This content was uploaded by our users and we assume good faith they have the permission to share this book. If you own the copyright to this book and it is wrongfully on our website, we offer a simple DMCA procedure to remove your content from our site. Start by pressing the button below!
-> CO | (excited state) COf + C (graphite) -> 2CO occurs at temperatures far below those where it is significant as a result of thermal activation. The energy deposition rate in, for instance, an advanced gas-cooled reactor core is —0.25 W/g (CO2) and this generates enough oxidising species to completely remove the graphite in a few years. The reaction is controlled in the reactors in the United KingSom and elsewhere by adjusting the levels of carbon monoxide in the coolant and also in the advanced gascooled reactors by adding methane (CH4) and hydrogen to the coolant at levels of less than 0.1% (compared to carbon monoxide levels of less than 3%). The reactions which take place in the gas phase and the mechanisms of protection of the graphite are well understood (Best etal., 1985; Kelly, 1985). It is found experimentally that the weight loss of a graphite volume
5.3 Solid Moderators
with uniform gas composition in the open porosity x = (dx/dd)0 df [exp (d/df) - 1] for d^lAdf (5-34)
x = (dx/dd)0[3d-130d']
for
d^l.ld'
where d is the dose to the gas in Wh • g~ *, (dx/dd)0 is the initial rate of weight loss, and d' is a constant equal to 18.58 x 103 (41/P) (T/673) Wh • g" 1 , P is the gas pressure. The values of (dx/dd)0 for AGR moderator graphite are shown in Fig. 5-14. The oxidation modifies most of the technically important properties such as Young's modulus, thermal conductivity, strength, diffusivity and permeability but does not effect the thermal expansion coefficient or Poisson's ratio. It is necessary to find appropriate rules to combine the property changes due to oxidation with those due to irradiation damage. The property changes due to oxidation have been examined by oxidising samples under carefully controlled conditions of coolant composition followed by
co Level
0.25 - 035 % 0.40-0.60 % 0.80 - 1.20 % xxx 1.80 - 2.20 % OOO
IE
0
500 Methane concentration ,vpm
Figure 5-14. Initial radiolytic oxidation rates for AGR moderator graphite based on gilsonite coke, vpm = Volume parts per million.
391
annealing out of the inevitable neutron irradiation effects which accompany the energy deposition. (Experiments can in principle be done in pure y-irradiation facilities, but these are always too limited in energy deposition rate in the gas compared to experiments in a nuclear reactor core which inevitably give a component of neutron damage.) Annealing temperature > 2000 °C is required to remove the effects of crystal lattice defects but above a critical dose at each irradiation temperature it is not possible to remove the crystal shape changes and this may cause complications in interpretation particularly for the mechanical properties. The UKAEA have studied a wide range of graphites and found that to a first approximation the changes in properties can be described by the same laws. These are: Young's modulus Thermal conductivity Strength Diffusivity
E — Eo exp (— 3.6 x) K= Ko exp (— 2.7 x) a = a0 exp (— 4 x) X = Xo + ( i - Xo) x2 (5-35)
where zero subscripts denote initial values, and x is the weight loss as in Eq. (5-34). The viscous flow coefficient and slip flow coefficient for transport of gases through graphite are not simple functions of weight loss, except when there are no inhibitors in the gas phase (Kelly et al., 1975). The theory of the changes in the diffusivity X, the slip flow coefficient K and the viscous flow coefficient B have been given (Johnson, 1981). The dependence of diffusivity on weight loss alone is found to be a result of its being dominated by gas flow through fine porosity where the inhibitors have least effect on the oxidation. The combination of these effects with those due to neutron damage is made using multiplicative rules. The changes in
392
5 Nuclear Reactor Moderator Materials
Young's modulus are given by: (5-36) 0
\£0/
£
where (E/E0)ox is given by Eq. (5-35) and (E/Eo){ is the irradiation damage effect alone. Similar rules apply to thermal conductivity and strength, the first two having been verified experimentally (Birch et al, 1990). The effect of oxidation on dimensional changes is more complicated, as expected, because there is no effect on the thermal expansion coefficient and consequently on the dimensional changes up to the beginning of shrinkage reversal where the relationship will break down. Irradiation experiments on isotropic graphite samples which were radiolytically pre-oxidised have shown that the shrinkage reversal is delayed in dose, although once it begins it has the same dose dependence (Kelly etal., 1976). However the effect on dimensional changes when the oxidation and irradiation damage occur simultaneously are very difficult to predict. Two recent studies have attempted to understand these effects, the first using substitutional doping with B 1 1 to enhance the rate of irradiation damage retained for a given displacement rate by a factor of three to give comparability to the ratio of energy deposition rate to atomic displacement rate found in the core of the advanced gascooled, graphite moderated reactors, while the second alternated specimens between two rigs, one with oxidising atmosphere and the other with inert gas to give the ratio correctly. The results of these experiments (Birch et al., 1990) show that multiplicative rules can be used to predict the simultaneous effects of radiation damage and radiolytic oxidation on Young's modulus and thermal conductivity, while oxidation does not modify the changes in thermal expansion coefficient. Because of
this effect the dimensional changes are also unaffected over the dose range where they are controlled by the thermal expansion coefficient. Unfortunately the experiments had to be terminated before the alternating set of specimens had achieved the dose required for shrinkage reversal. In the boron doped samples, the same volume expansion was found for samples in oxidising and inert atmospheres, but the oxidising samples in this condition have a higher thermal expansion coefficient than the non-oxidised samples and thus the turn around component of the dimensional changes has been reduced, as is found in pre-oxidised samples. The thermal oxidation of moderator graphite is important in a number of contexts. In the early, air-cooled reactors, the rate of reaction of the initially very pure graphite was very low at the normal graphite operating temperatures of less than 200 °C and the heat produced by the exothermic reaction C (graphite) + O 2 ^ 2 C O easily removed by the coolant flow. However the reactivity of the graphite is increased by irradiation damage and also by the deposition of catalytic material from the coolant to the graphite. The graphite in the British Production Pile at Windscale caught fire in 1957 during an operation to release Wigner energy in the graphite which raised the graphite temperatures to 400 °C. Investigation of the possible causes of this accident (see Arnold, 1992) revealed that one of them was associated with unstable heating of the graphite by the oxidation reaction. The heating rate of the graphite at temperatures where the oxidation takes place uniformly throughout the porosity can be written: H = H0 exp(-E/k Tg) (W • cm" 3 )
(5-37)
5.3 Solid Moderators
where E is the activation energy, k is Boltzmann's constant and Tg is the graphite temperature. If the heat transfer area/unit length of the fuel channel is A, and the associated graphite volume is V, then the heat loss to the coolant is
= A(Tg-Tc)h
(5-38)
where Tc is the coolant (air) temperature and h the heat transfer coefficient. Equation (5-37) in general has two common solutions with Eq. (5-38), in the lower temperature case an increase in Tg increases the heat transfer more than the heat output, but at the higher temperature the opposite is true and an increase in Tg produces a larger increase in heating rather than cooling, an unstable situation. In a similar way it is necessary to consider the effects of accidental air ingress into graphite moderated reactors with inert gas cooling. The chemical reactivity of the graphite, which includes contributions from very reactive carbonaceous deposits formed in the graphite pores, catalytic material deposited from the coolant and increases due to radiation damage, is monitored in the fixed moderator reactors by taking samples using specially designed machines (Wickham, 1989). The oxidation of the graphite at low temperatures may be dominated by variation of reactivity with depth in the graphite, but as temperature and reactivity increase the reaction becomes more and more limited to a surface layer. Similar considerations apply to reaction with small concentrations of water in the helium coolant of high temperature gas cooled reactors, the supply of oxidant to the interior of graphite components being limited by the diffusivity X, itself dependent upon the weight loss. Careful studies of the changes in properties of graphite subjected to uniform thermal oxidation (Rounth-
393
waite et al., 1967; Board and Squires, 1967) has demonstrated that the reduction in properties is greater for Young's modulus and strength than in the case of radiolytic oxidation for the same weight loss and also that the changes depend on the oxidising gas. The differences can be explained qualitatively in that radiolytic oxidation tends to expand existing porosity by eating away the pore walls, whereas in the thermal oxidation case the catalyst particles drill fine holes in the structure from the open pore surfaces and rapidly open closed porosity. At very high reactivities or temperatures the oxidation is limited by diffusion of the oxidant through the boundary layers at the graphite surface. The very complex behaviour of graphite under irradiation which requires very large experimental programmes for its evaluation, together with the effects of oxidation have led to: (a) The provision of means to monitor the irradiation effects actually occurring in the reactor cores and regular comparisons with the data base and design predictions for much of the reactor life - an expensive process. (b) The requirement of regulatory authorities that graphite behaviour must be experimentally re-evaluated if changes are made in the manufacturing process or manufacturing campaigns on the same basic material are widely separated in time. Recent agreements to share data between the U.S.A., Germany, Japan and the U.K., may help to alleviate requirement (b). Graphite has become important in the fusion reactor programme as a liner material providing first wall protection in TOKAMAK reactor designs. The data base established for fission reactors can be used provided that appropriate corrections are made for the greater proportion of high
394
5 Nuclear Reactor Moderator Materials
energy neutrons contributing to the damage process (Birch and Brocklehurst, 1987). However the greatest difficulty lies in the erosion/corrosion of the carbon liner by the plasma and the redeposition of carbon from the plasma, currently the subject of extensive study. 5.3.5 Zirconium Hydride
Metal hydrides are potentially efficientmoderators because the density of hydrogen atoms in many such materials far exceeds that found in liquid hydrogen. They are thus particularly suitable to thermal reactors which need to minimise core weight and volume. Zirconium hydride has for instance, been used in SNAP (Systems for Nuclear Auxiliary Power) reactors with liquid metal cooling. The much larger KNK reactor developed in Germany uses zirconium hydride as the moderator with sodium cooling. The future of such reactors might be as remotely based power plants (Polar regions or planets), moveable power plants or space systems. The hydrides are generally prepared by direct reaction of hydrogen with the metal at an elevated temperature with 800 °C being appropriate for zirconium, followed by cooling in a hydrogen atmosphere. A hydrogen pressure of one atmosphere is generally adequate but higher pressures are sometimes used. The preparation of the shapes necessary to form a suitable moderator structure may be achieved by two different methods, the direct hydriding of metal shapes or by reconstituting a monolithic body from hydride powder, although both have their disadvantages. The density of zirconium hydride is 14% less than that of the metal and hydride ductility is low even at the temperatures necessary for hydriding, thus a difficulty
arises and cracking can be expected in hydriding samples more than ~ 1/2 cm thick. It is possible to prepare larger moderator elements by hydriding very slowly in order to keep the hydrogen gradients to a minimum, and in this way sections of about 2.5 cm thickness can be hydrided (Vetrano, 1970). Improvements to this technique can be made by using material (metal) with fine grain size and random orientation. This can be induced by 0.3-0.5 wt.% of zirconium carbide which prevents grain growth during hydriding (Van Houten and Baxter, 1962). The alternative technique is to hydride zirconium powder directly and to form the required monolithic bodies by powder metallurgy techniques. Components prepared by these methods generally have poorer physical and mechanical properties than those of directly hydrided bodies. The principal difficulty resides in the fact that the temperatures required to produce sintering on a sufficiently short timescale are such as to produce dissociation of the hydride. The conditions are such that it is difficult to achieve densities of the hydride body greater than 80% of theoretical density. The properties of zirconium hydride are given in Table 5-3 and Table 5-4. The number of hydrogen atoms/cm3 of hydride can be calculated from: (H/M) Q 60.23 (5-39) ML where (H/M) is the hydrogen to metal atom ratio, Q is the hydride density (g-cm~ 3 ), M w is the molecular weight of hydride. ATH varies with temperature, and it is observed that a precipitous loss of hydrogen occurs at a characteristic temperature for each hydride; clearly this temperature must be avoided for operational use. The phase diagram is shown in Fig. 5-15 (Moore and Young, 1968). Nn =
395
5.3 Solid Moderators
Table 5-3. Properties of zirconium hydride. Property Structure Density (30 °C)
Value
Comments
f.c.t.a 5.610g-cm~ 3 7xl022at.-cm"3
ZrH,
Specific heat (298 K) 40.71-mor1 K ' 1 3 0 . 5 1 - m o r 1 K" 1
i5
ZrH 1>25 Thermal expansion coefficient Z r H 1 5 4 (20-850°C) ZrH 1 8 3 (20-550°C) Thermal conductivity
14.2 x l O ^ K " 1 9.15xlO" 6 K" 1 20W-m-1K'1
Insensitive to temperature or hydrogen content
Face-centred tetragonal.
