4.18.6.1 Joining of CFC to Heat Sink

CFCs, bonded either mechanically or otherwise to a metallic structure, are being used in most of the major fusion devices, including ITER,105,106

Tore Supra,111 Wendelstein 7-X,112 TFTR in United States,113 JT 60U, and JT60SU in Japan.110-113 Silicon-doped carbon is used as the first wall material for the Chinese reactor HT-7.114 See Table 3 for a description of ITER candidate composites (INOX Sepcarb NS-31, Sepcarb NB-31, and Dunlop Concept C1, for example).

As mentioned, CFCs will be used in ITER (see Table 5 for performance specifications), and specifically a 3D CFC for the divertor in the initial phases of the ITER project. The diverter, which is among the most technically challenging ITER com­ponents, is located at the bottom of the plasma cham­ber where the CFCs (and tungsten) are bonded to a copper alloy (CuCrZr). CFCs have been selected

for the lower part of the vertical targets for the initial phase of ITER operation (without tritium), and they must be able to resist ‘steady state’ heat flux up to 10MWm~2 for at least three thousand 400 s pulses and up to 20 MW m~2 in transient events.

As discussed earlier, the use of graphite-based materials (particularly CFCs) in a divertor is believed to be an advantage for the first phase of ITER operations because CFCs have good demonstrated performance in the currently operational plants (e. g., Tore Supra). Their primary competitor, tung­sten, also suffers from macroscopic cracking, melting, and possible melt layer loss, thus making the poten­tial damage to the divertor components more serious than that of CFC. As conditions of additional heating and off-normal events will be very likely during initial operation of ITER, CFC is the present refer­ence design solution for the lower part of the vertical targets for the ITER initial phase, when tritium retention-related issues are not relevant.

In the ITER design, CFC must be joined to the copper alloy CuCrZr-IG (ITER Grade)116-119 in order to transfer the heat loads. Two different designs being considered for this component are the CFC-bonded flat-tile (Figure 33(a)) and monoblock (Figure 33(b)). In order to join the mating surface of the CFC to the copper alloy heat sink, a pure copper interlayer (1-2 mm thick, oxygen free high conduc­tivity copper, 99.95%, CTE: 15.4 at RT, up 20.6 at 700 °C and up to 21.6 at 800 °C) is used to relieve, by plastic deformation, the thermal expansion derived between the CFC and copper alloy heat sink. The CTE of CFC can be found in Table 3, with that of the copper alloy stresses being (16-19) x 10~6K_1 at 700 °C).116 This joining design is shown schematically in Figure 34. Some alternatives to a pure copper interlayer have been proposed within the EU project ExtreMat. For example, an Mo interlayer (1-2 mm),

(b)

Figure 33 (a) Flat tile design. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3,

© 2010 The American Ceramic Society. (b) Monoblock design. Courtesy of J. Linke, FZJ, Germany.

a Cu/W fiber interlayer, and CFC monoblocks with a Cu/W fiber interlayer (Figure 35) have been prepared and tested up to 10MWm~2. Results (unpublished) indicate that this method is less promising compared to the use of a pure Cu thin layer. Ti-doped CFC have also been prepared and tested for flat tiles.117,118

Among the several possible options, the flat tile and monoblock configurations (see Figure 33(a) and 33(b)) with a pure copper interlayer have yielded the most promising results. In particular, the monoblock gives a more robust solution in com­parison with the flat tile for the vertical target and it is now considered as ITER reference geometry.119,120

The monoblock design requires drilled blocks of CFC into which a CuCrZr tube is inserted and joined; also necessary in the monoblock is a pure copper interlayer between the CFC and the copper alloy to relax interface stresses, which have been (modeled and) measured at ±45°, ±90°, ±135°. If 0° is considered to be the flux direction,121,122 the monoblock is preferred to the flat tile design. This design is also much easier to manufacture because of its better heat flux performances and because of its intrinsic ability to attach even in the presence of cracks at the interface of CFC and CuCrZr, caused by its preparation process.

Methods to join CFC to CuCrZr derive from techniques developed to join carbon-based compo­nents to metals. Brazing is a well-known joining process recommended for joining dissimilar

(a) (b)

11500 mm

Figure 35 Cross-section of a carbon fiber composite monoblock with a Cu/W fiber interlayer. Courtesy of EU-Extremat Project and Pintsuk, G. JFZ, Germany.

materials. If properly done, it results in good mechan­ical strength, high fatigue performance, and minimal thermal resistance at the interface.123-125 In the case of joining of CFC for nuclear applications, some additional restrictions must be considered. A primary consideration is that the joining materials should be low activation materials (LAMs), even if the volume occupied by these materials is negligible in respect of the total volume of the structural materials in the reactor.126 Also important is that the materials be irradiation-stable at the anticipated conditions. Furthermore, the use of pressureless joining
techniques is preferred as the parts to be joined are relatively large. Some joining materials are not allowed, for example, elements with high vapor pres­sure (e. g., zinc or cadmium), or those giving dangerous transmutation reactions. The ITER project mandates thermodynamic and mechanical stability up to at least 800 °C under vacuum for the joint, in order to satisfy requirements of Table 5; the joint must survive the thermal, mechanical, and neutron loads faced by the component, and it is expected to operate in a cyclic mode with an acceptable reliability and lifetime. The joining technology must also be compatible with the overall component manufacture process and in particular with the preservation of the thermomechanical properties of the precipitation — hardened CuCrZr alloy.127

