To unfault a 1/3[1010](1010) dislocation loop in alumina, we must propagate a 1/3 [0001] partial shear dislocation across the loop plane.17 This is described by the following dislocation reaction:

1[10l0] + ±[0001] ! i[10T1] (prismatic) [9

faulted loop partial shear unfaulted loop

This reaction is symmetric with that shown in eqn [8] for unfaulting a (0001) basal loop in alumina. The reaction in eqn [9] is shown graphically in Figure 4.

When we pass a 1/3 [0001] shear through a 1/3[1010](1010) dislocation loop, the cation planes beneath the loop assume new registries such that in eqn [6], ajpj and a2p2 commute as follows: a1p1 ! a2p2 ! a1p1. The anion layers beneath the loop are left unchanged (B! B, C! C). In addition, the Al-only cation layers are unchanged (p0 ! p0). Taking the faulted (0001) stacking sequence in eqn [2] and assuming that the planes to the right
are above the ones on the left, we perform the 1/3 [0001] partial shear operation as follows:

(P’S («AC) (P’S) (a.$2C) (P’S) (a, p,C) |(P’g|(a, frC) (P’S) (oftC) (faulted)

(a2P2C) (P’S) (a, p,C)

(P’S (a, P,C) (P’S («2P2C) (P’S (a, p,C) (P’S (012P2C) (P’S) (a, p,C) (unfaulted)

[10]

After propagating the partial shear through the loop, we are left with an unfaulted layer stacking sequence. The resultant dislocation loop, 1/3 [1011], is a perfect dislo­cation. The resultant 1/3 [1011] (0001) dislocation is a mixed dislocation, in the sense that the Burgers vector is canted (not normal) relative to the plane of the loop.

1.06.6.1 ISHI Studies and Thermal Stability of Nanofeatures in NFA MA957

In this section, we describe the status of developing a potentially transformational new class of materials, we call nanostructured ferritic alloys (NFA), with emphasis on He management for radiation tolerance as discussed in Section 1.06.3.2 , NFA manifest high tensile, creep and fatigue strengths, unique ther­mal stability, and remarkable irradiation tolerance. The outstanding characteristics of NFA result from the presence of an ultrahigh density of Y-Ti-O rich nanofeatures (NF). The multifunctional NF, which are remarkably thermally stable, impede dislocation climb and glide, enhance SIA-vacancy recombina­tion, and, perhaps most importantly, trap He in small, high-pressure gas bubbles. The bubbles reduce the amount of He reaching GBs, thus mitigating tough­ness loss at lower temperatures and potential degra­dation of creep rupture properties at higher temperature. He trapped in a high number density of small bubbles also mitigates many other manifesta­tions of irradiation effects, including void swelling.

As discussed in Section 1.06.2, ISHI in mixed spectrum fission reactor irradiations provides an attractive approach to assessing the effects of He-dpa synergisms. To reiterate, the basic idea is to use Ni (or B or Li)-bearing implanter layers to inject high-energy a-particles into an adjacent material that is simultaneously undergoing fast neutron-induced displacement damage. The a-particles can be pro­duced by two-step 58Ni(nth, g)59Ni(nth, a) thermal neu­tron (nth) reactions. A series of ISHI irradiation experiments have been carried out in HFIR. Micron — scale NiAl injector coatings were used to uniformly implant a-particles to a depth of «5-8 mm in TEM disks for a large matrix of alloys irradiated over a wide range of temperatures and dpa at controlled He/dpa ratios ranging from ^1 to 40appmdpa~ . Here, we compare the cavity structures in a 14Cr NFA, MA957, to those in an 8Cr FMS, F82H, following HFIR irradiation at 500 °C to 9 dpa and 380appm He. The experimental details are given

23,50,51

elsewhere.

