Extensive experimental investigations found3 that the ferritic steels, whose high temperature mechanical properties are far inferior to austenitic stainless steels, displayed excellent radiation resistance. The ferritic — martensitic steels (9-12% Cr) have, therefore, been chosen for clad and wrapper applications, in order to achieve the high burn-up of the fuel. This is based26-29 (Table 5) on their inherent low swelling behavior. The 9Cr-1Mo steel, modified 9Cr-1Mo (Grade 91), 9Cr-2Mo, and 12Cr-1MoVW (HT9) have low swelling rates at doses as high as 200 dpa. For example, HT9 shows 1% swelling at 693 K for 200 dpa. The threshold dose for swelling in ferritic steels is as high as nearly 200 dpa in contrast to 80 dpa for the present generation D9 austenitic stainless steel. It is established that the void swelling depends crucially on the structure of the matrix lattice, in which irradiation produces the excess defects.

Extensive basic studies have identified19,30-33 the following reasons as the origin of superior swelling resistance in ferritic steels:

1. The relaxation volume for interstitials, that is, the volume of the matrix in which distortion is introduced by creatingan interstitial, in bcc ferrite is larger19 than fcc austenite. For every interstitial introduced, the lattice distortion is high and hence the strain energy of the lattice. Hence, the bias toward attracting or accommodating interstitials in the bcc lattice is less. This leaves higher density of‘free’ interstitials in the bcc lattice than fcc lattice. As a result, recombina­tion probability with vacancies increases significantly and supersaturation of vacancy reduces. Conse­quently, the void nucleation and swelling is less.

2. The migration energy of vacancies in bcc iron is only 0.55 eV, against a high value in fcc austenite, 1.4 eV. Vacancies are more mobile in bcc than fcc, increasing the recombination probabilities in bcc ferrite. Another factor is the high binding energy between carbon and vacancy in bcc iron (0.85 eV), while it is only 0.36-0.41 eV in austenite. This leads19 to enhanced point defect recombina­tion in bcc than fcc, due to more trapping of vacancies by carbon or nitrogen.

3. In bcc iron, it is known30 that there is a strong inter­action between dislocations and interstitials solutes, forming atmospheres of solutes around dislocations.

The formation of ‘atmospheres’ around dislocations makes them more effective sinks for vacancies than interstitials, resulting in suppression of void growth,

V Ka і

(d)

Figure 5 Initial structure24 of normalized and tempered modified 9Cr-1Mo steel: (a) Monocarbides (MC) and M23C6 along lath boundaries in a carbon extraction replica of the sample and (b) Microdiffraction of fine particle marked B, confirming the crystal structure of MC. Energy dispersive analysis of X-rays (EDAX) identifying the MC particles (B) to be rich in

(c) V and (d) Nb.

provided the following conditions are satisfied: ‘atmo­spheres’ comprise of either oversized substitutional atoms or interstitials, dislocations have high binding energy with solutes, and concentration of solute atoms at the core of the dislocation exceeds a critical value. On the other hand, if‘atmosphere’ is made up of undersized atoms like Si or P, the voids can grow. The ‘atmosphere’ of interstitials reduces the dislocation bias for additional capture and inhibits dislocation climb, thus converting them to saturable sinks. Such a scenario would increase the recombination prob­abilities, suppressing the void growth.

These fundamental differences in the behavior of solutes and point defects in bcc lattice make ferritic steel far superior to austenitic steels, with respect to radiation damage.

The challenging task for materials scientists to use ferritic steels directly in fast reactor fuel assembly was with respect to enhancing the high temperature mechanical properties of the ferritic steels, especially high temperature creep life and irradiation creep resistance.

The first embrittlement correlation for the TTS of the Japanese RPV materials, JEAC 4201, was issued in 1991. Additional surveillance data have been com­piled since 1991 and in 2002 the Japanese electric power utilities started a project with CRIEPI to develop a new mechanistically based embrittlement correlation.126

Soneda and coworkers have adopted a two — step approach to developing a new correlation method.6 , ,7 In the first step, the microstructural

effects due to radiation damage are modeled, and the mechanical property changes engendered by such change are detailed. The microstructural changes, namely, the formation of solute atom clusters and MD features, due to irradiation are modeled using the following equations:

bulk chemical contents of Cu and Ni, DCu is the Cu diffusivity, f is the dose rate, t is the irradiation time, and tr is the relaxation time, respectively.

Equations [13 and 14] represent the time evolu­tion of solute atom clusters and the MD clusters, respectively (see Hiranumu eta/.126 for a full descrip­tion of the equations). In eqn [13], it is to be noted that solute atom clustering occurs with MD features as the nuclei. This process can occur without Cu atoms but is accelerated by their presence. In eqn [14], the formation of MD features is affected by the irradiation temperature and also the bulk Ni content. Equation [15] models the depletion of the matrix Cu content because of the formation and growth of Cu-enriched solute atom clusters. Note that the depletion of the matrix Cu reduces the formation rate of Cu-enriched solute atom clusters. Equation [19] gives an expression for the diffusivity of Cu atoms, which combines terms from both irradiation — induced vacancies and thermal vacancies.

Mechanical property changes are correlated with the microstructural changes using the following equations:

ATSC = Х1бРЦ~

= X1^Xf^+f PffiC [20]

f (CCT, Csc) = X11 CCu~ ^ + X12 [21]

CSC

g (CNi) = (1 + X13 (cNi)X14 У [22]

h(ft) = X9(1 + X10DSC)f t DSC ~ DCu [23]

ATmd = X17P Cmd [24]

AT = J (ATSc)2 + (ATmd)2 [25]

where A Tsc and ATmd are the contributions of solute atom clusters and MD features, which are calculated using eqns [20 and 24] as functions of Csc and CMD, respectively. In calculating the contri­bution ofsolute atom clusters, an empirical model, in which the TTS is proportional to the square root of the volume fraction of solute atom clusters, is used. The average volume per cluster, which is necessary for calculating the volume fraction, is modeled using eqns [21-23], which take into account the effect of chemical composition and the growth of the clusters during irradiation. The Greek characters in the above equations are coefficients, and were optimized using

the surveillance database of Japanese commercial reactors.126,127 The partitioning of the total embrit­tlement between that due to copper clusters and MD features is shown in Figure 17. It can be seen that the MD has a weak dependence on the Cu level of the steel.

Figure 18 shows a comparison between the calcu­lated and measured TTS. The standard deviation of the prediction error is smaller than that of the other correlation equations used in Japan and in the United States, as shown in Figure 10. When a plant-specific adjustment is applied to the initial transition temper­ature, the standard deviation of the prediction error becomes much smaller and is as low as 6 °C. A practical output of this approach is the develop­ment of a new embrittlement correlation method for Japanese RPV steels, and this method has been adopted in theJEAC 4201-2007. Thus, this study is a good example of how the understanding of a funda­mental mechanism can be applied in a real-world engineering application.

4.05.5.3 Summary

It is clear from the discussion above that there has been successful development of mechanistically based DDRs for both CMn and MnMoNi steels.

Different DDRs have been developed in different countries to describe the hardening and embrittle­ment of the various RPV steels. The inevitably approximate nature of the DDR expressions, the lim­ited variation of different parameters in each surveil­lance database, and the limited amount ofsurveillance data mean that the effects ofmany parameters must be implicit. Different irradiation and compositional var­iable ranges in different surveillance schemes may contribute significantly to the forms of the DDRs and the strength of different dependences. The lim­itations in the form of the DDRs and the R&D into outstanding issues are the subject of the next section.

Based on the aforementioned finding in ODS steels, the recrystallization processing was extensively stud­ied to change the substantially elongated grain struc­ture to the equi-axed grain structure. The Y2O3 content should be <0.25 mass % to attain the recrys­tallized structure in the ODS ferritic steels. The two types of 12Cr-ODS steels in the chemical composi­tion of 0.23Y2O3-2W-0.4Ti (A3) and 0.34Y2O3-3W- 0.4Ti (A15) were extruded at 1150 °C, followed by 60% cold rolling and annealed at 1200 °C for 1 h. From these heat-treated bars, the internal creep rup­ture specimens were machined and tested at 650 °C. Figure 1835 exhibits a comparison of the creep rupture strength of recrystallized (A3) and unrecrys — tallized (A15) 12Cr-ODS steels at 650 °C, where
von Mises’ equivalent stress was estimated for the internal hoop stress. The unrecrystallized specimen shows significant strength anisotropy in uni-axial and internal creep rupture strength, whereas the recrystal­lized specimens reveal decrease of anisotropy, where uni-axial creep rupture strength decreases and inter­nal strength approaches the uni-axial strength. These results demonstrate that the recrystallization process adequately improves the creep rupture strength in the internal hoop direction. Furthermore, softening by recrystallization makes it possible to manufacture cladding by cold-rolling processing.

