There are relatively few data in the literature concerning the postirradiation creep behavior of zir­conium alloys as pointed out by Peehs and Fleisch.149 Even in the thorough review by Franklin et a/.,134 very few results concerning the postirradiation creep are given. In the case of the SRA zirconium alloys142,150-155 or RXA Zy-2,142,156 several authors have shown that irradiation leads to a strong decrease of the creep rate (Figure 17). This phenomenon is attributed to the presence of a high density of irradi­ation defects that harden the material. However, according to Ito eta/.142 and Schaffler eta/.,1 2 irradi­ation does not seem to affect strongly the stress sensitivity coefficient of SRA Zy-4 (Zircaloy-4), at least for the high stress range. However, for low applied stress, Ito et a/.142 have shown that the stress sensitivity coefficient is lower after irradiation than before irradiation. They have also shown that irradi­ation has a weak effect on the creep activation energy of SRA Zy-4 for temperatures from 330 to 600 °C and for stresses from 77 to 384 MPa. Murty and Mahmood157 have suggested that the creep anisot­ropy of RXA Zy-2 is decreased by irradiation. According to these authors, this phenomenon is due to the activation of other slip systems than the pris­matic slip system after irradiation, such as the basal and the pyramidal slip systems.

Cappelaere et a/.154 and Limon and Lehmann15 have shown that for low applied stress, a ‘tertiary

creep’ occurs for SRA Zy-4, even though the creep strain level remains low. This phenomenon cannot be explained by the increase of the stress due to the thinning of the wall of the tube. This phenomenon is therefore interpreted as a result of the recovery of the irradiation defects during the creep test and also due to the beginning of the recrystallization that can occur for high-temperature creep tests. Tsai and Billone15 have come to the same conclusions by analyzing their own long-term creep tests. The recovery of irradiation loops during creep tests has been observed, using TEM, by several authors on SRA Zy-4(154) or RXA Zr-1% Nb-O alloy,124 but it is the recent work by Ribis etal.105 that gives the most detailed study of the microstructure evolution during creep tests of the above alloy. The microstructure is compared to that observed after postirradiation heat treatment or after creep of the nonirradiated mate­rial. In this study, it is clearly shown that in RXA zirconium alloys, the irradiation loops are progres­sively annealed during the creep test, as for a heat treatment without an applied stress, the magnitude of the recovery being similar in both cases. Moreover, these authors show that other mechanisms associated with the deformation occur. Indeed, it is noticed that for tests performed at 400 °C and for low applied stress (130 MPa), in addition to the recovery of loops, the microstructure observed after creep tests exhibits a high dislocation density, much higher than the dislocation density observed in the nonirradiated material deformed up to the same plastic strain. According to these authors, this phenomenon results

from the irradiation loops that act as obstacle to dislocation motion, especially in the prismatic planes, and limit their mean free path. This leads to an important multiplication of dislocations in order to accommodate the plastic strain. This high dislocation density can then lead to a significant hardening in addition to the hardening due to loops. This could explain that for long-term creep test performed at 400 °C under an applied stress of 130 MPa, although a significant recovery of the irradiation damage occurs, the creep strain remains limited. Additional hardening due to the high density of small p-Nb needles can also occur in the case of Zr-Nb alloys. For higher applied stress, higher than 200 MPa, these authors suggest that a sweeping of loops probably occurs. This mechanism is believed to be similar to the dislocation channeling mechanism that is observed for burst tests and tensile tests.113,124 This mechanism therefore allows the deformation of the material for high applied stress, despite the high loop density.

The initial microstructure of the steels evolves dur­ing service, due to high temperature and irradiation for prolonged times, leading to modification of defect structure and secondary phases. These changes harden the steel, leading to concomitant embrittle­ment, which is discussed below.

