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Category Archives: Comprehensive nuclear materials
Distribution of Fluence Within an Individual Moderator Brick
Having obtained the component mean fluence, temperature, and weight loss, the variation of these parameters throughout the particular component of interest is required.
The fluence reduces exponentially away from the fuel in the radial direction, but is influenced by surrounding fuel sources. The exact distribution is usually calculated using a reactor physics code for a 5 x 5 array pertinent to the area of interest. Figure 14 is an example for the Windscale Piles.
5% of the reactor heat which is generated within the graphite. Graphite weight loss variation within a component is more complex and is calculated by various empirical industry codes. If the axial variation in fluence, temperature, and weight loss along the brick length is deemed to be important, three-dimensional physics, temperature, and weight loss calculations will be required.
Radiation Effects
A fair amount of data is available for radiation response of vanadium alloys partly because they were candidates of cladding materials of LMFBR. For example, void swelling is known to be quite small if the alloy contains Ti. However, data are limited for V-4Cr-4Ti because this composition was decided as the reference one for fusion only recently. For this alloy, the feasibility issues of radiation effects are considered to be loss of ductility at lower temperature, embrittlement enhanced by transmutant helium at high temperature, and irradiation creep at intermediate to high temperature.
The mechanism of the loss of uniform elongation of vanadium alloys at relatively low temperature (<673 K) and low dose (~0.1 dpa) has been a longterm research subject. Microstructural observation after tensile tests showed that radiation-induced defect clusters were lost in layer structures and the defect-free zones were accompanied by dislocation channels as shown in Figure 16.37 This fact implies flow localization during deformation. Although the mechanism of the flow localization needs further investigation, it is inferred that interaction ofdislocations with
radiation-induced defect clusters, precipitates, or complexes of the two species is responsible. If the precipitates, most likely Ti-CON, play the role in this process, reduction of impurities in the matrix can improve the properties. Figure 17 compares the uniform elongation after irradiation for V-(4-5)Cr-(4-5)Ti alloys and those with doping of Al, Si, and Y. The significant increase in uniform elongation by the addition of Al, Si, and Y, which are known as getters of interstitial impurities such as O, N, and C in the matrix, suggests that the reduction of the interstitial impurities in solution enhances the radiation resistance.38 The effects of interstitial impurities on the formation of dislocation loops and precipitates were investigated by ion irradiations. Figure 18 shows temperature dependence of the densities of loops and precipitates.39 The loop density was not influenced by O level, but the precipitate density increased with O level below 973 K.
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Helium embrittlement is a critical issue, which is thought to determine the upper temperature limit for vanadium alloys. Past experimental evaluations of the helium effects involved uncertainties because controlled generation of helium during irradiation in a similar manner to that in fusion condition has been quite difficult. As a result, the past evaluation of the helium effects varied from weak to very strong.3 The Dynamic Helium Charging Experiment (DHCE) using fission reactors40 is one of the few potential neutron irradiation experiments with controlled variation of He/dpa ratio including typical fusion
conditions. DHCE is highly anticipated as a potential method to extend our understanding of the helium effects. However, for conclusive evaluation, a 14MeV neutron source is certainly necessary.
The irradiation creep tests have made progress recently, partly because of the progress in fabricating high quality pressurized creep tube specimens with reduced impurity levels. Figure 19 shows the normalized creep strain as a function of applied stress by irradiation in HFIR and JOYO in Li and Na
US-832665 698 K-Li (HFIR) NIFS-HEAT-2 698 K-Li (HFIR) NIFS-HEAT-2 731 K-Na (JOYO) о 2p
50 100 150 200
Applied stress (MPa)
Figure 19 Creep strain as a function of applied stress for V-4Cr-4Ti (US-832665 and NIFS-HEAT-2) irradiated in Li (HFIR) and Na (JOYO) environments. The creep strain was normalized as that at two displacements per atom. Reproduced from Fukumoto, K.; Narui, M.; Matsui, H.; etal. J. Nucl. Mater. 2009, 386-388, 575-578.
environments. The data also compare the performances of US and Japanese reference alloys.41 It was found that the creep strain rate exhibited a linear relationship with the effective stress up to 150 MPa at ^700 K and the differences with the environments and the heats are small.
Differences in Neutron Spectra
There are significant differences in neutron spectra for water-cooled, sodium-cooled, and other types of fission-based reactors. It should be noted that there is a conventional but slightly misleading practice to differentiate between ‘fast’ and ‘thermal’ reactors. Thermal reactors have a significant portion of their spectra composed of thermal neutrons. Thermalized neutrons have suffered enough collisions with the moderator material that they are in thermal equilibrium with the vibrations of the surrounding atoms. Efficient thermalization requires low-Z materials such as H, D, and C in the form of water, graphite, or hydrocarbons. At room temperature the mean energy of thermalized neutrons is 0.023 eV.