Table 5-4. Mechanical properties of zirconium hydride. Compound
ZrH 0 7 2
Temperature
Elongation
K
Young's modulus MPa
Tensile strength MPa
2.7 x 1022
300 700 1000 1200
0.85 x 10 5 0.77 0.40 0.30
157 155 19.3 11
0 1.5 49 72
0 1 93 100
4.0 xlO 2 2
300 700 1000 1200
0.81 x l O 5 0.61 0.44 0.17
91 131 19.3 22.7
0 1 53 22
0 1 100 61
The behaviour of zirconium hydride under irradiation has been studied using test elements from the inner and outer moderators of the KNK reactor (Diinner et al., 1978) and in irradiation experiments in the US GETR reactor at Vallecitos, the HFR reactor at Petten in Holland and the BR 2 reactor at Mol in Belgium (Paetz and Lucke, 1972). The experiments in the KNK study used a so-called materials test element in the first reactor core and samples from the moderator attached to a fuel element. The samples of ZrH 2 in the element contained either
Reduction in area
61.54 at.% H (x = 1.6) or 63 at.% H (x = 1.7), and the fuel element moderator was also initially at x = 1.7. The samples from the test element showed no change in dimensions and no reduction in mechanical integrity. Direct exposure to the sodium coolant for 2 Yi years showed that the compatibility under these conditions (420-480 °C) was very good, only a surface layer of ~100]im thickness showed some evidence for debonding of grains. Sectioning of the pellets and examination of the phases present indicated that at 480 °C some loss of hydrogen (to x = 1.63) had oc-
5 Nuclear Reactor Moderator Materials
396
at,%H 58 59 60 61 62 63
1000
64
65
66
/-
900 800 700
y 600 lMeV) varied between 0.48 xl0 2 1 n-cm~ 2 and 1.15 xlO 2 1 n-cm~ 2 in a time of 1.08 x 107 s. The second experiment (II) contained samples of 63.1, 62.0 and 61.0 at. % with a prior oxide coating at temperatures of 520 °C and 570 °C. The neutron dose was 1.3xl0 2 1 n-cm~ 2 (£ n >lMeV) in 10 7 s. In experiment III (in the HFR, Petten) the aluminium capsules containing samples with a hydrogen content of 63.1% were irradiated to 0.76xl0 2 1 n-cm" 2 (E n >lMeV) in 1.49xl0 7 s. The irradiation temperature was estimated to be between 100 °C and 350 °C. The further experiment (IV) exposed samples containing 62.4 at.% of hydrogen and one with 63.5 at.% hydrogen (oxide coated). The last sample was exposed at 580 °C and the others in two groups at 420 °C and 500 °C. The doses were 0.56 and 0.66 x 10 2 1 n-cm" 2 (£ n > 1 MeV) in times of 1.93 x 106 s and 6.45 x 106 s respectively. Analysis of the hydrogen content of the specimens showed that in experiment I the hydrogen content had increased at the cold end and decreased at the hot end, as was also observed in a parallel control experiment. Some hydriding to 63.4 at.% was observed in the 450 °C samples. The low temperature samples showed no change. Examination of the dimensional changes of the samples showed that a volume change (increase) only occurred if the irradiation took place in the s-phase and two-phase region. (5-zirconium hydride shows no di-
5.3 Solid Moderators
mensional changes after irradiation. Estimations of the amount of e-phase formed during irradiation were made. The volume increases for a given dose were found to be linearly related to the amount of e-phase, the largest value being 0.47% in pure ephase. A comparison of the density of unirradiated and irradiated zirconium hydride showed that 5- and e-hydride differed by 0.46% and thus the growth in volume was held not be explicable by a transformation from d- to the e-phase. Careful control experiments showed that these effects were not due to temperature. There were no other published data on the irradiation behaviour of the e-phase, although the lack of irradiation effects in 1 MeV) were examined. The original material had a density of 1.84 to 1.85 g-cm" 3 and contained ~ 1 % by weight of BeO. The process of cutting specimens revealed considerable embrittlement. Ells and Perryman annealed samples for one hour at increasing temperatures measuring density changes and collecting the gases released. The densities decreased slowly up to 700 °C and then more rapidly at temperatures from 700 °C to 1000°C reaching reductions of 20% at the higher temperatures. The gas contents varied from
5.3 Solid Moderators
5.2 cm3/(cm3 metal) to 24 cm3/(cm3 metal) from the low to the high dose part of the sample, the theoretical maximum being 23.2cm3/(cm3 metal). The grain boundaries were found metallographically to be decorated with large cavities after a few hours at 700 °C. Swelling continued slowly for hundreds of hours at a given annealing temperature. Simple models of the swelling which assumed that the gas pressure in bubbles overcame the yield stress to produce swelling underestimated the density decrease. Similar results were obtained by Rich et al. (1959) who also found that the irradiated micro-hardness of ~450 fell sharply between 600 °C and 900 °C to - 50. The bend strength fell linearly with annealing temperature above ~ 300 °C to low values at ~700°C. Rich and Walters (1961) studied the mechanical properties of beryllium irradiated at 350 °C and 600 °C. The material was hot pressed and extruded at 1050°C to a flat form. Tensile samples were cut from longitudinal and transverse directions and these were irradiated inside and outside hollow fuel elements in the Pluto materials testing reactor at Harwell. The samples at 350 °C received doses up to 1.6 x 10 1 7 n-cm~ 2 (fission) and 2.6 x 10 1 9 n-cm~ 2 (thermal) while those at 600 °C received doses of up to 4 . 8 x l 0 2 O n c m " 2 (fission) and 1.59 x 10 2 1 n-cm" 2 (thermal). The specimens did not show any changes in density greater than the measurement error of 0.2% at either 350 °C or 600 °C. Post irradiation annealing produced significant density reductions only for temperatures above 1000°C (1 h) for 35O°C and above 900°C (1 h) for 600°C material. Mechanical testing was carried out at temperatures of 20, 150, 300, 450 °C and 600 °C (minimum of two at each temperature). In the 350 °C sample the longitudinal yield stress was increased for temper-
399
atures below the irradiation temperature (by a factor of 1.6 at room temperature) but was essentially unchanged at and above the irradiation temperature. The longitudinal elongation was found to be reduced significantly below the irradiation temperature. In the samples irradiated at 600 °C there was no significant change in the longitudinal yield stress or the elongation at any test temperature up to 600 °C. The transverse properties of the material showed much smaller elongations at lower temperatures in the unirradiated state but there was no significant effect of irradiation at 350 °C or 600 °C. The dose dependence of the longitudinal yield stress measured at 20 °C or 150°C was found to be approximately logarithmic for fission neutron doses between 1018 and 2 x 10 2O n-cm" 2 . Annealing of these samples at temperatures between 300 °C and 1000 °C led to steady decreases of the yield stress. The United Kingdom Atomic Energy Authority carried out substantial irradiation programmes in the Hifar reactor in Australia and the Dounreay materials testing reactor with a view to using the material as a fuel cladding for the advanced gas-cooled reactor (AGR) at Windscale. In this work loss of ductility was significant for irradiation temperatures between 450 °C and 700 °C (Rhodes and Sumerling, 1961) due to the formation of helium bubbles on grain boundaries. The problems of ductility loss in high dose, high temperature irradiations led to the abandonment of its use as a fuel cladding. As we shall see, the corrosion of irradiated beryllium also renders its use difficult. Beryllium is not corroded at a significant rate in air below a temperature of 570 K, but at higher temperatures an adherent layer of beryllium oxide is formed. The oxidation resistance varies significantly with the manner of formation of the beryllium
400
5 Nuclear Reactor Moderator Materials
body, vacuum cast material being superior to that prepared by powder metallurgical methods (Higgins and Antill, 1961). Beryllium corrosion in water occurs at a faster rate than observed for zirconium, but is reduced at pH values greater than ~6.5. Gregg et al. (1960,1961) investigated the oxidation kinetics of electrolytic flake beryllium at temperatures between 500 °C and 750 °C in oxygen (Aylemore et al., 1960), carbon dioxide (Gregg et al., 1961) and carbon monoxide (Aylemore et al., 1961). The effect of addition of water vapour in these gases was also investigated (Aylemore et al., 1961; Higgins and Antill, 1961) investigated oxidation in the same gases over long times at temperatures from 500-1000 °C. In relation to the possible use of beryllium in carbon dioxide coolant, Bennet et al. (1961) studied oxidation of both unirradiated and irradiated material. The observed reactions were Be + CO 2 2Be + CO 2
BeO + CO 2BeO + C
The kinetics of oxidation in CO 2 are of two kinds which are dependent upon the temperature. At temperatures up to 850 °C the oxidation is protective, the rate of oxidation decreasing with time. At higher temperatures the oxidation first decreases with time and then increases rapidly. In the presence of water vapour this occurs at much lower temperatures. Oxidation of pre-irradiated materials showed little difference at 600 °C and 700 °C, there being little increase in volume with time at temperature. At 850 °C the oxidation was initially similar to non-irradiated material, with the accelerating oxidation rate being observed on both materials (normal breakaway oxidation). However when the large swelling sets in, in association with the formation of helium bubbles at grain bound-