Generally, a large number of alloying elements, W, Nb, Mo, Ta, V, or Ti are dissolved into the matrix of ferritic steels, some of them being larger than the iron atom. This could lead to the expansion of the unit cell of ferrite, making an element say, chromium undersized, with a positive binding energy with iron self-intersti­tial. Such a situation would lead to enrichment of chromium near the sink-like grain boundary. The reverse could happen if the size of the alloying ele­ments happen to be smaller than iron.

The w, G and a phases are all enriched in Si and Ni — elements which are known to segregate to in­terfaces during irradiation. With the exception of G phase, all the other phases and the a0 phase are rich in Cr. In those ferritic steels, where Cr is depleted near voids and at other interfaces which act as point defect sinks, it follows that in steels containing higher than 11 or 12% Cr, the chromium enrichment within the matrix may lead to local con­centrations exceeding those (>14%) at which a0 forms thermally. Further, enrichment of Cr may also result from the partial dissolution of chromium rich precipitates such as M23C6 during irradiation. In addition, RIS of phosphorus can also lead to the formation of phosphides in some of the steels. The irradiation-induced point defect clusters and loops may also facilitate and enhance nucleation of these phases. Although the relatively soft 8-ferrite im­proves the ductility and toughness of the 12Cr steel, the fracture could be initiated at the M23C6 pre­cipitates on the 8-ferrite-martensite interface. The presence of 8-ferrite, extensive precipitation and radiation-induced growth of M23C6 precipitates and formation of the embrittling intermetallic phases in the 12Cr-1MoVW steel in the temperature range 573-773 K are together responsible37 for the relative change in impact behavior of 9Cr-1MoVNb and 12Cr-1MoVW between 323 and 673 K.

Irradiation-induced microstructural changes are the factors that govern the creep and embrittlement behavior, which therefore, has to be minimized using appropriate chemistry and structure.

The DDRs for MnMoNi steels presented in the last section provide convincing examples of the applica­tion of fundamental insight to the prediction of changes in mechanical properties of operating RPVs due to radiation damage. Mechanistic understanding is continually developing as research continues and more data are obtained. Advances may lead to modifications in the form, or the values of, the fitting parameters. The major topics are the following:

• The effect of flux

• The role of Ni, Mn, and Si

• The possibility of new mechanisms at fluences

beyond the range for which there are data in the

current surveillance databases

There are two aspects of the effect of flux: first, the prediction of embrittlement at low fluxes and second, improvements in the general description of the effect of flux on embrittlement. It was described in the previous section that recent BWR data from the SSP capsules have greatly expanded the available BWR data, leading to an improved shift model. Carter et at}2 pointed out that, although this pro­vides a better description of BWR plate data, the model still tends to underpredict the embrittlement of BWR welds for measured AT41j greater than 60 °C. This suggests that there may be further improvements necessary in the description of embrit­tlement in the low flux range. Indeed, there may be general improvements in the description of flux. Odette considers that there is a systematic flux effect in the range of 0.8-8 x 1011 n cm-2 s-1 E > 1 MeV in the IVAR database which is not predicted by the EONY model.30 Further analysis of the IVAR data­base may lead to improvements in the description of the flux dependence of embrittlement at both low (surveillance) fluxes and high (MTR) fluxes.

The DDRs for MnMoNi steels discussed in the previous section really apply to only steels with Ni < 1.3 wt%. High Ni welds have been used in a limited number of civil PWRs, notably VVER 1000 reactors. High Ni welds were selected because vessel designers wished to take benefit from the greater hardenability and superior SOL properties (com­pared to lower Ni steels). At present the response of Cu-containing high Ni steels to irradiation doses of <60mdpa can be understood in terms of the framework of the understanding developed for Cu-containing MnMoNi steels with <1.2 wt% Ni, that is, hardening from MD and solute-enriched clusters (see, e. g., the work of Williams et at}2 ). A notable difference is that there is little evidence for a plateau in hardening from CECs as the fluence increases; rather a continuous increase with fluence is observed (for fluences up to ^60 mdpa).