The designation ‘fast’ reactor, as compared to ‘thermal’ reactor, refers to the portion of the neutron spectrum used to control the kinetics of ascent to full power for each type of reactor. As shown later, this practice incorrectly implies to many that fast reactors have ‘harder’ neutron spectra than do ‘softer’ thermal reactors. Actually, the opposite is true.
Examples of typical flux-spectral differences in fission-based reactors are shown in Figures 2-5. The local spectrum at any position is determined primarily by the fuel (U, Pu) and fuel type (metal, oxide, carbide, etc.), the coolant identity and density, the local balance of fuel/coolant/metal as well as the proximity to control rods, water traps, or core boundaries. Additionally, it is possible to modify the neutron spectra in a given irradiation capsule by including in it
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or enclosing it with a moderator or absorber. Metal hydrides are used in fast reactors to soften the spectrum, while in mixed-spectrum reactors the thermal — to-fast ratio can be strongly reduced by incorporating elements such as B, Hf, Gd, and Eu.
The most pronounced influence on neutron spectra in fission reactors arises from the choices of coolant and moderator, which are often the same material (e. g., water). Moving from heavy liquid metals such as lead or lead-bismuth to lighter metals such as sodium leads to less energetic or ‘softer’ spectra.
Use of light water for cooling serves as a much more effective moderator. Counterintuitively,
however, this leads to both more energetic and less energetic spectra at the same time, producing a two — peaked ‘fast’ and ‘thermal’ distribution separated by a wide energy gulf at lower fluxes.
Such two-peaked spectra are frequently called ‘mixed spectra.’ The ratio of the thermal and fast neutron fluxes in and near such reactors can vary significantly with position and also with time.4 Using heavy water, we obtain a somewhat less efficient moderator that does not absorb neutrons as easily as light water, but one that produces an even more pronounced two-peak spectral distribution where the thermal-to-fast neutron ratio can be very large.
These spectral differences lead to strong variations between various reactors in the neutron’s ability to displace atoms and to cause transmutation. Depending on the reactor size and its construction details there can also be significant variations in neutron spectra and ‘displacement effectiveness’ within a given reactor and its environs, especially where more energetic neutrons can leak out of the core. Examples of these variations of displacement effectiveness for fast reactors are shown in Figures 6 and 7. Compared to fission-derived spectra, there are even larger spectral differences in various fusion or spallation neutron devices.
The reader should note the emphasis placed here on flux-spectra rather than simply spectra. If we focus only on light water-cooled reactors for example, there are in general three regimes of neutron flux of relevance to this review. First, there are the relatively low fluxes typical of many experimental reactors that
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can produce doses of 10 dpa or less over a decade. Second, there are moderate flux reactors that are used to produce power that can introduce doses as high as 60-100 dpa maximum over a 30-40 year lifetime and finally, some high-flux thermal reactors that can produce 10-15 dpa year in stainless steels.
Most importantly, fast reactors also operate in the high-flux regime, producing 10-40 dpa year-1. Therefore, the largest amount of published high — dpa data on stainless steels has been generated in fast reactors. Some phenomena observed at high exposure, such as void swelling, have been found to be exceptionally sensitive to the dpa rate, while others are less sensitive (change in yield strength) or essentially insensitive (irradiation creep). These sensitivities will be covered in later sections.
For light water-cooled reactors, the various flux regimes need not necessarily involve large differences in neutron spectra, but only in flux. However, the very large dpa rates characteristic of fast reactors are associated with a significant difference in spectrum. This difference is a direct consequence of the fact that fast reactors were originally designed to breed the fissionable isotope 239Pu from the relatively nonfissile isotope 238U, which comprises 99.3% of natural uranium.
In order to maximize the breeding of 239Pu, it is necessary to minimize the unproductive capture of neutrons by elements other than uranium. One
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Figure 6 Displacement effectiveness values of dpa per 1022ncm-2 (E > 0.1 MeV) across the small core (30cm tall and ~30cm diameter) of the EBR-II fast reactor, showing effects of neutron leakage to soften the spectrum near the core axial boundaries. Near core center (Row 2) the spectrum and displacement effectiveness are dictated primarily by the use of metal fuel, producing a maximum of ~5.2 dpa per 1022 n cm-2 (E> 0.1 MeV). In mid-core Row 4 the radial leakage is just becoming significant. |
strategy used to accomplish this goal is to avoid thermalization of the reactor neutrons, which requires that no low atomic weight materials such as H2O, D2O, Be, or graphite be used as coolants or as moderators. For this purpose, sodium is an excellent coolant with a moderate atomic weight. The use of sodium results
in a neutron spectrum that is nominally single-peaked rather than the typical double-peaked (thermal and fast) neutron spectrum found in light water or heavy water reactors. The single-peaked fast reactor spectrum is significantly less energetic or softer, however, than that found in the fast peak of light water reactors. Depending on the fuel type (metal vs. oxide) the mean energy of fast reactor spectra varies from ^0.8 to ^0.5-0.4MeV while light water-cooled reactors have a fast neutron peak near ~1.2 MeV.
