Work began at an early stage to assess the thermo­mechanical properties of candidate insulating materi­als for fusion applications. In an attempt to determine the best combination of mechanical, thermophysical, and dielectric properties for the demanding H&CD applications, Al2O3 (both alumina and sapphire), AlN, Si3N4, BeO, and MgAl2O4 in numerous differ­ent grades were examined ‘as-received’ and following irradiation.142-149 At room temperature, the unirradi­ated thermal conductivity of a typical alumina is of the order of 30 W m-1 K~ , and that of BeO about 280Wm~1K~1. These values are sufficiently high
for IC and LH heating systems to ensure adequate cooling in most cases; however, the thermal conduc­tivity in ceramics is reduced because of increased phonon scattering, by the presence of point defects and to a lesser extent by extended defects or aggre­gates. Hence, one expects a reduction in thermal conductivity on irradiation, together with a notable influence of the irradiation temperature, that is, irradiation above temperatures at which the radiation- induced defects become mobile and can either recombine or aggregate should lead to a lower degra­dation of the thermal conductivity, while low — temperature irradiation should have a marked effect because of the increased point defect stability. The expected general behavior was confirmed by the early data (Figure 11), and indicated that a maxi­mum reduction to about one-third of the room temperature thermal conductivity value could be expected.142-145 This will occur for a neutron fluence value (dpa), which strongly depends on the irradia­tion temperature. For near room temperature irradi­ation (300 K), reduction to the lower saturation level was observed by about 1023 n m — (0.01 dpa), whereas at 600 K this lower saturation level was only reached following a fluence of above 1024nm-2. Within rea­sonable margins, these values applied for Al2O3, AlN, and MgAl2O4. Similar PIE results were obtained at a later date for reactor irradiations at different tem­peratures of a wide range of ceramic materials.150 Because of the importance of point defects in the reduction of thermal conductivity, it is reasonable to expect that postirradiation measurements may un­derestimate the effect due to possible postirradiation annealing. An attempt to measure thermal conductivity in situ during reactor irradiation, although unable to quantify the degradation, did highlight a very rapid decrease in thermal conductivity by <1022nm — (0.001 dpa) at the startup of irradiation, followed by

151

saturation.

Finally, one should mention the specific case of sapphire and CVD diamond, the original and the present reference materials for ECRH. For sapphire, the need for low-temperature (<100 K) operation to minimize dielectric loss also provided a gain in thermal conductivity (200Wm-1K-1 at 100 K, c. f. about 30 W m-1 K-1 at room temperature). However, in addition to the dielectric loss showing a very low neutron tolerance (<1020nm- ) at this low temper­ature,128 the high thermal conductivity was reduced by over two orders of magnitude also by 1020nm — (10-5dpa), because of the enhanced point defect stability.147,152 In the case of CVD diamond, the increase in the room temperature dielectric loss was still tolerable up to 1022nm-2 (10-3 dpa).134 Unfortunately, although the extremely high thermal conductivity at room temperature («1800 W m-1 K-1) already began to degrade by 1020nm-2 (10-5dpa), it was at the tolerance limit by 1021 nm-2 (Figure 12).134 Almost identical results were reported after electron irradiation to 3 x 10-6dpa where the thermal con­ductivity was reduced by about 9%, confirming the importance of point defects.153

The fracture toughness of ferritic steels has been characterized by numerous different parameters. It is not the purpose here to discuss history, so the main parameter is linear-elastic KIc and its use with respect to the elastic-plastic Kjc. The KIc parameter has, in the past, been one of the most commonly used parameters, also for structural steels, but its limita­tions for describing the transition behavior controlled by both cleavage and ductile cracking are widely recognized (discussed later in detail) today. Many high-alloyed quenched and tempered steels, which exhibit practically no plastic deformation, still have moderate fracture toughness, KIc, and can be used and no additional benefit is achieved by using an elastic-plastic parameter like KJc. For these steels, the measurement of fracture toughness at one or a few temperatures is all that is necessary. For low — alloyed structural steels, which typically exhibit a pronounced ductile-brittle transition and may be loaded in a wide temperature range, the situation is different. In this case, an elastic-plastic parameter is needed. An example of such an application is an irradiated RPV where safety and performance have to be demonstrated in accident and abnormal service conditions which are more severe (i. e., at lower tem­peratures) than in normal operation. The integrity analyses must be based on the material fracture toughness covering the entire transition tempera­ture range. In the mid and lower temperature por­tions of the transition curve, cleavage fracture has to be explicitly considered, and the parameter KJc and Master Curve analysis of data are the preferred methods to determine fracture toughness. It should be noted that application of the methodology is not just limited to ferritic steels, but the fracture is generally cleavage or a stress-controlled type of mechanism.