It is possible, however, that an additional high fluence embrittlement mechanism may operate in Ni and Mn-containing steels. Specifically, it has been suggested that at long times and or at high doses, Mn, Ni, and Si could form a new phase in RPV steel.47’130 This late-blooming phase (LBP) would produce an additional increment of hardening at high fluences, that is, late-onset embrittlement or anomalous hardening at high doses. If this is the case, then the DDRs described in the previous sec­tion may become nonconservative. Recently, there have been an increasing number of observations of MnNiSi clusters in irradiated low Cu steels (see, e. g., Auger et a/.131 and Soneda et a/.127), and there is intensive research aimed at establishing whether NiMnSi clusters represent segregation to small microstructural features (thereby lowering interfacial or strain energies) or represent precipitates of a distinct Ni-Si-Mn-enriched phase that is thermo­dynamically favored at RPV operating temperatures and RPV steel compositions. It is the latter possibility that gives rise to the possibility of a late-onset embrittlement.

Based on the aforementioned finding, cladding manufacturing tests were conducted using cold­rolling pilger mill in 12Cr-ODS ferritic steels with the limited Y2O3 content <0.25 mass % to induce the recrystallized structure, and their internal creep rup­ture properties were evaluated at 700 °C, not at 650 °C. The chemical composition of the manufac­tured cladding is listed in Table 1, where the speci­mens are denoted as F1 to F4. The four levels of Y2O3 contents were selected in the range of <0.25 mass %, and the titanium content ranged from 0.13 to 0.31 mass %. Cold rolling by a pilger mill was

Figure 19 Optical microstructure of the F1, F2, F3, and F4 specimens in the final claddings. Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.; Nakashima, H. J. Nucl. Sci. Technol. 2002, 39(8), 872-879.

repeated twice with a reduction ratio of about 50% per rolling. The intermediate heat-treatment to soften the cold-rolled cladding was performed at 1100 °C for 30 min, and the final heat-treatment was performed at 1150 °C for 0.5 h.

Figure 19 shows the optical microstructures of the manufactured claddings in the longitudinal and transverse directions.36 All of the specimens seem to be recrystallized. However, the extent of recrys­tallization depends on the yttria and titanium con­tents. In the transverse cross-section, the grain size becomes slightly finer with increasing yttria and titanium contents from F1 to F4. In the F4 speci­men, the elongated grains can be still seen and the aspect ratio in the longitudinal (L) and transverse (T) directions is large, whereas the aspect ratio of specimen F1 appears to be nearly unity. These findings show that F4 specimen with higher yttria and titanium contents did not attain the perfectly recrystallized grain structure by the annealing of 1150 °C for 0.5 h.

4.10.5.1 Wigner Energy

The release ofWigner energy (named after the phys­icist who first postulated its existence) was histori­cally the first problem of radiation damage in graphite to manifest itself. The lattice displacement processes previously described can cause an excess of energy in the graphite crystallites. The damage may comprise Frankel pairs or at lower temperatures the sp3 type bond previously discussed and observed by Urita et a/.20 When an interstitial carbon atom and a lattice vacancy recombine, or interplanar bonds are broken, their excess energy is given up as ‘stored energy.’ If sufficient damage has accumulated in the graphite, the release of this stored energy can result in a rapid rise in temperature. Stored energy accu­mulation was found to be particularly problematic in the early graphite-moderated reactors, which operated at relatively low temperatures. Figure 12 shows the rate of release of stored energy with

temperature, as a function of temperature, for graph­ite samples irradiated at 30 °C to low doses in the Hanford K reactor.32 The release curves are character­ized by a peak occurring at ^200 °C. This temperature has subsequently been associated with annealing of interplanar bonding involving interstitial atoms.20