Through-focus sequence TEM images were used to characterize the bubbles and voids, with care taken to avoid surface artifacts. Bubble-like features were generally not found in the unimplanted regions of either MA957 or F82H. As illustrated in a typical underfocused image in Figure 40(a), a high number density (Nb « 4.3 x 1023 m~3) of very small (average db « 1.2 nm) bubbles are observed in the NFA. The inserts in Figure 40(a) show examples of the decora­tion of larger features with cavities. Image overlap analysis suggests that most bubbles are associated with a similar number density («6.5 x 1023 m~3) of NFs.51,2 0 However, the degree of bubble-NF associ­ation has not yet been fully demonstrated and quan­tified. The boundary in MA957 in Figure 40(a) appears to be relatively cavity free, and there does not seem to be a large nearby NF-cavity denuded zone. Assuming equal partitioning of all the 380 appm He to 4.3 x 1023 m~3 bubbles («80 He atoms/bubble), g = 2Jm~2 is consistent with ty « 2g/kmkT« 0.6nm at 500 °C, where к is the real gas compressibility factor, which is in remarkable agreement with the measured average cavity size. Thus, we conclude that the He is primarily stored in near-equilibrium bubbles at a capillary pressure of 2g/r, « 6500 MPa in this case. A higher He content of2000 appm parti­tioned to the same number of bubbles («400 He atoms/bubble) increases гу to «1.1 nm, still far below the critical size for void formation, which is estimated to be well over 10 nm. The 9 dpa irradia­tion at 500 °C has no observable effect on the NFs.

As shown in Figure 40(b) and 40(c), a lower number density of (Nb « 5.3 x 1022 m~3) of somewhat larger (db « 2.1 nm) bubbles (the smaller population of cavities in this case) are observed in F82H, along with much larger faceted cavities; the larger cavities are likely voids. The smaller matrix bubbles in F82H are clearly formed on dislocations, as high­lighted by the black and white contrast insert in Figure 40(c). Figure 40(d) shows that the cavity size distribution is much narrower in the MA957, with a maximum diameter of less than 2.5 nm. In contrast, the largest diameters exceed 10 nm in F82H and it appears that a bimodal bubble-void

(d) ^bubble (nm)

Figure 40 (a) Underfocused TEM image of in situ He injected MA957; (b) and (c) underfocused TEM image of in situ He injected F82H; (d) cavity size distributions in MA957 versus F82H.

cavity size distribution is developing.51 Model-based extrapolation of these results suggests that significant swelling may develop at higher He and dpa.

The ISHI results suggest that NF are effective in trapping He in fine-scale bubbles at least up to 500 °C. The sink strength of 4.3 x 1023 m~3, 1.2 nm bubbles is Zb ~ 3.2 x 1015m~2, which is significantly higher than the typical total sink strength in FMS alloys (<«1015m~2). Presumably, Zb could be increased by an additional factor of at least two to three in alloys with larger numbers of NFs and bigger associated bubbles at higher He levels.

It is important to emphasize that both He bubbles and NFs are key to highly irradiation-tolerant alloys. The primary role of the NFs is to provide preferred and thermally stable sites for forming bubbles. In principle high densities of stable bubbles can seques­ter He up to high levels and serve as sinks that, in principle, mitigate all manifestations of displacement damage. Helium management schemes based on these principles are critical to developing fusion energy and may also play a role in fission applications intended to reach very high dpa levels. The effect of He bubbles on defect damage accumulation is being
investigated using the in situ implantation technique, including at higher He and dpa, for a wide range of irradiation temperatures and large alloy matrix.

The diffusion of the point defects created by the irradiation and their subsequent absorption at
dislocations and interfaces in the material is the most essential process that restores the material to its almost normal state. The adjective of‘almost nor­mal’ is anything but a casual remark here, but it hints at some subtle effects arising in connection with the long-range diffusion that constitute the root cause for the gradual changes that take place in crystalline materials exposed to continuous irradiation at ele­vated temperatures. If these effects were absent, then a steady state would be reached in the material sub­ject to continuous irradiation at a constant rate and temperature in which the rate of defect generation would be balanced by their absorption at sinks, mean­ing the above-mentioned dislocations and interfaces. As vacancies and self-interstitials are created as Frenkel pairs in equal numbers, they would also be absorbed in equal numbers at these sinks. At this point, the microstructure of these sinks would also be in a steady, unchanging state. While this steady state would be different from the initial microstruc­ture or the one reached at the same temperature but in the absence of irradiation, it would correspond to material properties that reached constant values.

The subtle effects alluded to in the above remarks arise from the interactions of the point defects with strain fields created both internally by the sinks and externally by applied loads and pressures on the materials that constitute the reactor compo­nents. The internal strain fields from sinks give rise to long-range forces that render the diffusion migration nonrandom, while the external strains induce aniso­tropic diffusion throughout the entire material. In the

next section, we derive the diffusion equations for cubic materials to clearly expose these two funda­mental effects.