aries the oxidation rate increases rapidly due to the associated increases in surface area and open pore volume. A similar effect was observed at 1000 °C in about 2 h, compared to 1000 h at 850 °C. At both temperatures the detailed oxidation rates correlated with detailed changes in the surface area. The relevance of these changes to reactor faults is obvious and a further reason for the rejection of beryllium as a cladding, and its preferred use as a reflector (low dose, low heating) or neutron multiplier. 5.3.7 Beryllium Oxide
Beryllium oxide (BeO), a refractory material, can be considered a potential moderator although the same helium producing reactions are important as in the case of beryllium metal, already described. The treatment of beryllium ores for extraction of the metal, either by the fluoride route (Kawecki, 1955) or by the sulfuric acid route (Schwensfeier, 1955) produce beryllium hydroxide which can be reduced to the oxide by heat-treating at 1800°C. The resulting product is not sufficiently pure for reactor use. Purification is carried out by further dissolution in sulfuric acid, and removal of aluminium by precipitation with ammonium sulfate and this is followed by crystallisation to beryllium sulfate, evaporation and cooling. The high purity sulfate thus obtained is calcined to oxide at 1150°C. Alternatively the freshly prepared beryllium hydroxide can be converted with acetic acid to beryllium acetate which is distilled at 400 °C and then decomposed at 600- 700 °C to give oxide powder. 5.3.7.1 Manufacture
BeO artifacts may be formed using most conventional fabrication techniques for ce-
401
5.3 Solid Moderators
ramies. A variety of binder materials such as resins or starches may be used as pre-firing binders. Where machining of sintered products is required (Beaver and Lillie, I960), a prefiring of the artifacts is carried out at temperatures from 1200-1500 °C, after which machining is possible with carbide tipped tools, followed by a final heat treatment at 1700-2000 °C. The densities of cold pressed and sintered materials may vary considerably, but high densities may be produced using high purity oxide and careful control of the pressing and heating conditions. Densities of 2.98 - 3.04 g-cm ~3, compared to the theoretical density of 3.03 g-cm" 3 were achieved in studies at Battelle Memorial Institute (Udy and Boulger, 1949) after a final sintering at 1400 °C. The crystallite size was found to increase approximately linearly with sintering temperature from 1100-1600 °C, but the bend strength falls steadily over the same sintering temperature range by a factor of about four. The manufacture of beryllium oxide was described by Elston and Labbe (1961) and Carniglia and Hove (1961) together with the properties of the consequent materials. The first of these authors employed sintering under stress in a graphite mould at a temperature of 1750°C and a pressure of 170kg-cm" 2 for 2 - 4 h to obtain high density. Similar methods were employed at the Atomics International Laboratory, but it was also found that MgO additions enhanced densification without causing deleterious effects such as excessive grain growth. Studies in the United Kingdom showed that the density increased rapidly with hotpressing above 1400 °C, reaching densities of ~3g-cm~ 2 above 1700°C. Grain growth occurs rapidly above ~1800°C and samples processed at higher temperatures were found to be weak and brittle.
The optimum hot pressed condition is apparently located at 1700-1800 °C. The properties of high density beryllia are shown in Table 5-7 below. The properties are dependent upon the initial density of the material in the usual way, e.g. the Young's Modulus and thermal conductivity increase with density, whereas the thermal expansion coefficient is insensitive to density, the strengths decrease with density (Beaver and Lillie, 1960; Carniglia and Hove, 1961; Elston and Labbe, 1961).
Table 5-7. Physical and mechanical properties of BeO. Property Crystal structure
Value Hexagonal
Lattice parameters
ao-2.69xl0~8cm co = 4.39xl0~ 8 cm
Theoretical density
3.035 g- cm" 3
Melting point
2550°C
Specific heat 173 K 273 K 373 K 673 K 1973 K 2073 K
O.OSOJg^K"1 0.92 1.29 1.76 2.06 2.21
Thermal conductivity 273 K 373 K 673 K 873 K 1073 K 1273 K 1473 K 1673 K 1873 K 2073 K
ZSlW-cm-^K-1 2.09 0.83 0.42 0.22 0.21 0.17 0.17 0.17 0.17
Thermal expansion coefficient 298- 373 K 298- 573 K 298= 873 K 298-1073 K 298-1273 K
5.5xlO~ 6 K~ 1 8.0 9.6 10.3 10.8
402
5 Nuclear Reactor Moderator Materials
The elastic modulus shows a sharp fall at temperatures above 1300°C and this is accompanied by a similar fall in the bend strength. BeO shows significant creep at a temperature of 155O°C. The critical characteristics required of a nuclear moderator relate to its behaviour under irradiation and its corrosion. 5.3.7.2 Radiation Effects Initial studies of the behaviour of beryllium oxide under irradiation (Elston and Labbe, 1961; Clarke and Williams, 1961; Clarke et al., 1961) showed that the bulk material expanded linearly with neutron dose (En > 1 MeV) for irradiations at temperatures ~100°C. These dimensional changes are accompanied by lattice parameter changes in both the c and a axis directions. For small dimensional changes in isotropic material the macroscopic length changes agree quite well with 1/3 (Ac/c + 2Aa/a). For small changes in dimensions (f§400ppm), small changes in Young's Modulus of both signs were observed, together with negligible effects on strength. The changes in Ac/c are about ten times greater than Aa/a (Elston and Labbe, 1961) and in this work it was found that these changes led to disintegration of the material for doses > 2 x 10 2O n-cm~ 2 (En > 1 MeV) at temperatures - 1 0 0 °C. At an irradiation temperature of 350 °C the changes in lattice parameters occurred more slowly, but with an increased ratio (Aa/a)/(Ac/c). The thermal conductivity showed substantial reductions with a reduced temperature dependence, which could largely be removed by annealing at 800 °C. Extensive studies of the macroscopic dimensional changes and the associated lattice parameter changes were made by later workers (Wilks, 1967). These measure-
ments were made using both powdered polycrystalline samples and single crystals, both of which yielded similar results. The results show a significant scatter but may be summarised as follows: (a) The expansions of the a- and olattice parameters are much smaller for a given dose at temperatures >400°C, than those obtained at temperatures below 150°C. (b) Irradiations at 700 °C show saturation of the oaxis lattice parameters at 2.2% after a dose of 4 x l 0 2 O n - c m ~ 2 (En > 1 MeV). (c) For irradiations ~ 1000 °C the expansion of the a parameter is 0.2%. The macroscopic dimensional changes observed can be generalised: (i) For a given dose the volume expansion decreases with increasing irradiation temperature. (ii) For doses > 2x 1020 n • cm" 2 (£ n > 1 MeV), at temperatures below 150°C large grain sized material expands more quickly than small grain size material. The rate of expansion increases at intermediate doses and then decreases at doses 10 x 10 2 O n-cm- 2 (En > 1 MeV). Comparison of the macroscopic and microscopic changes has been made by a number of authors. A number of factors have been identified which may contribute to the macroscopic volume expansion VM, these are: (a) The volume change determined by lattice parameter measurements Vx. (b) The volume expansion due to microcracking Vc. (c) The crystal expansion produced by dislocation loops VL too large to be major contributors to Vx. (d) The volume expansion Vg due to helium bubbles not observed in Vx. In single crystals VM = Vx for irradiations from -100-600°C, but for irradiations at 650-1100°C, VM is greater than Vx (at
5.3 Solid Moderators
least for doses up to 5 x l 0 2 1 n - c m 2 (En > 1 MeV)). In polycrystalline materials for irradiations at temperatures less than 150°C VM = VX. In the absence of microcracking VM = VX for irradiation temperatures up to 600 °C (Collins, 1965), Woolaston and Wilks (1964) found for irradiations at 1000°C VM>VX for doses up to 4 x 1020 n • cm " 2 (£ n > 1 MeV) in the absence of microcracking. Measurement of the volume of intergranular helium bubbles showed that both helium bubbles and dislocation loops contribute to VM. The loops observed by transmission electron microscopy generally lie on the basal plane and thus produce, as in graphite, an expansion in the c-axis crystal direction. Collins (1965) and Hickman and Pryor (1963) propose detailed models of the macroscopic volume changes in the absence of microcracking; the latter fitting the experimental data more accurately. The dimensional changes in the BeO crystallites lead to material integrity problems. Cracking has sometimes been observed in single crystals irradiated at low and high temperature, but it seems likely that these are associated with inclusions in the crystals. Microcracking in polycrystalline samples can be observed by increases in the open porosity, a decrease of the strain component of X-ray line broadening and most important technically, a sharp decrease in the bend strength leading to eventual disintegration. The results of many workers (see Wilks, 1967) show that irradiation induced microcracking in BeO depend upon the grain size, density, fabrication method and the irradiation temperature. At a given irradiation temperature microcracking and powdering occur at decreasing dose with increasing grain size, whereas for a given grain size the doses required increase with irradiation temper-
403
ature. Cold pressed and sintered material is more resistant to microcracking than hotpressed material, probably because the latter tends to be microcracked prior to irradiation. Lower densities lead to earlier microcracking at irradiation temperatures of about 100 °C. It is believed that the principal cause of microcracking and consequent disintegration at low temperatures of irradiation is the anisotropic crystallite dimensional change which produces strain at grain boundaries. Kingery (1960) and Carniglia and Hove (1961) analysed the stress system associated with the crystallite dimensional changes and concluded that the intergranular stress was proportional to the grain size. Nolle (1963) carried out a more rigorous analysis and concluded that the stress was independent of the grain size. Using this result, Clarke (1964 a and 1964 b) calculated the strain energy due to the anisotropic growth. The model predicted that microcracking would occur when the misfit strain s = [4Syf/(Ed)]1/2