When a cracked material is loaded, a plastic zone will develop at the crack tip. The size of this plastic zone depends on the crack tip loading (stress inten­sity) and the material yield strength. The radius of the plastic zone (rpl) can be expressed for mode I loading (tension loading perpendicular to crack plane) in a simple form as follows: ±( Kl

2я ffys

where KI is the stress-intensity factor and sys is mate­rial yield strength.

The plastic zone size is thus a measure of plasticity at the crack tip and can be used to assess the applica­bility of different fracture toughness parameters. In predominantly elastic cases, the plastic zone size can be very small, and the material may be analyzed using LEFM. The parameter KIc is generally determined without correction term for plasticity, when the plas­tic zone size is small in relation to the specimen dimensions. Otherwise, an elastic-plastic parameter such as the J-integral should be used.

For a compact specimen [C(T)] geometry and load­ing, the value of Kj can be calculated directly from the load (Pt) and the original crack length (a0/ W):

Pi

vWrnffi where W is specimen width, B is thickness, Bnet is net thickness, and function f (a0/W) is defined dependent on specimen crack depth.

For elastic-plastic conditions where a large plastic zone has developed and some stable crack growth can occur, one normally cannot use the LEFM parameter KIc and eqn [3] without at least some correction term(s). The use of EPFM is more practical to deter­mine directly the J-integral which takes into account the plastic component of the work done to the speci­men or component. The fracture toughness equiva­lence Kj, denoted as Kjc, can then be converted from the J-integral, which is first divided into elastic and plastic components:

Jc — Je + Jp [4]

where the elastic component of J is calculated from the elastic K (Ke) as follows:

J — [5]

where E is the elastic modulus and n is Poisson’s ratio.

The plastic component of J (Jp) is calculated from the total and elastic work using the measured load versus load line displacement or crack opening data. The value of Jc is converted to Kjc as follows:

K* 4 Jc rb [6]

Despite the limitations of LEFM as discussed earlier, one of the basic standards applied in the past in fracture toughness testing has been American Society for Testing and Materials (ASTM) E 399 which
defines the methodology for Kjc determination. This standard has recently been revised and issued as part of ASTM E 1820-08 (Annex 5),2 but the basic approach is essentially the same as in the previous versions of ASTM E 399. Because these standards, as well as a corresponding linear-elastic European stan­dard, BS 7448, are still being used to assess some ferritic steels, it is important to know how they differ from the Master Curve approach used in ASTM E 1921-08.3 In particular, the scope of the application requires discussion since it affects how Kjc data should be analyzed with respect to Kjc.

A common feature of LEFM KIc standards (such as the previous ASTM E 399) is that they express the fracture toughness as a single value, which should be a material property characterizing the resistance of a material to fracture. The value of KIc should be insensitive to specimen size, if the measured value fulfills the specified size criteria. When these condi­tions are met, the stress-strain condition at the crack tip has been thought to be predominantly plane strain, thus ensuring sufficient constraint to produce a minimum fracture toughness value for the material. The qualification to obtain a valid KIc measurement requires that relatively large test specimens have to be tested. The value of KIc is supposed to represent a lower limiting value of fracture toughness, but the method (ASTM E 1820, Annex 5) does not ade­quately cover ferritic steels where fracture by a cleav­age mechanism in the transition or lower shelf region has known statistical characteristics different from ductile initiation fracture toughness.2

Based on the current knowledge of fracture mechanics, confirmed by numerical finite element model analyses of local stress-strain fields, many of the arguments for the LEFM parameter KIc are not valid, especially for ferritic steels.4 First, for crack tip constraint, the KIc size requirement has been shown to be overly conservative, leading to highly oversized specimens. Also, it has been shown that KIc is not a size-insensitive parameter; a size adjustment, similar to that made for Master Curve specimens (discussed later in Section 4.14.1.3), should also be made to the values of KIc to correct for size effects and to make the data comparable with Kjc. Justification of this size effect argument has been demonstrated in many com­parisons made between the Kjc and the older KIc data measured with different size specimens.4 It is impor­tant to note that increasing the specimen size (both ligament and thickness) gradually diminishes the effect of the size adjustment, which means a reduced,
but still existing, dependence on specimen size even with large dimensions (discussed later in Section 4.14.1.3.3). One should also remember that the Mas­ter Curve size adjustment is valid only in the transi­tion region where cleavage fracture initiation is expected to occur, even after some stable crack exten­sion before cleavage fracture. Due to the size effect, the fracture toughness decreases in the transition region with increasing specimen size whenever cleav­age fracture is encountered. Examples of applying the statistical size adjustment are given in Section 4.14.3.