One consequence of attaining successful breeding conditions is that the spectrum-averaged crosssection for fission is reduced by a factor of 300-400 relative to that found in light water spectra. To reach a power density comparable to that of a light water power-producing reactor, the fast reactor utilizes two concurrent strategies: increases in fissile enrichment to levels in the order of 20% or more, and most importantly, an increase in neutron flux by one or two orders of magnitude.
Thus, for a given power density, the fast reactor will subject its structural materials to the punishing effects of neutron bombardment at a rate that is several orders of magnitude greater than that in light water reactors. At the same time, however, the softer ‘fast’ spectrum without thermalized neutrons leads to a significant reduction in transmutation compared to typical light water spectra, at least for stainless steels and nickel-base steels.
Void Swelling
4.04.2.1 Compositional Dependence of Void Swelling
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Nimonic PE16 was first identified as a low-swelling alloy in the early 1970s. Void swelling data derived from density measurements on fuel pin cladding materials from the Dounreay Fast Reactor (DFR) were reported by Bramman et al.1 and were complemented by electron microscope examinations described by Cawthorne etal.8 Swelling in STA PE16 was found to be lower than in heat-treated austenitic steels and comparable to cold-worked steels. Comparison of data for PE16 and FV548 (a Nb-stabilized austenitic steel) irradiated under identical conditions in DFR to a peak neutron fluence of ^6 x 1026nm~2 indicated that the lower swelling of PE16 was due to smaller void concentrations at irradiation temperatures up to ^550 °C and reduced void sizes at higher
temperatures. At around the same time, Hudson et al9 compared the swelling behavior of PE16, type 316 steel, and pure nickel, using 20MeV C2 ion irradiations in the Harwell VEC (variable energy cyclotron). The materials were implanted with 10 appm (atomic parts per million) ofhelium prior to ion bombardment to peak displacement doses >200 dpa (N/2) at 525 °C. Void swelling in 316 steel and nickel exceeded 10% at the highest doses examined, compared to ^0.5% in PE16. Void nucleation appeared to occur earlier in nickel (at ^0.1 dpa) than in PE16 or type 316 (~2 dpa), but the peak void concentration was higher by a factor of about 10 in the austenitic steel than in nickel or PE16.
Hudson et al.9 originally attributed the swelling resistance of PE16 to the presence of the coherent, ordered face-centered cubic, Ni3(Al, Ti) g0 precipitates, which were thought either to trap vacancies and interstitials at their surface, thereby enhancing point — defect recombination, or to inhibit the climb of dislocations, thereby preventing them from acting as preferential sinks for interstitial atoms. In support of the first ofthese two suggested mechanisms, Bullough and Perrin10 argued that the surface of a coherent precipitate would be a more effective trapping site than an incoherent one where the identity ofthe point defects would immediately be lost (and where, as a consequence, void nucleation was likely to occur). The efficiency of point defect trapping would be expected to be greater the higher the total surface area of the g0 precipitates, that is, to be inversely proportional to the precipitate size at constant volume fraction. On the other hand, the second mechanism proposed by Hudson etal. should be most effective at an intermediate particle size where dislocation pinning is strongest. Support for the latter process was provided by Williams and Fisher11 from HVEM (high-voltage electron microscope) irradiations of PE16 at a damage rate of about 10~2dpas_1 at 600 °C, where the swelling rate was higher at small (3 nm) and large (70 nm) g0 particle diameters than at intermediate sizes of about 20 nm.
However, it is now considered that any effect that the g0 precipitates may have on the swelling resistance of Nimonic PE16 is secondary to that of the matrix composition. The generally low-swelling behavior of Ni-based alloys compared to austenitic steels was shown by Johnston etal.12 following bombardment with 5 MeV Ni2+ ions at 625 °C. The damage rate in these experiments was 10~2dpas_1 and the displacement dose was originally estimated as 140 dpa but this was subsequently revised by Bates and Johnston13 to 116 dpa (based on displacement
energy Ed = 40 eV). In addition to precipitation- hardened alloys, including PE16 and Inconel 706, this experiment included nonhardenable high-Ni alloys, such as Inconel 600 and Hastelloy X, a range of commercial steels, and Fe-Cr-Ni ternary alloys containing 15% Cr and 15-35% Ni. The alloys were preimplanted with 15 appm helium prior to ion bombardment, and the irradiation temperature was chosen as being close to the peak swelling temperature for ion — irradiated austenitic steels. The extent of void swelling was determined by electron microscope examinations in low-swelling alloys, but was estimated from step — height measurements (comparing the surfaces of irradiated and nonirradiated regions) in high-swelling materials. As illustrated in Figure 1, the results showed negligible swelling (<0.1%) in PE16, Inconel 706, Hastelloy X, and the Fe-15Cr-35Ni ternary alloy, low swelling (<1%) in other high-Ni alloys, but high swelling (generally >20%) in austenitic steels. In commercial alloys containing ~18% Cr, minimum swelling occurred at Ni contents of about 40-45%. Although void diameters generally appeared to be smaller in the Ni-based alloys than in austenitic steels, the main factor accounting for reduced swelling was a much lower void concentration. In the ternary alloys, reducing the Ni content from 35% to 30% resulted in
Figure 1 Swelling versus nickel content of commercial alloys and ternary Fe-15Cr-Ni alloys bombarded with Ni2+ ions to a damage level of 116 dpa at 625 °C. Reproduced from Johnston, W. G.; Rosolowski, J. H.; Turkalo, A. M.; Lauritzen, T. J. Nucl. Mater.1974, 54, 24-40.