(5-40)
where d is the grain size, E is the Young's Modulus and y{ the grain boundary surface energy. If the misfit s is proportional to irradiation dose then Eq. (5-40) predicts that the dose at which microcracking occurs is proportional to d~1/2. Hickman (1966) obtained a linear relationship (Fig. 5-16) between the neutron dose required to produce microcracking against d~1/2 using the results from five publications concerning irradiations at 50-100°C. The slope of the line yielded yf ~ 1.2 J • m~ 2 for £ = 3x 10 11 N • m~ 2 . This value is in quite good agreement with direct measurement. According to Hickman the same relationship holds for irradiations up to 700 °C. The cause of microcracking and powdering which occurs at higher irradiation tern-
404
5 Nuclear Reactor Moderator Materials
0.5
0.1
1 2 3 4 5 Integrated Neutron flux (1020 nvt>1MeV)
Figure 5-16. Neutron dose required to produce microcracking during irradiation at 50-100°C as a function of grain size in beryllia (Hickman, 1966).
peratures is less certain since the lattice parameter changes are small, two obvious possibilities are anisotropic crystallite dimensional changes due to loops which are not reflected in X-ray lattice parameter measurements or the formation of bubbles on grain boundaries. A third possibility is due to differential granular thermal expansion on cooling from the irradiation temperature. The presence of both lattice defects and microcracking would be expected to reduce the thermal conductivity of BeO, which in its unirradiated state appears to be a lattice conductor. For irradiations at 100°C the conductivity decreases with increasing dose, for doses between 1019 and 4 x l 0 2 O n - c m - 2 (En > 1 MeV), while for fixed doses the reduction in conductivity, at least for small doses 10 2O n-cm - 2
(En > 1 MeV) is less the higher the irradiation temperature (Collins, 1965; Cooper, 1963). Keilholtz et al. (1964) have apparently observed a flux level effect in irradiations at 900 °C. Collins (1965) and Clarke (1963) both observed increases in the bend strength at low doses which appears to be due to the pinning of dislocations necessary for the formation of microcracks. However once the onset of microcracking has occurred the bend strength decreases very rapidly with increasing dose. Small changes in the elastic constants occur, compatible with the density reductions until a rapid fall takes place due to microcracking. There are no changes in the thermal expansion coefficients. A few measurements have been reported of the stored energy in irradiated beryllia. This is quite significant for irradiations at 100 °C, reaching values of up to - 380 J • g" 1 at a dose of 4 x 1020 n • cm" 2 (En > 1 MeV), but this is much reduced at higher temperatures (Roux et al., 1964; Heuer and Stolarski, 1966; Elston, 1963; Dupre et al., 1963; Barner, 1964). It is clear from the above that BeO is only usable as a solid moderator for a limited range of fast neutron doses at each temperature because of the onset of microcracking due to one or more mechanisms of internal distortion. Beryllia is a relatively stable and inert material, although it reacts with water vapour, fused alkalis and hydrofluoric acid at high temperatures. The reaction with water vapour is important for the firing and finishing of the beryllium oxide artifacts, but because of the potential damage to the ceramic and also because of the toxicity of the beryllium released to atmosphere, BeO should not be exposed to moisture at temperatures above ~1200°C.
5.4 Liquid Moderators
BeO dissolves readily in fused alkalis, alkali carbonates and pyrosulphates but reacts only slightly with sulfuric acid. There is a slight reaction with carbon, silicon and boron at high temperatures.
5.4 Liquid Moderators 5.4.1 Light and Heavy Water Moderators
The use of liquid moderators is dominated by light (H2O) and heavy (D2O) water, both of which have desirable properties as reactor coolants and radiation shields. Very large numbers of light water moderator pressurised water reactor (PWRs) and boiling water reactors (BWRs) have been built, while one country (Canada) has pursued the construction of heavy water moderated and cooled (CANDU) reactors. The use of light water as a reactor moderator and coolant has been pursued from the earliest times in nuclear power development. As a reactor coolant it possesses the virtues of high specific heat, high density and low viscosity. It is also easily available, cheap and chemically compatible with many materials which might be used in construction. A limitation of the use of water lies in the low boiling point at atmospheric pressure which would seriously restrict the thermal efficiency of a watercooled system. The required properties can be achieved by increasing the water pressure and this is done in every case where the water takes the role of the coolant: the two types of light water reactors, PWRs and BWRs differ only in the degree of overpressure employed. The systems must be operated at a pressure which will maintain the coolant/moderator as a liquid, which is a function of the core temperature. This saturation pressure increases rapidly from 1 bar at 100 °C to - 2 0 0 bar at 350 °C. Typ-
405
ical conditions in a PWR are 150 bar, 300 °C while for a BWR - 80 bar at 280 °C and while these relatively lower temperatures limit the reactor thermal efficiency, they also reduce the problems encountered with structural materials, particularly compared to gas-cooled reactors. The designers of virtually all commercial PWRs and BWRs have used light water as both coolant and moderator, although reactors have been designed and built using, for instance, graphite moderators with light water cooling (Russian RBMK design for example), or heavy water moderator with light water cooling (United Kingdom SGHWR for instance). Collier (1989) has given an up-to-date summary of the development of water reactors. In the most widely adopted system the coolant from the PWR core passes through a steam generating heat exchanger where steam is generated for supply to the turbine, thus avoiding radioactivity reaching this part of the plant. In the direct cycle BWR the steam passes to the turbine from the reactor core, thus requiring integrity of the whole steam circuit. The moderating power of light water is excellent, but the cross-section for neutron absorption is such that it is necessary to use enriched uranium to achieve system criticality. The slightly enriched uranium ( 2 % - 3 % U 235 ) necessary adds significantly to the associated system costs: the important neutron absorbing reactions are ^ i f a y ^ D and 2 D(n,y) 3 T for water of natural isotopic composition. In the case of heavy water moderators neutron losses are such that reactor criticality can be achieved without fuel enrichment, thus offsetting the increased cost of heavy water compared to light water. The induced radioactivity in the moderator/coolant is important from the point of view of dose to the operators, although as we shall see, a
406
5 Nuclear Reactor Moderator Materials
Table 5-8. Neutron reactions in oxygen. Isotope 16Q
99.76
16 16
17
Half-life s
Comments
O(n,p) 16 N O(n,a) 13 N
7.43 600
6.13, 7.10 MeV
O(n,p) 17 N O(n,a) 14 C
4.14 1.8 xlO 1 1
Reaction
Abundance
0.0374
O
17
17 18Q
0.204
18
O(n,y) 19 O
considerable part of the radioactivity is due to corrosion products in the coolant/ moderator. The most important direct reactions in the water are 16 O(n,p) 16 N and 18 O(n,y) 19 O for which the half-lives are 7.43 s and 29.4 s. Table 5-8 shows the important neutron reactions with oxygen. However as we shall see, the radiolysis of water by the energetic reactor radiations and control of the coolant chemistry are the most important features of the use of water moderators/coolants. 5.4.2 Properties of Light and Heavy Water Table 5-9. Physical properties of light and heavy water. Property
Light water H2O
Molecular weight, amu Density, g • cm ~ 3 Boiling point, K Freezing point, K Critical temperature, K Temperature of maximum density, K Heat of vapourisation, g Thermal conductivity, Viscosity, kg • m ~L - s Refractive index
Heavy water D2O
18.02 1.00 373.1 273.15 647.3
20.03 1.107 377.5 276.96 644.6
277.13
284.34
2232.9
2073.4
2.12 1.96 1.005 xlO~ 3 1.2514 xlO~ 3 1.3326 1.3283
29.43
5.4.3 Radiation Chemistry and Control of Water Moderators
The presence in the reactor core of highly energetic radiation, produces in the water moderator/coolant the phenomenon of radiolysis, similar to that already described for carbon dioxide used as a coolant with a graphite moderator. The water molecules may be excited, ionised or broken up by the reactor irradiation. The excess energy is transferred to other molecules over a timescale of a few atomic vibrations, and this is followed by reaction of the species created amongst themselves and with the normal water molecules. Extensive studies of the radiolysis of water have been carried out (Allen, 1961; Cohen, 1969; Burns et al, 1976 and 1980; Ibe and Uchida, 1983, for example). The process can be represented simply by (Ibe and Uchida, 1983). H 2 O -• H 2 , O 2 , H 2 O 2 ? H, OH, HO 2 H O 2 , O 2 , O , H + , OH e aq
at elevated temperatures (300-410°C). Table 5-10 gives some G-values for the production of primary radiolytic species at these temperatures. The energy deposition rates in the reactor core may be estimated from equations proposed by Cohen (1969).