As mentioned previously, ASTM E 1820 (Annex 5) actually invalidates the KIc determination if the value exhibits transition behavior indicative of some cleav­age fracture. If KIc values characterize only the upper shelf behavior of ferritic steels, no size effect typical of the transition behavior should exist even iflater cleav­age initiation occurs. Thus, KIc determinations per­formed as per the previous ASTM standard versions do not necessarily fulfill the requirement or follow the recommendations now in ASTM E 1820. This issue is significant, because it concerns an essential qualifica­tion criterion of KIc determination. Another related issue is that ASTM E 1820-08 does not specify selec­tion of the test temperature to make sure that there is no cleavage fracture in the test. No posttest measures are required or recommended to confirm that the test data are not affected by cleavage initiation. The ques­tion arises: which of the reported KIc values should be size-adjusted consistently with Kjc data and which should not be size corrected?

Another important aspect is the 95% secant requirement (Pmax/PQ< 1.1), which limits stable crack growth to very small amounts for typical struc­tural steels exhibiting a rising tearing resistance curve.5 The methodology for KIc determination according to the ASTM E 1820 standard is thus not applicable for high-tearing resistance or high-toughness steels, which is also mentioned in Annex 5. The size require­ment for KIc is essential, since it means that the speci­men (ligament) size is the only factor which can be affected in pursuing a valid test result, if the test has to be performed at a specific temperature. If the secant requirement cannot be met, it is possible that no value of KIc can be determined at that temperature. For all of these reasons, fracture toughness testing in the tran­sition region is recommended to be made following standard ASTM E 1921, which takes into account the statistical nature of cleavage fracture.

Direct fracture toughness determination for the reactor vessel surveillance programs of nuclear power plants (NPPs) was one of the first applications of the

Master Curve methodology. Compared to the conven­tional Charpy V-notch methods, the Master Curve concept represents a new approach which makes pos­sible direct fracture toughness determination with only a few relatively small specimens, which is a more efficient use of limited test material. Using the Master Curve method allows statistical confidence to be applied to the directly measured data. The traditional practice of estimating fracture toughness from Charpy data (using correlations) is more unreliable due to large uncertainties associated with the correlations and the subsequent safety margins to meet regulatory require­ments. However, the Charpy-based methods are still in use and will continue to be used until a large amount of surveillance capsule Master Curve data is available. Existing data from correlations is discussed in Section 4.14.5, and material reference curves are discussed in Section 4.14.4.2. Use of these Charpy-based methods may remain to be the only way of estimating static fracture toughness in some cases, but is not the pre­ferred approach for future surveillance programs.

Characteristic differences of thermal creep effects between homogeneous materials and pebble beds are:

• In pebble beds, the contact surfaces between the pebbles increase with time. Blanket relevant creep time periods of greatest interest are in the order of less than a day because stress relaxation effects are expected to be quite fast.104

• Pebble beds are loaded first with relatively small stress gradients ds/dt; therefore, it is not differen­tiated between instantaneous plastic deformations and conventional thermal creep effects (the ther­mal creep correlations given below include both effects).

Thermal creep strains are measured by UCTs by keeping the uniaxial stress s constant at a given temperature T. Figure 21110 shows creep strains for Li4SiO4 pebble beds for different values of s and T; the thermal creep strain occurring during the stress increase period has been taken into account in this representation.

The data are well fitted by a constant exponent n for the time dependence. A constant exponent n was also found for most of the Li2TiO3 pebble beds.95 For some batches, during the first several hours, the same exponent n was observed; however, there was a subsequent increase in the creep rates. The Li2TiO3 batches that showed this behavior were in general characterized by low sintering temperatures, large

porosities, and small grain sizes, (for details, see Reimann et a/.98). Eventually, impurities could play a role, too. The data are fairly well plotted in Arrhenius graphs by straight curves, as shown in Figure 22.95

For a selection of ceramic breeder material candidates, the influence of pebble-bed thermal conductivity as a function of, for example, packing factor and pebble-bed compression was studied by
means of UCTs. These were performed for the ceramics Li4SiO4, Li2TiO3, Li2ZrO3, and Li2O with temperatures up to 480 °C and pressures up to 8 MPa, with packing factors varying between 56% and 63.5%. Creep strains in the pebble beds are identified to be functions of temperature, stress, and time and are found to be of the form ecr = A(T)amtn where m and n are independent of temperature and need to be deduced experimentally.97 This research was extended for Li4SiO4 pebble beds with temperatures up to 850 °C and pressures up to 9 MPa and concludes that thermal creep effects are negligible at temperatures below 600 °Q96,97 Creep behavior is also determined by the pebble properties: as mentioned earlier, lower creep strains were found for Li2TiO3 with small grain sizes (<5 pm) and high

98

sintering temperatures.