The discovery of Fullerenes9 and carbon nanotubes,10 and other nanocarbon structures, has renewed inter­est in high-resolution microstructural studies of car­bon nanostructures and the defects within them.11 This in turn has given new insight to the nature of displacement damage and the deformation mechan­isms in irradiated graphite crystals. The binding energy of a carbon atom in the graphite lattice12 is about 7 eV. Impinging energetic particles such as fast neutrons, electrons, or ions can displace carbon atoms from their equilibrium positions. There have been many studies of the energy required to displace a carbon atom (£d), as reviewed by Kelly,13 Burchell,14 Banhart,11 and Telling and Heggie.15 The value of £d lies between 24 and 60 eV. The latter value has gained wide acceptance and use in displacement damage calculations, but a value of ^30 eV would be more appropriate. Moreover, as discussed by Banhart,1 Hehr et a/.,16 and Telling and Heggie,15 an angular dependence of the threshold energy for displacement would be expected. The value of £d in the crystallo­graphic c-axis is in the range 12-20 eV,11,17 while the in-plane value is much greater.

The primary atomic displacements, primary knock-on carbon atoms (PKAs), produced by ener­getic particle collisions produce further carbon atom displacements in a cascade effect. The cascade carbon atoms are referred to as secondary knock-on atoms (SKAs). The displaced SKAs tend to be clustered in small groups of 5-10 atoms and for most purposes it is satisfactory to treat the displacements as if they occur randomly. The total number of displaced car­bon atoms will depend upon the energy of the PKA, which is itself a function of the neutron energy spec­trum, and the neutron flux. Once displaced, the carbon atoms recoil through the graphite lattice, dis­placing other carbon atoms and leaving vacant lattice sites. However, not all of the carbon atoms remain displaced and the temperature of irradiation has a significant influence on the fate of the displaced atoms and lattice vacancies. The displaced carbon atoms easily diffuse between the graphite layer planes in two dimensions and a high proportion will recom­bine with lattice vacancies. Others will coalesce to form C2, C3, or C4 linear molecules. These in turn may form the nucleus of a dislocation loop — essentially a new graphite plane. Interstitial clusters may, on fur­ther irradiation, be destroyed by a fast neutron or carbon knock-on atom (irradiation annealing). Adja­cent lattice vacancies in the same graphitic layer are believed to collapse parallel to the layers, thereby forming sinks for other vacancies which are increas­ingly mobile above 600 ° C, and hence can no longer recombine and annihilate interstitials. The migration of interstitials along the crystallographic c-axis is discussed later.

Banhart11 observed typical basal plane defects in a graphite nanoparticles using high-resolution trans­mission electron microscopy (HRTEM). These defects can be understood as dislocation loops which form when displaced interstitial atoms cluster and form less mobile agglomerates. Other interstitials condense onto this agglomerate which grows into a disk, pushing the adjacent apart. Further agglomeration leads to the for­mation of a new lattice planes (Figure 4).

Other deformation mechanisms have been pro­posed for irradiated graphite. Wallace18 proposed a mechanism whereby interstitial atoms could facili­tate sp3 bonds between the atomic basal planes, this mechanism allowing the stored energy (discussed in Section 4.10.5.1) to be explained. Jenkins19 argued that the magnitude of the increase in shear

modulus (C44) with low dose irradiation could not be explained by interstitial clusters pinning dislocations, but that a few sp3 type covalent bonds between the planes could easily account for the observed changes. More recently, Telling and Heggie,15 in their ab-initio calculations of the energy of formation of the ‘spiro- interstitial,’ advocate this mechanism to explain the stored energy characteristics of displacement dam­aged graphite, particularly the large energy release peak seen at ~-473 K (discussed in Section 4.10.5.1). The first experimental evidence of the interlayer interstitial-vacancy (IV) pair defect with partial sp3 character in between bilayers of graphite was recently reported by Urita et at20 in their study of double­walled carbon nanotubes (DWNTs).

Jenkins19 invoked the formation of sp3 bonding to explain the c-axis growth observed as a result of displacement damage. If adjacent planes are pinned, one plane must buckle as the adjacent planes shrink due to vacancy shrinkage; buckled planes yield the c-axis expansion that cannot be explained by swelling from interstitial cluster alone. Telling and Heggie15 are very much in support of this position on the basis of their review of the literature and ab-initio simulations of the damage mechanisms in graphite. Their simulations showed how the spiro-interstitial (cross-link) essen­tially locked the planes together. Additionally, diva­cancies could lead to the formation of pentagons and heptagons in the basal planes causing the observed bending of graphene layers and c-axis swelling.11,21,22 The predicted c-axis crystal expansion via this mecha­nism is in closer agreement with the experimentally observed single crystal and highly oriented pyrolytic graphite (HOPG) dimensional change data.

The buckling of basal planes as a consequence of irradiation damage has been observed in HRTEM studies of irradiated HOPG by Tanabe21 and Koike and Pedraza.22 In their study, Koike and Pedraza22 observed 300% expansion of thin HOPG samples subject to electron irradiation in an in-situ transmis­sion electron microscope (TEM) study. Their exper­imental temperatures ranged from 238 to 939 K. They noted that the damaged microstructure showed retention of crystalline order up to 1 dpa (displace­ments per atom). At higher doses, they observed the lattice fringes break up in to segments 0.5-5 nm in length, with up to 15° rotation of the segments with respect to the original {0001} planes.

The evidence in favor of the formation of bonds between basal planes involving interstitials is consid­erable. However, such bonds are not stable at high temperature. As reported by numerous authors and
reviewers11’15’19’20 the sp3 like bond would be ex­pected to break and recombine with lattice vacancies with increasing temperature, such that at T>500K they no longer exist. Indeed’ the irradiated graphite stored energy annealing peak at ~-473 K, and the HRTEM observations of Urita eta/.20 demonstrate this clearly. Figure 5 shows a sequential series of HRTEM images illustrating the formation rates of interlayer defects at different temperatures with the same time scale (0-220 s) in DWNTs. The arrows indicate possi­ble interlayer defects. At T = 93 K (Figure 5(a)) the electron irradiation-induced defects are numer­ous’ and the nanotubes inside are quickly damaged because of complex defects. At 300 K (Figure 5(b)), the nanotubes are more resistive to the damage from electron irradiation, yet defects are still viable. At 573 K (Figure 5(c)), defect formation is rarely observed and the DWNTs are highly resistant to the electron beam irradiation presumably because of the ease of defect self-annihilation (annealing).

In an attempt to estimate the critical temperature for the annihilation of the IV defect pairs, a system­atic HRTEM study was undertaken at elevated temperatures by Urita et a/.20 The formation rate of the IV defects that showed sufficient contrast in the HRTEM is plotted in Figure 6. The reported numbers were considered to be an underestimate as single IV pairs may not have sufficient contrast to be
convincingly isolated from the noise level and thus may have been missed. However, the data was con­sidered satisfactory for indicating the formation rate as a function of temperature. The number of clusters of IV pairs found in a DWNT was averaged for several batches at every 50 K and normalized by the unit area. As observed in Figure 6, the defect for­mation rate displays a constant rate decline, with a threshold appearing at ^450-500 K. This threshold corresponds to the stored energy release peak (dis­cussed in Section 4.10.5.1) as shown by the dotted line in Figure 6. Evidentially, the irradiation dam­age resulting from higher temperature irradiations (above ^473-573 K) is different in nature from that occurring at lower irradiation temperatures.

Koike and Pedraza22 studied the dimensional change in HOPG caused by electron-irradiation-induced dis­placement damage. They observed in situ the growth c-axis of the HOPG crystals as a function of irradiation temperature at damage doses up to ~ 1.3 dpa. Increasing c-axis expansion with increasing dose was seen at all temperatures. The expansion rate was however signifi­cantly greater at temperatures ;S473 K (their data was at 298 and 419 K) compared to that at irradiation tem­peratures ^473 K (their data was at 553, 693, and 948 K). This observation supports the concept that separate irradiation damage mechanisms may exist at low irradiation temperatures (~T<473 K), that is,

buckling due to sp3 bonded cross linking of the basal planes via interstitials, and at more elevated irradiation temperatures (T ^ 473 K), where the buckling of planes is attributed to clustering of interstitials which induce the basal planes to bend, fragment, and then tilt. Koike and Pedraza22 also observed crystallographic a-axis shrinkage upon electron irradiation in-situ at several temperatures (419, 553, and 693 K). The shrinkage increased with dose at all irradiation temperatures, and the shrinkage rate reduced with increasing irradia­tion temperature. This behavior is attributed to buck­ling and breakage of the basal planes, with the amount of tilting and buckling decreasing with increasing tem­perature due to (1) a switch in mechanism as discussed above and (2) increased mobility of lattice vacancies above -~673 K.

