , , , , ,

AECL tests at NRU and JAEA tests at JMTR addressed the impact of neutron irradiation on pebble-bed properties, such as conductivity. In these cases, the constraints were modest: either higher
burnup with few pebbles in the heat flow direction or low volumetric heat loads and low lithium burnup.135-139 Verrall et a/.138 seem to the first to report on bubble formation in a ceramic breeder. They observed this in Li2O irradiated in NRU up to 1% lithium burn-up up, see Figure 45.

The effective thermal diffusivity of a Li2TiO3 pebble bed was studied in an in-pile irradiation experiment by Kawamura et a/.140 at the JMTR test reactor. The cylindrical assembly of the Li2TiO3 pebble-bed was instrumented with a number thermo­couples to determine the radial temperature profile. The derived thermal diffusivity as function of temperature is shown in Figure 36. The dimension of the pebble bed was 20 mm in diameter and 260 mm in length. The effective thermal diffusivity of the Li2TiO3 pebble bed was found to decrease with increasing irradiation temperature. This tendency remained up to a thermal neutron fluence of 1 x 1024nm~2, while the thermal conductivity at given temperatures also remained constant.

None of these experiments addressed the pebble — bed deformation behavior at the strong temperature gradients envisaged for breeding blankets. Out — of-pile testing is less representative as it requires the use of heater plates, with their specific impact on reduced bed thickness, and running a ceramic — heater interface at the highest temperature. Also, when material is irradiated in a stressed state, there is an additional phenomenon of irradiation — induced creep.

As a major step in the preparation of the European HCPB TBM program in ITER, an in-pile test of pebble-bed assemblies was defined. This experiment was designed to address the neutron-irradiation effects on the thermal-mechanical behavior of a breeder pebble bed at HCPB DEMO representative levels of temperature and defined thermal-mechanical loads.121,130,141 A schematic is given in Figure 37. The core of each test element is a horizontal cylindri­cal bed of ceramic breeder pebbles, either Li4SiO4 (OSi) or Li2TiO3 (MTi), with an outer diameter of about 45 mm and bed thickness of about 10 mm, sand­wiched between two beryllium pebble beds. The breeder and beryllium pebble beds are separated by Eurofer-97 steel plates. The heat flow is managed so as to have a radial temperature distribution in the ceramic breeder pebble bed as flat as reasonably possible. The test element design and test matrix required extensive pretesting, improved pebble-bed modeling, design curves for the material character­istics, and performance analyses allowing in-reactor operation in High Flux Reactor (HFR) Petten121,130,141 (Figure 38).

A specific pretest compaction procedure by press­ing and heating has been developed, in particular to increase the thermal conductivity of the beryllium pebble beds and provide conditions that would result in limited changes during in-pile operation.130,141 The compaction procedure consisted of a subsequent loading of the pressure plate of the total assembly to 3 MPa. The X-ray pictures combined with the

□

0 2 4 6 8 10 Thermal neutron fluence (xl023nm-2)

12

Figure 36 The effective pebble-bed conductivity derived from in-pile І_і2ТІОз pebble-bed irradiation in JMTR, as performed by Kawamura and coworkers.140 Reproduced from Kawamura, H.; Kikukawa, A.; Tsuchiya, K.; et al. Fusion Eng. Des.

2003, 69, 263-267.

Sealing plate: all Eurofer-97

Figure 37 Schematic of HCPB pebble-bed assembly (PBA) test-element for in-pile testing in the High Flux Reactor at Petten, The Netherlands. Each test-element has a cylindrical ceramic breeder section, either Li4SiO4 or Li2TiO3, in between two cylindrical shaped beryllium pebble beds and separated by Eurofer plates.

dimensional measurements taken during assembly allowed the determination of actual bed size and compaction values. An example of a postassembly X-ray picture is shown in Figure 39.

The design and safety requirements for in-pile operation required the development of a full-coupled thermomechanical model in the MARC finite-element code. In this way, the pressure buildup and stress relax­ation in the pebble beds could be simulated in detail to guide the required reactor startup profile.121,130,141

The two test elements with Li4SiO4 pebbles were irradiated at nominal temperatures of 600 and 800 °C in the breeder bed, to see any effects ofthermal creep. The other two test elements contained Li2TiO3 peb­bles with different grain sizes and were irradiated at the same temperature, the nominal temperature of 800 °C. The pebble beds were typically purged with a helium-hydrogen mixture of reference composition (0.1% H2). The gas purge entered at the lower beryl­lium bed and exited at the upper beryllium beds.