In Figure 12, the release rate exceeds the specific heat and therefore, under adiabatic conditions, the graphite would rise sharply in temperature. For am­bient temperature irradiations it was found9 that the stored energy could attain values up to 2720Jg which ifreleased adiabatically would cause a temper­ature rise of some 1300 °C. A simple experiment,8 in which samples irradiated at 30 °C were placed in a furnace at 200 °C and their temperature monitored, showed that when the samples attained a temperature of ^70 °C their temperature suddenly increased to a maximum of about 400 °C and then returned to 200 °C. In order to limit the total amount of stored energy in the early graphite reactors, it became nec­essary to periodically anneal the graphite. The gra­phite’s temperature was raised sufficiently, by nuclear heating or the use of inserted electrical heaters, to ‘trigger’ the release of stored energy. The release then self-propagated slowly through the core, raising the graphite moderator temperature and thereby par­tially annealing the graphite core. Indeed, Arnold33 reports that it was during such a reactor anneal that the Windscale (UK) reactor accident occurred in 1957. Rappeneau eta/.34 report a second release peak at very high temperatures (^1400 °C). They studied the energy release up to temperatures of 1800 °C of graphites irradiated in the reactors BR2 (Mol, Belgium) and HFR (Petten, Netherlands) at doses between 1000 and 4000 MWdT-1 and at tempera­tures between 70 and 250 °C. At these low irradiation temperatures, there is little or no vacancy mobility, so the resultant defect structures can only involve interstitials. On postirradiation annealing to high tem­peratures, the immobile single vacancies become increasingly mobile and perhaps their elimination and the thermal destruction of complex interstitial clusters or distorted and twisted basal planes contrib­ute to the high-temperature stored energy peak.

The accumulation of stored energy in graphite is both dose and irradiation temperature dependent. With increasingly higher irradiation temperatures, the total amount of stored energy and its peak rate of release diminish, such that above an irradiation temperature of ^300 °C stored energy ceases to be a problem. Accounts of stored energy in graphite can be found elsewhere.1,8,29,32

Figure 27 gives Raman spectra for unirradiated and irradiated graphite as well as for baked carbon.

In the spectral range shown, there is a prominent G-peak at 1580 cm-1 associated with the basal plane bond stretching of V axis sp2 atoms. The D-peak at 1350 cm-1 is associated with the breathing mode of sp2 atoms and disordered carbon structure. The sec­ond D-peak at 2700 cm-1 is indicative of the crystal­line structure of the graphite.

In Figure 28 the normalized positions of the G — and D-peaks, and the ratio of the peak intensities are compared for various graphites (unirradiated and irradiated). HOPG is obviously the most ordered structure followed by the PGA needle coke graphite and then the medium grained graphite grades. The most disordered materials are the baked carbon (NBG-18 baked) followed by irradiated BEPO (a UK test reactor) graphite.

Fluence (1020 ncm-2 EDND)
c-axis

Figure 29 qualitatively demonstrates that there is a relationship between Raman spectra and crystal structure disorder. The higher the disorder, the higher are the D — and G-peak wave number and 1(D/G) ratio.

In Figure 29(a), the crystal length La has been calculated from the full width at half maximum (FWHM) using the method proposed by Tuinstra and Koenig.56 Both figures demonstrate that Raman can be used to quantify the disorder in the graphite structure, either as manufactured or due to irradiation.

4.13.2.1 Historical Perspective on Concrete Longevity

Concrete, originally based on lime that hardened by atmospheric carbonation, has been utilized as a con­struction material for several thousand years. Cement has been around for at least 12 My when reactions occurred between limestone and oil shale during spon­taneous combustion in Israel to form a natural deposit of cement compounds.8 The oldest known concrete is from Yugoslavia and is about 7600 years old.9 Gypsum mortars were used by the Egyptians to fabricate the Great Pyramid at Giza about 2500 BC. The Romans were the first to use hydraulic limes and discovered the benefits of pozzolans. The survival of several ancient concrete structures (e. g., Pantheon in Rome, Figure 1) attests to the durability that concrete can attain.