5.4 Liquid Moderators
407
Table 5-10. G-values of primary radiolytic species at elevated temperatures. Species G-value (number/100 eV)
aq
0.4
H+ 0.4
H 0.3
H2 2.0
OH 0.07
O 2.0
-2.7a
The negative value implies a rate of destruction.
Energy deposition due to y-rays Ey = 9.0 x l(T 2 PM w /(F f FW)(W • cm" 3 ) Energy deposition due to neutrons £ n = 1.7xl0- 2 P/(F f F w )(W-cm- 3 ) where P is the reactor thermal power (W), M w is the mass fraction of water in the core, V{ is the average void fraction in the core and Vw is the volumetric space for coolant in the core. Estimation of the concentrations of radiolysis products in various parts of reactor circuits is complex, the concentrations tending to increase in the core and decay away in external parts of the circuit. Direct observation of concentrations at sampling points is limited and computer models such as AQUARY have been developed to make predictions (Ibe and Uchida, 1983) containing as many as 40 reactions, listed with their rate constants and activation energies in Table 5-11. The radiolysis of heavy water follows the same pattern as in light water. However, it is now essential to recover the deuterium released to use it to generate D 2 O for further use. The presence of neutrons in a heavy water moderator leads to tritium generation by the reaction 2 D (n, y) 3 T producing /^-activity of 0.018 MeV energy and long half-life. Activity levels of 40 Ci/dm3 may be generated (Ursu, 1982) so that it is necessary to submit the moderator and coolant to a tritium scrubbing process. It should also be noted that the reaction 2 D (y, n) XH depletes the deuterium. A separate source of isotopic pollution of heavy water is atmospheric humidity and leakage
or condensation of normal water. The heavy water surface should be covered with a pressurised inert dry gas. 5.4.3.1 Boiling Water Reactors
The radiolytic process in the water systems is in general undesirable since it creates potentially explosive gases in the circuit and also produces oxygenated water and various free radicals which enhance the corrosive action of water on circuit materials. The corrosive action on circuit materials leads to the problems of crud deposition on reactor fuel and excessive activity in reactor circuits leading to high operator dose rates. Further the presence of radiolytically generated oxygen in boiling water reactors is a contributor to the occurrence of stress corrosion cracking in stainless steel reactor circuit components. As a consequence of the radiolysis and gas stripping reactions in the reactor core, the reactor water may contain ~ 200 ppb of dissolved oxygen together with the appropriate stoichiometric ratio of dissolved hydrogen. This level of oxygen in association with internal stresses can produce intergranular stress corrosion cracking of stainless steels and it has been demonstrated in laboratory test that a reduction in dissolved oxygen reduces such effects. Programmes have therefore been instituted to investigate the possibility of reducing the level of radiolytically created dissolved oxygen in BWR systems. A number of candidate additions for boiling water reactor moderator/coolants
408
5 Nuclear Reactor Moderator Materials
Table 5-11. Reactions among radiolytic species (after Ibe and Uchida, 1983). No.
1 2 3
4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23
24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40
Reaction e-+H2O eaq + H +
e~ + H 2 O 2 H+H e~ + H O 2 eaq + O 2 2e a -+2H 2 O 2 OH HO 2 + OH O2+OH H + OH" H + eaq + H 2 O HO2"+e; +H2O
H+VoH" H + OH H 2 + OH H 2 O 2 + OH H + H2O2 H + O2 HO 2 O2+HO2 HO 2 + HO 2 2O2+2H2O2 H + HO 2 H + O2 O 2 + e a q + H2O2 H 2 O 2 + OH~ H2O O2+H + HO2+H2O H2O2 2O HO 2 + O OH + O H2O + O OH + O H2 + O H2O2 + O H+ O
-+ -+ -» -•
H + +OH~ H OH OH + OH"
- H2 -> H O 2
- o2 -> -• -» -> -> -^ -> -> -• -> -> -> -> -> -> -• -> -> -> -• -*
H 2 +2OH~ H2O2 O2 + H2O O 2 + OH~ ea-q + H 2 O H 2 + OH~ OH + 2 O H H2O H2O H + H2O HO 2 + H 2 O OH + H 2 O HO 2 H++O2 HO2+O2 O2 + H 2 O 2 O 2 + H 2 O 2 + 2OH~ H2O2 HO 2 HO 2 + O H " HO2+H2O2 H + OH~ HO 2 H 2 O 2 + OH" 2OH
- o2
-> -> -> -> -> -• ^
O 2 + OH H + O2 2OH HO 2 H + OH HO 2 + OH OH
Rate constant at 25 °C (mol L/s) 1.6 xlO 3 2.4 xlO 1 0 2.4 xlO 1 0 1.3 xlO 1 0 1.0 xlO 1 0 2.0 xlO 1 0 1.9 xlO 1 0 1.64 xlO 6 5.0 xlO 9 1.2 xlO 1 0 1.2 xlO 1 0 2.0 x 107 4.5 x 108 6.3 x 107 1.4378 xlO 1 1 2.0 xlO 1 0 3.4 xlO 7 2.7 x 107 9.0 x 107 1.9 xlO 1 0 8.0 xlO 5 1.5 xlO 7 2.7 x 106 5.6 xlO 3 2.0 xlO 1 0 2.0 xlO 1 0 1.8 xlO 8 5.0 x 108
Activation energy (kcal/mol) 3 3 3 3 3 3 3 3 2 3
3 3 3 3
3 3 4.6 3.4 4.5 3 3 4.5 4.5 4.5 3 3 4.5 4.5 a
a
10
5 xlO 1.022 xlO 4 7.7 x l O ' 4 2.2 xlO 1 0 2.0 xlO 1 0
3 3
1.0
16.8
1.9 xlO 3 2.0 xlO 1 0 4.8 x 103 1.3 xlO 6 2.0 xlO 1 0
9.8 3 8.3
3 3
7.28
4.2 3
Calculated from the reverse reaction rate constant k15 by the equation k29 = Kw/k15.
have been evaluated (Law et al., 1984) including ammonia, hydrazine, morpholine and hydrogen with the objective of suppressing the radiolysis and consequent oxygen production, but hydrogen was se-
lected since it was not expected to effect water chemistry, it did not require modifications to the reactor clean up systems, the condensate demineraliser or significant additions of new equipment. There were
5.4 Liquid Moderators
also studies of hydrogen injection available from various research reactors. It was expected that a feedwater hydrogen concentration of 1.6 ppm would reduce the oxygen level in the reactor water from ~ 200 to 20 ppb. The stress-corrosion cracking of steel is substantially reduced at levels below ~ 25 ppb. A test of these ideas was undertaken at the U.S. Dresden reactor in mid 1982. It had previously been observed empirically that the oxygen concentration in steam and the hydrogen concentration in reactor water were related as (Hammar et al, 1965) D H J ^ [OJ 8team = X (Power)2
where K is a constant. The hydrogen concentration in the feedwater was varied, levels of 5, 200, 400, 1000 ppb and 1800 ppb being established, the first three for two days respectively and the two high concentrations for 4 hours. Measurements were made of the hydrogen and oxygen concentrations in the recirculating water, feedwater and main steam. The presence of excess hydrogen in the water suppresses the water decomposition and the production of oxygen apparently through a chain reaction which rapidly eliminates OH, HO 2 and H 2 and O 2 , the oxygen precursors, i.e. H-, + OH H +o 2 H + HO 2 H + H2O2
-> -> -»
H2O + H HO 2 H2O2 H2O + H
These chain reactions eventually lead to reconversion of the stable radiolysis products H 2 O 2 and H 2 back to H 2 O. The excess H 2 in the coolant provides the initial hydrogen radicals for a chain reaction without net consumption of H. The experimental data showed an "apparent" steam/ water partitioning for O 2 ~90 compared to the Henry law predicted value of ~ 180 at 285 °C. This difference is usually ex-
409
plained as a result of gas rich steam from the separator and additional radiolysis in downcomer regions. In normal BWR operation the stoichiometric hydrogen and oxygen produced by radiolysis are recombined in the off gas recombination unit. When hydrogen is being added the H 2 /O 2 ratio reaching the recombination unit is altered so that there is insufficient oxygen to react with the total hydrogen. The off-gas system can be operated safely provided that the hydrogen concentration is below 4% by volume or the oxygen concentration is below 5% by volume. This is readily achieved during normal operation but in practice, to avoid a transient situation, oxygen is added upstream of the recombination unit to remove the excess hydrogen. The experiments showed that there were no longterm effects on the feedwater and reactor water conductivity. The reactor water pH increased slightly (7.3 to 7.7) and the reactor water conductivity increased (0.28 to 0.34 jus/cm). The concentrations of feedwater corrosion products generally decreased, but the concentrations of some corrosion products (Fe, Ni, Co) in soluble or insoluble fractions increased slightly. The concentrations of activated corrosion products and hence the specific coolant activity increased by 2 - 3 times. The main source of radioactivity in BWR coolant is 16 N during operation. Normally the nitrogen activity carried by the steam is a few percent of the total nitrogen activity generated in the core under normal chemistry conditions. The addition of hydrogen increased the steam activity, probably as a result of an increase of the cation species of 16 N at the expense of the anion species in water. The possibility of suppressing the volatile 16 N by addition of electron scavengers such as nitrous oxide is currently being considered.