Wettability is a key factor in the joining of CFC to the heat sink. It is not within the scope of this chapter to review this subject, and an overview can be found elsewhere.12 ,12 ,129 However, to restate the most salient point, the Cu interlayer cannot be obtained by directly casting copper on the CFC surface, because Cu does not wet CFC at all, the contact angle of molten copper on carbon substrate being about 1400.128,129 The poor wettability of CFC is related to the non­metallic character of its bonding, whereas the bonding electrons in copper are delocalized.123,124,128

Copper can be directly cast on CFC when the CFC surface is modified to form carbides by a solid state chemical reaction between the composite surface and elements such as Si, Al, Ti, Zr, Cr, Mo,

or W; some metals can form carbides with ‘metal-like’ behavior128 and are usually well wetted by molten metals. Several patents refer to joints between carbon-based materials and copper130-135: for exam­ple, pure Cr and Ti react with C to form carbides. Cr and Ti wet carbon-modified surfaces very well with a contact angle of 35-40° at 1775 °C and of 50-60° at 1740 °C, in Ar, respectively.128

A joining technique based on CFC surface modi­fication is the Active Metal Casting (AMC®) tech­nique originally developed in the eighties by the Austrian company Plansee for nonnuclear purposes. ‘Active’ indicates the activation of the CFC surface to allow it to be wetted by Cu. Physical vapor deposition (PVD) or chemical vapor deposition (CVD) of Ti coating on the CFC surface is followed by a high temperature treatment to form TiC, which improves the wettability ofCFC by molten copper. Active Metal Casting consists of casting a pure Cu layer onto a laser-textured and TiC-modified CFC surface.136,137 The laser texturing enhances Cu infiltration into the CFC, and the TiC-modified CFC surface improves the wetting. The special laser treatment of the CFC surface produces a large number of closed conical holes (diameter ^50-500 pm, depth 100-750 pm), thus increasing the joined area and providing addi­tional crack growth resistance. Due to the open poros­ity of the TiC-modified CFC and laser machining, the cast Cu penetrates into the CFC up to 2 mm. An example of a full-scale component produced by Plansee, Austria, is shown in Figure 36.

AMC® was successfully applied for both flat tile and monoblock geometry. However, AMC® technol­ogy requires laser machining of CFC surfaces, which might not be economically attractive for large-scale production. Laser-induced stresses in the joined area and cracks induced during the joining process have recently been measured and modeled.1 However,

Plansee has recently improved AMC® by using silicon and titanium to modify the CFC surface (TiSi-AMC) (Figure 37(c)).130 Ansaldo Ricerche — Genova, Italy, has proposed a joining technique based on a Cu-Ti-based (Cu ABA) commercial alloy, reinforced by 2D randomly oriented carbon fibers uniformly distributed in the brazing alloy. The joining is carried out at approximately 1000 °C. The Ti reacts with carbon to form a thin TiC layer that promotes wetting.131 Carbon fibers are expected to mitigate thermal expansion mismatch between CFC and the braze (Figure 38) and to react with titanium in the brazing alloy, resulting in beneficial thermal fatigue strength of the joint. This technique was

Plansee

Tungsten (upper part)

Blocks 33

Austenitic steel

Blocks 32

CFC — (lower part)

Blocks 1

Copper/steel
tube joint

Figure 36 Vertical target full-scale prototype manufactured by Plansee (high heat flux units) and Ansaldo Ricerche (support structure and integration). Reproduced from Missirlian, M., et al. J. Nucl. Mater. 2007, 367-370(2), 1330-1336.

successfully tested on CFC NB31-Cu flat-tile and monoblock joints. Several other solutions have been investigated to modify CFC surface, for example, by TiN or TiC, within the EU project ExtreMat.118

A method has been proposed based on the CFC surface modification by reaction of Cr, Mo, and W. Both Cr and Mo have been extensively used as active elements in brazing alloys for copper active brazing and in patents referring to nonnu­clear applications.130-135,138,139 As example W, Mo, and Cr powders were deposited on CFC (CFC NB31, Snecma Propulsion Solide, France) by the slurry technique: details related to the process can be found elsewhere.132-134 Cr-carbide-modified CFC appears to have yielded the best results (Figure 39). A 15-pm-thick carbide (Cr23C6, Cr7C3) layer has been identified by XRD on CFC; the CTE of the carbides lies between that of CFC and copper (reported above) (CTE of Cr7C3 is 10 x 10-6K-1).135