4.16.4.1 Expected In-Reactor Performance

As implied by Figures 21 and 22, Tables 1 and 2 and

in the earlier sections, permeation barriers can be used to reduce the effective permeation in laboratory

testing.175-179,183,195,196198-203,205 PRFs from labora­tory experiments have been reported to be many

175,183,195,196,201,202

thousands in certain barrier systems.

However, while the data available in the open litera­ture are quite limited, there is significant evidence that the effectiveness of the permeation barriers decreases in radiation environments. There were three sets of experiments21 — performed in the high flux reactor (HFR) Petten reactor in the Neth­erlands. In the first of these experiments21 , reported in 1991 and 1992, tritium was produced by the liquid breeder material Pb—17Li. Permeation of tritium through a bare 316 stainless steel layer was compared with that through an identical layer cov­ered with a 146-p. m thick aluminide coating. Over the temperature range 540-760 K, the barrier was reported to decrease the permeation by a factor of 80 compared to the bare metal, that is, PRF = 80. In the LIBRETTO-3 experiments,220 three different permeation barrier concepts were tested with the tritium again produced by the liquid breeder mate­rial. One irradiation capsule for tritium breeding was coated on the outside with a 6-8-p. m thick CVD layer of TiC. A second capsule was coated on the inside with a 0.5-1.5-p. m thick layer of TiC followed by a 2-3-p. m thick layer of Al2O3. The third barrier was an aluminide coating produced by the cementation process. The aluminum-rich layer was ^120-p. m thick with about 5 mm of Al2O3 on the outside. The single TiC layer reduced the tritium permeation by a factor of only 3.2, the TiC and Al2O3 layer reduced the permeation by 3.4, and the pack cementation aluminide coating reduced the permeation by a factor of 14.7. These are surprisingly small reductions PRFs compared to laboratory experiments. In a third set of experi- ments,221 the tritium production was achieved with the solid ceramic breeder materials. Both double-wall tubes and single-wall tubes with a permeation bar­rier were tested. The double wall configuration had an inner layer of copper. The permeation barrier on the other system was an aluminide coating with a thickness of 7 mm. The aluminide coating was reported to be 70 times more effective than the double-wall configuration in suppressing permeation. Unfortunately, different breeder materials were used for the two different experiments, and the results could have been strongly affected by the amount of tritium released from the ceramic as well as the form of release (T2 vs. T2O). The bottom line on the irradiation testing of barriers is that barriers do not perform as well in a reactor environment as expected from laboratory experiments: a PRF > 1000 has not been achieved in reactor environments.

Of the flux of particles that will impact the PFMs, the highest particle flux will be the ionized fuel itself. The energy of the impacting fuel ions on the various plasma-facing areas depends on many variables. For the divertor, where most of the interactions will occur, the majority of particles will have energies in the eV range. On the first wall where the interaction intensity is less, the charge exchange particles will mostly be in the keV range. For larger fusion devices of the future where dense plasmas and higher mag­netic fields result in the thermalization of the ener­getic helium (eqn [1]), those helium ions will have energies similar to the fuel ions. Electrons, which are in number density equilibrium with the plasma ions, also travel along the plasma field lines, albeit in the opposite direction. The high-energy neutrons pres­ent in the D+T reaction (14.1 MeV), or those for the D+D reaction (2.4 MeV) have mean free paths of several centimeters in graphite and so will not interact strongly with the first wall. However, these neutrons will be scattered and slowed down within and behind the first wall, resulting in a nearly isotro­pic flux of high-energy neutrons throughout the fusion device. The reaction of the plasma neutrons, ions, and electrons with graphite PFMs, which is discussed in some detail in the following sections, can have a wide range of effects. These effects include physical and chemical erosion of the first wall and thermomechanical property degradation of the bulk and surface material.

The discussion thus far has been limited to the operation of tokamaks in the quasi-steady state (long pulse). All present-day large tokamaks are pulsed machines with pulse lengths of seconds, where the plasma discharge consists of a rapid heating phase, a steady state, and a cool down phase. In this case, the heat flux is approximately uniform around the cir­cumference of the machine and scales with the machine power. However, a significant number of these plasma shots end in an abrupt and somewhat violent fashion referred to as disruption. When this occurs, the plasma rapidly becomes unstable and instantaneously ‘dumps’ its energy onto the PFC. This causes significantly larger heat loads than dur­ing normal operation, and in many cases, defines the design limits for these components.

