In addition to the impact of thermal loads and neu­trons, plasma wall interaction comprises the contact of the PFMs with hydrogen isotopes, the helium ash and impurities originating from eroded surfaces. The interaction of these different particles with the PFM leads to near-surface material modification and deg­radation in the nm and low-pm range. The extent of the damage depends on the energy of the particles, their fluence, the surface temperature of the PFM and the temperature gradients within the PFC during steady state and thermal shock loading. Furthermore, the material’s microstructure, composition, and pre­damage, for example, cracking, have an influence on the material’s performance. Knowledge of all these parameters, particularly with regard to operational conditions, is essential to determine the material’s lifetime due to erosion, the amount of dust formation, which influences the plasma performance, and the safety hazards due to H and He-retention.

4.17.4.4.1 He-irradiation

The incident energy of He ions impinging on the wall of fusion devices is expected to vary between eV and MeV. Energies in the eV range are representative for the majority of particles in a magnetic fusion device that have lost most of their energy due to the interaction with the plasma. In contrast, MeV He-ions, carrying more or less all of their initial energy, are typical for inertial confinement facilities. This broad range of energetic particles results in varying amounts of sputtered material (erosion rate

1.8 times higher than for deuterium242) and further interaction of the ions with the tungsten wall leads to additional material degradation. This starts with the formation of vacancies along the path of the ions and continues via vacancy clustering, bubble formation, blistering, and the formation of spongy structures. However, there exists a lower energy limit for He-penetration, which is related to the surface bar­rier potential that was theoretically calculated to be about 6 eV,243 and with slight variations this value was verified by experiments.149,244 For faster ions penetrating into the material, the generation of par­ticular defects depends on the ion energy, ion

fluence, and temperature.26,245,246 The correlation

between these parameters is summarized as follows:

The term chemically assisted physical sputtering refers to the transfer of energy from an incident particle to a molecule on the surface. The energy gained is sufficient to break any remaining bonds of the molecule to other atoms on the surface resulting in the release of the molecule, or a fragment of the molecule, from the surface. In the case of beryllium bombarded by deuterium plasma, the sputtering of beryllium deuteride was first recorded in JET71 dur­ing operation with a beryllium divertor plate. Since that time, a series of systematic investigations were performed in PISCES-B to quantify the magnitude of this erosion term.72,73

The results from PISCES-B show a surface tem­perature dependence of the sputtering rate72 of BeD molecules. The maximum in the BeD sputtering rate (at 175 °C) corresponds with the onset of thermal decomposition of BeD2 molecules73 from a standar­dized sample of BeD2 powder. Even at the maximum loss rate, the chemical sputtering remains smaller than the physical sputtering rate of beryllium atoms from the surface over the incident energy range examined (50-100 eV). Molecular dynamics simula­tions have predicted,74 and subsequent measure­ments have validated the prediction, that chemical sputtering can dominate physical sputtering of beryl­lium as the incident deuterium ion energy decreases below 50 eV.

A distinction should be made between chemical sputtering and chemically assisted physical sputtering. Chemical sputtering involves the formation and loss of volatile molecules from a surface. In the case of beryl­lium deuteride, the molecule decomposes into a deu­terium molecule and a beryllium atom before it becomes volatile, so at least to date there is no evidence for chemical sputtering of beryllium during deuterium particle bombardment. Documentation of the chemi­cally assisted physical sputtering of beryllium may be important for determining material migration patterns in confinement devices and the identification ofberyl — lium deuteride molecular formation in plasma- exposed surfaces may also help explain the hydrogenic retention properties of beryllium.

PH CuCrZr alloy is commercially available under several trade names, for example, Elbrodur® CuCrZr from KME Germany AG, Outokumpu Oy CuCrZr, Zollen CuCrZr, C18150®, Trefimetaux CuCrZr, MATTHEY 328® from Johnson Matthey Metals, and YZC® from Yamaha Co, Ltd. The chem­ical compositions of these alloys differ by a small amount, with Cr varying from 0.4 to 1.5% and Zr 0.03-0.25%. Low Cr content is to prevent the forma­tion of coarse Cr precipitates. The element, Zr,
improves the hardening by the enhancement of fine homogeneous precipitation and improves the ductility of the alloy by inhibiting intergranular fracture.8-10 CuCrZr-IG is the ITER grade with tighter specification for composition and heat treat­ment. CuCrZr alloys are available in different forms, for example, bars, tubes, wires, foils, sheets, and plates. Hot forming, brazing, and inert gas welding are applicable for component manufacturing.

