Various proposed graphite irradiation models, includ­ing the UKAEA irradiation creep models, require the pinning and structural terms to be separated. Origi­nally, the pinning and structure terms for Gilsocarbon were separated by ‘eye,’ which is subjective. Recently, more sophisticated statistical, pattern recognition and curve fitting methods have been used.60

Having separated out the pinning term, the struc­ture term is a function of irradiation temperature and dose, whereas the pinning term is only a function of temperature. The pinning term is assumed to not affect creep rate, whereas the structure term affects it. In some models, the structure term is also assumed to be a function of radiolytic weight loss and to be correlated to dimensional change.

Reinforced concrete structures almost from the time of construction can start to deteriorate in one form or the other as a result of exposure to the environment (e. g., temperature, moisture, and cyclic loadings).77 The rate of deterioration is dependent on the compo­nent’s structural design, materials selection, quality of construction, curing, and aggressiveness of envi­ronmental exposure. Termination of a component’s service life occurs when it no longer can meet its functional and performance requirements, it becomes obsolete, or the maintenance costs become excessive.

Primary mechanisms (factors) that, under unfavor­able conditions, can produce premature deterioration of reinforced concrete structures include those that impact either the concrete or steel reinforcing materi­als (i. e., mild steel reinforcement or post-tensioning systems). Degradation of the concrete can be caused by adverse performance of either its cement-paste matrix or aggregate materials under chemical or phys­ical attack. In practice, these processes may occur concurrently to reinforce each other. In nearly all physical and chemical processes influencing the dura­bility of concrete structures, dominant factors involved include the transport mechanisms within the pores and cracks and the presence of water. Chemical attack may occur in several forms: efflorescence or leaching; attack by sulfate, acids, or bases; delayed ettringite formation; and alkali-aggregate reactions. Physical attack involves the degradation of concrete due to external influences and generally involves cracking due to exceeding the tensile strength of concrete, or loss of surface material. Physical attack mechanisms for concrete include salt crystallization, freezing and thawing, thermal exposure/thermal cycling, abrasion/ erosion/cavitation, irradiation, fatigue or vibration, biological attack, and settlement. Degradation of mild steel reinforcing materials can occur as a result of corrosion, irradiation, elevated temperature, or fatigue effects, with corrosion being the most likely form of attack. Post-tensioning systems are susceptible to the same degradation mechanisms as mild steel reinforce­ment plus loss of prestressing force primarily due to tendon relaxation and concrete creep and shrinkage. Of these, corrosion and loss of prestressing force are the most pertinent from the perspective of NPP dura­bility. Additional information on durability of NPP reinforced concrete structures is available.34

In zirconium alloys, as in other metals, under irradia­tion both vacancies and SIAs (Frenkel pairs) are created within the cascade leading to an increase of the point-defect concentration with the irradiation dose. However, even at very low temperature, the Frenkel pair concentration saturates at values about 1% due to the mutual recombination of vacancies and SIAs.43 At higher temperatures, point defects migrate and can therefore disappear because of a large variety of defects/defects reactions. Three major mechanisms contribute to defect elimination: vacancy-SIA recombination, point-defect elimina­tion on defect sinks (dislocation, grain boundaries, free surface, etc.), and agglomeration in the form of vacancy dislocation loops and interstitial dislocation loops. It has to be noted that, because of the rapid migration of SIAs compared to the slow migration of vacancies, at steady state the vacancy concentration is
several orders of magnitude higher than the SIA concentration.

Because of the elimination of point defects on point-defect clusters, the clusters can grow under irradiation depending on their relative capture effi­ciency. In the case of cubic metals, since the relaxa­tion volume of SIAs is usually much larger than that of vacancies, edge dislocations eliminate SIAs with a higher efficiency than vacancies (positive bias toward SIAs). Assuming an isotropic diffusion of point defects, this phenomenon leads to a preferred absorp­tion of SIAs by dislocations, provided that there is another type of sink within the material. Because of this preferential absorption of SIAs, the intersti­tial loops tend to grow under irradiation and the vacancy loops tend to shrink.