In the past, plasma spraying was considered as a high deposition rate coating method, which could offer the potential for in situ repair of eroded or damaged Be surfaces. Development work was launched during the early phase of the ITER R&D Program in the mid-1990s.136 In the plasma spray process, a powder of the material to be depos­ited is fed into a small arc-driven plasma jet, and the resulting molten droplets are sprayed onto the target surface. Upon impact, the droplets flow out and quickly solidify to form the coating. With recent process improvements, high quality beryllium coat­ings ranging up to more than 1 cm in thickness have been successfully produced. Beryllium deposition rates up to 450 gh-1 have been demonstrated with 98% of the theoretical density in the as-deposited material. Several papers on the subject have been published.136-138 A summary of the main achieve­ments can be found in Table 4.

However, based on the results available, the initial idea of using plasma-sprayed beryllium for in situ (in tokamak) repair was abandoned for several rea­sons. First was the complexity of the process and requirements to control a large number of para­meters, which affect the quality ofthe plasma sprayed
coatings. Some of the most important parameters include plasma spray parameters such as (1) power, gas composition, gas flow-rate, nozzle geometry, feed, and spray distance; (2) characteristics of the feedstock materials, namely, particle size distribution, morphol­ogy, and flow characteristics; (3) deposit formation dynamics, that is, wetting and spreading behavior, cooling and solidification rates, heat transfer coeffi­cient, and degree of undercooling; (4) substrate conditions, where parameters such as roughness, temperature and thermal conductivity, and cleanli­ness play a strong role; (5) microstructure and properties of the deposit, namely, splat characteris­tics, grain morphology and texture, porosity, phase distribution, adhesion/cohesion, and physical and mechanical properties; and (6) process control, that is, particle velocity, gas velocity, particle and gas temperatures, and particle trajectories. Second, plasma-sprayed beryllium needs (1) inert gas pres­sure, (2) reclamation of the oversprayed powder (more than 10%), and (3) strict control of the sub­strate temperature. The higher the temperature the higher the quality of the plasma-sprayed coating, but unfortunately, an easy and reliable method to heat the first wall to allow in situ deposition was not found. Finally, tools to reliably measure the quality of the coating and its thickness are not available today and a strict control of the coating parameters is difficult to achieve.

Thus, it was concluded that plasma-sprayed beryl­lium for in situ repair is too speculative for ITER without further significant developments. Neverthe­less, this method still remains attractive and could be used for refurbishment of damaged components in

Table 4 Main achievements of ITER-relevant plasma-sprayed technology (summary of best results, not always achieved together)

Value/results Comments

Residual porosity (%) Thermal conductivity (WmK-1) Bond strength (MPa) Substrate temperature (°C)	~2 Up to 160 at RT 100-200 >450	Could be more than 5% Depends on temperature of substrate, maximum achieved at T~ 600-800 °C with addition of H Reasonable Very important for good strength, adhesion, and thermal conductivity. Keep in mind that CuCrZr temperature should not be higher than 500 °C for several hours due to overageing of CuCrZr Needed, but very difficult to do in situ Reasonable Reasonable It means that more than 10% of powder will be lost in chamber For first-wall conditions tested
Substrate preparation Deposition rate (kg IT1) Thickness (mm) Deposition efficiency (%) Thermal fatigue (MW m~2/ number of cycles)

hot cell, albeit it may be cheaper to replace a dam­aged component with a new one.

The irradiation behavior of copper and copper alloys has been extensively studied up to high doses (>100dpa) for irradiation temperatures of ^400- 500 °C.60 Most of the irradiation experiments of cop­per and copper alloys have been done in mixed spectrum or fast reactors, such as HFIR, Fast Flux Test Facility (FFTF), or EBR-II. It should be noted that the accumulation rate of helium in copper in fusion reactors is significantly higher than in fission reactors (~10 appm dpa-1 in fusion reactors vs. 0.2 appm dpa-1 in fast reactors).22 Attention must be paid to transmutation effects such as helium when the irradiation data of copper and copper alloys from fission reactors are applied for fusion reactor design.

4.20.5.1 Effect of Irradiation on Physical Properties of Copper and Copper Alloys

Neutron irradiation leads to the formation of trans­mutation products and of irradiation defects, dis­location loops, stacking fault tetrahedra (SFT), and voids. All these features result in reduction of electrical and thermal conductivities.36,37,61-63 At irradiation temperatures between 80 and 200 °C, the electrical resistivity is controlled by the forma­tion of dislocation loops and stacking fault tetra- hedra and transmutation products. The resistivity increase from radiation defects increases linearly with increasing dose up to ~0.1 dpa and saturates. The maximum measured resistivity increase at room temperature is about ~6%. At irradiation tempera­tures above ^200 °C, the conductivity change from extended radiation defects becomes less significant, and void swelling becomes important to the degrada­tion of the electrical conductivity.