1.02.6.1 Diffusion Mechanisms

Diffusion in ceramic materials is a process enabled by defects and controlled by their concentrations. Owing to the existence of separate sublattices, cat­ion and anion diffusion is restricted to taking place separately (i. e., without exchange of anions and cations), which is one of the main differences with respect to diffusion in other materials.22 Therefore, mechanistically, diffusion theory is applied in cera­mics by considering the anion and cation sublat­tices separately. Interestingly, it has recently been suggested23 that where there is more than one cat­ion sublattice, cations can move on an alternate sublattice through the formation of cation antisite defects. Finally, it can be the case that ion transport in one of the sublattices is more pronounced. For example, in oxygen fast ion conductors, oxygen self-diffusion is faster than cation diffusion by orders of magnitude.24-26

Figure 21 The interstitialcy mechanism of diffusion. The red and blue ions are lattice species, the blue ion with the red perimeter is initially an interstitial species but becomes a lattice species.

to an interstitial site. This mechanism is important for the diffusion of dopants such as boron in silicon.33 In hyperstoichiometric oxides, such as La2NiO4+d, it was recently predicted that oxygen diffuses predomi­nantly via an interstitialcy mechanism.26

Collective mechanisms involve the simultaneous transport of a number of atoms. They can be found in ion-conducting oxide glasses22 and have been pre­dicted during the annealing of radiation damage.34 Finally, in the interstitial-substitutional exchange mechanism, the impurities can occupy both substitu­tional and interstitial sites.22 One possibility for the interstitial atom is to migrate in the lattice until it encounters a vacant site, which it then occupies to become a substitutional impurity (dissociative mech­anism).22 Another possibility for the impurity inter­stitial atom is to migrate in the lattice until it displaces an atom from its normal crystallographic site, thus forming a substitutional impurity and a host interstitial atom (kick-out mechanism). The intersti­tial-substitutional mechanism has been encountered in zinc diffusion in silicon and gallium arsenide.[6]

Naturally, there are potential energy barriers hin­dering the motion of atoms in the lattice. The activa­tion energy associated with the barriers may be overcome by providing thermal energy to the system. The jump frequency o of a defect is given by[7]

f AGm

o =n exp( ~w

where DGm is the free energy required to transport the defect from an initial equilibrium position to a saddle point and n is the vibrational frequency. In real materials, the atomic transport may be locally affected by interactions with other defects especially if the defect concentration is high.35-37

M. L. Grossbeck

.

1.04.1 Introduction

The most commonly considered mechanical proper­ties of metals and alloys include strength, ductility, fatigue, fatigue crack growth, thermal and irradiation creep, and fracture toughness. All these properties are important in the design of a structure that is to experi­ence an irradiation environment. While determining the mechanical properties of irradiated materials, ten­sile properties, typically yield strength, ultimate tensile strength, uniform elongation, total elongation, and reduction of area are the most commonly considered because they are usually the simplest and the least costly to measure. In addition, the tensile properties can be used as an indicator of the other mechanical

properties. Space in a reactor or in an accelerator target is often so limited that the larger specimens required for fatigue and fracture toughness testing are not prac­tical; consequently, the number of specimens that can be irradiated is so small that a meaningful test matrix is not possible. Shear punch testing of 3-mm diame­ter disks, typically used as transmission microscopy specimens, was developed to address the problem of irradiation space. Although much information can be obtained from shear punch testing, the tensile test remains the most reliable indicator of strength and ductility. For these reasons, the tensile test is usually the first mechanical test used in determining the irradiated properties of new materials. This chapter addresses the tensile strength and ductility of alloys.

In situ He implantation (ISHI) in mixed spectrum fission reactors is a very attractive approach to assess the effects ofHe-dpa synergisms in almost any mate­rial that avoids most of the confounding effects of doping. The basic idea is to use an implanter layer, containing Ni, Li, B, or a fissionable isotope, to inject high-energy a-particles into an adjacent sample simultaneously undergoing neutron-induced dis­placement damage. Early work proposed implanting He using the decay of a thin layer of a-emitting isotope adjacent to the target specimen.47 However,
the isotope decay technique produces few dpa at a very high He/dpa. The first proposal ISHI in a mixed fast (dpa)-thermal (He) spectrum proposed using 235U triple fission reactions to inject «16MeV a-particles uniformly in steel specimens up to 50 mm thick; the 50 mm thickness permits tensile and creep testing as well as microstructural characterization and mechanism studies at fusion relevant dpa rates and He/dpa ratios.31 The triple fission technique was applied to implanting ferritic steel tensile specimen, albeit without complete success.48 A much more prac­tical approach is to use thin Ni-bearing implanter foils to uniformly deposit He up to a depth of ~8 mm in Fe in a thick specimen at controlled He/dpa ratios.49