A commercial brazing alloy Gemco® (87.75 wt% Cu, 12 wt% Ge and 0.25 wt% Ni; Wesgo Metals) has been used to braze Cr-modified CFC to Cu and

(c)

Figure 37 (a and b) Laser structuring of carbon fiber composite (CFC) in Active Metal Casting (AMC®) process and cross-section of the AMC® CFC-Cu joint. Courtesy of Chevet, G. Ph. D. Thesis 2010, University of Bordeaux, France.

to CuCrZr in a single step process,140 which is an advantage in comparison with other joining technol­ogies that require two steps: first joining CFC to Cu, and then CFC-Cu to CuCrZr. Flat-tile (a) and monoblock (b) mock-ups have been obtained by this technique and tested (Figure 40).

ENEA, Italy, has manufactured several actively cooled mock-ups of flat tile and monoblock type, by using different technologies; a new process (pat­ented) for the production of monoblocks is based on prebrazed casting and hot radial pressing (PBC+HRP) (Figure 41). The CFC surface modification is obtained by a titanium-copper-nickel commercial brazing alloy, which is followed by a Cu casting, then a radial diffusion bonding between the cooling tube and the CFC by pressurizing only the internal tube and keeping the joining zone in vacuum at the

120,141,142

required bonding temperature.

Complete manufacture and testing of this vertical target medium-scale mock-up (Figure 41) can be

(a)

31 mm

38 mm

(b)

Figure 39 Carbon fiber composite-Cu active brazing with Cu-Ti alloy and dispersed carbon fibers in the braze (b). Reproduced from Schedler, B.; Huber, T.; Eidenberger, E.; Scheu, C.; Pippan, R.; Clemens, H. Fusion Eng. Des.

2007, 82(15-24), 1786-1792. Reproduced from Ferraris, M., et al. In Ceramic Integration and Joining Technologies: From Macro to Nanoscale, 1st ed.; Singh, M., Ohji, T., Asthana, R., Mathur, S., Eds.; Wiley, 2011; Chapter 3,

© 2010 The American Ceramic Society.

considered as a success for both PBC and HRP pro­cesses, which can be an alternative to current techniques.

Several joining techniques are based on active braz­ing, which do not require any manner of CFC surface

modification. In this case, the active brazing alloys for CFC include elements such as Ti, Zr, Cr, and Si, which allow CFC wettability by the molten brazing alloy. A drawback of active brazing can be that active elements may form brittle intermetallics or com­pounds of low melting point. In one study, a TiCuNi brazing alloy produced by Wesgo Metals in the form of sheets (70Ti-15Cu-15Ni) was used to join CFC and silicon-doped CFC to pure copper, but the presence of Ni and Ti brittle intermetallics at the joint interface had a detrimental effect on thermal fatigue resistance tests of the joined component.

In Japan, a brazing technology was developed for ITER143 to join CFC to Cu by a NiCrP brazing alloy, followed by the joining of the CFC-Cu to the CuCrZr by low temperature HIP (500 °C). Recently,

a new brazing process was developed, based on the NiCuMn alloy, after metallization of the CFC surface. Several other brazing alloys have been devel­oped for CFC-Cu joints: Ag-based (63Ag-35Cu-2Ti, 59Ag-27Cu-13In-1Ti), or Cu-based brazing alloys (Cu-3Si-2Al-2Ti; Cu-Mn; Cu—Ti). Ag was discarded in view of nuclear transmutation-related issues.144

Depending on the actual energy flux on the Be PFCs in ITER during ELMs, melt damage may or may not occur. For Type I ELMs, which are compatible with the ITER divertor lifetime (-10 MJ ELMs16,18), the expected energy flux on the main chamber in ITER will be in the range of 2-3 MJ. The area of the wall over which this flux will be distributed is —30-60 m2, for a toroidally symmetric energy deposition. This leads to ELM energy fluxes —0.02-0.08 MJ m~2 on the main chamber wall, which will cause no Be melting at all. If toroidal asymmetries and/or poloi — dal structures dominate the ELM energy deposi­tion on the first wall, a substantial reduction of the first-wall effective area for energy deposition is expected (by a factor of —5). In this case, the ELM energy fluxes on the first wall would be 0.1-0.4MJm~2, which can cause up to 18 pm of
melting, lasting ^300 ms.209 Figure 21 shows the results of an analysis carried out with the code described in Raffray and Federici.21 , The erosion lifetime, expressed in number of ELMs or corresponding ITER full power pulses (approxi­mately 700 ELMs/pulses for a Be target initially 10 mm thick) is found to sharply decrease above a certain ELM energy threshold. Depending on the duration of the ELM event, the threshold energy density varies between 0.2 and 0.7 MJ m~2. For com­parison, the results of a W wall are also shown. More recently, analysis has been carried out using more sophisticated modeling tools and the results are described elsewhere.212