Jenkins19,23 also discussed the deformation of graphite crystals in terms of a unit c-axis dislocation (prismatic dislocation), that is, one in which the Bur­gers vector, b, is in the crystallographic c-direction. The c-axis migration of interstitials can take place by unit c-axis dislocations. The formation and growth of these, and other basal plane dislocation loops undoubtedly play a major role in graphite crystal deformation during irradiation.

Ouseph24 observed prismatic dislocation loops (both interstitial and vacancy) in unirradiated HOPG using scanning tunneling microscopy (STM). Their study allowed atomic resolution of the defect structures. Such defects had previously been observed as regions of intensity variations in TEM studies in the 1960s.25

Telling and Heggie’s15 first principle simulations have indicated a reduced energy of migration for a lattice vacancy compared to the previously estab­lished value. Therefore, they argue, the observed lim­ited growth of vacancy clusters at high temperatures (T >900K) indicates the presence of a barrier to further coalescence of vacancy clusters (i. e., vacancy traps). Telling and Heggie implicate a cross-planer metastable vacancy cluster in adjacent planes as the possible trap. The disk like growth ofvacancy clusters within a basal plane ultimately leads to a prismatic dislocation loop. TEM observations show that these loops appear to form at the edges of interstitial loops in neighboring planes in the regions of tensile stress.

The role of vacancies needs to be reexamined on the basis of the foregoing discussion. If the energy of migration is considerably lower than that previously considered, and there is a likelihood of vacancy traps, the vacancy and prismatic dislocation may well play a larger role in displacement damage induced in­crystal deformation. The diffusion of vacancy lines to the crystal edge essentially heals the damage, such that crystals can withstand massive vacancy damage and recover completely.

Regardless of the exact mechanism, the result of carbon atom displacements is crystallite dimensional change. Interstitial defects will cause crystallite growth perpendicular to the layer planes (c-axis direction), and relaxation in the plane due to coalescence of vacancies will cause a shrinkage parallel to the layer plane (a-axis direction). The damage mechanism and associated dimensional changes are illustrated (in sim­plified form) in Figure 7. As discussed above, this conventional view of c-axis expansion as being caused solely by the graphite lattice accommodating small interstitial aggregates is under some doubt, and despite the enormous amount of experimental and theoretical work on irradiation-induced defects in graphite, we are far from a widely accepted understanding. It is to be hoped that the availability of high-resolution microscopes will facilitate new damage and annealing studies of graphite leading to an improved under­standing of the defect structures and of crystal defor­mation under irradiation.

Dimensional changes can be very large, as demon­strated in studies on well-ordered graphite materials, such as HOPG that has frequently been used to study the neutron-irradiation-induced dimensional changes of the graphite crystallite.13,26 Price27 conducted a study of the neutron-irradiation-induced dimensional changes in pyrolytic graphite. Figure 8 shows the crystallite shrinkage in the й-direction for neutron doses up to 12 dpa for samples that were graphitized at a temperature of 2200-3300 °C prior to being irra­diated at 1300-1500 °C. The a-axis shrinkage in­creases linearly with dose for all of the samples, but the magnitude of the shrinkage at any given dose decreases with increasing graphitization temperature. Similar trends were noted for the c-axis expansion. The significant effect of graphitization temperature on irradiation-induced dimensional change accumula­tion can be attributed to thermally induced improve­ments in crystal perfection, thereby reducing the number of defect trapping sites in the lattice.

Nuclear graphites possess a polycrystalline struc­ture, usually with significant texture resulting from the method of forming during manufacture. Con­sequently, structural and dimensional changes in polycrystalline graphites are a function of the crys­tallite dimensional changes and the graphite’s texture. In polycrystalline graphite, thermal shrinkage cracks that occur during manufacture and that are prefer­entially aligned in the crystallographic й-direction will initially accommodate the c-direction expansion, so mainly a-direction contraction will be observed. The graphite thus undergoes net volume shrinkage. With increasing neutron dose (displacements), the incompatibility of crystallite dimensional changes leads to the generation of new porosity, and the volume shrinkage rate falls, eventually reaching zero. The graphite now begins to swell at an

increasing rate with increasing neutron dose. The graphite thus undergoes a volume change ‘turn­around’ into net growth that continues until the gen­eration of cracks and pores in the graphite, due to differential crystal strain, eventually causes total dis­integration of the graphite.

Irradiation-induced volume and dimensional change data for H-451 are shown28 in Figures 9-11. The effect of irradiation temperature on volume change is shown in Figure 9. The ‘turn-around’ from volume shrinkage to growth occurs at a lower fluence and

, there is less accommodating volume

Changes to C33 and C44 in HOPG and natural graph­ite crystals have been reported at 50, 650, and 1000 °C53 and at 150 °C.54 For HOPG, the 150 °C data indicated that C33 slightly reduced with increas­ing irradiation (Figure 24) and this was attributed to the increase in ‘C axis lattice spacing. However, there is no clear trend at the other temperatures (Figure 24).

In the case of shear, at a very low temperature of 50 °C there was a significant increase in C44, but at higher temperatures the increase was less (Figure 25). The data for natural crystal showed similar trends but there was significantly more scatter. The trend in the increase in C44 at the lower temperature would go towards explaining the increase in modulus in polycrystalline data at low fluence. However, it is surprising that the increase is only modest at the higher temperatures, although the maximum fast neutron fluence is very low and data is required at the intermediate temperatures.

Seldin and Nezbeda53 also measured the shear strength but unfortunately there is considerable scat­ter and no definite trend.

4.11.11.1 Thermal Conductivity

Taylor et al55 measured the change in thermal conductivity in HOPG with fast neutron
irradiation. The thermal conductivity along the basal planes (the ‘a’ direction) is much greater than the value perpendicular to the basal planes (the V direction). Taylor et al. also measured the change in thermal resistivity in irradiated graphite, and when this data is normalized, the data indicated that thermal resistivity temperature dependence changed with irradiation as given in Figure 26. This is the so-called ‘d’ curve that is used in the United Kingdom to predict thermal resistivity in irradiated graphite.

4.13.1 Introduction

As concrete ages, changes in its properties will occur as a result of continuing microstructural changes (i. e., slow hydration, crystallization of amorphous constituents, and reactions between cement paste and aggregates) as well as environmental influences. These changes do not have to be detrimental to the point where concrete will not be able to meet its functional and performance requirements; however, concrete can suffer undesirable changes with time because of improper specifications, violation ofspecifications, or adverse performance ofits cement paste matrix or aggregate constituents under physical or chemical attack. Additional information related to environmental effects on concrete is provided in molten core concrete interaction (Chapter 2.25, Core Concrete Interaction).

Portland cement concrete durability is defined as its ability to resist weathering action, chemical attack, abrasion, or any other process or deterioration.1 A durable concrete is one that retains its original form, quality, and serviceability in the working envi­ronment during its anticipated service life. The mate­rials and mix proportions specified and used should be such as to maintain concrete’s integrity and, if applicable, to protect embedded metal from corro — sion.2 The degree of exposure anticipated for the concrete during its service life, together with other relevant factors related to mix composition, work­manship, and design, should be considered.3 Guide­lines for production of durable concrete are available in national consensus codes and standards, such as American Concrete Institute (ACI) 3184, which have been developed over the years through knowledge acquired in testing laboratories and supplemented by field experience. Serviceability of concrete has been incorporated into the codes through strength requirements and limitations on service load conditions in the structure (e. g., allowable crack widths, limitations on midspan deflections ofbeams, and maximum service level stresses in prestressed members). Durability gen­erally has been included through items such as specifi­cations for maximum water-cement ratios, minimum cementitious materials contents, type cementitious material, requirements for entrained air, and minimum concrete cover over reinforcement. Requirements are frequently specifiedin terms of environmental exposure classes (e. g., chloride and aggressive ground environ­ments). Specifications in terms of service life require­ments (e. g., short <30years, normal 30-100years, and long >100 years) have only recently been developed, primarily through European standards.5

Water is the single most important factor con­trolling the degradation processes of concrete (i. e., the process of deterioration of concrete with time is generally dependent on the transport of a fluid through concrete), apart from mechanical deteriora­tion. The rate, extent, and effect of fluid transport are largely dependent on the concrete pore structure (i. e., size and distribution), presence of cracks, and micro­climate at the concrete surface. The primary mode of transport in uncracked concrete is through the cement paste pore structure (i. e., its permeability). The domi­nant mechanism controlling the rates ofwater penetra­tion into unsaturated or partially saturated concrete is absorption caused by capillary action of the concrete’s pore structure. To improve the durability of concrete, generally the capillary and pore size within the con­crete matrix should be reduced to a minimum.