However, in hcp zirconium, the point-defect diffusion is usually considered to be anisotropic although there is little experimental evidence of this phenomenon. From the experimental results, it is believed that vacancy migration is only slightly anisotropic but the SIA migration is believed to be significantly anisotropic, as shown by atomistic com­putations. This diffusional anisotropy difference (DAD) has a strong impact on capture efficiency of point defects by sinks.44 Indeed, assuming SIAs to have a higher mobility in the basal plane than along the (c) axis and that the vacancies have an isotropic diffusional behavior, it can be seen that grain bound­aries perpendicular to the basal plane absorb more SIAs than vacancies. On the other hand, grain bound­aries parallel to the basal plane absorb more vacancies

than SIAs. Similarly, a line dislocation parallel to the (c) axis absorbs more SIAs than vacancies and a line dislocation in the basal plane absorbs more vacancies than SIAs. As discussed by Woo,44 this geometrical effect due to the DAD can overwhelm the conven­tional bias caused by the point-defect/sink elastic interaction difference (EID). Thus, contrary to the implications of the conventional rate theory, edge dislocations in a-zirconium are not necessarily biased toward SIAs, and grain boundaries are no longer neutral sinks. As will be described in the following, this phenomenon can explain some anomalous irra­diation-induced microstructural features as well as the growth phenomenon of zirconium alloys.

Irradiation creep of austenitic steels can be envisaged as having four stages. These are the transient regime, the creep regime in the absence of swelling, swelling — enhanced creep, and creep disappearance. The first three contributions have traditionally been described using the following equation.1

2 = A[1 — exp(—dpa/t)] + B0 + DS.

a

The equivalent strain per unit equivalent stress, sometimes called the creep modulus B, is the sum of a transient contribution that saturates usually at a dpa or less, the creep compliance B0 in the absence of swelling, and stress-enhanced creep where the enhanced creep rate is proportional to the void swelling rate. Bubble swelling also accelerates irradi­ation creep, but the influence is expressed primarily in the early stages of creep.167

The coefficients A and t are empirical, experimen­tally determined constants that are very material- specific (composition, thermal-mechanical treatment and texture) and sometimes stress-state dependent. For many high-exposure applications, the transient can be ignored. Even more importantly, the transient term can be obscured if significant precipitation — related strains are developing concurrently. In gen­eral, the magnitude of the transient-induced strain increases with increasing stress, but the duration in time or dpa usually does not increase with stress.

The creep transient involves both the dose needed to establish equilibrium densities of point defects, but most importantly, it involves the dose required for establishment of the quasi-equilibrium dislocation density. The transient is most pronounced for cold — worked steels that start at higher dislocation density. As recombination-annihilation mechanisms reduce the dislocation density the instantaneous creep rate drops until the quasi-equilibrium dislocation density is reached and the steady-state creep rate B0 is estab­lished. For any austenitic steel it can be safely assumed that B0 is ~1 x 10—6 (MPa dpa)—1 and is effectively a ‘crystal constant’ similar to the 1% per dpa swelling rate of austenitic steels.

The creep compliance B0 is not only independent of composition but also starting state, dpa rate, and temperature over the range of reactor-relevant condi­tions. There appears to be one exception, however, in that the creep rate at temperatures somewhere below ^100 °C can increase significantly above B0, as shown at 60 °C by Grossbeck and Mansur for various austenitic steels.168,169 At very low temperatures, vacancies are relatively immobile and cannot cancel the climb of dislocations produced by more mobile interstitials. Examples of such behavior have been seen in other studies.170,171 A good example of acceler­ated creep at low temperatures is shown in Figure 66.