As illustrated in Figure 5(a)-5(c), there are at least three basic approaches to implanter design. Here we will refer to thin and thick, specifically meaning a specimen (ts) or implanter layer (ti) thick­ness that is less than or greater than the corre­sponding a-particle range, respectively. Ignoring easily treated difference in the a-particle range (Ra) and atom densities in the injector and specimens for simplicity, thick implanter layers on one side of a thick specimen produce linearly decreasing He con­centration (XHe) profiles, with the maximum concen­tration at the specimen surface that is one half the concentration in the bulk injector material, XHeo = XHei/2 (Figure 5(a)). If a thin specimen is implanted from both sides by thick layers, the He concentration

XHei Ri 2 Ra

У

X

R

(a)

(c)

is uniform and equal to one half that in the bulk injector material (Figure 5(b)). In contrast, a thin layer implants a uniform concentration of He to a depth of the Ra — ti. In this case, the He concentration in both the implantation layer and specimen is equal and lower than in the bulk (XHei) as XHes = tiXHei/ 2Ra (Figure 5(c)). Thus, the He/dpa ratio can be controlled by varying the concentration of the iso­tope that undergoes n, a reactions with thermal neu­trons, ti, and the thermal to fast flux ratio.

ISHI experiments were, and continue to be, carried out in HFIR using thin (0.8-4 gm) NiAl coating layers on TEM disks for a large matrix of Fe-based alloys for a wide range of dpa, He/dpa (<1-40 appm He/dpa), and irradiation temperatures. In this case, 4.8 MeV a-particles produce uniform He concentration to a depth of «5-8 gm (Figure 5(c)). Further details are given elsewhere.50 The first results of in situ implan­tation experiments in HFIR have been reported and are discussed in Section 1.06.6.23,51-53 The technique has also been used to implant SiC fibers irradiated in HFIR.50 More recently, the two-sided thick Ni implanter method was used to produce He/dpa ratios «25appm/dpa in thin areas of wedge-shaped specimen alloys irradiated in the advanced test reactor to «7 dpa over a range of high temperatures.54

1.01.8.1 Effective Medium Approach

The fate of the radiation-produced atomic defects, namely self-interstitials and vacancies, is mainly determined at elevated temperatures by their diffu­sion from the places where they were created to the sinks where they are absorbed or annihilated, as in the encounter of an interstitial and a vacancy. As there are many sinks within each grain of an irradiated material, the spatial distribution of the atomic defects requires the solution of a very complex diffusion problem. Clearly, some approximations must be sought to arrive at an acceptable solution. First, we assume that the rate of defect generation, that is, the rate of displacements, is constant, and the defect concentrations between the sinks no longer changes with time or adjust rapidly when the number of sinks and their arrangement changes. Then within the regions between sinks, the diffusion fluxes are stationary, that is,

j = ~dXD, J (Г)С (r)+F'(r)C(r) [128]

is independent of time.

The task is then to divide the solid into cells, each containing one individual sink, and to solve in each diffusion equations of the following type

V • j = P — recombination [129]

for each mobile defect. Here, P is the rate of defect production per unit volume generated by the radia­tion, and the other terms represent the rates of defect disappearance by recombination with other migrat­ing defects. On the outer boundary of this cell, the defect concentration C must then match the concen­tration in adjacent cells occupied by other sinks, and its gradient must vanish. This cellular approach has been pioneered by Bullough and collaborators,51 but the drift term, the second term in eqn [128], is omit­ted when solving the diffusion equation. Its effect is subsequently taken into account by changing the actual sink boundary into another, effective boundary at which the interaction energy between the sink and the approaching defect becomes of the order of kT, k being the Boltzmann constant.

An alternate approach52,53 is to view a particular sink as embedded in an effective medium that main­tains an average concentration of mobile defects at a distance far from this particular sink, and to neglect production and losses of mobile defects nearby. In this approximation, the r. h.s. of eqn [129] is set to zero, and the outer boundary condition far from the sink is that C approaches a constant value C that remains to be deter­mined later from average rate equations. The diffusion equation is now solved with and without the drift term and the resulting defect current to the sink is evaluated. The ratio of the currents with and without the drift defines the sink bias factor. It is thereby possible to define unambiguously bias factors for each type of sink, and with these determine the net bias for a given microstructure. It is this embedding approach that we follow here to evaluate the bias of a sink.