Fusion neutrons produce a significant amount of gaseous and solid transmutation products in copper. The major solid transmutation products include Ni, Zn, and Co. The calculated transmutation rates for copper in fusion first wall at 1 MW-year m-2 are 190appmdpa-1 Ni, 90appmdpa-1 Zn, and 7 appm dpa-1 Co.2 Ni is the main transmutation element that affects the thermal conductivity of copper. It should be noted that water-cooled fission reactors would produce significantly higher transmutation rates of copper to Ni and Zn (up to ^5000 and 2000 appm dpa-1, respectively) because of thermal neutron
reactions. The data from fission reactor irradiation experiments must be treated with care when they are applied for fusion design.

A typical pressurized thermal shock (PTS) experi­ment can be conducted by loading a thick-wall pres­sure vessel with an embedded crack by an external load and a severe thermal transient. The aim of a PTS experiment is to load the test vessel so that unstable crack initiation occurs. In connection with these large-scale PTS tests, a material test program is typically performed to produce the necessary mate­rial property data. If the thermal transient includes large temperature changes, any crack initiation after the peak temperature may be affected by the WPS effect. WPS involves the material loaded at the peak temperature during the thermal transient to a stress-intensity level that is higher than the critical fracture toughness at lower temperatures during the transient. Because the material was preloaded to a higher toughness level, the critical fracture toughness at a lower temperature of the transient is higher than without the preloading. The WPS effect is complex and is related to the crack tip stress state and is expected to be most pronounced in short transients where strain ageing will not diminish the effect.

4.14.3.2.1 PTS test by Framatome

In the 1980s, Framatome performed a thermal shock pressure test for a thick-wall pressure vessel, includ­ing a very long but shallow crack.29 The crack length (2c) in the 230-mm-thick cylinder was 1000mm and the depth (a) was 17 mm. The vessel material (A508 Cl. 3) was characterized by KIc tests conducted with large, 75 and 100-mm thick, C(T) specimens. The aim was to characterize the material directly ahead of the crack front. The data from these characterization tests have been reanalyzed using the Master Curve
method.30 The analysis was performed by making the size adjustment of the material test data to the 1000 mm crack length, corresponding to that of the vessel. After this correction, the material characteri­zation data and the crack tip load data for the test vessel for the observed first, second, and third initia­tions should fall on the same Master Curve. The comparison in Figure 24 shows that the vessel initia­tions occurred within the 5% and 95% fracture probability bounds estimated for the material, that is, two initiations fell almost on the mean Master Curve and the third close to the 95% probability level. In this case, the result indicates no WPS effect between the initiations, although some effect can be expected due to some warm prestressing during the decreasing temperature of the transient. The C(T) fracture toughness and the PTS test data generally coincide for all three initiations.

Ceramic breeder materials offer a wide range of possibilities for the development of fusion energy based on the deuterium-tritium fuel cycle. Currently, most ceramic breeder R&D is focused on the lithium orthosilicate and metatitanate systems in the form of pebble beds.

This chapter started with blanket designs, material requirements, manufacturing routes, pebble and pebble-bed thermomechanics, tritium production and release properties, neutron-irradiation behavior, chemistry, and modeling.

One of the most important missions of ITER is to provide a test bed for breeding blanket modules, the so-called test blanket modules (TBMs). However, because ITER testing is a cost-intensive exercise and is most likely the only opportunity to test a fusion

blanket component prior to a DEMO, the questions of the choice of materials for the ITER TBM and the definition of a set of requirements (and the related qualification program) to ensure safety, reliability, and test performances become particularly impor­tant. Accordingly, Ying et a/.19 proposed a roadmap outlining the necessary development steps for quali­fying and accepting the pebbles for ITER and fusion applications (see Figure 65).

For each development step, a set of criteria is presented as a means for initial screening before proceeding to the next evaluation tests to reduce development costs. However, it is important to rec­ognize that ITER conditions (neutron fluence about

1. 5—2 orders of magnitude lower) are far from suffi­cient to qualify any specific breeder material to be used in DEMO. Thus, parallel with ITER and subsequent to ITER testing, tests such as HICU or in fusion relevant neutron sources such as International Fusion Materials Irradiation Facility (IFMIF) for any candidate ceramic breeders under typical reactor blanket conditions with relevant nuclear environment are necessary for this purpose.

Though a significant R&D effort on ceramic breeder development has already been made and a vast amount of data on material performance have been obtained, the knowledge to date on the limiting factors in blanket designs for long-term operation is still modest. These limitations are addressed here:

The quantity of W needed for the PFCs in a fusion device such as ITER or DEMO represents only a small fraction of the yearly production and the world’s reserves120 and its production can be easily satisfied by existing industrial capabilities. The same point is valid for stellarators and even more for iner­tial fusion devices, which only work with thin coat­ings. However, the issue of component lifetime has to be taken into account. Depending on the component lifetime, the recycling rate, and the storage time until a hands on level is achieved (see Section 4.17.3.2.6), the operation of numerous power plants may require an amount of tungsten that exceeds what is currently available from the market.