410
5 Nuclear Reactor Moderator Materials
A further recent development of BWR chemistry (Wood, 1990) involves the addition of metallic ions to control radiation fields. Work in the United States has shown that injection of 5 ppb of zinc reduces the build-up of radiation due to 60 Co by up to 50%, the dose rate coming into equilibrium in an operating time of about two effective full power years. The injection of zinc appears to work in several ways, reducing 60 Co activity in the coolant, making surfaces less receptive to cobalt pick-up, and competing directly with cobalt for active sites. It was found that a burst of radioactivity due to 63 Zn occurred at shut-downs, but this has been avoided in later operation. According to Wood (1990) six plants in the USA and Switzerland are now using this process. The latest Japanese BWRs have developed the use of iron additions to modify the composition of deposits on fuel surfaces. The plants have high capacity cleanup systems which are very effective in removing iron based impurities. An unexpected result of these additions was the observation that corrosion products were predominantly nickel based and that the resultant crud released activity from the reactor core easily, the addition of iron giving relatively more stable nickel ferrite than nickel oxide. The result however has been a substantial reduction in operator exposure. 5.4.3.2 Pressurised Water Reactors
The chemistry of pressurised water reactors differs considerably from that in boiling water reactors because of the use of soluble boric acid in moderator/coolant is general and longstanding in commercial reactors to control the reactivity of the core. Thus the concentration of boric acid (H3BO3) tends to be high after reactor
shut-downs following fuel loading and to fall with the reactivity during the fuel cycle. The presence of the boric acid leads to increased corrosion of some circuit materials and the need for a variety of additional control systems. In order to minimise the transport of soluble corrosion products and hence minimise the build-up of radioactivity deposited out of the core it has become the general practice at PWRs to add a buffering alkali (0.7-2 ppm) 7 LiOH to optimise the pH value at temperature in the range 5.6-7.5. In the Russian VVER pressurised water reactors potassium hydroxide and ammonia are also used to buffer the boric acid. Some 7Li is created in both cases by the 10 B(n,a) 7 Li reaction also. The intention is to determine the optimum pH with the appropriate lithium hydroxide concentration. Experimentally it is agreed that deposition on fuel (crud) is significantly reduced by raising the pH value above 6.9, an industry standard value which has been maintained for ten years. Studies of reactor deposits in the United Kingdom and reactor coolant studies in West Germany indicate that minimum solubility occurs between pH values of 7.2 and 7.4. Studies in the three Ringhals reactors in Sweden and Millstone 3 PWR in the U.S.A., show that the dose rates around the circuit are reduced by raising the pH level to 7.4 with a maximum lithium concentration of 3.5 ppm. French and German studies have respectively shown similar improvements with pH values of 7.2 and 7.4 respectively, the latter associated with a lithium concentration of 2 ppm. Studies of the optimum value of pH are still continuing because it is still uncertain as to how far the lithium concentration can be increased without producing increased corrosion of the Zircaloy fuel cladding or stress corrosion cracking of the Zircaloy
5.4 Liquid Moderators
fuel cladding or Inconel 600 alloy tubing in steam generators. The Ringhals and Millstone 3 plants referred to above have mode to pH values of 7.2; changes in the circuit radiation fields as a consequence of this change are currently being evaluated. 5.4.4 Transport of Corrosion Products in Water Reactor Systems
It has been found that after a water reactor circuit has been in operation for about one year it may contain a corrosion product inventory of several tens of kilograms. Operational experience indicates that under steady state conditions water reactor coolants are generally carrying less than ten grams at any moment in time, e.g., only ~ 0 . i % of the corrosion product is subject to mass transfer at any time. This result implies that deposition is a dominant and rapid process in the coolant circuit. The fuel surface area in a typical large PWR is ~ 6 x l 0 3 m 2 and given a deposit level of ^0.1 mg • cm ~2 and a neutron flux of 10 1 3 -10 1 4 n • c m ^ s " 1 , this will produce a large inventory of radioactive material. The transport of this activity to the external reactor circuits gives rise to problems of dose to station operations. The important sources of radiation are shown in Table 5-12. It is observed that the deposition of fuel surface crud is generally greater in BWRs
Table 5-12. Important corrosion activation products. Parent nuclide
Radionuclide
Co-59 Ni-58 Fe-58 Fe-54 Cr-50
Co-60 Co-58 Fe-59 Mn-54 Cr-51
Energy (MeV)
Half-life
1.2 0.8 1.1 0.8 0.3
5.3 y 71 d 45 d 313 d 28 d
411
than PWRs (Ivars and Elkert, 1980), by perhaps an order of magnitude and it is less regular in PWRs. Deposition levels of up to 5mg-cm~ 2 have been observed in BWRs corresponding to several tens of kilograms of in-core deposit. In PWRs levels of say 0.05-0.10mg • cm" 2 are more common but levels of up to 9mg-cm" 2 have been observed. The basic problem is due to a complex series of transport processes, both particulate and soluble species are involved in releases from deposit both in and out of the reactor core. The major cause of operator dosage is 60 Co, although 53 Co may be important also in the first few years of reactor operation. The BWR and PWR systems are now considered separately.
5.4.4.1 Crud Transport in Boiling Water Reactors
The stages in the transport of crud in BWRs may be described as follows: (a) Metallic impurities in ionic, colloidal and insoluble oxide forms are released into the condensate and feedwater from corroding ferrous alloy (and in older plants from brass and cupro-nickel components) in components and pipes. A proportion of these impurities is transported to the feedwater depending upon the efficiency of the feedwater/condensate purification plant. It is important to minimise anionic impurity levels to keep corrosion levels. (b) Metal ions are released into the water from corroding surfaces, principally stainless steel, in the primary circuit. The general chemical impurity level is dependent upon the coolant purification rate and the efficiency of the coolant plant is therefore crucially important. (c) Particulates and colloidal species are created when the saturation concentra-
412
5 Nuclear Reactor Moderator Materials
tions for particular species are exceeded (generally oxides or hydrates). This may be a local rather than general phenomenon in a circuit. (d) Soluble ionic species may be adsorbed onto crud particles in the water. An example is the adsorption of dissolved cobalt by ion-exchange or absorption-description on Zircalloy surfaces in the core. (e) Crud and other species which cannot be filtered from the coolant deposit on fuel surfaces. Important variables in the deposition process are coolant iron concentration, heat flux, nucleate boiling, surface condition and Reynolds number. The crud is basically iron based but other species such as copper, nickel, cobalt, manganese are present. Colloidal species and small oxide particles deposit preferentially on boiling surfaces compared to larger more crystalline particulates. (f) The initial deposited material is loosely held on the surface but as the deposit thickness increases the binding at the surface becomes stronger. This can lead in extreme cases to fuel cladding failure. The increased residence time of the inner layer leads to increased radioactivation. (g) There are many mechanisms leading to the release of radioactive species from the core including dissolution, ion-exchange, wear, erosion and spalling due to coolant induced shear forces. Correlations have been established (Comley, 1985) between conditions in fuel channels and surface deposit levels. The peak deposition tends to occur between the position where boiling begins and coolant velocity consequently increases and produces erosion. Deposition decreases with increasing fluid wall shear stress. (h) The released radioactive species are in forms both capable and incapable of removal by filtration and conditions in the coolant, including small temperature dif-
ferences, O 2 and H 2 concentrations may produce some activity redistribution. (i) Deposition of the radioactive material and other corrosion products on surfaces outside the reactor core apparently depends upon the corrosion of the surfaces in these locations. Modelling of this process assumes incorporation of 60 Co in growing oxide films on the surface and hence the contamination rate is governed by the corrosion kinetics of stainless steel. Release of 60 Co from the oxide film is governed by normal solid state diffusion and hence is relatively slow. The corrosion rate of the surface depends on coolant purity and hence this is a critical parameter in deposition. (j) The net result of these processes is that the out-of-reactor deposits consists of double layers, the inner layer due to corrosion of the base metal and the outer layer deposited from the coolant. It is probable that Fe 2 O 3 from the oxidising fuel surface environment is transformed to Fe 3 O 4 in the outer deposits, particularly on coolant circuitry where steam separation occurs and dissolved oxygen levels are low ( l - 3 x l 0 2 p p b ) . The conditions of BWR coolant, that is neutral and oxygenated at ~ 280 °C lead to solubilities of iron of a few ppb, cobalt a small fraction of a ppb, with copper and chromium (as a chromate) soluble at rather higher levels, but the high specific activity of cobalt is significant even at very low levels. The transport phenomena in BWRs are described by a variety of computer models (see Comley, 1985 for references). 5.4.4.2 Crud Transport in Pressurised Water Reactors