From the JET Be divertor experience, we expect that only a very small part of the melt layer produced during each ELM will be mobilized (typically <5%) and may lead to a Be influx into the plasma. Larger ELM energy fluxes onto the Be wall in ITER are indeed possible and would lead to serious pro­blems for the use of Be as main plasma PFC in ITER, both because of lifetime issues and because of plasma contamination. However, for the argu­ments explained above, a regime with repetitive ELM energy loads which are not compatible with the lifetime of the Be main chamber wall in ITER is not compatible with the ITER divertor lifetime either and will not be the reference regime of ITER operation. The development of techniques that can either eliminate or greatly reduce ELM energy losses
without significantly degrading confinement have therefore been recognized to be critically important for successful operation of ITER and have stimulated further worldwide research on ELMs mitigation.213

Development of an insulator coating is the criti­cal feasibility issue for self-cooled Li blankets. A very limited number of ceramics are stable under long-term exposure to Li at high temperatures. The present candidates are Er2O3, Y2O3, and, under lim­ited conditions, AlN. In addition to concept verifica­tion studies using several physical coatings, chemical or reactive coatings have been explored as a potential means to cover large and complex surfaces.

Considering the very low tolerable crack fraction (<10~6), in situ healing and two-layer coating with metallic overlayer are promising candidates. Recent loop tests with high impurity control have demon­strated that a two-layer coating with a vanadium overlayer is stable in flowing lithium. Thus, verifica­tion of in situ coatings, including the healing function, and two-layer coatings for large and complex surfaces would be the next necessary step in development.

Er2O3 is also of interest as a candidate for the required tritium permeation barrier coating.43 Thus, collaborative effort to develop this material for appli­cation in fusion blankets, either for electrical insula­tion or for tritium permeation reduction, seems to be an efficient development strategy.

Studies on the effects of radiation on the coating (resistivity and mechanical properties) are limited and
further research is necessary. In particular, the perma­nent effects of radiation need to be assessed by con­trolled irradiation experiments. In addition to the use of fission neutrons and charged particles, the opportu­nity to use the International Thermonuclear Experi­mental Reactor-Test Blanket Module (ITER-TBM) and International Fusion Materials Irradiation Facility (IFMIF) is anticipated for integrated coating function tests and high fluence 14MeV neutron irradiation tests, respectively.

In-service inspection programs have the primary goal of ensuring that the NPP structures have sufficient structural margins to continue to perform in a reli­able and safe manner.86,87 A secondary goal is to identify environmental stressors or aging factor effects before they reach sufficient intensity to poten­tially degrade structural components. Routine obser­vation, general visual inspections, leakage-rate tests, and destructive and nondestructive examinations are techniques available to identify areas of NPPs that have experienced degradation.

Determination of the existing performance char­acteristics and extent and causes of any observed distress is accomplished through a structural condi­tion assessment that routinely initiates with a general visual inspection to identify suspect areas followed by application of destructive or nondestructive exami­nations to quantify the extent and significance of any observed degradation. Basic components of a condi­tion assessment include (1) a review of ‘as-built’ drawings and other information pertaining to the original design and construction so that information, such as accessibility and position and orientation of embedded steel reinforcing and plates in concrete, is known prior to the site visit; (2) detailed visual exam­ination of structure to document easily obtained information on instances that can result from or lead to structural distress (e. g., crack mapping); (3) determination of the need for additional surveys or application of destructive or nondestructive test­ing methods; (4) analysis of results; and (5) prepara­tion of a report presenting conclusions and recommendations. More detailed information on guidelines on conduct of surveys of existing civil engineering buildings is available.88-91

Some general guidance on assessment of NPP deg­radation is also available.92-95 However, NPP rein­forced concrete structures present special challenges for development ofacceptance criteria because oftheir massive size, limited accessibility in certain areas, sto­chastic nature of past and future loads, randomness in strength, uncertainty in material changes due to aging and possibly degradation, and somewhat qualitative nature of some nondestructive evaluation techniques. Improved guidelines and criteria to aid in the inter­pretation of condition assessment results, including development of probability-based degradation accep­tance limits, are required. (Some information on probability-based crack acceptance limits for beams and shear walls considering loss of steel area and concrete spalling is available.96)

The first ceramic breeder blanket concepts used blocks or plates; later, pellet designs emerged. The larger the breeder component, the greater the threat to its integrity because of cracking or fragmentation due to thermal stresses and irradiation-induced swelling and embrittlement. This is the major reason for favoring pebble-bed concepts.

4.15.3.2.1 Pellets or blocks

Pellet or block fabrication makes use of proven technologies in the ceramic industry. Pressing and sintering of ceramic powders is a widely used and cost-effective industrial process. Pellets and rectan­gular blocks can be manufactured up to some centi­meters in size with excellent material homogeneity and controlled density. Thus, LiAlO2, Li2ZrO3, and Li2TiO3 pellets meeting dimensional, microstruc­tural, and purity characteristics were produced by Pechiney in collaboration with CEA.52 Similar results were obtained by ENEA, SCK/CEN, UKAEA- Springfields, and US laboratories.65,73-75

Kapychev et a/.30 fabricated pellets of Li4SiO4, metasilicate (Li2SiO3), and aluminate (LiAlO2), with a diameter of about 10 mm and heights of 5, 10, and 14 mm.