The PBA has been irradiated for 294 FPD (full power days), achieving lithium burnups of 1.5—2.2% for Li4SiO4 and 2.8—2.9% for Li2TiO3, without enriching in 6Li. The dose in the Eurofer-97 parts

ranged from 2 to 3 dpa.142

Extensive analyses of the in-pile data using FEM calculations showed that the maximum

temperatures in the Li4SiO4 pebble beds of the top and bottom test-elements are about 600 and 800 °C, respectively. The average temperatures in the Li4SiO4 beds are about 550 and 740 °C, respectively.1 2

In postirradiation examinations of both Li4SiO4 samples, a little sintering and a significant amount of cracking or fragmentation have been observed. No significant difference between the lower and higher temperature case was found. Most of the evidence of cracking and fragmentation in the Li4SiO4 pebbles is observed toward the middle of the bed (highest temperature, highest deformation). This is visible in scanning electron microscopy images (Figure 42). There is some evidence of grain growth. Reactions of Li4SiO4 pebbles with Eurofer were found to be very small.

The maximum temperatures in the Li2TiO3 pebble beds of both the two test-elements in the middle are about 780 °C. The average temperatures in the Li2TiO3 beds are about 690 and 720 °C, respectively.142

In both test-elements, Li2TiO3 pebbles showed a significant amount of sintering and necking, which was found most significant in the test-element #3. The average temperature of this test-element was higher by about 35 °C. Almost no fracture or frag­mentation was seen. There appeared to be a small reaction layer, distributed uniformly along the Euro — fer (Figures 43-44).

In the high burnup irradiation EXOTIC-7 (see details in Section 4.15.5.1.2), the pellet stacks and pebble beds were found to be essentially intact by neutron radiography analyses after irradiation, and except for one capsule containing Li2ZrO3 pellets, three out of five were found intact after unloading. Fragmentation of the 0.1—0.2 mm Li4SiO4
pebbles was also observed but was very difficult to quantify33’143-145 (Figure 46). The pores observed in the images are related to the pebble fabrication method rather than to neutron irradiation.

Chikhray etal. irradiated Li2TiO3 + 5 mol% TiO2 in the Kazakhstan water water research reactor (WWRK) reactor to lithium burnups of about 20% at 760 and 920 K; see details in the next section and Chikhray et a/.82 Pebble crush tests showed reduc­tion of strength’ whereas microhardness tests also revealed ingrowth of soft phases. X-ray diffraction measurements showed traces of LiTi2O4’ LiTiO2’ and Li4Ti5O12; see Chikhray eta/.82 for more details.

In an IEA-framed international collaboration’ European Li4SiO4 and Li2TiO3 and Japanese Li2TiO3’ reference pebble materials were tested in a high fluence irradiation project at the HFR in Petten, named high neutron fluence irradiation of pebble stacks for fusion (HICU).146-148’194 The neutronic analyses as reported by Fischer and coworkers149’150 demonstrated that relevant nuclear irradiation para­meters such as the displacement damage accumula­tion’ the lithium burnup’ and the damage production function W(T) are met with the selected neutron shielding and 6Li enrichments chosen. This project is conceived to irradiate ceramic breeder pebble stacks at high temperatures under blanket prototypical ratios of fast neutron damage (dpa) and lithium burnup.

Compared with the PBA experiment, the pebble stacks are smaller’ but capsule dimensions are up to about 10 mm; X-ray tomography was used for detailed mapping of pebble location prior to irradiation.148’151

4.17.3.1 Fabrication and Microstructure

Tungsten and tungsten alloys are commercially available in many forms, for example, as bulk rods, plates and discs, or thin coatings on various kinds of substrates. For each of these tungsten products, opti­mized production routes exist involving mainly pow­der metallurgical techniques for bulk materials and PVD and chemical vapor deposition (CVD) as well as plasma spraying (PS) for coatings. Each of these processes has its own advantages and disadvantages as well as an individual influence on the material’s microstructure and subsequently the material prop­erties. In addition to the fabrication method, the raw materials, the alloying elements and dopants/impu — rities, pre — and postthermomechanical treatments, and the final shape/geometry have a strong impact on the achieved microstructure.

Focusing on the powder metallurgy fabrication route, tungsten powder is obtained from ammonium paratungstate ((NH4)2WO4), tungsten oxide (WO3), and tungsten blue oxide (WO3_x) by hydrogen reduc­tion at temperatures in the range of 700-1100 °C. Vari­ous grain sizes can be produced depending on the reduction temperature and the hydrogen dew-point. The purity of the metal powder obtained is >99.97%. In the manufacture of doped or alloyed tungsten pro­ducts, the dopants or alloying elements are either introduced into the raw materials before reduction or they can be added to the metal powder after reduction.

Following the reduction stage, the powder is sieved and homogenized. The initial densification of the powder in various plate and rod geometries takes place predominantly through die pressing and cold isostatic pressing. The pressed compacts are subsequently sintered at temperatures between 2000 and 2500 °C (2273-2773 K), mostly using furnaces with hydrogen flow. This increases the density and the strength of the pressed blanks.40

After sintering, the products have a rather low density of about 80% of the theoretical value and poor mechanical properties. To increase density and improve mechanical properties, the sintered products are subject to a mechanical treatment such as rolling, forging, or swaging at temperatures up to 1600 °C. Intermediate annealing, leading to recovery and recrystallization, is necessary to maintain sufficient workability. The working temperature can be reduced as the degree of deformation increases. In this way, forged parts such as rods and discs as well as sheets and foils are produced.40

The final step, that is, the mechanical treatment, changes the microstructure from isotropic with grain sizes determined by the initially used powder size into anisotropic. Depending on the deformation method, the grains may show either:

• an elongated, needle-like structure along the deformation direction for radially forged rods and uniaxially rolled plates (see Figure 1(a)), or

• a flat disc-shaped structure for axially forged discs

or blanks and cross-rolled plates (see Figure 1(b)).