Much of the description of transport processes given above for BWRs is applicable to PWRs, but there are some basic dif-
5.4 Liquid Moderators
ferences. The closed circuit of the PWR means that there are no external sources of corrosion products, but there are major differences in the materials of construction such as the large areas of high nickel alloys in the steam generators. In general there are also larger temperature differences around the PWR coolant circuit (280 °C to 320 °C) than in the BWR case (274 °C to 280 °C). This difference in temperature can determine the solubility of metals carried by the coolant and hence the potential for deposition in the core. The coolant chemistry of PWR cases is also a variable due to the use of co-ordinated boron/lithium levels to control the pH and hence minimise the in-core deposits. The solubility of corrosion products in PWR water is clearly an important part of the transport and given the importance of minimising deposits and the consequent radiation dose to the reactor operators, extensive theoretical and experimental studies have been conducted over more than twenty-five years (see the BNES Conferences on "Water Chemistry of Nuclear Reactor Systems"). The consensus of the studies is that it is the operating pH value of the core that is the dominant parameter, as noted already, the higher alkalinities limiting the in-core deposition. However other factors have been considered besides the variable coolant chemistry, Darras (1980) for instance cites redox transitions at low levels of dissolved H 2 (or the presence of O2), hydro-thermic transitions which cause precipitation of spinels from the soluble metals and localised boiling effects which may be involved in the two previous effects. It is also the case that the presence of the reactor radiation fields may have an effect on the deposition of fresh corrosion products. In laboratory experiments, for instance, the deposition rate of iron was
413
found to be enhanced by ionising radiation and pH value (Cohen, 1969). The behaviour of depositing and transporting species in reactor are difficult to predict, not least because of the inability to obtain representative in-core samples during operation. There is however good reason to suppose that pre-conditioning surfaces before full power operation can reduce activity build-up in the short term and perhaps for longer times. Differences in circuit materials and design details, as well as operation are important. 5.4.4.3 Cnid Transport in Pressurised Heavy Water Reactors (CANDU)
Many of the aspects of corrosion product generation, release, transport and deposition in PWRs apply to the pressurised heavy water reactors. In spite of the very significant design differences very similar levels of crud and radioactivity are observed. An extensive programme has been carried out by Atomic Energy, Canada (AECL) in support of the CANDU reactors. Le Surf (1978) has summarised much of this work and concluded that the important processes are: (a) Surfaces outside the reactor core corrode. (b) Corrosion products enter the coolant circuit as ions and particulates. (c) The corrosion products are transported to the reactor core by the coolant and are deposited there. (d) Activated corrosion products are released from the core and transported to the exterior. (e) The activated ions exchange with ions from the corroding surfaces rendering them radioactive. These processes are clearly very similar to those discussed under the PWR section,
414
5 Nuclear Reactor Moderator Materials
in the absence of the effects of boric acid in the coolant. A considerable effort has been expended in attempting to model these processes for all important reactor systems. Comley (1985) has summarised much of this work and also gives more detailed observations of deposition and activity levels in individual reactors. It is worth briefly considering how new plants can be designed and operated to reduce the problems associated with crud generation and circuit activity, the latter in particular being important as allowable doses to operators are reduced. The first priority should be the exclusion of high cobalt alloys and the reduction of cobalt impurities in any structural material where this is economically possible. It is also known that a large amount of corrosion product is released in the commissioning phase, and removal or limitation of these products prior to reactor criticality would significantly reduce the inventory of active materials. The third important possibility is the pre-treatment (conditioning) of surfaces before or during commissioning. For instance new Japanese BWRs have pre-oxidised primary circuit surfaces to give a mature oxide film which is less able to absorb cobalt than a growing film. In the use of PWRs the chemically reducing coolant can be beneficially treated with hydrogen during reactor commissioning. This modifies the oxide films apparently reducing the release processes. In a number of PWRs, particularly in France areas requiring inspection and maintenance in the circuits, are being electro-polished to reduce the micro-surface area, this is expected to reduce both the surface corrosion ratio and activity retention. In existing reactors there will be further studies of the role of pH/boron concentration/lithium concentration and the effects
of higher (>2.2ppm) lithium concentrations on Zircalloy fuel clad corrosion and Inconel-600/690 stress corrosion cracking. There will, because of the preponderance of water reactors, be continuing vigorous development programmes to understand the circuit chemistry and their consequences, particularly with regard to reducing the dosage for operators.
5.5 References Allen, A. O. (1961), The Radiation Chemistry of Water and Aqueous Solutions. New Jersey: Van Nostrand. Amelinckx, S., Delavignette, P. (1966), Chemistry and Physics of Carbon 1,1. Arnold, L. (1992), Windscale 1957, Analysis of a Nuclear Accident. New York: Macmillan. Aylemore, D. W., Gregg, S. I, Jepson, W. B. (1960), J. Nuclear Materials 2, 169. Aylemore, D. W, Gregg, S. X, Jepson, W. B. (1961), J. Nuclear Materials 3, 190. Barner, J. O. (1964), General Atomic Report GA-5123. Barnes, R. S. (1961), United Kingdom Atomic Energy Report, AERE R3769. Beaver, W. W., Lillie, D. W. (1960), Reactor Handbook, Materials, Chapter 44, 879.1 Bell, J. C , Bridge, H., Cottrell, A. H., Greenough, G. B., Simmons, J. H. W, Reynolds, W. N. (1962), Phil. Trans. Roy. Soc. A254, 361. Bennet, M. I, Crick, N. W., Blythe, P. C , Antill, J. E. (1961), UKAEA Report, AERE-R3783. Best, J. G., Stephen, W. X, Wickham, A. J. (1985), Prog. Nuclear Energy, Vol. 16, No. 2, 127. Binkele, L. (1972), High TemperaturesI High Pressures 4, 401. Binkele, J. (1978), Non-Equilibrium Thermodynamics 3, 257. Birch, M., Brocklehurst, J. E. (1987), A review of the behaviour of graphite under the conditions appropriate for protection of the first wall of a fusion reactor. AEA Report ND-R-1434(S). Birch, M., Schofield, P., Brocklehurst, X E., Kelly, B. X, Harper, A. H., Prior, H. (1990), Extended Abstracts "Carbon" 90, 242. Blakslee, O. L., Proctor, D. G., Seldin, E. X, Spence, G. B., Weng, T. X (1970), / Applied Physics 41, 3373. 1 This reference contains a large amount of material concerning the manufacture of beryllium and the consequent material mechanical properties.
5.5 References
Board, J. A., Squires, R. I (1967), Proc. Second SCI Conference on Industrial Carbons and Graphites. London: SCI, p. 289. Bridge, H., Kelly, B. T., Gray, B. S. (1963 a), in: Proc. Fifth Biennial Conference on Carbon. Oxford: Pergamon Press, p. 289. Bridge, H., Kelly, B. T., Gray, B. S., Sorensen, H. (1963 b), in: Proc. Int. Conference Radiation Damage in Reactor Materials. Vienna: IAEA, p. 531. Broeklehurst, J. E., Gilchrist, K. E., Kelly, B. T. (1981), in: Proc. 15th Biennial Conference on Carbon. Carbon Society, p. 546. Burns, W. G., Moore, P. B. (1976), Radiation Effects 30, 233. Burns, W. G., Moore, P. B. (1978), Water Chemistry of Nuclear Reactor Systems 2. London: BNES. Burns, W. G., Marsh, W. R., Kimber, J. (1980), United Kingdom Atomic Energy Authority Report, AERER-9516. Carniglia, S. C , Hove, J. E. (1961), J. Nuclear Materials 4, 165. Clarke, F. J. P. (1963), in: Prog, in Nuclear Energy, Series IV, Vol. 5, p. 221. Clarke, F. J. P. (1964 a), Ada Met. 12, 139. Clarke, F. J. P. (1964 b), J. Nuclear Materials 11, 117. Clarke, F. J. P., Williams, J. J. (1961), J. Nuclear Materials 4, 143. Clarke, F. J. P., Tappin, C , Ghosh, T. K. (1961), /. Nuclear Materials 4, 125. Cohen, P. (1969), Water Coolant Technology of Power Reactors. New York: AEC Monograph, Gordon and Breach Science Publishers. Collier, J. G. (1989), in: Light Water Reactors, Vol. 1, Nuclear Power Technology: Marshal, W. (Ed.). Oxford, p. 208. Collins, C. G. (1965), US General Electric Report, GEMP 106A, 10A, 12A, 14A, 16A, 18A. Cooper, M. K., Palmer, A. R., Stolarski, G. Z. A. (1963), J. Nuclear Materials 9, 320. Darras, R. (1980), Chemistry and Radioactive build-up in the primary circuits of PWR plants, CEA-R5072. Diinner, P. (1978), Interatom Report 54, 3145. Dupre, M., Elston, J., Sicard, L. C. R. (1963), Acad. ScL, Paris 256, 90. Ells, C. E., Perryman, E. C. W. (1959), /. Nuclear Materials 1,13. Elston, J. (1963), J. Nuclear Materials 8, 268. Elston, X, Labbe, C. (1961), / Nuclear Materials 4, 143. Genthon, J. P. (1976), IAEA Specialists Meeting on Radiation Damage Units, Harwell. Glasstone, S., Edlund, M. C. (1950), The Elements of Nuclear Reactor Theory. New York: Macmillan. Grasshof, P. (1978), Interatom Report No. 54.3145. Gray, B. S., Thorne, R. P. (1968), /. Brit. Nuclear Eng. Soc. 7, 91. Gregg, S. X, Hussey, R. I , Jepson, W B. (1960), /. Nuclear Materials 2, 225.