4.15.3.2.2 Pebbles

For TBR considerations, the density of the pebbles should be high and enable a dense packing to achieve a high lithium density. Further comments on pebble shapes are given in a later section. The presently used or developed processes are as follows:

1. A melting-spraying process was used at KIT (formerly FZK), in collaboration with Schott Glas — werke, for the production of 0.25-0.63 mm Li4SiO4 and Li4SiO4-SiO2 pebbles76 (see Figure 11). After annealing, spherical pebbles of 95-96% of theoretical density (TD) exhibiting satisfactory mechanical strength were obtained. Long-term annealing experiments on various candidates ceramic breeder materials were performed by Piazza et a/.77 An alternative route avoiding use of carbonate and using hydroxide was developed by Knitter et a/.,78 with slightly lower density. The reference composition is Li4SiO4 + 2.5 wt SiO2, resulting in a two-phase Li4SiO4 + Li6Si2O7 struc­ture in as-melted condition, and Li4SiO4 + Li2SiO3 after heat treatment (see Figure 12).

A melting-dropping process was investigated by Tsuchiya et a/.79 in collaboration with Mitsubishi to produce 1-mm Li2O spheres.

2. Sol-gel type processes were developed at Japan Atomic Energy Agency (JAEA), with Nuclear Fuel Industries, to produce 1 mm Li2O and 1.6 mm Li2TiO3 pebbles80 (see Figure 13). Similarly, Muis and coworkers81 at Energy Research Centre of The Netherlands (ECN) produced 0.5-1.0mm Li2TiO3

pebbles. In these cases, the pebble densities were <80% TD. Further work led to pebbles with Li2TiO3 + 5% TiO2 composition82 (see Figure 14).

Wu et a/.83 started the development of a sol-gel type process for Li4SiO4 and achieved 75% of TD for 1.2 mm diameter pebbles.

3. A process consisting of extrusion, spheronization, and sintering has, for several years, been used by AECL to produce 1.2 mm LiAlO2, Li2ZrO3, and Li2TiO3 pebbles in collaboration with Cera­mics Kingston.3 Material densities are in the 80-90% TD range.

Similar process trials were made by Lulewicz and Roux53 at CEA, with Pechiney, to produce 1 mm Li2ZrO3 pebbles, and later work concerned Li2TiO384,193 (Figure 15).

4. An agglomeration-sintering process has been used by JAEA, in collaboration with Kawasaki Industries,

Figure 13 Manufacturing process for Li2TiO3 by sol-gel method at Japan Atomic Energy Agency. Reproduced from Tsuchiya, K.; Kawamura, H.; Uchida, M.; Casadio, S.; Alvani, C.; Ito, Y. Fusion Eng. Des. 2003, 69, 449-453.

ECN SEI 15.0kV x50 100ptm WD36mm

Figure 15 Scanning electron micrographs of L12T1O3 pebbles produced by an extrusion-spheronization method (see text).

for producing 1 mm Li2O, Li4SiO4, and Li2ZrO3 pebbles. Pebble densities in the 90% TD range were obtained.85 This process has also been inves­tigated at CEA for producing 1 mm Li2TiO3 peb­bles. Pebble density of 90% TD and good mechanical strength were obtained.84

5. Zhu et al. developed a wet process for fabrication of Li3TaO4 ceramic pebbles. Typical pebble dia­meters are about 0.7-1.0mm, and the density achieved is over 90% TD, with crush loads more than 40 N.56 X-ray diffraction (XRD) patterns showed 99% of p-Li3TaO4 and traces of LiTaO3.

The necessity to recycle ceramic breeders after service imposes specific requirements on pebble manufacturing technologies. This reprocessing aspect may become a significant driver in fusion power economics on a longer term. See Sections 4.15.7 and 4.15.8.8 for further discussion.

Zirconium alloys are described more fully in Chapter 2.07, Zirconium Alloys: Properties and Character­istics. They are used in fusion reactors partly because of their corrosion resistance in aqueous environments and low neutron cross-sections.157 However, zirco­nium readily forms embrittling hydride precipitates. Zirconium alloys oxidize and the surface ZrO2 may be an effective permeation barrier’ preventing both hydrogen release and formation of detrimental
hydrides. Andrieu eta/.158 demonstrated that the rate of tritium release of zircaloy-4 (Zry4) decreased sub­stantially upon oxide formation in tritiated water.

Zirconium has multiple phases at temperatures of interest: for example, a-, p-, and g-Zr coexist in equilibrium at 833 K. Most solubility and diffusivity studies have been conducted on the single-phase a-Zr generally at 773 K and below (Figure 16). Above this temperature, zirconium alloys dissolve up to 50 at.% hydrogen and this solubility decreases rapidly with decreasing temperature, causing hydride precipitates within the alloys. The solubility has been found to vary slightly with the alloying content. Yamanaka et a/.159 note that the solubility in the p-phase decreases with alloying additions, while the solubility in the a-phase increases with alloying additions.