In addition to bulk materials, research and develop­ment is also directed on tungsten coatings. One possi­bility would be the plasma-spraying process, in which powders are injected into a plasma flame, melted, and accelerated toward the (heated) substrate. The depos­ited layers are splat-cooled, leading to a flat disc­shaped microstructure. Depending on the atmospheric conditions, the result may be layers with high porosity and oxygen content (water stabilized and atmospheric plasma spraying, APS, see Figure 2(a))41,42 or low porosity and good thermal contact (low-pressure or vacuum plasma spraying, LPPS/VPS).26,43-47

In contrast, PVD and CVD coatings show a columnar structure perpendicular to the coated sub­strate with grain sizes in the range of the coating thickness (see Figure 2(b)). PVD coatings, which are also used as thin intermediate layers below a plasma-sprayed tungsten top layer,48 are deposits of tungsten vapor on the substrate surface, which is in the source’s line of sight.43,49 CVD coatings are reactions of a W-containing gaseous phase and have the ability to coat complex geometries.6,50-52 In both cases, a high density (~100%) of the coatings is achieved.

The coated substrate can be graphite as used for AUG (PS),12,13 CFC as used for the ITER-like wall

(a)

project in JET (PVD),21,53 55 or copper and steel as it might be used for first wall applications in future fusion devices (PS, PVD, CVD).44,49,50,56-60

As mentioned earlier, the first wall materials in next — generation machines will receive many tens of dpa. At low doses (<0.01 dpa), there are essentially no mechanical property changes expected in graphite materials (see Chapter 4.10, Radiation Effects in Graphite). However, even at these low doses, thermal conductivity and stored energy are of concern, spe­cially for low irradiation temperatures (<400 °C). For displacement levels >0.01 dpa, significant property changes occur, including strength, elastic modulus, specific heat, coefficient of thermal expansion (CTE), Poisson’s ratio (v), and thermal conductivity. In addi­tion, the dimensional stability under irradiation is important because the induced stresses may be sig­nificant and may need very tight tolerances at the plasma edge. It has been shown in fission neutron experiments that specific heat Cp6 and v7 are not greatly affected by irradiation. Moreover, only mod­erate changes in the CTE occur, but the magnitude and nature of the CTE change is highly dependent on the type of graphite.6,8-10

The irradiation-induced property changes for graphite and composites that have received the most study by the fusion community deal with dimensional stability, strength, elastic modulus, thermal conduc­tivity, and hydrogen retention. A large body of data exists on the thermophysical changes in graphites, coming mainly from graphite-moderated nuclear reac­tor development programs. A smaller body of research
exists on CFCs, mainly from the same source, but with some additional data from fusion research. These data suggest that CFCs have very similar irradiation behav­ior compared to graphite. In Chapter 4.10, Radiation Effects in Graphite, Burchell discusses radiation dam­age mechanisms in graphite, and some of the specific property changes that occur in fission reactor appli­cations. Because they are of special significance to fusion energy, the radiation effects in CFCs in gen­eral and the radiation-induced degradation in thermal conductivity in graphite and CFCs in particular will be focused on in the remainder of this section. How­ever, it is first important to contrast nuclear graphite (essentially a form of purified structural graphite) with that of graphite composites. For the purposes of dis­cussing graphite materials for fusion applications, the term composites is applied specifically to continuous fiber composites, typically woven, and infiltrated with pitch or some other resin that is graphitized to form a highly crystalline graphite matrix. The fibers compris­ing these composites are, as compared with most forms of graphite, highly crystalline and of comparatively high strength, elastic modulus, and thermal conduc­tivity. The fibers themselves are typically either poly­acrylonitrile (PAN) or Pitch derived. In general, one would select the PAN-based fiber, which is some­what less expensive, if the application required higher strength while the Pitch-based fibers would produce a product with superior elastic modulus and thermal conductivity.