CuCrZr alloys are used in the conventional aged condition. The reference ITER heat treatment in­cludes solution annealing at 980-1000 °C for 1 h, water quench, and aging at 450-480 °C for 2-4 h.1 Typical microstructure of the prime-aged CuCrZr is shown in Figure 1(a). The alloy contains an equiaxed grain structure and uniformly distributed fine Guinier-Preston (GP) zones exhibiting primarily black dot contrasts and a small number of precipitates with lobe-lobe contrast. The number density of precipitates is on the order 1022m~3, with a mean diameter of ^3 nm. A low density of micron-size Cr particles and grain boundary precipitate-free zones were also observed.12-18 CuCrZr is susceptible to overaging and recrystallization during prolonged exposure at elevated temperatures. Overaging of CuCrZr causes significant coarsening of grain struc­ture and fine precipitates. Li et al.14 reported a lower number density (~1.9 x 1022m~ ) of larger (~9nm in diameter) precipitates with a mixture of coherent and incoherent particles after CuCrZr was hot iso­static pressing (HIP) treated at 1040 °C for 2 h at 140 MPa followed by solutionizing at 980 °C for 0.5 h with a slow cooling rate of 50-80 °Cmin-1 between 980 and 500 °C, and final aging at 560 °C for 2h (Figure 1(b)). The average grain size was >500 pm in comparison with ^27 mm grain size in the prime-aged alloy.

4.14.2.1 Standard Test and Analysis Procedure (ASTM E1921)

4.14.2.1.1 Test specimens used

ASTM E1921 specifies three specimen types for testing: standard compact tension (C(T)), disk­shaped compact tension (DC(T)), and single-edge notched bend (SE(B)). The type shall be selected based on the form and dimensions of the end-product or component (plate, forging, bar, etc.). The C(T)

Kjc(med) (MPaVm, B = 25 mm)

60 80 100 120 140 160 180 200 220

0.27 0.24 0.21 0.18 0.15 0.12 0.09 0.06 0.03 0.00

-60 -50 -40 -30 -20-10 0 10 20 30 40 50 60
T-T0 f C)

Figure 7 The number of specimens required by ASTM E 1921 for T0 determination (n = 1/number of specimens) and how it depends on temperature and median KJc.

specimen is commonly used, however, in the nuclear power industry, the SE(B) specimen with a square cross-section is preferred since that is a form that can be used directly from the Charpy V-notch test speci­mens contained in most current RPV surveillance programs. The standard test method does not directly specify the minimum specimen size, but only an upper limit for the measuring capacity (Kjc(limit)) depending on the size of the ligament (b0 = W — a0) (see eqn [14]). A value exceeding this limit may be included in the analysis as censored data, even if the test did not occur by cleavage fracture initiation. Those values shall be lowered to the limit value and included as censored (with reduced weight as indicated in eqn [16]) data. It should be noted that censored values shall not be included in determining the minimum number of specimens required for the T0 determination (discussed in Section 4.14.2.1.3). When very small-size specimens are used, more specimens need to be tested for a valid analysis. The effect of test temperature (T — T0) or the corresponding median Kjc on the required minimum number of valid test values is given in Figure 7, where parameter n is a weight factor (ASTM E1921) representing the inverse of the number of specimens needed for overall validity. When using subsize specimens, the expected number of speci­mens required for a valid analysis is, however, larger. The estimated numbers of specimens needed with different size single-edge notched bend specimens are given in Table 1.15

Two SE(B) specimen configurations are specified: square (W = B) or rectangular (W = 2B) cross-sections.

Due to the same ligament size (b0), the W = 2Bgeom­etry has the same measuring capacity as the C(T) specimen with the same thickness. Commonly used configurations are, for example, in RPV surveillance programs, the Charpy-size (10 mm square cross­section) and the half Charpy-size (5 x 10 mm rect­angular cross-section), which also have the same measuring capacity (equal b0). For typical medium strength (quenched and tempered) structural steels, these specimens have been found to give almost identical results.15 Side grooving of specimens is optional (total side-grooved depth may not exceed 0.25B), but recommended to increase stress triaxial — ity near the specimen surfaces.15 The side-grooved half Charpy-size design, with dimensions, is shown in Figure 8. It should be noted that the total speci­men thickness (not net thickness) shall always be used in the Master Curve analysis independent of the side-grooving. Testing with miniature specimens are discussed in Section 4.14.2.2.