The swelling-creep coupling coefficient D was originally assumed also to be a crystal constant at ^0.6 x 10—2MPa,—1 but as discussed later, it was found that D declines with increasing swelling to approximately one-third of this value or even to zero, depending on the stress state, stress history, and swelling history. This decline is an expression of the fourth stage of irradiation creep, variously designated as creep cessation, creep disappearance, or creep damping.

One feature of irradiation creep that distinguishes it as different from thermal creep is that it varies with stress to power 1.0 rather than a higher power typical of thermal creep. This linearity is shown in Figure 67. The creep equation presented above also predicts

that the creep rate is proportional to dpa both before and after swelling begins.

As discussed later, some important characteristics of creep have been redefined in the past two decades, especially for the creep compliance. B0 is known to be generally independent of alloy composition, thermal-mechanical treatment, irradiation tempera­ture, and dpa rate, but swelling is known to be very sensitive to all of these variables. This means that the irradiation creep modulus B quickly assumes all of the parametric sensitivities of void swelling. When the swelling rate reaches only 0.017% per dpa the swelling-enhanced contribution equals the B0 contri­bution, effectively doubling the creep rate.

There are a number of consequences of the cou­pling between swelling, creep, and precipitation — related strains.

1. The onset of swelling can be detected by a jump in creep modulus B long before measurable swelling — induced changes in dimension can be detected, and often before microscopy confirms the pres­ence of voids.

2. Attempts to measure B0 in the presence of low and sometimes undetectable levels of voids or bubbles will lead to misleading values, usually higher than ~1 x 10-6 (MPa dpa)-1.

3. Any local stress gradient generated by a swelling gradient will be reduced to a very low level by a local gradient of creep exactly matched to that of swelling.

4. Attempts to measure B0 in the presence of precipitate-related strains will lead to mislead­ingly different values, either too large, too small, and even negative values.

5. Whenever the stress state is generated solely by swelling, the coupling between creep and swelling guarantees that the system cannot operate at a stress level higher than D-1 or 160 MPa.1,16

The effect of radiation damage on ferritic RPV steels46-50 has for some time been considered in terms of

• The formation of matrix damage (MD), that is, defect clusters and dislocation loops. It is well established that in low copper steels the shift in impact or yield strength properties depends on Vdose.

• The irradiation enhanced formation of copper- enriched clusters (CECs). (CEC are also referred to as CRPs (Cu-rich precipitates) as they were originally assumed to be strictly Cu-rich rather than simply Cu-enriched.) It has been demon­strated that, in many low-to-medium Ni steels and alloys, the yield strength change due to copper precipitation rises to a plateau value that is then unchanged by subsequent irradiation.

• The irradiation induced/enhanced grain boundary segregation of embrittling elements such as P.

The first two mechanisms contribute to embrittle­ment by increasing the steels’ hardness as illustrated in Figure 6(a). The third mechanism induces embrit­tlement without hardening. The latter mechanism is not necessarily found in all RPV steels under operating conditions. Indeed, for MnMoNi steels irradiated in surveillance schemes in Western LWRs, the observed embrittlement is associated with the first two mechanisms; that is, the total shift in the ductile to brittle transition, as measured at the Charpy 41J level, is

A T41J = A T41J(CRP) + A T41J(MD) [2]

where CRP = Cu-rich precipitate and MD = matrix damage, or equivalently the increases in yield strength, Affj, is given by

Asy = Asy(CRP) + Asy(MD) [3]

The first two mechanisms serve to harden the mate­rial and increase the yield strength sy, while the third mechanism causes a drop in the fracture strength, sF. The effects of these changes on the fracture behavior are illustrated in Figure 6(b), where the temperature dependence of the yield stress, sy, and the fracture stress, sF, are plotted. It can be seen that the effect of

irradiation in causing hardening or a change in sF is to cause a change in the transition temperature. In the figure, ATT is caused by an increase in sy, while segregation of P to grain boundaries can lower the fracture stress and result in a shift ATT2. If both mechanisms are operative, then a combined shift of ATT3 occurs.