As observed in Sections 4.18.2.2 and 4.18.2.3,

the composite materials, due to their typically higher strength and elastic modulus, have a superior performance in terms of thermal stress and thermal shock. Another key advantage of these materials stems from the fact that they tend to fail in a less abrupt manner than seen for graphite or ceramics in general due to the presence of the reinforcing fibers, which bridge evolving crack fronts. This can be seen by casual inspection of Figure 6, which compares the nuclear graphite Poco AXF-5Q(historically used in TFTR and for other nuclear applications) and the FMI-222 balanced weave, 3D CFC. From Figure 6, and by a comparison of the graphite and composite data of Table 2, it is clear that the FMI-222 CFC material has both higher bend strength and higher elastic modulus (greater slope) as compared to this Poco graphite. Moreover, it is clear from Table 2 that other engineering properties of importance, such as strength and thermal conductivity, are superior for the CFC. These superior properties are primarily attributable to the exceptional quality of graphite

fiber. Unlike nuclear graphite, which is on the order of 20% porosity with a relatively imperfect, heavily faulted, inhomogeneous amalgam of filler particles (such as coke) and graphitized binder (such as pitch; see Section 4.10.2 in Chapter 4.10, Radiation Effects in Graphite for a discussion of graphite man­ufacture), graphite fibers, while somewhat different depending on the starting material (PAN, Pitch, Rayon, etc.), are extremely uniform, and highly crys­talline with density that can approach theoretical den­sity. This leads to exceptional properties. For example, the PAN-based T-300 fiber has a tensile strength of 3.66 GPa, slightly higher than the 2.41 GPa strength of the P120 fiber of the FMI-222 composite of Table 2, or more than 40 times that of the Poco AXF-5Q graphite. Similarly, the elastic moduli of T-300 and P-120 fibers are 21 and 75 times the elastic modulus of the Poco AXF-5Qgraphite. In the case of the P-120 fiber, which has been graphitized at a very high temperature, very long, defect-free basal planes oriented along the axis of the fiber result in excep­tional 1D thermal conductivity (640 W m-1 K, twice that of copper). This property is the primary reason for the twofold increase in ambient thermal conduc­tivity of the FMI-222 composite as compared to the Poco AXF-5Qgraphite. Clearly from this example of thermal conductivity, the architecture (fiber weave or loading) will determine the composite properties.

Examples of practical fusion CFCs are the mate­rials chosen for consideration and application by the ITER project. Table 3 provides the nonirradiated thermophysical property data for selected CFCs

of the ITER project. A review of these properties emphasizes the anisotropic nature of the composite system, which is engineered through selection of the fiber type and route to matrix infiltration, fiber architecture, and final heat treatment of the system. All the materials for this application have been

engineered with a preferred thermal conductivity direction (the x direction in the table), and in order to maximize thermal conductivity, the composites will tend to have a higher volume fraction offibers in that direction and the fibers will be ofthe higher conduc­tivity pitch-based type. In the directions normal to

this preferred thermal conductivity direction, for strength, cost, and fabricability reasons PAN-based fibers are typically chosen. The composite INOX Sepcarb NS31 underwent a final processing step of 10 ± 2% liquid silicon infiltration. This silicon reacted with carbon-producing SiC, which is thought to mitigate chemical erosion and tritium retention while enhancing oxidation resistance.

Also observed from Figure 6 is the clear differ­ence in the shape of the load-displacement curves for the two materials. Clearly, the composite material has significant nonelastic behavior, which is attributed to the progressive load transfer from the composite to the high-strength fiber as the matrix becomes exten­sively microcracked. This contrasts with the graphite material, which undergoes abrupt failure when the load exceeds some critical stress adequate to propa­gate a crack through the test article. This added toughness of the composite is another key attribute to the systems that make it particularly attractive for fusion applications where disruption (shock) loading tends to produce interconnected cracking in materi­als leading to loss of material mass.

An extensive development programme performed especially in Europe has enabled the production of very good Be/CuCrZr alloy joints by HIPping. The progress on the fabrication of Be/CuCrZr joints in Europe is described in Lorenzetto et a/.155,161,162 and Sherock et a/.163 The HIP joining temperatures ranged from 540 to 580 °C. Different interlayers such as Cr, Ti, and Cu were tested and reported. The selection of the joining conditions used for the fabri­cation of representative first-wall mock-ups was done on the basis of mechanical test results performed with guillotine shear test specimens. The best results were obtained with Ti and Cu interlayers at 580 °C, in which the shear strength exceeded the yield strength of the parent materials as creep of the CuCrZr part or rupture of the Beryllium part was observed. Per­formance achieved with representative first-wall mock-ups exceed the present ITER design require­ments, namely, 30 000 cycles at 0.5 MW m~ peak heat flux plus transient events up to 1.4 MW m — for about 1000 cycles (see Section 4.19.5.1.2).

A neutron irradiation programme is still in progress to complete the full characterization with irradiated mock-ups (see Section 4.19.5.2).

In the frame of the ITER Programme, in addition to Europe, the countries that in the past were inter­ested to supply the ITER first wall were China, Korea, the Russian Federation, and the United States. The underlying development work which has recently been performed in these countries to estab­lish the necessary fabrication capability is summar­ized elsewhere (see, e. g., Sherock et a/.,163 Nishi et a/.,164 Hong et a/.,165 Youchison et a/.,166 Lee eta/.,167 Park era/.,168 Liu,169 and Chen170).