The solubility of hydrogen in ZrO2, regardless of the crystal structure (10-4 to 10-5mol hydrogen per mol oxide), is much lower than in the base metal and is even lower than that in Al2O3. a-ZrO2 exhibits a solubility almost an order of magnitude lower than p-ZrO2.160

Greger et a/.161 have reviewed hydrogen diffusion in zirconium. The diffusivities reported in studies they cite and in others is plotted in Figure 17. At 623 K, the diffusivity of hydrogen in zirconium is 10 m s, while the diffusivity in

-19 -20 2 -1 158,163,165

ZrO2 is only 10 to 10 m s. ’’ Austin

106

‘as

105

CO

E о E,

>.

104

о w

1000

Temperature, 1000/T (K-1)

et al.163 were able to measure the diffusivity in both a — and р-phases by measuring the activity, due to tritium, in tomographic slices of samples. The diffu — sivity values do not have a very strong dependence on crystallographic orientation or on alloy composition.

some of these phases. However, these quantities appear to be relatively large for the zirconium-matrix material. Further, autoradiography shows depletion in some iron-rich precipitates and at 623 K, the diffusivity in ZrFe2 is 2.5 x 10-11 m2 s-1, slower than in bare zirconium.168 The permeability values through hydrides might be larger because of the high solubil­ity of hydrogen isotopes in the hydride phase. How­ever, the volume fraction of hydrides tends to be small and the activation energy has been shown to be independent of the presence of the hydride.169

Zirconium alloys that lack an oxide layer are not useful in hydrogen environments that exceed the solubility of hydrogen in zirconium, because of hydride formation. At relatively low use temperatures (<623 K), and in aqueous or otherwise oxidizing environments, zirconium oxide is able to grow and is an effective barrier against the permeation of hydro­gen. Above this temperature, the integrity of the oxide layer cannot be maintained and the effective perme­ation of hydrogen isotopes is increased substantially.

The impact of the material’s microstructure is related to the amount of intrinsic defects at which He can be trapped and therefore determines the amount of He-retention.2 0 SC tungsten contains fewer defects than powder metallurgically (PM) produced tungsten (grain boundaries) and plasma-sprayed tungsten (grain boundaries and porosity), which is directly related to the thickness of porous/spongy structures (porosity about 90%,256 see Figure 10) that form depending

on energy and temperature.257,258 However, investiga­tions at 1650 K have shown that at such high tempera­tures there exists no difference between SC-W and PM-W, even for ion energies as low as 25 eV. The main trap sites at such high temperatures are thermal vacan­cies while intrinsic defects play a minor role.149,251

With regard to the material’s lifetime under He-exposure, the migration of He-bubbles toward the surface and the formation of pores and porous/ spongy structures seem to prevent the rupture and exfoliation that can accompany blistering. This is important as the exfoliation of blisters creates dust,19 which limits the plasma performance. For an antici­pated flux of 2 x 1018 He+ m~2 s-1 at 850 °C in iner­tial confinement devices, this may lead to a removal of 20mmyear-1 from the wall.109 Therefore, one way to increase the material’s lifetime might be to operate it at higher temperature.253 Another approach would be to develop advanced microengineered materials that have typical feature sizes less than the classical helium migration distance (20 nm).109 However, bubbles, holes, and porous/spongy structures signifi­cantly influence the material’s performance by reducing its thermal conductivity in the near-surface layer. This might play an important role when deter­mining the erosion and melt formation under com­bined He-irradiation and transient thermal loads, which will be shortly addressed in Section 4.17.4.4.3.

4.19.3.2.1 Implantation

The use of beryllium as a plasma-facing material in tokamaks has prompted many experimental studies of retention and emission of hydrogen implanted into beryllium-like metals from ion-beams or plasmas. References and discussions of these studies can be found in reviews.82- 5 Here, we review those studies which are relevant to H retention in Be in a fusion plasma environment. This section is mainly excerpted from Federici et all Two basic parameters for under­standing H retention are the hydrogen diffusivity and
solubility. Studies of solubility and diffusivity are reviewed in Causey and Venhaus85 and Serra et al.86 Figures 487-90 and 587,91,92 show experimental values for hydrogen solubility and diffusivity in W and Be. For Be there are significant differences between results from various studies. These differences may be due to effects of traps and surface oxide layers. The presence of bulk traps tends to increase the measured values of solubility and to decrease the mea­sured values of diffusivity (see Federici et al.7), especially under conditions where the concentration

ІХІ0-7

ІХІ0-8

d” 1 xio-9 E,

■f ixio-10

=5

Q 10-11

10-12 10-13

of hydrogen in solution is smaller than the concen­tration of traps. For this reason, studies done on materials of higher purity and crystalline perfection, and at higher temperatures and with higher concen­trations of hydrogen in solution, tend to give more reliable results. The porosity and oxide inclusions present in beryllium produced by powder metallurgy are also likely to lead to inconsistent results in mea­surements of hydrogen solubility and diffusivity. In the Be experiments, the effects of traps were not characterized and may be dominant. One firm con­clusion is that the solubility of hydrogen is very low in both Be and W.