an increase in overall swelling from <0.1% to ~12%, although it was noted that the 35% Ni alloy showed a localized swelling of ^5% in a region close to a grain boundary. Additional experiments reported by Johnston et a/.12 indicated that the peak swelling temperature for PE 16 irradiated with 5 MeV Ni2 ions was 675 °C, but even then, swelling at 116 dpa remained below 0.2%.
Swelling data for a wider range of pure Fe-Cr-Ni austenitic alloys, with Cr contents up to 30% and Ni up to 100%, following Ni ion bombardment to 116 dpa at 675 °C, were reported by Bates and Johnston.1 These results showed a strong dependence on both Cr and Ni, with the swelling increasing with increasing levels of Cr but being minimized at Ni contents of about 45-60%. Examination of the dose dependence of swelling in ternary alloys containing 15% Cr and 20-45% Ni showed that the incubation dose required for the onset of swelling increased with increasing Ni content. Furthermore, although high-swelling rates of the order of 1% per dpa were attained in 20-35% Ni alloys, the swelling rate of the 45% Ni alloy remained low even at doses above 250 dpa.
Following their earlier C2+ ion irradiation experiments, Hudson and coworkers moved to the use of
46.5 MeV Ni6+ ions to investigate void swelling behavior. This was considered preferable because the recoil spectra of high-energy Ni ions provided a better simulation of fast neutron damage, and because carbon implantation encouraged the formation of carbides which acted as void nucleation sites. A summary of some of the Ni ion irradiation work carried out by the Harwell group was given by Makin et a/.14 No significant differences in the swelling behavior of Nimonic PE16 were evident between ST or aged conditions. Peak swelling in Ni6+ ion- irradiated PE16 (preimplanted with 10appm He) occurred at 625 °C, where a swelling of ~1.5% was recorded at 120dpa(N/2). Void concentrations in PE16 were reported to be lower by a factor of about 5 than in similarly irradiated type 316 and 321 austenitic steels.
A drawback of charged particle irradiation experiments for evaluating void swelling is that the evolution of other microstructural features may differ significantly from that during neutron irradiation (see also Chapter 1.07, Radiation Damage Using Ion Beams). In the case of Nimonic PE16, for example, the precipitation and/or redistribution of the g0 phase during long-term neutron exposure might be expected to influence swelling behavior. In order to simulate swelling in a more appropriate microstructure, Bajaj eta/.15 examined the effect of4MeV Ni2 ion irradiation on PE16, which had been preconditioned by exposure to neutrons in Experimental Breeder Reactor-II (EBR-II). Reactor-conditioned samples had been exposed to neutron fluences in the range of 3-6 x 1026nm~2 (E > 0.1 MeV) at temperatures from 454 to 593 °C. Swelling rates during Ni ion irradiations at 675 °C were higher by a factor ofabout five in reactor-conditioned material than in a nonconditioned sample. The increased swelling rate was attributed to changes in the matrix composition resulting from an increased volume fraction of g0 in the reactor-conditioned material.
Early attempts to account for the effects of matrix composition on void swelling focused on the stability of the austenite phase. Harries16 suggested that the swelling behavior of austenitic steels and nickel — based alloys could be rationalized in terms of their Ni and Cr equivalent contents (i. e., the relative austenite and ferrite stabilizing effects of their constituent elements), with the composition of high — swelling alloys then falling into the (g + s) phase field in the Fe-Cr-Ni ternary phase diagram. Harries postulated that the partitioning of solute elements into the sigma phase would have a detrimental effect on the swelling resistance of austenite. Watkin17 took a similar approach, but found that an improved correlation could be obtained using the concept of electron vacancy numbers rather than Ni and Cr equivalents. The average electron vacancy number, Nv of the matrix is calculated from the atomic fractions of its constituents, with allowance being made for the precipitation of carbides and g0 (or g00, etc.), and has been widely used to predict the susceptibility of nickel-based alloys to the formation of intermetallic phases.18Nv was calculated from:
Nv = 0.66Ni + 1.70Co + 2.66Fe + 3.66Mn + 4.66(Cr + Mo)
Watkin found that void swelling in a range of alloys with Ni contents up to ^43%, which were irradiated in DFR to a peak dose of 30 dpa at 600 °C, remained low for Nv below about 2.5 (corresponding to low susceptibility to s phase formation), but increased approximately linearly at higher Nv. However, as was clearly argued by Bates and Johnston,13 correlations based on sigma-forming tendency could not account for the minimum in swelling observed at about 45% Ni, since higher Ni contents should continue to be beneficial.