4.21.1 Introduction

The use of a ceramic coating for electrical insulation is a key technology for fusion blanket systems using liquid metals as breeding and coolant material and solid metals as the structural material. Particularly for the blanket system using liquid lithium and vanadium alloys (Li—V blankets), coating develop­ment is a major feasibility issue (see also Chapter 4.12, Vanadium for Nuclear Systems for vanadium alloy and liquid Li blankets). Overviews of the coating development for liquid lithium blankets are available in recent publications.1-3 It should, however, be noted that, with the development of a more fundamental understanding of coating behavior and blanket design, there has been a paradigm shift in coating development. This chapter describes the present status of insulator coating R&D, in addition to a historical overview of its development.

The Master Curve concept, developed originally for quasistatic loading condition, has proven to be useful for dynamic KJd tests conducted at a high loading rate. Assuming that only the value of T0 is rising (along with the material yield strength) with increasing loading rate, dK/dt, the dynamic fracture toughness, Kjd and associated T0 should, in principle, be esti­mated from the value of T0 measured by static tests. This dependence was empirically evaluated in 1997 by Wallin using a large dataset consisting of dynamic and static fracture toughness data measured for vari­ous structural steels (yield strength ranging from about 200 to nearly 1000 MPa). The Master Curve method was applied to both the static and dynamic loading rates.44 Based on these results, as well as an IAEA round-robin exercise (report to be published in the IAEA Report Series within the framework of the Technical Working Group on Life Management of Nuclear Power Plants), the Master Curve approach
appears to be fully applicable to dynamic fracture toughness measurements conducted in the ductile — to-brittle transition region.

4.14.5.2 T0 Versus Charpy V-notch Transition Temperatures

In the case of dynamic loading with notched speci­mens, the correspondence with the fracture tough­ness test and T0 is complicated due to several uncertainties associated with the Charpy V-notch impact test. First, the loading situation is very differ­ent in the dynamic loading of a notched specimen compared to the quasistatic loading of a fatigue pre­cracked specimen. Due to the differences in the load­ing conditions, the measured Charpy energy includes a significant proportion of both crack initiation and propagation, and often some energy associated with crack arrest; whereas, the quasistatic SE(B) test in the transition region characterizes mainly crack initiation conditions. Additionally, the inherent data scatter and curve fitting required to obtain Charpy V-notch parameters increase the uncertainty of estimating and correlating the transition temperatures.

Correlations between the Charpy V-notch tem­peratures T28j and T41j versus T0 are presented in Sattari-Far and Wallin.11 The correlations, which are based on data from over 200 pressure vessel steels, are currently being reassessed in more detail, but in applications where the direct estimation of T0 is not possible, the correlations can be used as indicated below (s is standard deviation):

T0 = T28J — 19 °C (s = 22 °C) [42]

T0 = T41J — 26 °C (s = 25 °C) [43]

4.14.2 Summary and Conclusions

The Master Curve methodology has been described as an advanced, direct technique of determining the fracture toughness of ferritic structural steels. The application of the methodology has increased during the last decade and spread worldwide extending beyond the initial applications associated with NPP surveillance and integrity assessment programs. Today, the methodology is well known and increas­ingly accepted by safety authorities as a standar­dized method for application in safety assessments. Methods based on conventional approaches, such as the Charpy V-notch test, are still widely used, and probably will be used in parallel into the foreseeable future. Once a sufficient amount of reference data have been measured using the Master Curve method, it will gain even further acceptance. Also, further understanding of the limits of applicability for differ­ent steels will be obtained. Over this transfer period, the correlations developed between the different methods should play a significant role, providing support for properly analyzing data and encouraging the use of the more advanced methods. Note that some correlations, like those proposed for estimating crack arrest toughness from Charpy V-notch tests have brought new applications for the instrumented Charpy test. The overall trend in fracture mechanics testing is toward characterization and methods which allow the use of small and/or moderate-size speci­mens simulating the true loading conditions and accounting for the expected micromechanisms of fracture. The Master Curve approach and its implementation in the ASTM E 1921 methodology have proven to be a valuable and powerful analysis tool for a wide variety of applications involving ferritic steels.

Tritium diffusion in metals is simply the process of atomic tritium moving or hopping through a crystal lattice. Tritium tends to diffuse relatively rapidly through most materials and its diffusion can be measured at relatively low temperatures. Diffusivity, D, is a thermodynamic parameter, and therefore, fol­lows the conventional Arrhenius-type dependence on temperature:

D = Do exp(-ED/RT) [2]

where D0 is a constant and ED is the activation energy of diffusion. Measuring tritium diffusion is nontrivial because of the availability of tritium. Therefore, hydrogen and deuterium are often used as surrogates. From the classic rate theory, it is commonly inferred that the ratio of diffusivities of hydrogen isotopes is equivalent to the inverse ratio of the square root of the masses of the isotopes:

where m is the mass of the respective isotope, and the subscripts tritium and hydrogen refer to tritium and hydrogen, respectively. When this approximation is invoked, the activation energy for diffusion is gener­ally assumed to be independent of the mass of the isotope. Diffusion data at subambient temperatures do not support eqn [3] for a number of metals;2 however, at elevated temperatures, the inverse square root dependence on mass generally provides a rea­sonable approximation (especially for face-centered cubic (fcc) structural metals).3-9 Although eqn [3] provides a good engineering estimate of the relative diffusivity of hydrogen and its isotopes, more advanced theories have been applied to explain experimental data; for example, quantum corrections and anharmo — nic effects can account for experimentally observed differences of diffusivity of isotopes compared to the predictions of eqn [3].3, For the purposes of this

report, we assume that eqn [3] is a good approximation for the diffusion of hydrogen isotopes (as well as for permeation) unless otherwise noted, and we normal­ize reported values and relationships of diffusivity (and permeability) to protium.