4.17.3.2 Tungsten Grades

Within current R&D programs for the selection and characterization of candidate grades of W and W alloys for fusion applications, many materials pro­duced according to the schemes outlined above were investigated. These are discussed in the following sec­tion, which introduces some of their characteristics. The manifold production processes described below for pure W are also applicable to W alloys.

Pure tungsten (undoped)

• Sintered W is the most readily available and cheapest grade with a grain size that depends on the initially used W powder. However, it is characterized by high porosity, low recrys­tallization temperature (1000-1200 °C), and low strength at elevated temperature.96 The option of improving the sinterability by add­ing small amounts of activators (Ni, Fe)121 increases the radiological hazard due to addi­tional activation products that have to be taken

119

into account.

• Forged or swaged W offers an increased density and a refined microstructure compared to sin­tered material, resulting in higher ductility and mechanical strength. Forging and swaging are therefore the industrial production processes that are typically applied not only for pure tungsten but also for most kinds of tungsten alloys (see below). This grade of W is manu­factured in block shape or more commonly in the form of rods with different diameters (<90 mm)40 showing an anisotropic micro­structure12 with elongated grains along the axial direction and an increasing grain size and porosity with increasing rod diameter. Thus, increasing rod diameter leads to a decrease in mechanical strength and ductility. For the production of monoblock tiles, such as those planned for ITER, rods with a minimum diameter of 30-35 mm are necessary.

• Rolled Wis applied in the form of plates or foils with thicknesses from 0.02 to 20 mm.40,123,124 It offers a densified but layered microstructure that is strongly anisotropic, with flat disc-shaped grains parallel to the rolled surface affecting the mechanical properties (see Section 4.17.3.2.3) and resulting in the risk of delamination.

• Double-forged W is in the form of blanks with a diameter of 140 mm and a height of 45 mm. The double-forging process, first in the radial and then in the axial direction, provides a more

isotropic microstructure than it is generated by single forging. This material should act as a reference grade for establishing a reliable mate­rials database for finite element calculations.82

• SC W provides higher ductility than polycrys­talline W, higher thermal conductivity, lower neutron embrittlement, higher thermal fatigue resistance, and a more stable structure at ele­vated temperatures. The disadvantages are high cost and low industrial availability.96,125,126

• Metal injection molded (MIM)-W127-129 provides a dense and isotropic microstructure with grain sizes on the order of the powder particle sizes used. A final densification by hot isostatic pressing (HIP) at temperatures >2000 °C leads to an improvement of the mechanical proper­ties; recrystallization and grain growth do not play a role. Furthermore, the production pro­cess offers the possibility of net shaping.

• Spark plasma sintered (SPS)-W and resistance sin­tering under ultra-high pressure}20 132 The mate­rial is characterized by a short fabrication time of only a few minutes keeping the initial fine microstructure determined by the powders used. The finer the grain size, the higher the microhardness and the bending strength but also the lower the achievable density. The applica­tion of alternatively uni-, two-, or three — directional orthogonally applied forces for the material’s densification during the process leads to internal stresses, which have an influence on the recrystallization behavior. Recrystallization and grain growth occur at 1500 °C. Depend­ing on the amount of porosity, the finer the initial grain size of tungsten, the smaller is the grain growth.

• Severe plastically deformed W (and W alloys, see below) with ultra-fine grains in the nm range are produced by either high-pressure torsion at 400 °C84,133 or by the equal-channel angular extrusion or pressure (ECAE or ECAP) pro­cess at high temperatures (1000-1200 °C).134 The material shows stable, that is, deformation — independent, recrystallization temperatures and exhibits considerably enhanced ductility and

fracture toughness.61,85,86,135,136

• Plasma-sprayed Winvolves, in general, application of VPS, more precisely also called low-pressure plasma spraying (LPPS), which provides a significantly reduced oxygen content and improved thermophysical properties com­pared to atmospheric (APS) or water-stabilized plasma spraying.42 However, LPPS-W is typi­cally characterized by a lower thermal con­ductivity (up to 60% of bulk tungsten is reported67) and a lower strength than bulk W particularly when deposited on large sur­faces. The recrystallization temperature is similar to pure W.48,137 Although the thick­ness of the plasma-sprayed coatings required for fusion applications are flexible, coat­ings with 200 pm or thicker are commonly produced.26,43,67,138,139 Furthermore, PS is the only production method that offers the possibil­ity to produce and repair W components.57,60,96