Although the coefficient of permeability for con­crete depends primarily on the water-cement ratio and maximum aggregate size, it is influenced by the curing temperature, drying, cementitious materials content, and addition of chemical or mineral admixtures as well as the tortuosity of the path of flow. Concrete compressive strength has traditionally been utilized as an acceptance test for concrete, but it typically is not a good indicator of durability. Many structures have been fabricated with concretes having adequate 28-day compressive strength only to lose their func­tionality because they were facing an environment for which they had not been designed or because the concrete had not been placed or cured correctly.6

The safety-related concrete structures in nuclear power plants (NPPs) are designed to withstand load­ings from a number of low-probability external and internal events, such as earthquake, tornado, and loss — of-coolant accident. Consequently, they are robust and not subjected to high enough stresses during normal operation to cause appreciable degradation. In general, this has been the case, as the performance of reinforced concrete structures in NPPs has been good. (Operating experience is discussed in the next section.) However, as the NPPs age, degradation incidences start to occur at an increasing rate, pri­marily due to environmental-related factors. One — fourth of all containments in the United States have experienced corrosion, and nearly half of the con­crete containments have reported degradation related to either the reinforced concrete or post-tensioning

system.7 Although the vast majority of these struc­tures will continue to meet their functional and per­formance requirements during their initial licensing period (i. e., nominally 40 years), it is reasonable to assume that with the increasing age of the operating reactors there will be isolated examples where the structures may not exhibit the desired durability without some form of intervention.

Currently, the United States has 104 NPP units licensed for commercial operation, which provide about 20% of the electricity supply. As all but one of the construction permits for existing NPPs in the United States were issued prior to 1978, the focus for the existing plants has shifted from design to condition assessment. Here, the aim is to demonstrate that struc­tural margins ofthe plants have not eroded or will not erode during the desired service life due to aging or environmental effects. One of the key factors to maintaining adequate structural margins to protect public health and safety in the unlikely event of an accident is implementation of effective inspection and maintenance programs. An inspection program is important for identifying and characterizing any deg­radation that may be present in a timely manner. Once degradation has been identified, or its potential to occur established, a maintenance program is imple­mented to repair the degradation and arrest (as far as possible) the mechanism(s) causing the degradation. Proper maintenance is essential to the safety of NPP structures, and a clear link exists between effective maintenance and safety. Uncertainty in condition assessment can be assessed using probabilistic methods, which are also an essential ingredient of risk-informed management decisions concerning continued service of the NPP structures.

Long before the onset of significant phase evolution or void swelling is observed, the first manifestation of the radiation-induced microstructural/microchemical evolution appears in changes ofthe mechanical prop­erties. As shown in Figure 18 the stress-strain dia­grams of stainless steels begin to change significantly even at very low dpa levels. The strength of the alloy increases, the elongation decreases, and there is a progressive decrease in work-hardening. This behav­ior is dependent somewhat on test temperature but is not very sensitive to neutron spectrum.

Movement of dislocations in metals during deformation following irradiation is impeded by the microstructural components produced by radiation (dislocations, dislocation loops, voids, bubbles, preci­pitates) and therefore the strength of annealed steel

increases. The strength increase usually saturates at relatively low exposure levels (<10dpa) as shown in Figure 19, reflecting a similar saturation of micro­structural densities. Since the concentration of most radiation-induced microstructural components decreases with increasing temperature above ^300 °C, one would expect that the saturation strength would also decrease with increasing temperature, as is shown in Figures 20-23.

Irradiation of cold-worked steels also leads to strengthening at lower temperatures but softening can occur at higher temperatures if the saturation strength level at a given temperature is below the starting strength, as seen in Figure 21. Most impor­tantly, both annealed and cold-worked steels con­verge to the same saturation level when irradiated at the same dpa rate and temperature as seen in Figure 22.78

Similar convergence behavior has been observed in the evolution of microhardness.79 Note also that radiation-induced changes in strength are roughly independent of composition within the annealed 300 series stainless steels, especially at lower irradia­tion temperatures, as shown by Figure 19.

Such convergence behavior has been observed many times, but there are exceptions; for example, cold-worked steels converge in their notch tensile strength, but not to the level reached by annealed steels.80 Such behavior is usually observed in steels that twin heavily during deformation and were irradiated at low temperatures that resist recrystallization. Twin boundaries are not easily erased by displacements, so their hardening contribution persists.

Concurrent with an increase in radiation-induced hardening is a loss of ductility,81-83 as shown in Figures 23 and 24.

The concept of saturation or persistence of mechanical properties, especially with respect to tem­perature, applies to the most recent irradiation temperature, as demonstrated by comparing isother­mal and nonisothermal histories. In Figure 25 the mechanical properties of three model alloys are seen to converge during isothermal irradiation without being affected by composition, He/dpa ratio, and mechanical starting state.84 In Figure 26, however, an early detour in temperature led to differences from isothermal behavior, but these differences dis­appeared when the intended isothermal temperature was reestablished.84

Previous saturation states are soon forgotten, usu­ally by ^5 dpa, but only if the hardening components are easily erased and replaced at the new tempera­ture. If hardening arises primarily from dislocation loops and dislocations, this condition is easily met. If the primary hardening arises from a fine density of voids and especially bubbles produced at lower temperatures, then the microstructural memory can­not be easily erased, even at much higher tempera­tures. An example is shown in Figure 27 where a series of Fe-Cr-Ni ternary austenitic alloys were irradiated at 400 and 500 °C in ORR at high He/dpa ratios

(27-58 appmdpa-1) and 395 and 450 °C in EBR-II at very low He/dpa ratios (0.7-1.2 appmdpa-1).85

Note that there are very significant differences in hardening observed between the two experiments and that the differences arose primarily from a very large difference in cavity density, a difference that was too large to be explained in terms of helium content alone. It was later shown that the ORR experiment suffered a very large number (237 over 2 years) of unrecognized negative temperature set­backs of 1-2 h, with decreases varying from 50 to 500 °C.86 Even though the total dpa accumulated during these setbacks was only ~1% of the total dose, the frequent bloom of high densities of small Frank loops at lower temperatures provided a very large periodic increase in nucleation sites for helium bubbles on the new Frank loops that significantly strengthened the matrix. The loops could subse­quently dissolve but the bubbles could not.

In addition to temperature, the most prominent irradiation variable is the dpa rate and it is known that the microstructural densities, especially Frank loops and voids, are known to increase in concentra­tion as the dpa rate increases. Various radiation-stable phases such as %’ are also known to be flux-sensitive, while other phases such as carbides and intermetallics are more time-sensitive.1

Thus, it is not surprising that some sensitivity to dpa rate might be observed in strength properties, as

Test temperature = Irradiation temperature

suggested by the behavior shown in Figure 28 where both the transient rate of strength rise and saturation strength appear to increase with increasing dpa rate. Unfortunately, this figure does not represent a single variable comparison, and by itself is not sufficiently convincing evidence of flux sensitivity. The data shown in Figure 29 is much closer to a single variable comparison, indicating that the transient rise may or not be somewhat flux-sensitive, depending on the details of the microstructural evolution of each alloy. The authors of this study used microscopy to confirm the microstructural origins of the observed differ­ences of behavior as a function of dpa rate.

More recently, Chatani and coworkers showed that at relatively low irradiation temperatures char­acteristic of boiling water reactors, the radiation — induced increments in strength of 304 stainless steel increased by the 1/4 power of the increase in dpa rate.87 It was demonstrated that the black-spot micro­structure dominated the strengthening. It was also shown that the concentration of black spots varied with the square root of the flux as expected, and it is known that hardening varies with the square root of the loop density, thereby producing a fourth-root dependence. Thus, in the absence of any significant microchemical or phase stability contributions, it

appears that radiation-induced strengthening is affected by dpa rate but not very strongly.

The loss of ductility proceeds in several stages, first involving convergence of the yield and ultimate strengths as shown in Figures 29 and 30, such that a loss of work-hardening occurs and very little uniform elongation is attained. As the irradiation pro­ceeds, there is a progressive tendency toward flow localization followed by necking. As seen in Figure 31 the failure surface shows this evolution with increas­ing dose.

The flat faces observed at highest exposure in Figure 31 are often referred to as ‘channel fracture’ but they are not cleavage faces. They are the result of intense flow localization, resulting from the first moving dislocations clearing a path of radiation — produced obstacles, especially Frank loops, and thereby softening the alloy along that path. It is not possible to remove the voids by channeling but the distorted
voids provide a microstructural record of the flow localization as shown in Figure 32. Linkage of the elongated voids is thought to contribute to the failure.

Such a failure surface might best be characterized as ‘quasi-embrittlement’, which is a suppression of uniform deformation, differentiating it from true embrittlement, which involves the complete sup­pression of the steel’s ability for plastic deformation. This distinction is made because under some con­ditions quasi-embrittlement can evolve into true embrittlement.