4.16.2.2 Solubility

The solubility (K) represents equilibrium between the diatomic tritium molecule and tritium atoms in a metal according to the following reaction:

1 /2T2 $ T [4]

The solubility, like diffusivity, generally follows the classic exponential dependence of thermodynamic parameters:

K = K exp(-AHs/RT) [5]

where K0 is a constant and AHs is the standard enthalpy of dissolution of tritium (also called the heat of solution), which is the enthalpy associated with the reaction expressed in eqn [4]. A word of caution: the enthalpy of dissolution is sometimes reported per mole of gas (i. e., with regard to the reaction T2 $ 2T as in Caskey11), which is twice the value of AHs as defined here. Assuming a dilute solution of dissolved tritium and ideal gas behavior, the chemical equilibrium between the diatomic gas and atomic tritium dissolved in a metal (eqn [4]) is expressed as

1/2 (mTT + RT ln ^TI) = m0 + RTlnо [6]

Ptt

where c0 is the equilibrium concentration of tritium dissolved in the metal lattice in the absence of stress, тіт is the chemical potential of the diatomic gas at a reference partial pressure of pli, and ml is the chemical potential of tritium in the metal at infinite dilution. This relationship is the theoretical origin of Sievert’s law:

co = K (pii)1/2 [7]

where to a first approximation, the solubility is equiv­alent for all isotopes of hydrogen.

It is important to distinguish between solubility and concentration: solubility is a thermodynamic property of the material, while the concentration is a dependent variable that depends on system conditions (including whether equilibrium has been attained). For example, once dissolved in a metal lattice, atomic tritium can interact with elastic stress fields: hydrostatic tension dilates the lattice and increases the concentration of tritium that can dissolve in the metal, while hydrostatic compression decreases the concentration. The relationship that describes this effect in the absence of a tritium flux12-14 is written as

cL = co exp(~RT) [8]

where cL is the concentration of tritium in the lattice subjected to a hydrostatic stress (a = ay-/3), and VT is the partial molar volume of tritium in the lattice. For steels, the partial molar volume of hydrogen is ^2 cm3 mol-1,15 which can be assumed to first order to be the same for tritium. For most systems, the increase of tritium concentration will be relatively small for ordi­nary applied stresses, particularly at elevated tempera­tures; for example, hydrostatic tension near 400 MPa at 673 K results in a ^15% increase in concentration. On the other hand, internal stresses near defects or other stress concentrators can substantially increase the local concentration near the defect. It is unlikely that local concentrations will significantly contribute to elevated tritium inventory in the material, but locally elevated concentrations of hydrogen isotopes become sites for initiating and propagating hydrogen — assisted fracture in structural metals.

In contrast to disruptions, ELMs occur during nor­mal operation in the H-mode and are characterized as instabilities caused by the steep temperature and density gradients at the plasma edge, which deposit a significant amount of energy at a high repetition rate.181’182 In particular’ it is the expected high repe­tition rate for ELMs during the lifetime of the PFC (>1 million of events at a frequency of 1-25 Hz183) that, although yet unexplored, will impose high demands on the PFMs.

While it is the desire of plasma physicists to oper­ate in H-mode regimes with high-energy ELM depo­sition (> 1 MJ m~2) the response of bulk tungsten’ tungsten coatings, and tungsten alloys to such loading conditions, that is, surface melting, melt motion, material erosion, and vaporization, is

detrimental. To obtain further insight into material behavior under these conditions’ modeling of experi­mental conditions was carried out.9’167’168’190-195 It has been shown that with regard to melt motion/ erosion, the results of the different facilities cannot be directly compared196 and none of the testing facil­ities used provides identical conditions to those that will occur in a tokamak. However’ mitigation techni­ques have been explored for reducing the applied ELM energy, which, in general, can only be done at the expense of a higher repetition rate.183 The extent to which the ELMs have to be mitigated depends on the melt formation at tile edges due to the shallow plasma impact, which was experimentally found to be between 0.4 and 0.6 MJ m~2 for pure forged tung­sten.189’197 On the other hand’ the effect of crack formation during ELMs on the lifetime behavior of the PFCs has to be taken into account. As mentioned before, this behavior is yet unexplored at high repe­tition rates.