There are different types of models to describe the behavior of pebble beds (see, e. g., Reimann et a/.97): finite-element models based on continuum mechanics, also called finite-element modeling (FEM), and the so-called discrete-element models, DEM, for description of mechanics at micromecha­nics level (i. e., individual pebble-pebble interactions). The development of computational tools at these two different length scales (macro — and microscales) allows for a better description of the thermomecha­nics of the pebble beds.

Figure 28 Bi-axial pebble-bed deformation tests: experimental results and calculations with ABAQUS code system. Reproduced from Hermsmeyer, S.; Reimann, J. Fusion Eng. Des., 2002, 61-61, 367-373.

4.15.4.5.1 Continuum models

The macroscopic behavior of pebble beds is de­scribed by constitutive equations commonly used in soil mechanics, considering the granular material as a continuum material that can undergo reversible elastic deformations, inelastic volume compaction (consolidation), and pressure-dependent shear fail­ure. To account for these properties, different models have been developed, which are implemented in structural computational programs.91,115,117,120-125

Pebble-bed data (see sections above) had to be implemented, and user-defined subroutines had to be written, for example, for the thermal creep laws.115

The codes were first validated with fairly simple experiments.104,123,126,127 Figures 28 and 29 show examples of results from the biaxial experiment126 and the SCATOLA experiment for calculational benchmarking127 (Figure 29).

These codes were later validated with more com­plex mock-up experiments such as HELICA and HEXCALIBUR91,108,114,128,129 and, as outlined below, aimed to be validated by the pebble bed assembly (PBA) experiment.130 Figures 30(a)117 and 30(b)91 show a comparison between measured and calculated HELICA temperatures.

Figure 30 (a) Measured temperatures in the HELICAexperiment, reprinted from Dell’Orco, G.; di Maio, P. A.; Giamusso, R.; Tincani, A.; Vella, H. Fusion Eng. Des., 2007, 82, 2366-2374. (b) Calculated temperatures for the HELICA experiment, reprinted with courtesy from Gan, Y. Ph. D. Thesis, Universitat Karlsruhe, 2008.

1st cycle (PD=68.29%) 2nd cycle (PD=60.95%) 3rd cycle (PD=61.10%) Loading process Unloading process

6

30 25 20

ro CL u> c T5 ro _o "ro і

5

ГО CL

U)

4

~o ro _o "го I

Loading process

Unloading process

3

2

Axial strain (%)

(a)

Figure 31 Stress-strain behavior of granular materials in (a) a rectangular box under uniaxial compaction and (b) packing density effect. Reproduced from An, Z.; Ying, A.; Abdou, M. Fusion Eng. Des. 2007, 82, 2233-2238.

The final goal of these codes is to determine the thermomechanical behavior of pebble beds in TBMs for ITER and DEMO blanket modules. At present, the codes are set up for this task. Up to now, only small portions of a DEMO blanket have been modeled. In these calculations, maximum stresses (very loca­lized) of 6 and 2 MPa were obtained for the ceramic breeder and beryllium pebble beds, respectively. Because of thermal creep, these values were re­duced after 2 h to 75% and 25%, respectively, of the initial value.

G. Pintsuk

Forschungszentrum Julich, Julich, Germany © 2012 Elsevier Ltd. All rights reserved.