The tendency toward quasi-embrittlement grows with increasing swelling but the alloy is actually softening with increasing swelling rather than hard­ening. As shown in Figure 33 brittle fracture (defined as strength reduction with zero plasticity) of a Fe-18Cr-10Ni-Ti stainless wrapper in BOR-60 at 72 dpa maximum was observed at positions where peak swelling occurs.88 Some decrease of strength is

observed with increasing irradiation temperature, but the primary strength reduction for specimens tested at the irradiation temperature arises from the magnitude of swelling. Testing at temperatures below the irradia­tion temperature (e. g., 20 °C) demonstrates the same dependence on swelling and irradiation temperature, but the strength and plasticity values are higher. As expected, the strengths for tests conducted at 800 °C are uniformly much lower than that observed at lower temperatures, but there is an absence of any relation­ship between strength and swelling at this temperature.

As shown in Figure 34 failure surfaces at high swelling levels exhibit transgranular cup-cone mor­phology where failure proceeded by micropore coa­lescence arising from stress concentration between deforming voids.88 Similar fracture morphology has been observed in studies on other stainless steels.1

Although voids and bubbles initially serve to harden the microstructure,78 large swelling levels allow previously second-order void effects to become dominant.1,88,89 One of these second-order effects is the strong decrease of elastic moduli at high swelling levels. All three elastic moduli decrease initially at ^2% per each percent of void swelling.90-93 At >10% swelling this leads to significant reduction in strength.

As a consequence, the slope of the elastic region (Young’s modulus) of the stress-strain curve decreases, and more importantly, the barrier strengths of all sinks decrease as the shear modulus likewise decreases. Therefore, the yield and ultimate strengths decrease with increasing swelling, even though the elongation strongly decreases. Similar behavior has also been observed in pure copper.94

The nature of the void-related failure changes from quasi-embrittlement to true embrittlement for tests at or near room temperature, demonstrating another example of a late-term second-order process growing to first-order importance at higher swelling levels.

Hamilton and coworkers observed that above ~10% swelling the previously established saturation strength level of 316 stainless steel suddenly increased very strongly in room temperature tensile tests.95 Similar results were observed in Russian steels 96,97 As shown in Figures 35 and 36 the failure surfaces in such tests had rotated from the expected 45° (relative to the stress axis) to 90° as swelling approached 10%, indi­cating complete brittle failure, as also indicated by the fully transgranular nature of the failure surface. Con­currently, the ductility vanished and the tearing mod­ulus plunged to zero, indicating no resistance to crack propagation. Once a crack has initiated it then propa­gates completely and instantly through the specimen.

Neustroev and coworkers observed such failures in Russian steels that are subject to greater amounts of precipitation and determined that the critical micro­structural condition was not defined solely by the level of swelling, but by the obstacle-to-obstacle distance of the void-precipitate ensemble, indicating that stress concentration between obstacles was one contributing factor.96 However, it was the progressive segregation of nickel to increasing amounts of void surface and the concurrent rejection of chromium from the sur­faces that precipitated the rather abrupt change in failure behavior.1,95 This late-term void-induced micro­chemical evolution induces a martensite instability in the matrix, as evidenced by the failure surface being completely coated with alpha-martensite.95

The abrupt jump in strength just before failure observed by Hamilton and coworkers is the result of a stress-induced blossoming of a high density of small, thin, epsilon-martensite platelets, as seen in Figure 37. These platelets are essentially stacking faults that form under stress as a result of the influ­ence of both falling nickel level and low deforma­tion temperature to decrease the stacking fault energy of the matrix.1 When encountered by the advancing crack tip, the epsilon-martensite is con­verted to alpha-martensite in the strain field ahead of the crack, providing a very brittle path for further cracking.

The correlation between void swelling and both quasi-embrittlement and true embrittlement is observed not only in slow tensile tests (Figures 36, 38, and 39) but also in Charpy impact tests as shown

in Figure 39. Figures 40-44 present examples of swelling-induced failures in components experien­cing a wide range of physical insults. The example of Porollo et al. in Figure 44 (top) is particularly noteworthy in that it results from significant swelling at 335 °C, a temperature earlier thought not to produce significant amounts of swelling.

If there are no physical insults experienced by the component during irradiation, the continued segregation of nickel to void surfaces and the con­current rejection of chromium can lead to strong changes in composition in the matrix during irradia­tion, pushing the matrix toward ferrite rather than martensite at higher temperatures, especially for steels with nickel content of <10%. In some observations voids encased in austenite shells have been observed to exist in a pure ferrite matrix.98,99 To date, however, no significant component failure has been reported to result from this particular late-term instability.

Finally, there appears to be another late-term phase instability developing at lower irradiation tem­peratures that involves martensite but does not appear to be due to void swelling. Gusev et al. have shown that for irradiation temperatures below -~350 °C a growing tendency for stress-induced mar­tensite formation is occurring in Russian austenitic steels at doses in the range of 25-55 dpa when tested at room temperature.100-102 Surprisingly, this insta­bility results in a restoration of engineering ductility to preirradiation levels. However, the ductility is

regained not because the steel has softened, but because it becomes exceptionally strong and hard­ened during deformation. As a consequence, the steel has lost the ability to neck.

Instead of necking, a moving wave of deformation is initiated at the first attempted necking point. The wave front then travels nearly the full length of the gage section. Initially, there is a local deforma­tion in the order of 40-45%, but as the wave moves forward it leaves a relatively uniform local defor­mation in its wake. Everywhere behind the wave front there is measured 30-35% volume percent of martensite, as shown in Figures 45 and 46. The martensite is not only a product of the wave, but also the cause of the wave. Deformation-induced martensite resists further necking and forces the deformation to be displaced to the adjacent lesser deformed material. The mechanisms that cause the late-term onset of martensite instability have not yet been determined.

A property of important engineering interest is the fracture toughness Jc. While the fracture toughness of various unirradiated stainless steels can be quite

Figure 35 Fractographs of failure surfaces of 20% cold-worked 316 specimens cut from an FFTF duct at high exposure. Reproduced from Hamilton, M. L.; Huang, F. H.; Yang, W. J. S.; Garner, F. A. In Effects of Radiation on Materials: 13th International Symposium (Part II) Influence of Radiation on Material Properties; ASTM STP 956; 1987; pp 245-270. Note change of fracture mode from channel fracture when tested at 205 and 460 °C to brittle fracture when tested at 20 °C.

different, it appears that all austenitic steels studied undergo the same general evolution in toughness during irradiation. Mills has shown that three regimes of evolution occur.103,104 The first regime involves a low-dose threshold exposure range (< 1 dpa) where there is essentially no loss of toughness, and the second regime involves an intermediate exposure range (1-10 dpa) where toughness decreases rapidly with exposure, producing an order of magnitude reduction in Jc and two orders of magnitude degra­dation in tearing modulus. Finally, a saturation regime is reached, in which increasing exposure does not produce a further reduction in toughness. This saturation occurs well before any of the void-induced instabilities discussed above can occur. As shown in

Figure 47, the saturation level is remarkably inde­pendent of the original toughness level.

Welds in austenitic alloys were shown by Mills to exhibit lower initial toughness values and lower satu­ration toughness levels as well. The fracture toughness level is sensitive to the test temperature, however, as shown in Figure 48. At high test temperatures, the fracture mode changes from transgranular to inter­granular in nature, reflecting the effect oftest tempera­ture on both matrix strength and also the influence of helium embrittlement at grain boundaries.105 The level ofhelium needed to promote high temperature embrit­tlement is not very high, however, and can easily be reached after moderate neutron exposure in fast reactor-irradiated alloys with the lowest nickel level.

Early models ofprecipitate stability under irradiation were based on the ideas of Nelson et a/.,76 who suggested that precipitates would evolve to an equi­librium size determined by competing processes affecting their growth, via the radiation enhanced and/or thermal diffusion of solutes, and their simul­taneous dissolution due to damage arising in collision cascades. Two dissolution mechanisms were sug­gested: recoil dissolution due to the displacement of atoms from the precipitate into the matrix, and dis­ordering dissolution of ordered phases such as g, with the latter predicted to be the more effective. The model predicted that fine precipitates would continue to grow to some equilibrium size (depen­dent on temperature, dose rate, and solute levels), but that precipitates greater than this size would shrink. Experimental evidence for the dissolution of large preexisting Ni3Al g precipitates in heavy-ion- irradiated Ni-Al alloys was shown by Nelson et a/.76

These ideas were developed further and applied to g precipitates in ion and neutron-irradiated alloys by Baron et a/.77 The model developed by Baron et a/. indicated that, at a given particle size, a higher solute supersaturation was required under irradiation than in a purely thermal environment. The model appeared to be consistent with the observed coarsen­ing behavior of g0 precipitates during irradiation, though no evidence for the shrinkage oflarge particles was presented. For example, data for PE16 irradiated at fluences up to 7.5 x 1026nm—2 at 560 °C, which were reported by Chang and Baron,78 only examined the growth of g0 particles up to a maximum radius of — 15 nm under conditions where the predicted maximum equilibrium radius was — 35 nm.