It is important to note that Ni and Mn are known to strongly influence hardening in steels containing low levels of Cu and also CEC hardening in Cu-containing steels. In Cu-containing steels, satura­tion of the cluster hardening has been demonstrated in steels containing up to ^1wt% Ni. At steel Ni levels above ~1.5 wt% (and with Mn 1.2-1.7 wt%>), cluster hardening has not been observed to saturate. The precise Cu, Ni, and Mn levels at which the plateau is suppressed have not been fully char­acterized, and are the subject of current research. Similarly, the exact influence of Ni and Mn on the embrittlement of low Cu steels has not been fully established and is again a subject of ongoing research. (The term standard MnMoNi steels is used to refer to steels with typical Mn levels (<1.5wt%) and Ni levels < ~ 1-1.2 wt%. The limit on the level of Ni is the subject of ongoing debate
with some workers preferring a limit of 1.0 wt%, with others promoting a higher limit.) Indeed, as will be described in Section 4.05.6, there is concern about whether at high doses (typical of those achieved in plant with extended lives) there may be deviations from the simple framework established above.

Irradiation creep is defined as the difference in dimensional changes between a stressed and an unstressed sample irradiated under identical condi­tions. Irradiation creep is important for structural materials for nuclear services as it is a major contributor to the dimensional instability of irra­diated materials at temperatures where thermal creep is negligible. However, studies on irradiation creep of SiC(-based materials) have so far been very limited, although it is of high importance for the behavior of the SiC TRISO shell.

Price published the result of the irradiation creep study on CVD SiC in 19 7 7.59 In this work, elastically bent strip samples of CVD SiC were irra­diated in a fission reactor, and the steady-state creep compliance was estimated to be in the order of 10-38 (Padpam-2 (E> 0.18 MeV))-1 at 1053— 1403 K. Scholz and coworkers measured the transient creep deformation of SCS-6 CVD SiC-based fiber, which was torsionally loaded under penetrating pro­ton or deuteron beam irradiation.70-73 They reported several important observations including the linear stress and flux dependency of the tangential primary creep rate at 873 K, and the negative temperature dependence of primary creep strain at the same dose. Recently, Katoh etal. determined the bend stress relaxation (BSR) creep in Rohm and Haas CVD SiC and Hoya monocrystalline 3C-SiC during irradia­tion in HFIR andJMTR at 673-1353 K.74 The results reported for CVD SiC are summarized in Table 1.

In the BSR irradiation creep experiment by Katoh et al, the creep strain for CVD SiC exhibited a weak temperature dependence at <0.7 dpa in the ^673-^ 1303 K temperature range, whereas a major transition at higher doses likely exists between 1223 and 1353 K. Below 1223 K, the creep strain appeared highly nonlinear with neutron fluence because of the

early domination of the transient irradiation creep. The transient creep is speculatively caused by the rapid development of defect clusters and the structural relaxation of as-grown defects during early stages of irradiation damage accumulation. At 1 353 K, irradi­ation creep mechanisms, which are common to metals, are likely operating.

In metals, steady-state irradiation creep rates are generally proportional to the applied stress and neutron (or other projectiles) flux, f,75,76 and there­fore, irradiation creep compliance, B, has been con­veniently introduced75:

eic = s(Bf + DS)

where S is void swelling and D is a coefficient of swelling-creep coupling. Ignoring the swelling — creep coupling term (valid in the saturable swelling regime), preliminary estimations of the steady-state irradiation creep compliance of CVD SiC were given to be 2.7 ± 2.6 x 10-7 and 1.5 ± 0.8 x 10-6 (MPa dpa)-1 at ^873-^1223 K and 1353 K, respectively. If linear-averaged, creep compliances of 1-2 x 10-6 (MPa dpa)-1 were obtained for doses of 0.6-0.7 dpa at all temperatures. Monocrystalline 3C-SiC samples exhibited a significantly smaller transient creep strain by 0.7 dpa and a greater subsequent deformation when loaded along <011 > direction.