A better understanding of the swelling behavior of Fe — and Ni-based alloys resulted from a series of fast neutron irradiation experiments which were carried out in EBR-II in the early 1980s. Irradiation temperatures in these experiments ranged from about 400 to 650 °C. Initial data for a range of commercial alloys, including ferritic and austenitic steels, as well as nickel-based alloys, were reported by Bates and Powell19 and Powell eta/.,20 with higher dose data (up to a peak fluence (E > 0.1 MeV) of ^25 x 1026nm~2, corresponding to 125 dpa) being reported by Gelles21 and Garner and Gelles.22 Swelling data for Fe-Cr-Ni ternary alloys, irradiated in EBR-II to a peak fluence of 22 x 1026nm~2 (^110dpa), were presented by Garner and Brager23 The extent of void swelling in these experiments was determined by density change measurements. In general, alloys with nickel contents in the range of 40-50% exhibited the lowest swelling. Swelling in commercial nickel-based alloys was generally lower in ST than in aged conditions, this being attributed to the beneficial (though temporary) effect of minor elements remaining in solution and being able to interact with point defects19; subsequent precipitation during irradiation would be expected to reduce this benefit and the resulting densification, though small, would also effectively reduce the measured swelling. Swelling data for a number of ST alloys, which were irradiated in the AA-1 rig in EBR-II, are shown in Figure 2; data are shown for two withdrawals, at peak fluences of 14.7 x 1026nm~2 and 25.3 x 1026nm~2, with measurements for Inconel 600 and Inconel 625 reported at both fluence levels, data for Nimonic PE16 and Inconel 706 at the lower level, and data for Incoloy 800 and Hastelloy X at the higher level. The nickel contents of the alloys range from about 34% in Inco- loy 800 to 75% in Inconel 600. Swelling remained relatively low in the three Inconel alloys and in PE16. However, both Incoloy 800 and Hastelloy X exhibited high swelling at some temperatures, with swelling in the latter reaching ^80% at 593 °C. The reason for such high swelling in neutron-irradiated Hastelloy X (nickel content ^48%) is unclear, but it was noted that densification up to 3% occurred at the lower irradiation temperatures — indicating microstructural instability and possibly signaling changes in the composition of the matrix which may have affected the swelling behavior. (Note that Hastelloy X was identified as a low-swelling alloy in the Ni2 ion irradiation experiments described by Johnston eta/.12)
Some data for different heat-treated conditions of PE16 at the higher fluence level were reported by
Garner and Gelles,22 and are compared for irradiations at 538 °C (more or less corresponding to the peak swelling temperature for PE16 in the AA-1 experiment) with lower fluence data from Bates and Powell19 in Figure 3. The heat-treated conditions indicated in Figure 3 are ST (ST 4h at 1080 °C), A1 (ST and aged 16 h at 705 °C), A2 (ST and aged 1 h at 890 °C plus 8h at 750 °C), and OA (ST and aged 24 h at 840 °C). Note that the silicon content of the PE16 used in these experiments was much lower at 0.01% than the level of ^0.2% typically found in UK heats of the alloy. Overall, the data appear to show little effect of initial heat treatment on the swelling of PE16, except that the OA condition exhibited the most swelling (5.2%) at the higher fluence.