The first attempt to determine the dislocation bias by solving the diffusion equation with drift appears to have been made by Foreman.54 He employed a cellu­lar approach retaining only the defect production term. Furthermore, anticipating small bias values, the drift term was treated by perturbation theory and numerous approximations were introduced in the derivation. The intent was to obtain rough esti­mates; nevertheless, Foreman concluded that the bias was larger than the empirical estimate. Shortly there­after, Heald55 employed the embedding approach and used the solution of Ham56 in the form presented by Margvalashvili and Saralidze.57 We shall return to this solution below, as it is the only analytical one

known. However, this solution applies only when the mobile defect is modeled as a center of dilatation (CD). For more general interactions, Wolfer and Askin58 developed a rigorous perturbation theory that includes also the interaction induced by exter­nally applied stresses. The latter was shown to result in radiation-induced creep. They also compared the bias obtained from the perturbation theory carried out to second order with the bias from Ham’s solu­tion, and demonstrated that both results agree only for weak centers of dilatation. Vacancies can be con­sidered as such weak centers, but interstitials cannot. As a result, perturbation theory to second order is in general insufficient to evaluate the dislocation bias.

He cavities (T >> Stage V)

Irradiation at temperatures near or above 0.5 TM typ­ically results in only minor microstructural changes due to the strong influence of thermodynamic equi­librium processes, unless significant amounts of impu­rity atoms such as helium are introduced by nuclear transmutation reactions or by accelerator implantation. When helium is present, cavities are nucleated in the grain interior and along grain boundaries. The cavity size increases and the density decreases rapidly with increasing temperature. Figure 15 compares the helium cavity density for various implantation and neutron irradiation conditions in austenitic stainless

steels as a function of temperature.118,119 The tem­perature dependence of the cavity density is dis­tinguished by two different regimes. At very high temperatures, the cavity density is controlled by gas dissociation mechanisms with a corresponding high activation energy, and at lower temperatures by gas or bubble diffusion kinetics.118 The cavity density decreases by nearly two orders of magnitude for every 100 K increase in irradiation temperature in this very high temperature regime. The helium cavity densities in materials irradiated at low temperatures (near room temperature) and then annealed at high temperature are typically much higher than in materials irradiated at high temperature, due to excessive cavity nucleation that occurs at low temperature. In the absence of applied stress, the helium-filled cavities tend to nucle­ate rather homogeneously in the grain interiors and along grain boundaries. If the helium generation and displacement damage occurs in the presence of an applied tensile stress, the helium cavities are preferen­tially nucleated along grain boundaries and may cause grain boundary embrittlement.12

1.05.2.1 Introduction to Radiation Damage in Alumina and Spinel

a-Al2O3 and MgAl2O4 are two of the most important engineering ceramics. They are both highly refrac­tory oxides and are used as dielectrics in electrical applications (capacitors, etc.). Both a-Al2O3 and MgAl2O4 have been proposed as potential insulat­ing and optical ceramics for application in fusion reactors.1-3 In a magnetically confined fusion device, these applications include (1) insulators for lightly shielded magnetic coils; (2) windows for radio­frequency heating systems; (3) ceramics for structural applications; (4) insulators for neutral beam injectors; (5) current breaks; and (6) direct converter insula­tors.3 Such devices in a fusion reactor environment will experience extreme environmental conditions, including intense radiation fields, high heat fluxes and heat gradients, and high mechanical and electri­cal stresses. A special concern is that under these extreme environments, ceramics such as a-Al2O3 and MgAl2O4 must be mechanically stable and resis­tant to swelling and concomitant microcracking.