Typical investigations on various grades of W82’187’189’198 coatings21’54’186’199 and alloys146’157’187 were in the range of 10-100 repetitions. In a few cases up to 1000 repetitions and in single experiments even on the order of tens of thousands of repetitions have been obtained depending on the testing facility used. As the repetition rate is still rather low compared to the expected millions of events the main interest of

(a)

(b) these investigations was the qualification of different W grades and alloys (see Section 4.17.3.3) with regard to their damage and cracking thresholds. The characterization was done as a function of the main parameters described in Section 4.17.4.1, that is, microstructure, power density, and base temperature.

The results obtained so far showed that crack formation20 vanishes above a certain base tem­perature (see Figure 6).82,157,198 This temperature decreases with increasing material ductility, indicat­ing that the use of W alloys or fine-grained W is preferred. In the case of an anisotropic microstructure, this effect strongly depends on the material’s orienta­tion. Better results are obtained for grain orientations parallel to the loaded surface (see Section 4.17.4.1), yielding differences in the threshold temperature compared to the orthogonal direction of up to several hundred K (cf. Figure 6(a) and 6(b)). Recrystallization leads to a slight homogenization of the material’s
microstructure and therefore the mechanical proper­ties; however, there is no full convergency of the orientation-dependent thresholds.82

Despite the fact that for the currently limited number of applied pulses no crack formation was observed above a material and orientation-dependent temperature, the material is still damaged by plastic deformation and surface roughening. The evolution of this plastic deformation and of the related material hardening as a function of the applied number of loads is still unclear and has to be investigated. However, there are also heat load levels (at least up to Tbase < 800 °C), at which no visual material degradation could be determined and the future goal will be to investigate if these damage thresholds are still valid for high repetition rates, at higher base temperatures, and particularly in combination with neutron irradiation (see Section 4.17.4.3) and plasma wall interaction (see Section 4.17.4.4).

All the information given above on the effect of ELMs is also directly transferable to the short tran­sient events expected for inertial fusion applications and has been verified by IFE-related tests on dif­ferent W-based materials.201-204 There are coating parameters of high interest besides those mentioned above; these include the manufacturing-induced residual stresses at the surface, which are dependent on the used substrate, and the coating thickness. As mentioned in Section 4.17.4.1.1, the applied loading conditions and therefore the pulse length determine the heat penetration depth.163 As a result, the tem­perature and stress gradient induced under IFE applications should be similar to those in X-ray anodes (see Section 4.17.2). In case of thin coatings, residual and induced stresses might affect the coating to substrate interface and could lead to interfacial crack formation and delamination. This leads to minimum requirements for coating thicknesses that depend on the applied loading conditions.54 For example, in industrially produced X-ray anodes, W-Re coatings are typically used with a thickness of 200-700 pm205,206 to provide better mechanical and thermal-shock properties compared to pure W.204 However, the first experience on the influence of ELMs on coatings under real plasma operational con­ditions will be gained in the ITER-like wall project in JET, which involves testing relatively thin PVD- tungsten coatings (14-20 pm) on a CFC substrate that provides a strong and anisotropic CTE difference.19,142

The behavior of this material under the above outlined transient heat loads is of course a key factor for the lifetime assessment of PFCs. However, the
results obtained for pure thermal shock testing might underestimate the material damage and by this over­estimate its lifetime. Only a combination of thermal shock, thermal fatigue (see Section 4.17.4.2), neu­tron irradiation (see Section 4.17.4.3), and plasma wall interaction (see Section 4.17.4.4) will be able to give appropriate answers for the selection of suitable grades of W.

Reliable mechanical tests on CFCs joined to heat sinks are still an issue. As ASTM tests to measure the shear strength of CFCs joined to metals are not available, several laboratories have independently developed tests for CFC to metal joints, making interlab comparison of results almost impossible. Joints obtained by AMC® have been extensively tested,146 in particular for shear strength, with data ranging from 20 to 60 MPa for prepared samples. At 600 °C, the shear strength dramatically decreases to ^20 MPa. The shear strength and tensile strength of the improved AMC® (TiSi-AMC) joint are in the range of 54—73 MPa and of 39—64MPa, respectively.130,146

Monoblocks obtained by AMC® have been measured after HHF tests147: apparent shear strength has been measured in the range of 30—60 MPa. Some cracks have been found at ±45°, ±90°, and ±135°, considering 0° as the flux direction, leading to the de­tachment of CFC from the Cu layer before testing.148 The Cr-modified CFC—Cu joined samples148-152 have been measured by single-lap test (adapted from ASTM C 1292, C 1425) and off-set single lap test (adapted from ASTM D 905) at room temperature. Independent of the CFC surface machining and different casting process, results obtained133,138 for Cr-modified CFC on more than 50 samples yielded average values of apparent shear strength ranging from 26 to 32 MPa. The average shear strength is in any case higher than the interlaminar shear strength of the CFC (15 MPa).