4.17.1 Introduction

Until the mid-1990s, only few fusion devices used high-Z elements in plasma-facing materials (PFMs).1 These devices either operated at high plasma cur­rents and high plasma densities such as Alcator C-Mod2 and Frascati tokamak upgrade (FTU)3,4 or used high-Z materials only as test limiters such as Tokamak EXperiment for Technology Oriented Research (TEXTOR).5-9

Since then, high Z refractory metals have been attracting growing interest as candidates for PFMs because of their resistance against erosion and the need for low erosion and stability against neutron irradiation.10 Considerable effort has been made to study the behavior of high Z impurities in the core and edge plasmas, erosion/redeposition processes at the limiter/divertor surfaces, hydrogen isotope retention, and on material development and testing. In particular, the modification of ASDEX-upgrade (AUG) into a fully tungsten machine,11-17 which was achieved in 2007, provided positive answers to critical questions on the reliability of tokamak operation with high-Z plasma-facing components (PFCs) and the compatibility with standard and advanced H-mode scenarios and with the available heating methods.10

Among the challenges, for tokamak devices, that still remain are the strong increase of the W source and W concentration resulting from ion cyclotron reso­nance heating (ICRH) and the need for rigorous mod­eling to support the extrapolation of current results to ITER conditions. Clearly, not all questions posed by ITER can be answered by AUG only. For example, the effects of material mixing with Be, the melt behav­ior under transients, or the change of the hydrogen retention due to damage by high-energy neutron irra — diation18 cannot be addressed in AUG. Answers to some of these issues may be provided by the ITER — like wall project in Joint European Torus (JET), which is installing a bulk tungsten component for the strike point and physical vapor deposition (PVD)-W-coated carbon fiber composite (CFC) tiles for the remaining parts of the divertor.19-21 The remaining questions have to be answered by dedicated experiments in other plasma devices or can only be assessed by mod­eling. However, the results obtained so far do not exclude the use of W in ITER as a standard PFM.10 Further investigations related to future fusion power plants such as demonstration fusion reactor (DEMO) have to focus on the minimization of plasma heat loads to the PFCs to increase their lifetime. In particular, transient heat loads caused by instabilities significantly decrease the operation domain ofPFCs, due to thermal stresses and consequent enhanced erosion.22 There­fore, it is also important to mitigate all instabilities, such as edge localized modes (ELMs), that cause sig­nificant plasma transient heat losses.23 Plasma scenar­ios need to be developed, such that the conditions for achieving the required fusion yield are maintained in steady state, while at the same time sustaining tolerable heat loads on the PFCs. The above-mentioned upgrades to the JET24 and AUG15 will allow further optimization of the plasma scenarios under these con­ditions, in particular with DEMO relevant tungsten PFCs.25 These investigations will show how the iden­tified deficiencies of W can be overcome or how they have to be dealt with.

In addition to the application of tungsten in ITER and in potential future tokamak devices such as DEMO,26-29 tungsten also became an interesting alter­native for the divertor of stellarators, for example, Advanced Reactor Innovation Evaluation Studies — Compact Stellerator (ARIES-CS),30 and as a first
wall material for inertial fusion devices.31 Due to similar demands on the PFMs during the operation of all these devices, similar problems have to be solved for each application.

The disruption, or collapse, of the fusion plasma causes a potentially intense thermal load to the PFC of all large fusion devices. As discussed later in Section 4.18.4, such events will cause very high thermal stresses and significant material erosion. As these events are transient in nature, the ability of the PFC to withstand the disruption depends on the material’s ability to conduct, and its ability to absorb the deposited heat, before reaching a temperature or stress limit. For comparative purposes, a disruption figure of merit takes this into account:

Su(Cp pK )1/2
a E

where su is the ultimate tensile strength, Cp the specific heat, and p the density.

Figure 4 reports this disruption figure of merit for the materials in Figure 3. Consistent with the results of the thermal FoMth, high-quality, high-thermal conductivity composites and fine-grained graphites perform better than standard and larger grained gra­phites, and exhibit an order of magnitude better FoMd than beryllium and tungsten. As discussed later in Section 4.18.4, the erosion of graphite and beryllium is very high and dictates the use of thick tiles in high flux areas. This is in contrast to tungsten, which has a relatively low erosion yield, potentially allowing an armor thickness of only a few millimeters. Because the FoMs are essentially calculated on a per unit tile thickness, comparing tungsten with graphite can be somewhat misleading. However, because graphite

and beryllium are erosion-limited, the FoMs and the melting temperatures are useful evaluation tools. While the sublimation temperature of graphite 3350 ° C) is comparable to the melting point of tungsten (^3400 ° C), it is clear that beryllium, which has a melting point of 1300 °C, is at a distinct disadvantage. Removal of beryllium, as well as other metallic PFCs, by melting has been seen in several large experiments.