Many studies have been done on the retention and emission of H implanted into materials to provide data needed to predict H retention in fusion reactor environments. Figure 6 shows the retention of 1 keV deuterium implanted into Be at 300 K versus inci­dent fluence, measured by thermal desorption.9 D retention in Be was close to 100% at low fluences but saturated at high fluences. Earlier nuclear reac­tion analysis (NRA) measurements of D retained in Be within ~1 pm of the surface gave very similar results.94 This saturation behavior indicates that D implanted into Be at 300 K does not diffuse, but accumulates until it reaches a limiting concentration of ^0.3-0.4 D/Be within the implantation zone. At high fluences, the implanted zone becomes porous allowing additional implanted D to escape. This
saturation mechanism is confirmed by electron micros­copy, which shows bubbles and porosity in the implan­tation zone after high fluence H implantation.95 Saturation of retention by the same mechanism is observed for D implanted into stainless steel at 150 K where the D is not mobile.96 H retention in Be increases with increasing ion energy and decreases with increasing sample temperature.84,97 The retention of 1 keV deuterium implanted into W and Mo at 300 K98 is also shown for comparison in Figure 6.

Figure 784 shows retention of deuterium and tri­tium as a function of incident particle fluence from a set of high fluence experiments in which Be specimens were exposed to laboratory ion-beams (Idaho National Engineering and Environmental Laboratory (INEEL), University of Toronto Institute for Aerospace Studies (UTIAS)), linear plasma devices (Sandia National Laboratory (SNL)/Los Alamos National Laboratory — Tritium plasma experiment (LANL-TPE), University of California, San Diego-Plasma Interaction with Surface Components Experimental Station B (UCSD — PISCES-B)), a tokamak divertor plasma (DIII — D-DIMES), and a neutral beam (NB-JET). In some of these studies carbon deposition or formation of carbide or oxide surface layers occurred, which is likely to affect D retention. The figure shows the D retention in Be observed under a wide range of exposure conditions. The high fluence saturated con­centration tends to be lower at higher temperatures.

It must be noted that this phenomenon is very important because it implies that tritium inventories and permeation due to implantation in beryllium for ITER PFC applications should be significantly lower than was previously estimated using classical recombination-limited release at the plasma surface. A first attempt to model this saturation by allowing the recombination coefficient to become exponen­tially large as the mobile atom concentration near the plasma-facing surface approaches a critical value was made by Longhurst et al.99 For Be, calcula­tions suggest that the critical concentration is related to the yield strength using Sieverts’ law of solubility. On the basis of the results of these calculations, it can be concluded that the inventory of tritium in the beryllium first wall of a device such as ITER, because of implantation, diffusion, trapping, and neutron — induced transmutation, will be of the order of 100 g rather than the kilogram quantities estimated previ­ously,100,101 and most of that will result from neutron — induced transmutations in the Be itself and from trapping in neutron-induced traps. Current predic­tions of tritium inventory in ITER are briefly dis­cussed in Section 4.19.6.2.1.

Fusion neutrons will create vacancies and intersti­tials in plasma-facing materials. For metals at reactor wall temperatures, these defects will be mobile and will annihilate at sinks (e. g., surfaces or grain bound­aries), recombine, or agglomerate into defect clusters. Vacancy agglomeration may also lead to the formation
of voids. In beryllium, neutron-induced nuclear reac­tions produce helium and tritium, which may be trapped at defects or precipitate as gas bubbles. These defects, resulting from neutron irradiation, will increase the retention of hydrogen, by increasing the concentration of sites where diffusing hydrogen can precipitate as gas or become trapped as atoms. The effect of neutron irradiation on hydrogen reten­tion in metals is complex, but, in principle, this can be modeled, provided the material parameters are known, such as hydrogen diffusivity, solubility, trap binding energy, and defect microstructure produced by the neutron irradiation. For many metals, most of these parameters are known well enough to attempt such modeling. For beryllium, however, uncertainties in solubility, diffusivity, and trapping of hydrogen make such modeling of hydrogen retention difficult.

The problem of production of helium and tritium by nuclear transmutation in beryllium itself is dis­cussed in Section 4.19.4.4.5.

DS copper alloys contain a fine dispersion of nanometer-sized oxide particles such as alumina, zir — conia, hafnia, or chromia in the copper matrix, giving rise to high-strength and thermal stability of the alloys. This class of copper alloys can be manufactured by either conventional powder metallurgy or internal oxi­dation. Their properties strongly depend on the type, dimension, and volume fraction of the dispersed phase and processing techniques. Unlike PH copper alloys, the addition of finely dispersed oxide particles into the copper matrix can prevent recrystallization of the matrix and consequent softening even after exposure to temperatures approaching the melting point of the copper matrix. In addition, the oxide particles are insoluble in the solid state, and are essentially immune to coarsening because of their high melting point and high thermodynamic stability. This extends the useful temperature range of a DS alloy far beyond that possible for conventional PH alloys.