Over the last 30 years, many radiation damage experiments have been performed on a-Al2O3 and MgAl2O4 under high-temperature conditions by a number of different research teams. Figure 1 shows the results of one such study, where the high tem­perature, neutron irradiation damage responses of a-Al2O3 and MgAl2O4 are compared. The plot in Figure 1 was adapted from Figure 1 in an article by Kinoshita and Zinkle,4 based on experimental data obtained by Clinard et a/. and Garner et a/.5-

The neutron (n) fluence on the lower abscissa in Figure 1 refers to fission or fast neutrons, that is, neutrons with energies greater than ~0.1 MeV. Figure 1 also shows the equivalent displacement damage dose on the upper abscissa, in units of

displacements per atom (dpa). These dpa estimates are based on the approximate equivalence (for cera­mics) of 1 dpa per 1025nm~2 (En > 0.1 MeV).9

Figure 1 shows a stark contrast between the radiation damage behavior, particularly the volume swelling behavior, between a-Al2O3 and MgAl2O4. Specifically, MgAl2O4 spinel exhibits no swelling in the temperature range 658-1100 K, for neutron fluences ranging from ^3 x 1026 to 2.5 x 1027nm~2 (3-200 dpa). On the contrary, a-Al2O3 irradiated at temperatures between 925-1100 K exhibits signifi­cant volume swelling, ranging from ~1 to 5% over a fluence range of 1 x 1025 to 3 x 1026nm~2 (1-30 dpa). The purpose of the following discussion is to reveal the reasons for the tremendous dispar­ity in radiation-induced volume swelling between alumina and spinel.

Figure 2 shows a bright-field (BF) transmission electron microscopy (TEM) image that reveals the microstructure of a-Al2O3 following fast neutron irra­diation at T= 1050 K to a fluence of 3 x 1025nm~2

(^3dpa). The micrograph reveals a high density of small voids (2-10 nm diameter), arranged in rows along the c-axis of the hexagonal unit cell for the a-Al2O3. When voids are arrayed in special crystallo­graphic arrangements, as in Figure 2, the overall structure is referred to as a void lattice. Figure 2 shows the underlying explanation for the pronounced volume swelling of a-Al2O3 shown in Figure 1, namely the formation of a void lattice with increasing neutron radiation dose. This phenomenon is well known in many irradiated materials, both metals and ceramics, and is referred to as void swelling. Susceptibility to void swelling is a very undesirable material trait and basically disqualifies such a mate­rial from use in extreme environments (in this case, high temperature and high neutron radiation fields). It should be noted that TEM micrographs (not shown here) obtained from MgAl2O4 spinel irradiated under similar conditions to those in Figure 2 show no evidence of voids of any size.

The most recent source ofsignificant information on He effects on bulk microstructures and properties is from SPN irradiations of both AuSS and FMS to maximum doses of «20dpa and «1800 appm He at temperatures up to «450 °C.

1.06.4.1 Microstructural Changes

The microstructures of FMS alloys irradiated in STIP at temperatures below about 400 °C are domi­nated by defect clusters and bubbles, characterized by their respective number densities (Nd/b) and dia­meters (dd/b). Table 2 summarizes some recent TEM observations on FMS and AuSS after neutron and SPN irradiations. Figure 24 shows the development of defect structures in FMS F82H and T91 irra­diated in STIP-I209 at irradiation temperatures up to «360 °C. Note that the temperatures, dpa, and He levels (dose) are correlated with one another due to their mutual dependence on the proton flux. The microstructures in both FMS alloys are similar and are composed of small (1-3 nm) defects (likely small SIA cluster-dislocation loops, Figures 24(a)-24(d)), along with a lower number of larger dislocation loops (>3nm). Figure 25(a) plots the loop data for F82H as a function of irradiation temperature for the STIP-I and — II irradiations, along with some neutron data, including from an ISHI experiment. The STIP-I irradiation conditions were «10-12 dpa with He levels that increase with temperature from 560 to 1115 appm. The STIP-II data were at «10-19 dpa with between 750 and 1790 appm He. Note that the STIP-I irradiation initially ran at a much lower irradiation temperature while the STIP-II irradiation was more isothermal. Ignoring these confounding factors, the overall trends show that the dd increases with temperature while Nd decreases, after apparently peaking at «250 °C. The neutron data point at 10 dpa and 310 °C, with a low He content, following irradiation in the Peten High

Flux Reactor (HFR), is similar to the STIP-II data at higher He levels. The neutron data 400 °C from an ISHI study with 90 appm He at 3.8 dpa215 appears to be more consistent with the overall data trends than the corresponding STIP-II data point; however, this may be because of the lower dose in this latter case (Figure 25(b)).