Although it is clear that the swelling behavior of austenitic alloys is largely dependent on nickel content, there is ample evidence to show that minor solute additions can have significant effects. Much of the work on minor solutes has focused on steels similar to type 316, but some data are available for higher nickel alloys. For example, Mazey and Hanks24 used
46.5 MeV Ni6+ ion irradiations to examine the effects of Si, Ti, and Al additions on the swelling response of model alloys with base compositions approximating that of the matrix phase in PE16. Solute additions of ^0.25% Si or 1.2% Ti reduced swelling, but the addition of ~1.2% Al (in the absence of Si or Ti) markedly increased it. The beneficial effect of Si was believed to arise from its high diffusivity in solution (this is discussed further in Section 4.04.2.2), whereas that of Ti appeared to be related to the formation of Z phase (hexagonal-structured Ni3Ti). The addition of Al resulted in an increase in the concentration of voids, the surfaces of which were coated in a thin layer of the g0 phase (Ni3Al). A beneficial effect of Si on the swelling response of modified Incoloy DS alloys under Ni6+ ion irradiation was also reported by Mazey et a/.25 However, it should be noted that high Si contents can give rise to the formation of radiation — induced phases which are enriched with Ni and Si, such as the Ni3Si form of g0 and the silicide G-phase (M6Ni16Si7, where M is usually Ti, Nb, or Mn). G-phase particles are generally found in association with large voids and their formation may therefore give rise to an increase in the swelling rate.26,27
Swelling data derived from density measurements for neutron irradiated, modified Incoloy DS alloys, with Si contents ranging from 0.19 to 2.05% (compared to a specified level of 1.9—2.6% in the commercial alloy), are compared with data for a ‘PE16 matrix
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alloy’ and Nimonic PE16 in Figure 4. The materials were all in ST condition apart from PE16 which was in an STA condition (aged 4h at 750 °C). The alloys were irradiated in the UK-1 rig in EBR-II to fluences in the range of 9-16 x 1026nm~2 (E> 0.1 MeV) at temperatures of ^390-640 °C. These data are previously unpublished except those for STA PE16 (heat DAA 766) which were reported by Boothby.28 Swelling in the modified Incoloy DS alloys generally decreased with increasing Si content. The 0.19% Si
alloy exhibited high swelling at all temperatures with indications of swelling peaks at about 440 and 640 °C. Increased Si levels tended to suppress the high temperature swelling peak and reduce the magnitude of swelling at lower temperatures. The PE16 matrix alloy containing 0.24% Si exhibited a high temperature swelling peak but moderate swelling below ^550 °C, suggesting a beneficial effect of Mo (this being the main compositional difference between the PE16 matrix alloy and the modified Incoloy DS alloys)
Figure 4 Void swelling data derived from density measurements for Nimonic PE16, a PE16 matrix alloy, and modified Incoloy DS alloys, irradiated in the UK-1 rig in Experimental Breeder Reactor-II. Unpublished data from Boothby, R. M.; Cattle, G. C. Void Swelling in EBR-2 Irradiated Nimonic PE16 and Incoloy DS; FPSG/P(90)10, with permission from AEA Technology Plc.
at lower temperatures. However, swelling in the PE16 matrix alloy remained significantly higher at all temperatures than in STA Nimonic PE16 (containing 0.15% Si), indicating a significant benefit of the g0 forming elements Al and Ti.
Irradiation-Induced Swelling of Ta and Ta-Base Alloys
Swelling data for Ta and its alloys are limited to a few studies.19 Void formation in pure Ta was experimentally observed through TEM examination of material irradiated to 2.5 x 1022ncm-2 (E> 0.1 MeV) at temperatures between 673 and 1273 K.46 An empirical estimation of the bulk swelling taken from microstructural void size density data of that study is shown in Figure 6. Void concentrations in the material were highest at the peak swelling temperature and decreased with higher irradiation temperature with an associated increase in cavity size. Ordering of the voids at the peak
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Figure 6 Swelling data for pure Ta measured through microstructural void density measurements by Wiffen46 and from immersion density measurements by Bates and Pitner.47 |
swelling condition was reported to occur along the {110} planes in the bcc structure. A subject of considerable theoretical debate, the mechanisms of void ordering that have appeared in bcc and fcc metals have been examined,4 — 0 since the first reported occurrence in irradiated Mo.51 Disordered void structures in the microstructure of the higher temperature irradiated Ta appear as the size of the voids increase, though some rafting, or grouping, was reported.46
The swelling data of Wiffen46 derived from microstructural analysis correlate well with the immersion density data of Bates and Pitner47 (Figure 6), from which an empirical equation for percent swelling as a function of temperature, T (K), and fluence, F (in units of 1022ncm-2, E > 0.1 MeV), was developed, which is as follows:
D — = (F)°’4{1.69 exp[—(0.018T — 16.347)2/a]}
where a _ 14.87 + 44.57 exp[0.09(T — 1338.71)]
“_ 1 + exp[0.09(T — 1338.71)] [1]
The broader width of the swelling peak as a function of irradiation temperature for the calculation represented by eqn [1] compared to the microstructural data of Wiffen46 is believed to be associated with errors in the accurate irradiation temperature of these early measurements. Experimental evidence of decreased swelling at higher fluences was reported by Murgatroyd et al. and attributed to the transmutation of Ta to W, resulting in a shift in the lattice constant. Similar effects have been more closely examined in Mo and TZM alloys, and attributed to impurity segregation at void surfaces leading to shrinkage of the voids.53
Swelling measurements in Ta-10W and T-111 alloys are limited specifically to work by Wiffen, from which a later summary was given.19 For irradiations at 723 and 873 K to a fluence of 1.9 x 1022ncm-2 (E > 0.1 MeV), no swelling in T-111 was observed, though a possible densification of up to 0.36% may have occurred as evidenced in length measurements. In companion irradiations to that of pure Ta already discussed, involving irradiations to 4.4 x 1022ncm — (E > 0.1 MeV) at temperatures between 698 and 1323 K,46 samples of Ta-10W were included with postirradiation examination involving TEM analysis. The microstructure of the irradiated Ta-10W contained fewer voids than the companion Ta samples, with a lower swelling assumed in the Ta-10W alloy but with values not accurately quantifiable.19
LBE Compatibility
Molten LBE has a high solubility of nickel, iron, and chromium, which are the most important alloy elements
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in austenitic stainless steels. Thus, nickel super alloys and austenitic stainless steels cannot be used as the structural materials for LBE-cooled systems, especially at temperatures >500 ° C. Ferritic steels have been considered more appropriate for LBE application.