However, detailed examinations of g structures in neutron-irradiated Nimonic PE16 which were made by Gelles79 found no evidence to indicate that irra­diation-induced dissolution mechanisms limited the particle size. Microstructural examination of PE16, originally in ST, STA, and OA conditions, irradiated in EBR-II to -27 dpa (5.4 x 1026nm—2, E> 0.1 MeV) at 600 ° C, revealed that preexisting g dispersions in aged material were maintained but continued to coarsen even in the OA condition, and that a fine dispersion formed in ST material. Coarsening of the g0 particles in the OA material was accompanied by the formation of fine background precipitates in some regions. Further in-reactor precipitation of g0 also occurred at point defect sinks, including void surfaces and dislocations, in all the heat-treated con­ditions. Additional examinations by Gelles80 of ST PE16, irradiated to —30-50 dpa at temperatures in the range of 430-650 °C, indicated that g coarsening was controlled by radiation-enhanced diffusion below 600 °C with an activation energy that (in agreement with theoretical predictions for a process governed by point defect recombination) was approx­imately a quarter of that for thermal diffusion.

As described in Section 4.04.5.1 in relation to irradiation embrittlement effects, Yang81 examined an identically irradiated set of ST PE16 samples as Gelles, focusing on the precipitation of g at grain boundaries. Similar g structures to those described by Gelles and Yang were also observed by Boothby2 in the aged conditions of EBR-II-irradiated PE16, though at higher irradiation temperatures (>540 °C for the STA condition, and >600 °C for the OA condition), where doses were in the range 66-74 dpa, the spherical g precipitates which formed during thermal aging were almost entirely replaced by ‘skel­etal’ forms nucleated at point defect sinks. Figure 11 shows an example of the g distribution, imaged in dark field, in STA PE16 irradiated to 69 dpa at 570 °C; although small spherical precipitates were retained in a narrow region adjacent to the grain boundary, a much coarser dispersion is evident at the boundary itself and within the bulk of the grain.

The mechanical property performance of pure Mo is strongly controlled by the grain size, oxygen, nitrogen, and carbon concentrations as well as alloy additions. This is true for unirradiated as well as irradiated properties.71,76,83,84 The sensitivity to embrittlement at irradiation temperatures <873 K can be mitigated through a reduction in oxygen and nitrogen while keeping the carbon-to-oxygen ratio high to reduce the segregation of oxygen and nitrogen to the grain boundaries. A reduction in the grain size can further increase the number of sinks and reduce the mean distance that irradiation-induced defects must travel at temperatures at which mobility is limited.

Irradiated mechanical properties of wrought LCAC-Mo in both the recrystallized and stress — relieved conditions have been examined over several decades.81,82,84-86,99-106 In general, LCAC-Mo un­dergoes significant increases in tensile strength through radiation hardening <873 K, which produces reductions in ductility and high DBTT values. A summary of tensile properties as a function of irra­diation temperature and dose is shown in Figure 14. Irradiated stress-relieved LCAC-Mo shows less radia­tion embrittlement than as-crystallized materials at temperatures <1208 K.1 0 However, at higher irradia­tion temperatures, recrystallization of the stress — relieved material occurs, leading to large changes in the microstructure and less desirable properties.

Increases in hardness of 56% and 112% for stress — relieved and recrystallized LCAC-Mo, respectively, are reported following irradiation to 1.2 dpa at 543 K.100 The increase in hardness decreases slightly following 878 K irradiations to 2.4 dpa, but returns to values near those for the unirradiated material for irradiation at temperatures >1208 K. Materials irra­diated in the stress-relieved condition at tempera­tures >1173 K can result in softening compared to unirradiated materials because ofrecrystallization.107 A comparison of the changes in irradiated material hardness as a function of dose and temperature is plotted in Figure 15(a) for LCAC-Mo, TZM, and ODS-Mo. The two latter alloys are discussed in detail later. In the three materials investigated by Cockeram et a/.,107 the largest increase in hardness was measured for irradiations at 873 K, counter to what is observed in tensile properties for irradiation between 573 and 873 K.81,85 However, tensile failure in the lower temperature irradiated samples generally occurs before the samples yield because of the eleva­tion ofthe flow stress above the fracture stress for these test conditions. Recovery of the hardness, measured through 1 h anneals at increasing temperatures, is plot­ted in Figure 15(b) for the ^573 K irradiated material. The start of recovery is near 800 K with full recovery of the LCAC-Mo hardness by 1253 K. For LCAC-Mo irradiated at a higher temperature of 873 K, recovery of hardness begins near 1163 K and is completed near 1463 K.

Substantial increases in the DBTT to values >773 K are observed for irradiated LCAC-Mo over a range of fluences for irradiation temperatures <873 k.26,82,85,103,105,108 A summary of DBTT values for LCAC-Mo is presented in Figure 16, along with data from high-purity grain-refined LCAC-Mo, TZM, and ODS-Mo, which is discussed next. In general, recovery in the DBTT of LCAC-Mo is not observed until irradiation temperatures above 873­973 K, depending on material conditions. A reduced sensitivity to low-temperature irradiation embrittle­ment of LCAC-Mo is observed in materials with reduced levels of impurities. A high-purity form of LCAC-Mo (HP-LCAC-Mo) was developed through the use of 1873 K hydrogen atmosphere annealing of LCAC-Mo plates prior to further arc casting, extru­sion, and rolling into sheet stock.82 Levels of oxygen and nitrogen were < 4 wppm each, with carbon at 20wppm. Average grain diameters of 1.3 and 452 mm lengths were produced, and represent a considerable change in aspect ratio compared to LCAC-Mo values of 4-5 mm diameter and 78-172 mm lengths.100 The DBTT of 573 K irradiated HP-LCAC-Mo showed no increase over the unirradiated value (123 K) for irradiations up to 0.11 dpa, and increased to 723 K by 1.29 dpa. For 873 K irradiations, the DBTT remained below 223 K up to 1.29 dpa.

The majority of the irradiated mechanical prop­erty database for TZM is limited to displacement damage <5 dpa and irradiation temperatures <1000K, with much of the testing conducted at tem­peratures below that used in the irradiation. The data­base is also sparse, with little connectivity between

different experimental examinations. Furthermore, significant differences are observed in the nonirradi — ated tensile values of TZM because of variations in material processing leading to differences in grain size, impurity level, distribution of the strengthening phase, and differences in testing procedure. In general,

I

1100

900

Figure 15 (a) Change in hardness as a function of neutron dose for LCAC-Mo, TZM, and ODS-Mo irradiated between 573 and 1173 K up to 13.1 dpa. (b) Recovery of hardness as a function of isochronal annealing temperature for material irradiated ~573K. Adapted from Cockeram, B. V.; Smith, R. W.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2009, 393, 12-21.

displacement damage to the TZM alloy produces a large increase in the yield strength of the material with a corresponding drop in total elongation, with the level of change increasing with dose and/or lower irradiation temperatures. Some recovery of properties begins to be observed at test conditions above 973 K for materials irradiated at lower temperatures. A compila­tion of earlier and more recent tensile data for irra­diated TZM is provided in Figure 17. Irradiation <1.2 dpa and temperatures >800 K showed little effec­tive strengthening above the unirradiated values, for tensile tests below the irradiation temperature, 85,109,110 though higher displacement doses resulted in a signifi­cant increase.81 Total elongation in material irradiated < 1000 K was limited, with uniform elongation values <1%.m Plastic instability following yielding was also observed in irradiated TZM as well as LCAC-Mo. The plastic instability stress, defined as the maximum true stress in which plastic necking occurs, is strongly dependent on test temperature but nearly independent of fluence and irradiation temperature.81

Increases in the DBTT for irradiated TZM are significant and do not diminish until irradiation tem­peratures slightly higher than that of LCAC-Mo,81 but comparisons may be difficult because of differ­ences in materials and test methods. The DBTT value for unirradiated TZM is ^200 K and increases with increasing irradiation damage. An increase of 230 K in the DBTT over unirradiated values was reported for TZM irradiated to 2-4.8 dpa at 644-661 K.111 DBTTs
of 750 K for Mo-0.5Ti irradiated to ^16 dpa at 727 K26 and 1073 K for TZM irradiated to 12.3 dpa at 573 K81,85 were also reported. A summary of DBTT as a function ofirradiation temperature for LCAC-Mo and TZM is shown in Figure 16 for irradiation doses <50 dpa.8 Excessive embrittlement is observed for irradiations at <773 K. A reduction in the DBTT for LCAC-Mo appears near 873 K, while the DBTT of TZM remains high. Recovery of DBTT to near unir­radiated values occurs for irradiations at >1073 K.81, The stage V recovery temperature for vacancy diffu­sion in Mo is ~-873 K; however, the kinetics for micro­structural changes to occur are still relatively slow at this temperature. Therefore, embrittlement issues can be present until 1073 K.