As shown in Figure 26, the He bubble populations in the two FMS irradiated in STIP-I to 10 dpa at 295 °C are also similar. Due to the resolution limits, conventional TEM cannot image bubbles smaller than about 1 nm. Thus, bubbles are visible only at He concentrations and temperatures above «500 appm and «170 °C, respectively, in the STIP database.66 Figure 25(b) shows that 4> increases and Nb decreases with increasing temperature in the STIP — I and — II irradiations. The neutron data at 500 °C, from an ISHI study with 380 appm He at 9 dpa, is consistent with the STIP data trends.215 The smallest bubbles found in the SPNI studies, with an average diameter of «0.7 nm, were observed in F82H from the STIP-I irradiation to 9.9 dpa/560 appm He at 175 °C. Bubbles were not observed at lower tempera­tures and He levels.17,66,209,216,217 In this case, the He is presumably located in a very high concentration of subvisible He-V clusters, which may be overpressur­ized. At higher temperatures of 350 and 400 °C in the STIP-II irradiation, the apparent sizes of the bubbles are much larger than those for STIP-I.

Figure 27 shows the cavity structure in two F82H samples irradiated in STIP-II (left) and STIP-III
(right) to similar dose and He concentration at nomi­nal temperatures of 400 ± 50 and 440 ± 50 °C. Some of the cavities in the 400 °C STIP-II irradiation have transitioned from bubbles to larger voids, forming a bimodal size distribution, while there is a larger Nb of smaller bubbles in the 440 °C STIP-III case, with a monotonic size distribution. These differences are believed to be due to the fact that the STIP-III irradiation also ran at lower temperatures of about 200 °C during the initial phase of the irradiation up to «0.3 dpa and «30 appm. It is believed that the high Nb («5.1 x 1023 m~3) nucleated during this transient and then grew at higher temperatures and He levels. Thus, due to the large Nb, voids did not form in this case. In contrast, voids formed in the STIP-II irradi­ation, since it was more isothermal, resulting in a lower Nb («2.4 x 1023m~3) and closer to the peak swelling temperature for FMS of 400 °С.1 Large voids that are associated with precipitates are shown in Figure 28(a) for a STIP-II irradiation at 400 ± 50 °С. The corresponding microstructure for irradiations at lower irradiation temperatures of «350 ± 50 °С is again composed of a high density of small bubbles. Thus, the increase in the average sizes of the cavities in Figure 25 at 350 and 400 °С in STIP-II is due to the bimodal mix of bubbles and voids.

The strings of bubbles seen in Figure 28(b) also indicate a strong association between bubbles and dis­locations. Further, neither bubble denuded zones near GBs nor a significant number of grain boundary

bubbles have been found in STIP samples investigated to date. This suggests that grain boundary bubbles may be too small to image.

SPN irradiations also produce defect clusters, faulted Frank loops, and bubbles in AuSS (316L

60

50

40

10

Figure 25 Irradiation temperature dependences of (a) loop size and density and (b) bubble size and density of F82H specimens irradiated in STIP with SPN, in HFR with fission neutrons, and in HIFR with in situ He implantation (ISHI).

Figure 26 He bubble structures of (a) F82H irradiated to 9.7 dpa/670appm He and (b)T91 irradiated to 10.1 dpa/725appm He at 295°C. The scale shown in (b) is the same for (a). Reproduced from Jia, X.; Dai, Y. J. Nucl. Mater. 2003, 318, 207.

average sizes between 2.5 and 3.9 nm. In contrast, neutron irradiations of AuSS produce a lower density of larger loops. For example, mixed spectrum reactor irradiations at 400 °C, which produce smaller but significant amounts of He compared with the SPNI case, result in «30 times fewer and «6 times larger loops. The differences are even larger for fast reactor irradiations with much lower He levels. Thus, it appears that high He results in significant refinement of the loop structures in AuSS.

Figure 29 shows small 1-2 nm bubbles in 316LN AuSS for both 285 °C irradiations to 9.3 dpa and 705appm He at 285 °C and 19.4 dpa and 1800 appm He at 425 °C. Unlike the case of FMS irradiated at «400 °C, no large voids were observed in this case. The refinement of the bubble structures in SPNI with high He levels is even more profound. The mixed spectrum reactor irradiations spectrally tailored to produce «11 appm He/dpa produced 60 times fewer cavities with a bimodal distribution of bubbles and voids compared with the SPNI case with «7 times more He. The average cavity diameter is about 2.4 times larger in the mixed spectrum
neutron case. These observations are highly consis­tent with the concepts described in Section 1.06.3 with regard to higher He and Nb suppressing void formation. Again, however, the likely effects of the temperature history confound quantitative interpre­tations of both the loop and cavity microstructures in the AuSS as well as FMS.