To better define the irradiation creep behavior of SiC and the underlying physical mechanisms, it will be essential to further examine the stress depen­dence, dose dependence, effect of crystallographic orientation, microstructures of the crept samples, and the coupling between irradiation creep and swelling.

In addition to cracking, defects such as lack of fusion between weld beads or the weld bead and the side­wall, variable penetration, or second-phase inclusions can degrade weld quality. Lack of fusion defects is a notable concern when welding high-alloy nickel — based materials, which have notably ‘sluggish’ weld pools and are difficult to wet and tie into adjacent material (Figure 14(a)).

Inclusion-type defects are another concern and can be grouped into at least two types: (1) those that result from alloying additions or slag and (2) those that form via reaction with the environment. An example of the first type is given in Figure 14(b), which shows an unmelted iron-niobium Laves phase that is an intentional alloying addition to the flux coating of a shielded metal arc electrode. While alloying in this manner is a cost-effective way to tailor the composition of the electrode, it can lead to brittle second phases that also affect the local composition. In this case, the Nb-rich Laves phase is a strong melting point suppressant which can lead to either solidification — or liquation-type cracking.

The second type of inclusion is generally oxide or nitride-type particles that form via reaction with air. The corrosion-resistant alloys used in nuclear power systems (i. e., Fe-based stainless and Ni-based alloys) are especially prone to oxide-type defects as the nature of their corrosion resistance depends on the formation of stable, tenacious oxide films. An example of an aluminum-titanium-rich eutectic — type oxide that formed in a poorly shielded Alloy 690 fusion weld is shown in Figure 14(c). Another consideration of this oxide formation is that whatever metallurgical effect these alloying elements have is lost if they oxidize prior to solidification (e. g., the grain nucleating effect of Ti(C, N)-type particles).

Recent research shows that control of oxygen is critical to the weld puddle flow and wetting in nickel — based filler metals.10 In practice, this often translates into careful wire drawing practice so as to minimize the extent of embedded oxides or wire drawing lubri­cants into the filler metal. Figure 15 shows variability in the bead contour and tie-in of two filler metals welded under identical conditions, which was later traced to wire cleanliness. Additionally, separate test­ing shows that ^100 wt ppm levels of oxides can have macroscopic detriment on the regular flow and con­tour of Ni-30Cr-type filler metals.10

4.11.7.1 Introduction

In carbon dioxide (CO2)-cooled reactors, two types of oxidation can occur. The first is thermal oxidation which is purely a chemical reaction between graph­ite and CO2. This reaction is endothermic and is negligible below about 625 °C and is not impor­tant up to 675 °C. The second is radiolytic oxidation that occurs when CO2 is decomposed by ionizing radiation (radiolysis) to form CO and an active oxi­dizing species, which attacks the graphite. Radiolytic oxidation occurs predominantly within the graphite open porosity.

4.11.7.2 Ionizing Radiation

Ionizing irradiation can be defined as that part of a radiation field capable ofionization (charge separation) in CO2 either directly or indirectly. This leads to the creation of reactive species, which may react with the carbon atoms at the surfaces (external and more importantly internal) of the graphite components.

4.11.7.2.1 Energy deposition

Historically, ‘energy deposition’ has been used for a surrogate for ionizing irradiation, most probably because it is easy to measure using calorimetry and can be estimated from the reactor power. Energy deposition, sometimes referred as ‘dose rate,’ in the units of W/g of graphite, is a measure of the total energy absorbed in the gas in unit time from the scattering of g-radiation and fast neutrons.