Performance calculations for graphite and CFCs have been conducted in both laboratory test stands and operating tokamaks. Some experimental data generated using electron beam simulation are given in Figure 5. Here, the power is deposited by a ras — tered electron beam for approximately one second up to surface heat loads of 11 MW m~ . The samples were 2.5 x 2.5 cm tiles, 1 cm in thickness, facing the beam. Each sample had a large notch machined into
one edge (the highest stressed area) to serve as a stress intensifier. It was noted that, without the notch, the graphites did not crack. Figure 5 gives the maximum heat flux of which each material was tested and whether cracking of the tile occurred. The data indicate that CFC materials and graphites with a higher thermal conductivity and high density are superior. No cracking occurred in either the three composites studied, or the two FMI graphites, at the maximum power density applied. The superior per­formance of the composite materials agrees with the performance of CFCs in the large tokamaks such as TFTR and JT-60U. The reason for the superior performance of the CFCs and the graphites is most likely their low thermal expansion coefficient, high thermal conductivity, and high strength. In addition, the presence of the fibers in the CFCs may serve to

blunt and arrest cracks, thus increasing toughness. All monolithic graphites shown in Figure 5, with the exception of the two FMI-HDFG materials, cracked. It is interesting to note that this graphite possessed

the highest FoMd, even higher than that of the composites. However, strict correlation of improved performance with increased FoMd was not seen, although a loose correlation was noted. As pointed

Ф CL >>  О LL О О Ф CL Ф 0

12

Power density (kW cm-2)

Figure 5 The performance of several grades of graphite and graphite composites subject to thermal shock loading. Redrawn from Croessmann, C. D.; Gilbertson, N. B.; Watson, R. D.; Whitley, J. B. Fusion Technol. 1989, 127-135.

out by Watson,4 the CTE may be the most dominant property, with the lowest CTE graphites showing the best resistance to thermal shock.

Finally, it should be noted that there are many issues regarding the selection of carbon materials as PFCs other than simply their thermal shock behavior. The issues of radiation damage, erosion, and hydro­gen retention are the three leading issues/drawbacks to the use of graphite as a PFC and they are discussed in the following sections.

4.19.5.1 Joining Technologies and High Heat Flux Durability of the Be/Cu Joints

One of the most critical aspects of the design of the ITER first wall is the attachment of the Be tiles to the actively cooled copper alloy substrate. The primary threat to this attachment comes from the heat flux. However, secondary effects also arise from the fact that there will be thermal gradients as well as mechanical loads from disruptions, which will cause distortion. These effects cannot easily be simulated experimentally. It should, therefore, be the responsi­bility of the design to minimize these effects.

At the time of writing this chapter, the first wall of ITER was undergoing a major redesign. Because of the lack of more detailed information on the new design, most of the considerations presented here, albeit general, are on the basis of the design of actively cooled ITER first-wall panels, which consist of stainless steel tubes in a copper heat sink with 10 mm thick beryllium tiles bonded to the plasma­facing surface. This technology was deemed to be adequate for handling the previously assumed surface heat flux of 0.5 MW m, and could potentially with­stand 3 MW m~2(155) for a limited number of cycles. However, because of fatigue it would be incompatible with areas subjected to higher power densities during each pulse. In such cases, other options for the heat exhaust technology are being considered,156 using thinner Be tiles.

Dislocation channels are frequently observed during postirradiation deformation of copper and copper alloys.102,103 Greenfield and Wilsdorf104 were the first who observed an area free of irradiation defects in the middle of a slip-line cluster by TEM in a neutron-irradiated copper single crystal. Extensive studies were conducted to establish the correlation between the deformation behavior and the slip-line structure in neutron-irradiated copper single crys — tals.104-107 Sharp108-110 studied the deformation and dislocation channels in neutron-irradiated copper single crystals in detail, and established a direct cor­relation between the surface slip steps and dislocation channels. The channels are nearly free of irradiation — produced defects, and operate parallel to the primary {111} slip plane. The cleared channels are formed by cooperative localized motion of glide dislocations that interact with and annihilate the preexisting radi­ation defect clusters. The channel characteristics
have strong dependence on irradiation dose and test temperatures. The channel width decreases and the slip step height increases with increasing irradiation dose, and the channel width and the slip step height decrease with decreasing deformation temperature. Howe111 confirmed that the channel width, spacing, the slip step height, and the average shear per slip band increase with increasing test temperature in the temperature range of 4-473 K. The reduction in channel width was considered to be a consequence of impeded cross-slip.10 ,111