Several DS copper alloys have been evaluated for fusion applications, for example, GlidCop® Al15, Al25, Al60, and MAGT 0.2. Both GlidCop® and MAGT class alloys are strengthened by Al2O3 particles, produced by internal oxidation. GlidCop® Al25 and MAGT-0.2 have been studied extensively because of their balanced strength, thermal conduc­tivity, and ductility. GlidCop® Al25 (0.25 wt% Al) is produced by OMG America. CuAl25-IG is the ITER grade with the optimized fabrication process for improved ductility and reduced anisotropy. The microstructure of the CuAl25 alloy is characterized by elongated grain structure along the extrusion or rolling direction and a high density (average of 3.27 x 1022m~3) of dispersed Al2O3 particles with a mean diameter of 6-9 nm. The distribution of alumina particles can be highly heterogeneous, with some grains free of strengthening particles. A low number density of micron-size a-Al2O3 particles exists at grain boundaries. The density of dis­locations in the as-wrought condition can be as high as ~1.5 x 1015m~2.15-18,22-24 Most of the oxide particles in GlidCop alloys are triangular platelets with the remainder in the form of circular or irregular-shaped disks.25

MAGT 0.2 is a Russian alloy produced by SPEZS — PLAV Company. It contains 0.17% Al, 0.05% Hf, and 0.09% Ti in the form of oxide particles.25,26 GlidCop contains Al-oxide particles only, while in MAGT alloy, there are Al-, Ti-, and Hf-oxide particles, and mixed Al — and Ti-oxide particles. A majority of the oxide particles in MAGT 0.2 are spherical in shape with a small fraction in the form of circular disks, with an average particle size of 6 nm.25,26

A lower bound curve can be constructed to corre­spond to a lower limiting fracture probability, which normally is set to 5% or 2%, taking into account the uncertainty ofthe T0 determination. The uncertainty of determining T0 depends on the number of speci­mens used to establish T0. The uncertainty (A T0) is defined from a normal distribution with two vari­ables, the test temperature, and the number of speci­mens used for the T0 determination, as follows:

AT0 = — Z [19]

r

where b= 18-20 °C, depending on the value of T — T0, r is the number of valid test data used to determine T0, and Z is the confidence level (e. g., Z85% = 1.44).

The median KJc is used to determine the value of and the uncertainty of T0 according to ASTM E1921. When Kjc(med) is equal to or greater than 83 MPa Vm, b = 18 °C. Alternatively, b = 20 can be used for all values of Kjc(med) not less than the mini­mum of 58 MPa Vm.

KJc(0.02) = 24.1 + 290 exp{°.°19(T — T0(margin))}

[23]

Kfc(0.05) = 252 + 366 exp{0.019(T — T0(margin)) }

[24]

When the dataset consists of several test tempera­tures, the median KJc is obtained from the basic rela­tionship (eqn [18]) as follows:

Jmd) = — E f30 + 70 exp[0.019(T,- — T0)]} [25]

r i=1

where r is the number of valid test data.

The 2% lower bound curve is sometimes used as a criterion to determine if the material should be analyzed by taking into account possible material inhomogeneity. This further analysis can be done using the SINTAP procedure (discussed in Section 4.14.2.5), which ensures that a conservative lower bound definition is obtained regardless of possible low fracture toughness values.16,17 If there are values below the 2% curve, the data preferably should be analyzed using the not yet standardized multimodal procedure.17 An example of the 2% curve and the effect of the basic SINTAP analysis are shown in Figure 10.

4.14.2.1.3 Limits of applicability

The applied cleavage fracture model strictly applies only to the transition region of the material, although the model also includes a term to take into account the conditional probability of crack propagation. This term is needed because in and near the lower shelf, the fracture event is mainly controlled by crack
propagation and therefore cannot be correctly described by only the crack initiation term. The situation is complicated by the fact that the lower shelf is not always close to the assumed, theoretically estimated, constant value of 20 MPa Vm (although it usually is). It is also assumed (Section 4.14.1.3.2) that the crack initiators are randomly distributed and that no global interaction exists between the crack initiators. The material is also assumed to be macro — scopically homogeneous. The method applies for transgranular, cleavage fracture events, although it may be used, with caution, for intergranular-type fracture especially when the fracture event is pre­dominantly stress controlled (as is typical in the lower transition area). Figure 11 shows scanning electron microscope views of both transgranular cleavage and intergranular fracture mode surfaces. The scope of applicability and the limitations of the method are also discussed in Sections 4.14.2.5 and 4.14.3.1 and the application to structures in Section 4.14.3.2.