• CVD W provides a microstructure with a columnar grain structure parallel to the sur­face, high thermal conductivity similar to bulk

W, and a very high density and purity.6,140,141

Thicknesses up to 10 mm were produced,67 but its high cost is a significant drawback for prac­tical applications.52,96

• PVD W provides a featureless structure that is extremely dense and pore free. In contrast to plasma sprayed and similar to CVD-coatings, the deposition rates are low. Economic and process-related restrictions generally limit the deposited W thickness to 10-50 pm.13,54,55,67,142

• W foam for Inertial Fusion Experiment (IFE) applications provides structural flexibility dur­ing quasivolumetric loading. The material is microengineered with a relative density of ~21% and can be simultaneously optimized for stiffness, strength, thermal conductivity, and active surface area.143

Tungsten alloys

• Oxide dispersion strengthened W alloys such as W-La2O3, W-Y2O3, and W-CeO2 with oxide additions <2% are processed by powder met­allurgy methods similar to pure W.40 The insol­uble dispersoids, which are influenced in shape and distribution by the thermomechanical treatments during the production process,73,144 improve the grain boundary strength and machinability and play an important role in controlling recrystallization and the morphol­ogy of the recrystallized grains.68 This results in a higher recrystallization temperature by 100-350 K by suppression of secondary grain growth (i. e., grain boundary migration), lower grain size, higher strength after recrystalliza­tion, and better machinability than sintered W even at RT. This permits fabrication at lower costs.67 The size of the dispersoids in commercially available alloys is ~10 pm; how­ever, research on mechanically alloyed materi­als using submicron dispersoids is currently being performed.145 However, the presence of oxide particles with a melting temperature below those of tungsten has a negative effect

146,147

on the erosion resistance.

• W—3—5% Re is, compared to sintered pure W, characterized by a higher recrystallization tem­perature and strength even after recrystalliza — tion,148 better machinability, and improved ductility at low temperatures.67 The addition of Re, which has a high solubility in W, how­ever, reduces thermal conductivity, increases embrittlement after neutron irradiation, and significantly increases the cost and safety con­cerns because of the high Re activation under neutron irradiation.96

• W—1—2% Mo (TY and Ti) cast alloy. The addi­tion of Mo and the reactive elements Y and Ti, which reduce the amount of free oxygen and carbon and form obstacles to grain growth, improves the mechanical properties compared

to large grained pure cast W.67,73

• W—TiC produced by mechanical alloying and slow deformation techniques provides, similar to all other W alloys, higher strength and recrys­tallization temperature, better machinability, and improved ductility compared to pure W with superplastic behavior at temperatures of 1400-1700 °C.89 The addition of Ti-carbide par­ticles stabilizes the grains during the material’s production process. This generates an isotropic grain structure and has the additional effect of keeping a fine grain structure even in the recrystallized condition, but the alloy is more expensive. After recrystallization, the finer dispersoids of TiC particles improve the low- temperature impact toughness of refractory alloys following low-dose neutron irradia­tion.71,87-89,149-154 Other carbides, for example, ZrC155,156 or HfC (in combination with Re and Mo),96 can be used instead of TiC.

• K-doped W is a nonsag material that contains a maximum of 40 ppm of potassium.40 Originally known from the lighting industry, it provides high creep strength due to its aligned pore structure, high recrystallization temperature >1600 °C, and good machinability.68,77,78,157

• W—Si—Cr as a ternary or even by the addition of another element as a quarternary alloy is a

newly developed and not yet optimized mate­rial that is being investigated as a wall protec­tion material due to its favorable oxidation resistance, preventing excessive material ero­sion in case of accidental air ingress.158,159

• Severe plastically deformed W alloys offer, similar to pure tungsten (see above), significantly improved fracture toughness and ductility.61,84 The addition of alloying elements to the starting material (any developmental or commercial produced W alloy), such as Re or dispersoids, leads to an increasing stability of the grains and therefore a higher recrystallization tempera­ture and less grain growth.133

Any of the bulk materials mentioned above could be used and are being investigated in its cold-worked, stress-relieved, or recrystallized state. The latter is of particular interest due to in situ recrystallization of surface near regions during operation.10

In spite of the fact that a large variety of tungsten grades and alloys already exist, the attempts to fur­ther optimize these materials are ongoing. The fabri­cation and successful testing of He-cooled divertor mock-ups for DEMO and ARIES-CS102,160 under a heat flux of 10 MW m~2 are important driving forces for the present development of W alloys with improved performance in the fusion environment.25 However, R&D has to address many different issues related to the performance of the material when exposed to thermal loads, neutron irradiation, and the plasma; these will be discussed in the following section.