Exposure of 9Cr-ODS steels to an LBE environment at 530 °C was carried out in the DELTA Loop of the Los Alamos National Laboratory. The molten
alloy flow velocity in the loop is 1.2 m s-1, and oxygen sensors were used to measure and maintain an oxygen concentration of about 1 x 10~6wt%. Samples were exposed for 200, 400, and 600 h, in order to study the early stages of oxide formation and growth. A crosssectional view and the distribution of elements are shown in Figure 32.54 In a short time, the 9Cr-ODS steel formed a protective duplex oxide layer consisting
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Time to rupture (h)
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of an outer magnetite (Fe3O4) layer and an inner Fe-Cr spinel ((Fe, Cr)3O4) layer, which is sometimes accompanied by an O-enriched and Fe-depleted diffusion zone at the oxide-bulk interface. Over time, the outer magnetite layer is removed and the underlying spinel layer serves to mitigate more catastrophic corrosion degradation such as dissolution and liquid metal attack along the grain boundaries. Very thin oxides are not particularly protective in regard to loss of metal, as manifested by the thick diffusion zones associated with them. Furukawa pointed out that at temperatures above 600 °C, the thickness of the oxide layer diminishes with increasing temperature. This behavior can be ascribed to a change in the stable form of iron oxide from magnetite to wustite at 570 °C. Beyond this temperature, dissolution attack was observed on some portions of 9Cr-ODS steel, and the oxide layer’s adhesion to the material began to weaken.55
It has been reported that the addition of aluminum to steel effectively prevents LBE corrosion. Figure 33 shows the appearance of ODS steel specimens after a corrosion test in LBE for 1 x 104h at 650 ° C.43 The 18wt% Cr-ODS steel without the addition of Al dissolved markedly into LBE, while those ODS specimens containing 4wt% Al almost completely maintained their shape even in Al-added 14Cr — and 16Cr-ODS steels, indicating a very high resistance to LBE corrosion. It is noteworthy that this corrosion resistance was independent of Cr concentration from 13 to 19 wt% in Al-added ODS steels. From the distribution of elements across the cladding surface, we deduce that LBE corrosion can be prevented by the formation of an Al enriched film.56 It was demonstrated that Al-added 16Cr-ODS steel (16Cr- 2W-4Al-0.1Ti-0.35Y2O3) has superior corrosion resistance at 650 °C for 5000 h.
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18Cr 19Cr-4Al Figure 33 The appearance of Al added high Cr-ODS steel specimens after corrosion test in LBE for 1 x 104h at 923K (DO: 1 x 10~6wt%). Reproduced from Kimura, A.; Kasada, R.; Iwata, N.; etal. In Proceedings of ICAPP ’09, Tokyo, Japan, May 10-14, 2009; Paper 9220. |
Deficiencies in Current Creep Models at High Neutron Doses
The poor performance of the Kelly and Burchell model (eqn [25]) at predicting the high temperature (900 °C) and high dose 6 MPa tensile creep data suggests that the model requires further revision.50,7i H-451 graphite irradiated at 900 °C goes through dimensional change turn-around in the dose range 1.3—1.5 x i022ncm~2 [E>50] (—8.8-10.2 dpa). This behavior is understood to be associated with the
generation of new porosity due to the increasing mismatch of crystal strains. The Kelly-Burchell model accounts for this new porosity only to the extent to which it affects the CTE of the graphite, through changes in the aligned porosity.
Gray70 observed that at 550 °C the creep rate was approximately linear. However, at 800 °C he reported a marked nonlinearity in the creep rate and the changes in CTE were significant. Indeed, for the two high density graphites (H-327 and AXF-8Q) Gray reports that the 900 °C creep strain rate reverses. Gray postulated a creep strain limit to explain this behavior, such that a back stress would develop and cause the creep rate to reduce. Other workers have shown that a back stress does not develop.62 However, Gray further argued that a creep strain limit is improbable as this cannot explain the observed reversal of creep strain rate. Note that a reversal of the creep rate is clearly seen in the 900 °C tensile creep strain data reported here for H-451 (Figure 23). Also, a creep strain limit would require that tensile stress would modify the onset of pore generation behavior in the same way as compressive stress, because the direction of the external stress should be immaterial.70 More recent data52 and the behavior reported by Burchell71 show that this is not the case. Gray70 suggests that a more plausible explanation of his creep data is the onset of rapid expansion accelerated by creep strain; that is, net pore generation begins earlier under the influence of a tensile applied stress. Indeed, it has been observed52 that compressive creep appears to delay the turnaround behavior and tensile creep accelerates the turnaround behavior (Figure 18).