For a typical Magnox reactor, energy deposition is composed of approximately the following components:

• 36% from the neutrons

• 58% from the gamma

• 6% from the interaction of graphite atoms within the moderator

Of these, it is only the last two that directly con­tribute to ionization of the carbon dioxide gas, mainly through Compton scattering. These ratios will be slightly different in an AGR.

An assumption is made that the dose rate received by the graphite is the same as that absorbed by carbon dioxide within the pores of the graphite and that a fraction k of the fission energy from the fuel causes heating in the moderator. For a typical Magnox reac­tor, k is ^5.6% of the thermal power. The unit G_c is defined as the number of carbon atoms gasified by the oxidizing species produced by the absorption of 100 eV of energy in the CO2 contained within the graphite pores; G_c for pure CO2 = 3.

of corrosion with liquid sodium. Vanadium alloys attracted attention in the 1980s again for use in fusion reactors because of their ‘low activation’ properties. At present, vanadium alloys are considered as one of the three promising candidate low activation structural materials for fusion reactors with reduced activation ferritic/martensitic (RAFM) steels and SiC/SiC composites. Overviews of vanadium alloys for fusion reactor applications are available in the recent proceedings papers of ICFRM (International Conference on Fusion Reactor Materials).2-6 This chapter highlights the recent progress in the devel­opment of vanadium alloys mainly for application in fusion nuclear systems.

A heat treatment performed at a temperature higher than the irradiation temperature on vari­ous zirconium alloys results in a recovery of the radiation-induced hardening90,138 (Figure 16). This recovery can also be measured using microhardness tests.101,102,105,139-142 The recovery of the hardening is always associated with the recovery of the ductility and the fracture properties.138

Howe and Thomas90 have shown that in a cold — worked zirconium alloy most of the recovery occur­ring between 280 and 450 °C appears to be the annealing out of radiation damage rather than cold work. In the case of strongly cold-worked zirconium alloys such as SRA Zy-4, radiation hardening recov­ery is also observed. The hardness of the material can even become lower than the initial hardness of the SRA Zy-4(105) owing to the recovery of the disloca­tions, in addition to the recovery of the loops.

Some authors,101,140,143 on the basis of various experimental results, have suggested that there is an interaction between oxygen and irradiation — produced dislocation loops, which increases the dislocation-defect barrier interaction. During the recovery, this phenomenon can lead to an additional hardening, as shown by Snowden and Veevers.140

Several authors48,101,105,141,144,145 have shown that during a heat treatment performed on a RXA zirco­nium alloy, the (a) loop density strongly decreases and the loop size increases. This decrease of the obstacle density to dislocation motion has been clearly correlated to the decrease of the radiation — induced hardening.101,105

Concerning the nature of the loops, Kelly and Blake48 have studied 240 loops in a zirconium alloy sample heat-treated at 490 °C during 1 h after irradi­ation up to a fluence of 1.4 x 1024nm~2. These authors show that, although the initial microstructure is composed of both interstitial and vacancy loops in equal amount, after the heat treatment, two-thirds of the analyzed loops are vacancy loops and only one — third are interstitial loops. This implies that the interstitial loops undergo a more rapid recovery than the vacancy loops. These observations have been recently confirmed by Ribis et al.,105 who stud­ied the evolution of the proportion of the vacancy loops and interstitial loops with heat treatment for various temperatures. These authors have shown that after 960 h at 450 °C, only large vacancy loops in low density are observed.

In the literature, several mechanisms are proposed in order to explain the irradiation damage recovery. The most commonly agreed mechanism is based on bulk diffusion of vacancies during the recovery and their exchange between loops of various size.105,146-148 Indeed, the smaller vacancy loops emit vacancies that diffuse toward larger vacancy loops, which absorb more vacancies than they emit, leading to a growth of the larger loops at the expense of the smaller loops. On the other hand, interstitial loops always absorb vacancies whatever their size, since the vacancies are in supersaturation during the heat treatment, exp­laining the rapid disappearance of the interstitial

105,146

loops.