A detailed study involving an examination of sam­ples obtained from several ancient concrete struc­tures utilizing physical and chemical techniques concluded that these structures survived primarily because of careful selection of materials and con­struction, mild climatic conditions, and the lack of steel reinforcement.9 These structures, however, were not fabricated using current ‘hydraulic Portland cement,’ as it did not exist until about 1824. Some information, however, was presented in Mallinson and Davies9 relative to samples that were obtained for testing from several structures fabricated in the mid — to late 1800s. It was concluded that the durabil­ity of these structures was not only due to high cement contents but also due to the relatively slow cement-setting times and high construction quality. These Portland cements differ somewhat from the Portland cements used to fabricate NPP concrete structures in that the formulations have changed as well as the fineness of the cement. Also, modern con­cretes have incorporated admixtures to improve work­ability, modify hardening or setting characteristics, aid in curing, and enhance the performance or durability.

One of the most challenging engineering conse­quences of neutron irradiation is the development of dimensional instability, whereby a structural com­ponent can shrink or grow in volume and where it can be distorted in shape, often with both processes occurring at the same time. There are two major categories of such changes: conservative of volume and nonconservative of volume. A distinction can also be made between processes that distribute the resulting strains isotropically or anisotropically. Additionally, a further distinction can be made con­cerning whether the process to the first-order is stress-driven or not, or whether it is stress-sensitive to the second-order.

Depending on the crystal structure there are a variety of such distortion processes, some more prominent than others in a given crystal system. For austenitic stainless steels the phenomenon of radiation-induced growth (volume-conservative, an­isotropic distribution ofstrains in the absence ofstress) is not an issue, whereas for hexagonal close packed alloys based on zirconium and rhenium growth is often a dominant process.9,106 Austenitic steels also

Swelling (%) Swelling (%)

Figure 39 (Left) Correlation between ductility loss and swelling in several heats of irradiated Ti-modified steels in PHENIX. At ~5% swelling the total and uniform elongations converge and by ~10% no ductility remains. (Right) Correlation of swelling and embrittlement in Charpy impact tests of cold-worked Ti-modified 316 steel irradiated in PHENIX.

Reproduced from Fissolo, A.; Cauvin, R.; Hugot, J. P. Levy, V. In Effects of Radiation on Materials: 14th International Symposium; STP 1046; 1990; Vol. II; pp 700-713.

are not very prone to significant transmutation — induced changes in lattice parameter as sometimes observed in alloys based on rhenium and vana­dium.106107 See also Chapter 4.01, Radiation Effects in Zirconium Alloys.

Stainless steels experience three general cate­gories of radiation-induced strain processes. These are precipitation-related strains, void swelling, and
irradiation creep. In general, these three processes are not fully independent but are interrelated and often synergistic.

The effects of fast neutron irradiation on the tensile properties of several precipitation-hardened nickel — based alloys were investigated in the 1970s and 1980s.

Figure 11 Dark field, transmission electron micrograph, illustrating the distribution of g precipitates in solution treated and aged Nimonic PE16 irradiated in Experimental Breeder Reactor-II to 69 dpa at 570 °C. Unpublished data from Boothby, R. M. The Microstructure of EBR-II Irradiated Nimonic PE16; AEA TRS 2002 (FPSG/P(90)23), with permission from AEA Technology Plc.