Dislocation channels were also observed in neutron-irradiated copper single crystals under cyclic straining. , 3 The width and average spacing of channels changed with the number of cycles, in contrast to channels formed during tensile strain­ing where the width and spacing of channels were constant over a large range of strains.108

Dislocation channels are formed in neutron — irradiated copper alloys as well. Sharp114 observed the channeling effect in three different copper alloys neutron irradiated at ambient temperature, that is, Cu-0.8% Co, Cu-Al2O3, and Cu-4% Al single crystals. The channel spacing in the copper alloys were 1.2—1.5 pm, about half that observed in neutron-irradiated copper single crystals (2.3 pm). The channel width in Cu-0.8% Co alloy is similar to that for irradiated copper crystal (0.16 pm), and the chan­nels have the uniform width along the length. The presence of the second-phase particles in Cu-0.8% Co alloy has little effect on channeling. In the DS Cu-Al2O3 alloy, the channels are wider (0.24 pm) and
more irregular in width. The channel width can vary by a factor of 2 within a few microns along the length of a channel. A high density of dislocations surrounding the particles within the channel was observed in Cu — Al2O3, indicating great difficulty of dislocations in bypassing the (nondeforming) second-phase particles. In the single-phase Cu-4% Al alloy, however, no dislocation channels were observed.

Edwards13,40,64,115 studied thoroughly the deformed microstructure in neutron-irradiated CuCrZr alloys, and compared with the deformation microstructure in neutron-irradiated OFHC-Cu (Figure 19). Disloca­tion channels were observed during postirradiation deformation of the CuCrZr alloy neutron irradiated to 0.2-0.3 dpa at 100 °C. Channels were formed even before the upper yield point, and continued through­out the tensile deformation process. Some channels are completely free of defect clusters, and others contain a sizeable population of defect clusters. The width of cleared channels varied between about 100 and 250 nm. The channel formation is more pro­nounced in a higher-dose specimen than in a lower — dose specimen. In comparison with OFHC-Cu, CuCrZr showed little difference in deformation mode and channel characteristics in terms of width and size. While the channels in the OFHC-Cu were free of defects and dislocation debris, the channels in the CuCrZr alloy contained a small fraction of defects and precipitates. When the irradiated CuCrZr was annealed and deformed, deformation occurs in a much more homogeneous fashion, and no well-defined channels were observed.

The formation of dislocation channels in pure copper was investigated by in situ straining experi­ments on ion-irradiated copper in an electron micro — scope.116,117 Postirradiation straining of the thin foils of polycrystalline copper irradiated with 200 keV Kr
ions to about 2 x 10-4 to 0.02 dpa at room tem­perature showed that defect-free channels nucleate at grain boundaries, or in the vicinity of cracks, sug­gesting that grain boundaries and crack tips are nucleation sites for channels.117 Cross-slips were found to be responsible for channel widening and defect removal within the channel. Edwards et al.(A studied the initiation and propagation of dis­location channels in neutron-irradiated OFHC-Cu (Figure 20) and CuCrZr alloy in an interrupted tensile test. TEM observations suggested that chan­nels are initiated at boundaries, large inclusions, or existing channels. Channels are formed by interac­tions of newly formed dislocations with irradiation defects on the glide plane. Once formed, the channels propagate rapidly in the grain interior until they intercept another boundary, interface, or channel. Despite significant efforts, the exact mechanism of channel formation and evolution still remains unre­solved, and a clear connection between the slip pro­cesses, dislocation channeling, and localized flow in neutron-irradiated metals is still lacking.