Very little fracture toughness data exist for irradiated TZM. For precracked compact tension specimens irradiated to 0.29-0.35 dpa at 313-695 K, a 4MPaVm decrease in fracture toughness (15- 20MPaVm at Trm, unirradiated84,109) was observed up to the irradiation temperature.109 Kitsunai etal.112 examined the impact toughness of irradiated TZM and alloys, incorporating 0.1-1 wt% TiC additions to Mo. The Mo-TiC alloys showed dramatically improved toughness levels over TZM with increasing TiC concentration in samples irradiated to 0.08 dpa between temperatures of 573 and 773 K. Shown in Figure 18 is the shift in DBTT to lower tempera­tures for the irradiated TiC-containing Mo over the TZM alloy. For the 1 wt% containing sample,

the DBTT remained unchanged with irradiation, despite a ^50% increase in Vickers microhardness. More surprisingly, the Mo-1%TiC sample increased in toughness following 0.8 dpa irradiation attributed to grain boundary strengthening by (Ti, Mo)C radiation-enhanced precipitates.

Developed to improve the low-temperature duc­tility and weld characteristics of unalloyed Mo, the Mo-Re alloys have gained considerable attention over the past decade for use in nuclear applications. Single-phase solid solution a-Mo phase field extends
up to ^42 wt% Re, above which the a-MoRe2 phase precipitates. At higher Re concentrations, the W-MoRe3 phase is present. However, the exact phase boundaries are not well delineated at temperatures below 1773 K113,114 mainly because of the slow kinet­ics in phase development.115

The Mo-Re alloys show a hardening response to irradiation stronger than that of the pure metal and TZM following irradiation.116-118 The hardening response of Mo-Re alloys ranging in composition from 2 to 13 as well as 41 wt% Re following

irradiation up to 20 dpa at temperatures between 681 and 1072 Kwas examined by Nemoto etal.116 A linear increase in hardness with Re concentration was ob­served for the unirradiated controls as well as samples irradiated at temperatures <874 K. For samples irra­diated at 1072 K, little variation was observed with increasing Re content, though hardness values remained nearly double that of the unirradiated mate­rial. The dependence of hardness on irradiation tem­perature and fluence for Mo-5Re and Mo-41Re in comparison with LCAC-Mo is presented in Figure 19.

The high degree of radiation hardening at tem­peratures <1000 K exhibited by the Mo-Re alloys is further reflected in low ductility and reported embrittlement. Tensile elongation values of <0.3% were reported for Mo-5Re irradiated to 0.16 dpa and tested at the irradiation temperature of 320 K,11 though higher total elongations of 8% were reported for Mo-5Re fast reactor irradiated to 0.29 dpa and tested near the irradiation temperature of ^723 K.1 Unlike LCAC-Mo and TZM, the little deformation occurring in Mo-Re is mostly uniform at tempera­tures below 1000 K,11 ,119 while a small degree of work softening has been observed at higher tempera — tures.120 No evidence of dislocation channeling was found during microstructural examination of tensile tested Mo-5Re irradiated to 0.16 dpa at 373 K.118

For comparable irradiation fluences, dislocation loop concentrations are approximately two times higher than TZM and four to six times higher than Mo, while dislocation loop diameters are smaller in the Mo-5Re alloy.118 Void development begins to appear in Mo-Re alloys above 623 K118 and shows a slight increase in size with Re concentration (with corresponding decrease in number density) up to 10wt% with no further increase for the 41% Re alloy.116 Irradiation-induced void swelling in Mo with Re concentrations <13wt% is 0.5—1.5% for ^21 dpa irradiated material at 681-1072 K, while swelling for 41 wt% Re was near 0.1%.116

Radiation-induced precipitation has been reported by Nemoto eta/.116 in Mo—(2—41)Re alloys irradiated to 21 dpa between 681 and 1072 K, and by Edwards et a/.121 in Mo—41Re irradiated 28—96 dpa at 743— 1003 K. An initial formation of hcp-structured precipi­tates with a thin plate-like morphology consisting of solid solution Re and Os was observed, appearing with the {110}Mo//{0001}Re, <111>Mo//<2110>Re orien­tation relationship.121 The precipitation of these plates on dislocation loops resulted in the high density ofplates observed, which dominates the microstructure. On fur­ther irradiation to higher doses or higher temperatures, these plates develop and coarsen into the w-phase. This nonequilibrium phase development in Re-lean alloys was originally observed by Erck and Rehn12 in Mo—(27—30)Re irradiated by 1.8 MeV He ions at 1023— 1348 K. The a-MoRe2 phase was also reported appear­ing in all the Mo—Re samples examined by Nemoto and coworkers,116 but was suppressed in the stress-relieved specimens compared to the recrystallized materials.

While a limited (<1%) amount of ductility was reported in Mo—5Re alloys irradiated at tempera­tures <1000K to 34dpa,109,118,119 embrittlement of Mo—(1—20)Re irradiated 723—1073 K in a fast reactor up to 5 dpa and Mo—(13 and 47)Re irradiated 373— 673 K in a mixed spectrum reactor to 2 dpa has been reported by Fabritsiev and Pokrovsky.62 The reduc­tion in tensile strength, in many cases below the unir­radiated values with no plasticity occurring in the samples, was attributed to the hardening of the mate­rial by radiation-created defects along with RIS of oxygen, nitrogen, and transmuted impurities to the grain boundaries. The oxygen and nitrogen content in the embrittled alloys was reported to be near 70 appm.

Irradiation hardening above 900 MPa was also observed in Mo—41Re and Mo—47.5Re samples irra­diated to 1.46 dpa at temperatures >1073 K.120 Failure of Mo—41Re samples irradiated to 1.46 either prior to yielding or after ^5% elongation upon reaching 1600 MPa was observed at 1073 K (Tirr = Ttest). Exam­ples of the tensile curves for the two alloys in the irradiated, 1100 h aged and as-annealed condition tested at 1073 K is shown in Figure 20. Radiation

Mo (Nemoto et a/.116 18-21 dpa) ■ Mo-5Re (Nemoto et a/.116 18-21 dpa)

A Mo-10Re (Nemoto et a/.116 18-21 dpa) x Mo-41Re (Nemoto et a/.116 18-21 dpa)

♦ Mo-5Re (Hasegawa et a/.117 6.8-34 dpa) —O— Mo-5Re (Hasegawa et a/.117 6.8-34 dpa)

—□— Mo-41Re (Hasegawa et a/.117 6.8-34 dpa)

Figure 19 Vickers hardness as a function of neutron irradiation temperature and dose for LCAC-Mo, Mo-5Re, and Mo-41 Re alloys. Displacement damage levels are provided in the key. Reproduced from Nemoto, Y.; Hasegawa, A.; Satou, M.; Abe, K.; Hiraoka, Y. J. Nucl. Mater. 2004, 324, 62-70; Hasegawa, A.; Ueda, K.; Satou, M.; Abe, K. J. Nucl. Mater. 1998, 258-263, 902-906.

hardening to levels over twice the as-annealed condi­tion was observed for the alloys irradiated at 1223 and 1373 K; however, total elongation was between 4 and 12%. Analysis of the fractured surfaces of these samples revealed intergranular failure, with the sever­ity increasing with irradiation temperature. A compari­son of mechanical property data of Mo-Re samples from the sources discussed is shown in Figure 21.

The degree of RIS influencing the properties of Mo-Re alloys varies with temperature, dose rate, and total fluence. At temperatures <0.3 Tm, the recombi­nation of vacancies and interstitials generated by displacement damage dominates because of the lim­ited defect mobility, and therefore RIS is not a factor. At temperatures >0.5 Tm, a reduced driving force for segregation occurs because of the high thermal defect concentrations. At intermediate temperatures (^850-1430 K for Mo-Re), the radiation generated point defects diffuse to defect sinks such as grain
boundaries or dislocations. Any preferential coupling of vacancies or interstitial defects fluxes with solute atoms, including transmuted species, will create enrichment at the defect sinks. This is observed in the nucleation of Re-rich phases in the microstruc­tures of neutron-irradiated samples116,121 and the degradation in mechanical properties and transition to intergranular fracture in higher Re concentration alloys.62,120 Further information on RIS can be found in Chapter 1.18, Radiation-Induced Segregation.

Through modeling and experimental work, Erck and Rehn123 showed that the degree of segregation per dpa reaches a maximum for Mo-30 at.% Re (^45 wt% Re) near 1223 K and that for Mo-7 at.% Re (~13wt% Re) near 1473 K. While the Mo-5Re alloys irradiated up to 20 dpa show some limited ductility,109,118,119 the maximum irradiation tempera­tures were <1073 K and are therefore at or below the lower temperature limit expected for RIS.