The materials examined included a number of g/g’ — hardened alloys, such as the Inconel alloys 706 and 718 and the developmental alloys D68 and 7818, as well as g’-hardened alloys similar to Nimonic PE16. Earlier work by Broomfield eta/.82 on thermal reactor irradiated materials indicated that PE16 was more susceptible to irradiation embrittlement at elevated test temperatures than austenitic steels. Broomfield83 found that thermal neutron irradiated PE16 was most severely embrittled in low strain tests at ^550-650 °C, and attributed this to an increased tendency for intergranular failure arising from the effects of helium generated from the 10B(n, a)7Li reaction. Boron itself is considered to have a benefi­cial effect on (unirradiated) creep rupture life, as it segregates to grain boundaries and inhibits intergran­ular cracking, and additions of a few 10s of ppm are therefore, generally made to nickel-based alloys, including PE16.84 Nickel is also a major source of helium in neutron-irradiated alloys, with the two — stage 58Ni(n, g)59Ni(n, a)56Fe reaction becoming the dominant source at high thermal neutron fluences, and nickel threshold reactions accounting for the greater part of helium production in fast neutron spectra.85 For example, the rate of helium generation in fast reactor irradiated PE16 was estimated by Boothby28 to be ~1.2appm per dpa, with about 85% of the helium being generated from nickel threshold reactions (see also Chapter 1.06, The Effects of Helium in Irradiated Structural Alloys). Nevertheless, other factors, including irradiation — induced strengthening and grain boundary segrega­tion and precipitation effects, have been implicated in the embrittlement of fast neutron irradiated nickel-based alloys.

4.06.5.1 Introduction and Irradiated Properties Database for W and W Alloys

Despite the recurring interest in the use of tungsten as a structural material for very-high-temperature applications and for use as plasma facing components in fusion devices, the database on irradiated properties is very limited and based primarily on fast neutron irradiation experiments. Similar to all bcc materials, tungsten is susceptible to low-temperature embrit­tlement at T< 0.3 Tm (Tm = 3695 K) for fluences >1 x 1024nm~2, which makes this material even more limited in toughness and ductility.

Improvements in the unirradiated mechanical properties of tungsten are observed with the addi­tion of Re, which is found to increase ductility at elevated temperatures125 as well as fracture tough — ness126 through the reduction in the DBTT. How­ever, as discussed in this section, the gains in performance through added Re content do not nec­essarily hold for irradiated materials.

4.06.5.2 Irradiation-Induced Swelling and Physical Property Changes in W and

W Alloys

Early investigations into the behavior of irradiated tungsten59,125,127-129 examined the defect formation and recovery defects. Much of this initial work was through the examination of electrical resistivity fol­lowing irradiation. Increases in electrical resistivity of pure annealed tungsten of up to 24% following irradiation to 2 x 1022ncm~2 in a fast reactor and ~14% in a mixed spectrum reactor to 1021ncm~ have been reported.59

Keys etal.127,128 examined the recovery of neutron — irradiated tungsten through isochronal resistivity studies following irradiation at 343 K to dose levels of 1.5 x 1021ncm~2 (E > 1 MeV). The beginning of saturation in resistivity observed in their studies appears just below 102°ncm~2 and correlates with the work by Lacefield et a/.130 on the appearance of defect clusters by 2.4 x 1019ncm~ identified through TEM examination. The work by Keys et a/.127,128 identified distinct stages of recovery such as self-interstitial migration occurring near 0.15 Tm, with a less pronounced recovery at 0.22 Tm through divacancy and impurity migration, which was followed by vacancy migration above 0.31 Tm. The residual resistivity, not recovered following anneals above 0.4 Tm after irradiation to fluences >3.3 x 1019 n cm~2, was due to the development of Re in the tungsten from transmutation reactions with thermal neutrons.

Very little data exist on irradiation-induced swell­ing in W and its alloys. Data on pure W are restricted to two reported series of experiments concerning the temperature dependence of swelling. The irradiation — induced swelling measured by Matolich eta/.131 and Wiffen19 using immersion density methods is shown in Figure 23. It should be noted that there is an order of magnitude difference in fluences between the two stud­ies. No other systematic examination of the swelling dependence on fluence and temperature is available.