The use of a Master fracture toughness Curve is not a new concept. The ASME Boiler and Pressure Vessel Code, Section III, Appendix G, has used a lower bound static fracture toughness curve that is a ‘Mas­ter Curve’ indexed using the reference temperature

RTndt.35 Currently, the ASME KIc and KIR curves, indexed to the RTNDT of the material, describe the fracture toughness of the RPV and its lower bound variance with temperature. These curves were adopted in the early 1970s as a lower bound repre­sentation to a relatively small set of linear-elastic fracture toughness (KIc) and linear-elastic arrest toughness (KIa) values for 11 heats of RPV steel.36 The use of RTndt to normalize temperature was intended to account for the heat-to-heat differences in fracture toughness transition temperature, thereby collapsing the fracture toughness data onto a single curve. However, RTndt is not always successful in this regard, often providing a very conservative characterization of fracture toughness. RTndt is the material/heat-specific ASME Code-defined temperature per Section III, NB-2300,37 based on a combination of drop-weight nil-ductility transition temperature (NDT) and Charpy V-notch tests (>68J) for nonirradiated materials; or for irradiated materials, RTndt is the nonirradiated RTndt (IRT) plus the shift in the 41J Charpy V-notch temperature to account for irradiation (assumed to be ARTndt). Appendix A to Section XI of the ASME Code38 uses the same lower bound Master Curve as Section III,

Appendix G, for crack arrest toughness, and another Master Curve (again indexed using RTndt) for static initiation fracture toughness. This approach has been used now for over 30 years in the US nuclear industry and has been shown to be very reliable in that there have been no vessel failures, albeit very conservative for most materials.

Kim Wallin’s direct fracture toughness Master Curve provides a more complete representation of the material fracture toughness. The Master Curve reference temperature T0 can be used in an analogous manner as RTndT to index the position of the Master Curve. The obvious advantages of this approach are:

• The index temperature itself is based on measured fracture toughness rather than Charpy V-notch and drop-weight tests.

• The Wallin Master Curve has a well-described statistical shape that allows for better-defined direct use in either deterministic or probabilistic analyses.

• Direct measurement of irradiated fracture tough­ness eliminates the need to add a shift to an initial value for many applications and it provides, to some extent, the possibility to extrapolate outside the already characterized fluence area.

ASME Code Cases N-62926 and N-63128 were pub­lished in 1998 and utilize the ASTM E1921 test method for determining T0. These Code Cases per­mit the use of a Master Curve — based index tempera­ture (RTt0 = T0 + 19.4 °C) as an alternative to RTndT. Code Case N-629 is for Section XI applica­tions for both irradiated and nonirradiated RPV steels; Code Case N-631 is essentially the same Code Case, but it is for Section III design applica­tions for only nonirradiated RPV steels. These Code Cases allow the determination of RTt0 when T0 is measured. Application to RPV integrity requires the knowledge of uncertainties associated with the use of a measured RT t0 in place of RTndT. The use of Master Curve in the United States has been limited to a few examples for which the Nuclear Regulatory Commis­sion (NRC) has written a safety evaluation (SE):

• The first use was indirect in that Master Curve data were used to justify a lower value of non­irradiated RTndT for some Linde 80 welds rather than defining a value of RT t0 9; this application for the Zion RPVs was actually submitted and approved before ASTM E 1921-97 or the ASME Code Cases were finalized.

• The key approved use of Master Curve was for the Kewaunee RPV. The NRC did not accept the utility submittal approach,40 but modified it to reflect their interpretation of approximating the current Charpy shift-based regulatory approach.41 This interpretation resulted in the use of a deter­ministic Margin term that was larger than that used for the Charpy data application. However, there was still enough beneficial gain using the Master Curve approach for determining an irra­diated RT t0 (over the current regulatory approach) to allow the utility to move forward to replace steam generators and pursue license renewal.

• The most recent SE was issued for the plants that have vessels containing Linde 80 welds. The Bab­cock and Wilcox (B&W) fabricated welds were used in all of the B&W design vessels and some Westinghouse design vessels fabricated by B&W. The initial RTndT for these weld metals has always been uncertain since these weld metals tend to have low upper shelf levels that often fall below 68J after irradiation. Since 68J is part of the transition temperature definition for RTndT, these welds may be unduly affected by the Charpy 68J temperature requirements. Therefore, the B&W Owners Group developed a program to bet­ter define the initial nonirradiated RTndT using the Master Curve and the RT t0 approach in Code Cases N-629 and N-631. Their approach utilized Charpy V-notch testing to get the 41 J transition temperature change for assessing the effects of radiation embrittlement in the same manner as currently used for RTndT. Irradiated surveillance program materials were evaluated using the Mas­ter Curve to compare with the predictive method of initial RTt0 + AT41j. The methodology was accepted by the NRC but requires explicit margins to be applied.42