In discussing possible explanations for his creep strain and CTE observations, Gray70 noted that changes in the graphite pore structure that manifested themselves in changes in CTE did not appear to influence the creep strain at higher doses. The classical explanation of the changes in CTE invokes the closure of aligned porosity in the graphite crystallites. Further crystallite strain can be accommodated only by fracture. A result of this fracture is net generation of porosity resulting in a bulk expansion of the graphite. A requirement of this model is that the CTE should increase monotonically from the start of irradiation. A more marked increase in CTE would be seen when the graphite enters the expansion phase (i. e., all accommodating porosity filled). The observed CTE behavior, reported previously50 and in Gray’s70 work, does not display this second increase in CTE; thus, the depletion of (aligned) accommodation porosity is not responsible for the early beginning of expansion behavior.
The observation by Gray70 and Kennedy63 that creep occurs at near constant volume (up to moderate fluence) indicates that creep is not accompanied by a net reduction ofporosity compared to unstressed graphite, but this does not preclude that stress may decrease pore dimensions in the direction of the applied stress and increase them in the other, that is, a reorientation of the pore structure. Pore reorientation could effectively occur as the result of a mechanism of pore generation where an increasing fraction of the new pores are not well-aligned with the crystallites basal planes (and thus they would not manifest themselves in the CTE behavior) or accompanied with the closure of pores aligned with the basal planes.
Kelly and Foreman53 report that their proposed creep mechanism would be expected to break down at high doses and temperatures, and thus deviations from the linear creep law (eqn [12]) are expected. They suggest that this is due to (1) incompatibility of crystal strains causing additional internal stress and an increasing crystal creep rate, (2) destruction of interstitial pins by diffusion of vacancies (thermal annealing of vacancies in addition to irradiation annealing), and (3) pore generation due to incompatibility of crystal strains.
It is likely that pore generation can manifest itself in two ways: (1) changes in CTE with creep strain — thus, pores aligned parallel to the crystallite basal planes are affected by creep strain — and (2) at high doses, pore generation or perhaps pore reorientation, under the influence of applied and internal stress that must be accounted for in the prediction of high neutron dose creep behavior.
Brocklehurst and Brown62 report on the annealing behavior of specimens that had been subjected to irradiation under constant stress and sustained up to 1% creep strain. They observed that the increase in creep strain with dose was identical in compression and tension up to 1% creep strain, and that the CTE was significantly affected in opposite directions by compressive and tensile creep strains. Irradiation annealing of the crept specimens caused only a small recovery in the creep strain, and therefore provided no evidence for a back stress in the creep process, which has implications for the in-crystal creep mechanism. Thermal annealing also produced a small recovery of the creep strain at temperatures below 1600 °C, presumably because of the thermal removal of the irradiation-induced defects responsible for dislocation pinning. Higher temperature annealing produced a further substantial recovery of creep strain. Most significantly, Brocklehurst and Brown62 reported the complete annealing of the creep induced changes in CTE, in contrast to the total creep strain, where a large fraction of the total creep strain is irrecoverable and has no effect on the thermal expansion coefficient. Brocklehurst and Brown62 discuss two interpretations of their results, but report that neither is satisfactory. One interpretation requires a distinction between changes in porosity that affect the CTE and changes in porosity affecting the elastic deformation under external loads, that is, two distinct modes ofpore structure changes due to creep in broad agreement with the mechanism discussed earlier.
The modified Simmons model29,30,67 for dimensional changes (eqn [18]) and that for dimensional changes of a crept specimen (eqn [20]) both have pore generation terms which are currently neglected. It now appears necessary to modify the current Kelly-Burchell creep model (eqn [25]) to account for this effect of creep strain on this phenomena; that is, we need to evaluate and take account of the terms Fx and F’x as well as include the term (F’x — Fx) in eqn [25]. Such a term should account for pore generation and/or reorientation caused by fracture when incompatibilities in crystallite strains become exces — sive.71 Clearly, further work is needed in the area of irradiation-induced creep of graphite.
Predicting the Thermal Conductivity of Irradiated Graphite for Reactor Core Assessments
A method derived by Kelly71 is often used in the prediction ofirradiated graphite thermal conductivity
in the assessment of UK reactor cores. The thermal resistivity, 1/K(T), induced because of irradiation as a function of temperature, T can be given by the difference between the thermal resistivity due to fast neutron damage, and the unirradiated thermal resistivity:
1 _ 1 1
K(T)_ Kjt) ~ K0(T)
However,
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