The Mo—(41 and 47.5)Re alloys irradiated at 1073— 1373 K120 showed indications of RIS even at relatively low damage levels, part of which may have been a contribution of a thermal aging component, which in the unirradiated as-aged Mo-41Re and Mo—47.5Re showed increases in Re at the grain boundaries, lead­ing to precipitation of s — and w-phases at the grain boundaries in the 47.5Re containing alloy.115 Utilizing Mo-Re alloys with a more moderate Re content may improve the irradiation performance of these alloys, especially when considering higher doses and/or lon­ger irradiation times at temperatures at which thermal precipitation effects may further compound RIS influ­ences on mechanical properties.

Additional information on the fracture toughness data for Mo-Re alloys is also needed. Preliminary data by Scibetta and coworkers109 on precracked com­pact tension specimens of Mo-5Re showed reduc­tions in fracture toughness from unirradiated values of 17-23 MPa Vm at room temperature and 623 K, to ^11MPaVm for 0.35 dpa irradiated at 313 K, and 15MPaVm for 0.29 dpa irradiated at 643 K. These low irradiated values, at which no ductile crack growth was observed in the specimens, are a concern.

Recent work examining the irradiated properties of wrought, commercially available, ODS-Mo contain­ing lanthanum oxide particles has shown promising results.81,82,124 The fabrication methods produce a microstructure consisting of elongated grains with appreciable texturing and alignment of the oxide par­ticles. The high degree of working associated with fabrication produces a <2 pm grain size, which is stabilized from growth by the ODS particles. Irradia­tion of the ODS-Mo up to 13.1 dpa at 567 K and 883-882 K produced an increase in yield strength of 57-173%,124 while irradiation at 1143-1273 K produced a 10-34% increase. The irradiated tensile properties of ODS-Mo as a function of irradiation temperature and dose from the work of Cockeram and coworkers124 are shown in Figure 22. The increases in radiation strength are comparable to the higher limits for LCAC-Mo. The most striking result of the ODS work is the improvement in the DBTT for the irradiated samples.82,124 For 567 K irradiation to 12.3 dpa, the DBTT is 1073 K and is comparable to that of LCAC-Mo and TZM (Figure 16). How­ever, the DBTT for ODS-Mo irradiated to 13.1 dpa at 833-882 K is ^298 K, while that of LCAC-Mo is 573 and 973 K for TZM. For irradiation to 13.1 dpa at 1143-1209 K, the DBTT is 173 K, unchanged from the nonirradiated material, while those of LCAC-Mo and TZM are between 223-273 K.

The reduced susceptibility to irradiation embrit­tlement of ODS-Mo is in part due to the grain size reduction and presence ofthe oxide particles. Reduc­ing the distance of possible defect sinks such as grain boundaries and offering additional sites such as the oxide/matrix interface are particularly critical at lower irradiation temperatures at which defect mobility is limited. In addition, Cockeram and coworkers81,12 describe the fine but elongated grain structure as enhancing the plain strain condition acting on each plane of the lamina-shaped grains formed during the fracture process, inducing larger plastic deformation in irradiation-hardened material. This is also true for the HP-LCAC-Mo containing the high aspect ratio grains compared to other forms of LCAC-Mo produced. Currently, no fracture data on the irradiated

ODS-Mo are available. Unirradiated fracture tough­ness values are between 23 and 38 MPaVm, depending on the grain orientation tested.83

The 5 mm-wide ring-tensile specimens with a 1.5 mm-wide gauge section were prepared from the cladding of 12Cr-ODS steels (F94, F95, and 1DS) and 9Cr-ODS steels (M93).67 This type of specimen makes it possible to test mechanical properties in the hoop direction of the cladding. These ring-tensile samples were irradiated in the experimental fast reac­tor JOYO using the material irradiation rig at tem­peratures between 400 and 534 °C to fast neutron fluences ranging from 5.0 x 1025 to 3.0 x 1026nm~2 (E > 0.1 MeV). The yield strength of the irradiated samples as a function of test temperature is shown in Figure 39, together with that of the unirradiated ones.67 After irradiation, the yield strength of irra­diated F94, F95, and M93 cladding, is modestly higher (<10%) than that of the unirradiated ones at all test temperatures, due to irradiation hardening. Figure 40 plots uniform elongation before and after irradiation as a function of test temperature.67 Uniform elongation for unirradiated F94 and F95 cladding is almost the same at all test temperatures, and that of M93 is lower in relation to strength. Uniform elongation in the hoop direction for all three claddings is more than 3% at these test temperatures, though that of 1DS was particularly low (<1%) due to its microstructural anisotropy, as shown in Figure 17. Figure 40 indicates that there is no significant degradation in uniform elongations for F94, F95, and M93, due to irradiation. This indicates that the microstructural improvement by recrystallization or a-g-phase transformation is quite effective in maintaining well-balanced mechani­cal properties for ODS steel cladding, especially those of strength and ductility, not only for as-received con­ditions but also following irradiation.

In-pile creep rupture tests were conducted in JOYO using the Material Testing Rig with Tempera­ture Control (MARICO-2) as a new irradiation test device.68 The test specimens were prepared from the claddings of 9Cr-ODS steel (Mm14) and 12Cr-ODS steel (F14). Both end-plugs of the specimens were joined by means of pressurized resistance welding (PRW). The hoop stress was set by adjusting the pres­sure of the enclosed helium gas. To identify the rup­ture of time and specimens, a unique blend of stable

Table 3 Historical survey of yttrium-titanium-oxides reported to be stable under radiation Author Material Irradiation particle (dpa) Temperature (°C) Dose (dpa) Dose rate (dpa s 1) Result Pareige et al.74 12YWT 150 keV Fe 300 0.7 1.9 x 10~4 Stable dispersions Asano et al.75 MA957 1 MeV He (+) 450 150 2 x 10~3 Stable oxides 4 MeV Ni 650 Hide etal.76 MA957 42 keV He 475 200 1.0-1.4 x 10~2 Stable oxides at 25°C(+) 525 200 keV C~ 575 625 Hide etal.76 MA957 220 keV He at 25 °C 475 150 3.0 x 10~3 Stable oxides (+)3 MeV Ni+ 525 Little77 DT2203YO5 52 MeV Cr6 + (+) 475 50 3.0 x 10~4 Stable oxides 4 MeV He Saito et al.61 13Cr-0.5TiO2 -0.2Y2O3 1 MeV electron 400 12 2.2 x 10~3 Stable oxides 500 Kinoshita et al.78 13Cr-ODS (+) Nb, V, Zr 1 MeV electron 350 15 2 x 10~3 Stable oxides 450 Akasaka et al.79 9Cr and 12Cr-ODS JOYO 330 7.0 Not reporteda Stable oxides 400 2.5 450 14.0 500 15.0 Mathon etal.80 MA957 Thermal neutrons 325 0.8, 2.0 1 x 1014 ncm~2 Stable dispersions (OSIRIS) 3.5, 5.5 (E > 1 MeV) Monnet etal.63 DY EM10+Y2O3EM10 + MgO 1 MeV Helium 400 0.05 Not reported No change in oxide particles

Kimura et al.81 (13-19)Cr-4Al-ODS 300 20 1 x 10~4 No reported change in oxide size -500 aTypical fast reactor displacement rates in the driver fuel portion of the core are 1 x 10 6dpas 1. Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26-37.

aTypical fast reactor displacement rates in the driver fuel portion of the core are about 1 x 10 6dpas 1 Source: Reproduced from Allen, T. R.; Gan, J.; Cole, J. I.; et al. J. Nucl. Mater. 2008, 375, 26-37.
• F94 — F94 unirrad. □ F95 — F95 unirrad. A M93 M93 unirrad. О 1DS — 1DS unirrad.
8
sP O’
6
4
2
(3.56)
(0.45)

xenon and krypton tag gases was enclosed. The irradi­ation temperatures were 700, 725, and 750 °C, and the hoop stress ranged from 45 to 155 MPa. The maximum neutron dose reached 20 dpa. It was confirmed that in­pile creep rupture time is located within the out-of­pile data band, and there is no degradation in creep strength due to irradiation.68

MA957 and MA956 were irradiated in Fast Flux Test Facility (FFTF)-Materials Open Test Assembly (MOTA) at 420 °C up to 200 dpa.69 No voids were seen in this area, but precipitates did appear, which were expected to be a0. The results regarding the radiation damage resistance of ODS steels were highly encourag­ing. Evidence was apparent in both MA956 and MA957

of a0 precipitation, and in regions where recrystalliza­tion occurred before irradiation in MA957, a few voids were slightly observed. Gelles69 pointed out that these could be overcome by employing suitable alloy design and that ODS steel microstructures, when properly manufactured to provide a uniform oxide dispersoid in a structure, appear to be completely resistant to radiation damage at doses as high as 200 dpa.