The shear strength of the CFC—Cu joints (flat-tile geometry) obtained by using a commercial Gemco® brazing alloy to braze CFC to pure copper was 34 ± 4 MPa, measured by single lap tests at room temperature. This is comparable to the values obtained by other joining processes and higher than the intrinsic CFC shear strength.140

Mechanical tests on the monoblock braze require specific designs: some of them are adapted by ASTM D 4562-01 ‘‘Standard Test Method for Shear Strength of Adhesives Using Pin-and-Collar Speci­men’’, as in the compression test used by Plansee AG, but the joint is not stressed in uniform pure shear state.

Reliable nondestructive tests (NDTs) are also extremely important for nuclear fusion components, especially for high heat flux PFCs. NDTs on CFC—Cu joints are complex because of the different response of CFC and copper to the physical excitations used to test the component. The aim of NDT is to identify and localize defects in the joined components before submitting them to high heat flux tests or actual appli­cation. It is also important to identify the maximum acceptable defect size, as a function of its position, defined as the largest defect that is stable under spe­cific loads in the fusion device.139,149 Several techni­ques are used for NDT150: X-ray microradiography, X-ray microtomography, ultrasonic inspection,151,153 lock-in thermography,1 2 and transient infrared ther­mography (SATIR). SATIR (Figure 42) is the French acronym for infra red acquisition and data proces­sing device: it is a dedicated facility developed in Cadarache-France at CEA. SATIR consists of record­ing the surface temperature evolution of the compo­nent with an infrared device during the circulation of hot (^95 °C) and cold (~5 °C) water through the cooling channel of the component. The transient ther­mal response is compared to a ‘defect-free’ component; defects such as debonding of CFC tiles from heat sinks are detected by a slower temperature surface response (Figure 42).

Ultrasonic inspection has been applied to the flat tile and the monoblock design. Defects on joints between materials having very different acoustic impedance (e. g., copper and CFC) result in the gen­eration of a high reflected echo, making defect detectability more difficult.141 Lock-in thermogra­phy consists of applying a series of heat flashes on the CFC. The main advantage of this technique is that there is no need for an active cooling of the component and it can also be used as an inspection method during the manufacturing process.142,152

Several nondestructive tests have been performed on the Cr-modified CFC—Cu joined samples150 not only to test performance, but also to verify and compare the reliability of these tests on a CFC—Cu interface. However, the present conclusion is that nondestructive tests of joints should be validated by destructive tests such as morphological evidence of the detected defect and mechanical testing.

Thermal quench of a full-performance ITER plasma, with ~-350 MJ of thermal energy will result in signif­icant transient heat loads causing vaporization and melting especially of divertor material. The erosion lifetime due to these events will depend on mitigating effects resulting from vapor shielding, redeposition of the eroded materials, and melt layer behavior. Dis­ruptions could also result in significant Be erosion due to vaporization and possible loss of the melt layer. The evaporated and melt layer thicknesses are of the order of ~10 mm and ^50 mm, respectively, for the radiation energy density of 1 MJ m~ over 1 ms expected for the first wall under the assumption of no vapor shielding.214 Figure 22 compares the pre­dicted melted layer thickness for a thermal quench time of 0.1 and 1 ms for beryllium and tungsten. For a disruption energy density of 1 MJ m~ , we see that about 50 mm of Be are melted, as compared to 60 mm of tungsten. This result occurs even though beryllium melts at 1283 °C, whereas the melting point of tung­sten is 3410 °C. The explanation for this result is that, under very intense energy deposition, a nearly instantaneous thermal balance is established between

the energy deposited by the plasma and cooling by vaporization of beryllium. The vaporization temper­ature of beryllium is variously reported as 2480­2979 °C, as compared to over 5630 °C for tungsten. Similarly, the latent heats of melting and vaporization of Be are also lower than the corresponding tungsten values. This explanation is consistent with the results in Figure 22(b), which shows the amount of material that is vaporized for a thermal quench time of (1) 0.1 ms and (2) 1 ms. At an energy density of 1 MJ m~2 and for a time of 0.1 ms the thickness of vaporized material is 10 pm for beryllium and 2.5 pm for tungsten. It must be noticed that vapor shielding is not included in these calculations and that the results therefore should be considered conservative.

High-pressure noble-gas-jet injection, for exam­ple, of neon and argon, has shown to be a simple and robust method to mitigate the deleterious effects of disruptions in tokamaks.215 The gas jet penetrates the central plasma at its sonic velocity. The deposited species dissipate >95% of the plasma energy by radiation and substantially reduce mechanical stress on the vessel caused by poloidal halo currents. Nevertheless, there remains some concern that even mitigated disruptions could damage the Be wall
in ITER. Preliminary calculations show that even during a mitigated disruption in which the plasma energy is intentionally dissipated by radiation in ^1 ms by disruption mitigation techniques, the entire first wall of beryllium can melt to a depth of roughly 20-50 mm.212,216 The fate of this melted layer is uncertain. If the melt layer resolidifies, it provides a means of removing the oxide layer and creating a clean Be layer for oxygen gettering. On the other hand, if significant j xB forces associated with the plasma termination mobilize the melt layer within the vessel, it will likely lead to operational difficulties.

Another 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 to depths only where the thermal stress level is above the yield stress.