Though swelling data for W-Re alloys are also lim­ited, work by Matolich eta/.131 for W-25Re irradiated to 5.5 x 1022ncm~2 revealed no significant amount of swelling. The data are also shown in Figure 23. Microstructural examination of W-Re alloys with concentrations of 5%, 11%, and 25%Re showed no cavity formation for fluences between 4.3

Figure 23 Irradiation-induced swelling measured through immersion density methods of W and W-25Re by Matolich etal.131 and Wiffen.19 and 6.1 x 1021 ncm-2 (E > 0.1 MeV) at temperatures between 873 and 1773 K,1 2 while void cavities have been experimentally observed in irradiated pure W over similar fluences.133,134 The effects of increas­ing Re or Os content in W were experimentally shown to decrease the density and radius of dislocation loops and voids in 0.15 dpa proton and neutron — irradiated material by He eta/.135 This reduction in size and number density is the result of the restricted mobility of the radiation-induced defects by the lat­tice dilations from the Re and Os solute.

The transmutation of W to Re and Re to Os during irradiation can have an effect on microstructural, physical, and mechanical properties of the material. The transmutation of the material, which results in the shifting of solute concentrations to higher levels, may result in precipitation in alloys that are nomi­nally in a single-phase region. One example of micro­structural and physical property changes because of irradiation is the decalibration of type-C (W—3%Re/ W-25%Re or W-5%Re/W-26%Re) thermocou­ples, such as that used in fuel element centerline temperature measurements. Reviews of early experi­mental work on W/Re thermocouples and on the dependence of decalibration on the neutron fluence have previously been discussed.136,137 While displa — cive neutron damage may result in material changes such as vacancy clusters or dislocation loops, the maximum theoretical changes expected in emf out­put of the thermocouple is ~1 p. V °C-1,138 whereas the changes associated with transmutation effects can result in more significant decreases. A -300°C drift in temperature following 6000 h irradiation under 2.7 x 10 n cm-2 thermal and 8 x 1021 n cm — fast fluence was reported for a W-3%Re/W-25%Re couple.139

These changes can also be significant in fast reac­tor irradiations. Experimental work by Williams eta/.132 showed that for a W-5%Re/W-25%Re ther­mocouple irradiated to 6.1 x 1021 n cm-2 fast fluence at 1173 K, the precipitation-induced changes in the Seebeck coefficient were -6.6 and -0.02 mV °C-1 for the 5 and 25% Re alloys, respectively. Calculated final compositions following 6.1 x 1021ncm — (14 MeV) irradiation of W—5(wt%)Re produce W-5.130Re-

0. 021Os-0.150Ta alloy, while a W-26Re alloy trans­mutes to W-25.955Re-0.107Os-0.117Ta.140

Postirradiation examination of the microstructure of the irradiated 5, 11, and 25% Re alloys in Williams eta/.132 revealed w-phase precipitation at irradiation temperatures above 1373 K, though unidentifiable precipitation was apparent in the alloys at 1173 K.

The development of the w-phase over the equilibrium а-phase in irradiated samples, but not in the unirradi­ated annealed samples, is the result of irradiation — induced solute segregation to defect sinks. The development of the w-phase was also reported in microstructural studies of W-26Re irradiated up to 11 dpa at temperatures between 646 and 1073 K.141

It should be pointed out that the change or tem­perature shift under irradiation is proportional to the degree of localized transmutation and local tempera­ture gradients and therefore dependent on the pro­files of the temperature and irradiation fields to which the thermocouple is exposed. Therefore, experimental work typically involves the irradiation of the entire cable, while in reactor applications, significant variations in temperature and fluence may result. The changes in thermoelectric power (D) as a function of irradiation fluence can be mod­eled by the following:140

D = 0 for 0 < f < 0.25 x 1021 n cm-2

D = 100[1 — e°’067(°’25-f)] for 0.25 < f < 1 x 1021 n cm-2

D = 100[1 — e°’104(°’52-f)] for f > 1 x 1021 n cm-2 [2]

Though significant radiation-induced decalibration may occur in fission reactors, this effect may not be readily observed in fusion reactors where the thermal flux is much lower. In addition, typical end-of-life estimates of total neutron fluence of <1021ncm-2 suggest that transmutation effects resulting in decali­bration is not an issue.