415
Gregg, S. X, Hussey, R. X, Jepson, W. B. (1961), J. Nuclear Materials 4, 46. Hammar, L., Rose, R., Allison, G. M. (1965), in: Proc. Int. Conf Peaceful Uses of Atomic Energy, 3rd Geneva Conference, Vol. 9. United Nations, p. 408. Heuer, P. M., Stolarski, G. Z. A. (1966), /. Nuclear Materials 19, 70. Hickman, B. S. (1966), in: Studies in Radiation Effects, Vol. 1. New York: Gordon and Breach, p. 72. Hickman, B. S., Pryor, A. W. X (1963), J. Nuclear Materials 14, 1101. Higgins, J. K., Antill, X E. (1961), J. Nuclear Materials 4, 190. IAEA Specialist Meeting on Radiation Damage Units for Graphite (1972), Seattle. Ibe, E., Uchida, S. (1983), Water Chemistry of Nuclear Reactor Systems 3. London: BNES. Ivars, R., Elkert, X (1980), in: Experience of Water Chemistry and Radiation Levels in Swedish BWRs, Water Chemistry II. London: BNES, paper 49. Johnson, P. A. V. (1981), J. Nuclear Energy 20, 231. Kawecki, H. C. (1955), The Metal Beryllium. Cleveland: Am. Soc. for Metals, pp. 63-70. Keilholtz, G. W, Lee, X E., Moore, R. E. (1964), /. Nuclear Materials 11,235; J. Nuclear Materials 13, 87. Kelly, B. T. (1964), Phil. Mag. 9, 721. Kelly, B. T. (1978), Prog. Nuclear Energy, Vol. 2, No. 4, 1. Kelly, B. T. (1981), Physics of Graphite. Applied Science. Kelly, B. T. (1985), Prog. Nuclear Energy, Vol. 16, No. 1, 73. Kelly, B. T., Broeklehurst, X E. (1977), in: Proc. Pet ten Conference on Irradiation Creep in Nuclear Materials: J. Nuclear Materials 65, 79. Kelly, B. T., Broeklehurst, J. E. (1979), in: Proc. Fifth SCI Conference on Industrial Carbons and Graphites. London: SCI, p. 892. Kelly, B. T., Martin, W. H., Nettley, P. T. (1966a), Phil. Trans. Roy. Soc, A260, 37. Kelly, B. X, Martin, W. H., Nettley, P. T. (1966b), Phil. Trans. Roy. Soc. A260, 51. Kelly, B. X, Ashton, B. W, Lind, R., Labaton, V. (1975), in: Proc. Twelth Biennial Carbon Conference. American Carbon Society, p. 319. Kelly, B. X, Broeklehurst, X E., Martin, W H., Ashton, B. W. (1976), in: Proc. Fourth SCI Conference on Industrial Carbons and Graphites. London: SCI, p. 429. Kennedy, C. R., Eatherly, W. P. (1981), in: 15th Biennial Carbon Conference. American Carbon Soc, p. 552. Kennedy, C. R., Cundy, M., Kleist, G. (1988), Carbon 88. Newcastle U.X: Univ. Press. Kinchin, G. H., Pease, R. S. (1955), Rep. on Progress in Physics 18, 1. Kingery, W. D. (1960), Introduction to Ceramics. New York: John Wiley.
416
5 Nuclear Reactor Moderator Materials
Kretchman, H. F. (1957), The story of Gilsonite. American Gilsonite Co. Law, R. I , Lin, C. C , Cowan, R. (1984), L. Water Chemistry of Nuclear Reactor Systems 3. London: BNES. Le Surf, I E. (1978), Some Aspects of Primary and Secondary Water Chemistry in CANDU Reactors, AECL-6364. Linhard, J., Nielsen, V., Scharff, M., Thomsen, P. V. (1963), Mat. Fys. Medd. Dan. Vid. Selsk, Vol. 33, No. 10. Lux, I., Pazsit, I. (1981), Annals of Nuclear Energy 8, 319. Mantell, E. (1968), Carbon and Graphite Handbook. Interscience. Martin, W. H., Price, A. M. (1967), /. Nuclear Energy 21, 359. Moore, K. E., Young, W. A. (1968), J. Nuclear Materials 27, 316. Morgan, W. C. (1974), /. Nuclear Materials 51, 209. Nightingale, R. S. (Ed.) (1962), Nuclear Graphite. London: Academic Press. Nolle, H. (1963), N. Nuclear Materials 14, 1101. Norgett, M., Robinson, M., Torrens, I. M. (1975), Nuclear Eng. and Design 33, 50. Paetz, P., Lucke, K. (1972), J. Nuclear Materials 43,13. Price, R. J. (1974), Carbon 12, 159. Reynolds, W. N., Thrower, P. A. (1965), Phil. Mag. 12, 573. Rhodes, D., Sumerling, R. (1961), Internal UKAEA Document. Rich, J. B., Walters, G. P. (1961), United Kingdom Atomic Energy Authority Report, AERE-R-3684. Rich, J. B., Redding, G. B., Barnes, R. S. (1959), J. Nuclear Materials 1, 96. Robinson, M. T. (1969), Nuclear Fusion Reactors, Culham, p. 364. Rounthwaite, C , Lyons, G. A., Snowden, R. A. (1967), Proc. Second SCI Conference, Industrial Carbons and Graphites. London: SCI. Roux, A., Richard, M., Eyraud, L., Elston, J. (1964), Le Jounel de Physique 125, Supp. 3, p. 51. Ruland, W. (1968), Chemistry and Physics of Carbon 4, 1: Walker, P. L., Jr. (Ed.). New York: Marcel Dekker. Schwensfeier, C. W (1955), in: The Metal Beryllium. Cleveland: Am. Soc. for Metals, pp. 71-101. Seldin, E. J., Nezbeda, C. W. (1970), /. Applied Physics 41, 3383. Simmons, J. H. W. (1959), in: Proc. Third Biennial Carbon Conference. Oxford: Pergamon Press, p. 559. Simmons, J. H. W. (1965), Irradiation Damage in Graphite. Oxford: Pergamon Press. Simmons, J. H. W, Reynolds, W. N. (1962), in: Uranium and Graphite. Institute of Metals Monograph No. 27, p. 75. Spence, G. B., Seldin, E. I (1970), /. of Applied Physics 41, 3389.
Summers, L., Walker, D. C. B., Kelly, B. T. (1966), Phil. Mag. 14, 317. Taylor, R., Brown, R., Gilchrist, K., Hall, E., Hodds, A., Morris, F. (1967), Carbon, Vol. 5, p. 519. Taylor, R., Kelly, B. T, Gilchrist, K. E. (1969), /. Phys. Chem. Solids 30, 2251. Thompson, M. W, Wright, S. B. (1965), /. Nuclear Materials 16, 146. Udy, M. C , Boulger, F. W (1949), USAEC Report BMI-T-18. Battelle Memorial Institute. Ursu, I. (1982), Physics and Technology of Nuclear Materials. Oxford: Pergamon Press. Van Houten, R., Baxter, W G. (1962), Trans. Am. Nuclear Soc. 5, 488. Vetrano, J. B. (1970), Nuclear Engineering and Design 14, 390. Wickham, A. J. (1989), Private Communication. Wilks, R. S. (1967), United Kingdom Atomic Energy Authority Report, AERE-5596. Wood, C. J. (1987), Progress in Nuclear Energy, Vol. 19, No. 3, p. 241. Wood, C. J. (1990), Nuclear Engineering International, Feb. 1990, p. 30. Woolaston, H. J., Wilks, J. S. (1964), /. Nuclear Materials 11, 265; 12, 305. Wright, S. B. (1962), in: Radiation Damage in Solids, Vol. II. Vienna: IAEA, p. 239. Yoshikawa, H. H. (1964), Nuclear Sci. Eng. 19, 461. Zijp, W L., Rieffe, H. C. (1972), Reactor Centrum Nederland Report, RCN-161.
General Reading Billington, D. S., Crawford, J. H. (1961), Radiation Damage in Solids. London: Oxford University Press. Chadderton, L. T. (1965), Radiation Demage in Crystals. London: Methuen Ltd. Gittus, J. H. (1978), Irradiation Damage in Crystalline Solids. London: Applied Science. Kelly, B. T. (1966), Radiation Damage to Solids. Oxford: Pergamon Press. Kinchin, G. H., Pease, R. S. (1955), in: Reports on Progressive Physics, Vol. XVIII, p. 1. Thompson, M. W. (1969), Defects and Radiation Damage in Solids. Cambridge: Univ. Press.
Graphite Engle, G. B., Eatherly, W. P. (1972), High Temperature-High Pressures, Vol. 4, p. 119. Le Groupe Francaise d'Etudes des Carbones (1965), Les Carbones, Vols. 1 and 2. Paris: Masson et Cie S A. Kelly, B. T. (1981), Physics of Graphite. London: Applied Science Publishers.
5.5 References
Mantell, E. (1968), Carbon and Graphite Handbook. New York: Interscience. Nightingale, R. (Ed.) (1982), Nuclear Graphite. London: Academic Press. Reynolds, W. N. (1966), in: Chemistry and Physics of Carbon, Vol. 2: Walker, P. L., Jr. (Eds.). New York: Marcel Dekker, p. 121. Reynolds, W. N. (1986), Physical Properties of Graphite. New York, Amsterdam: Elsevier Press. Simmons, J. H. W. (1975), Radiation Damage in Graphite. Oxford: Pergamon Press. Ubbelodhe, A. R., Lewis, F. A. (1960), Graphite and its Crystal Compounds. Oxford: Univ. Press.
417
Light and Heavy Water Allen, A. O. (1961), The Radiation Chemistry of Water and Aquesus Solution. New Jersey: Von Nos-
trand. Cohen, P. (1969), Water Coolant Technology of Power Reactors. New York: AEC Monograph Gordon and Breach Science Publishers. Comley, G. C. W. (1985), in: Progress in Nuclear Energy, Vol. 16, No. 1, p. 41. BNES Water Chemistry Conference No. 1 (1977). BNES Water Chemistry Conference No. 2 (1986). BNES Water Chemistry Conference No. 5 (1989).
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors Frank A. Garner Pacific Northwest Laboratory, Richland, WA, U.S.A.
List of 6.1 6.2 6.3 6.4 6.5 6.6 6.6.1 6.6.2 6.6.3 6.6.3.1 6.6.3.2 6.6.3.3 6.6.3.4 6.6.3.5 6.6.3.6 6.6.3.7 6.6.3.8 6.6.4 6.6.4.1 6.6.4.2 6.6.4.3 6.6.4.4 6.6.5
Symbols and Abbreviations Introduction Influence of Reactor Environment Phase Stability During Irradiation Dislocation Evolution During Irradiation Transient Versus Steady State Behavior Dimensional Stability of Irradiated Steels Strains Arising from Precipitation and Dislocation Evolution Strains Arising from Void Swelling: An Overview Variables Which Influence the Swelling of Steels Crystal Structure Base Composition Solute Additions Thermomechanical Treatment Displacement Rate Temperature Temperature History Stress Strains Arising from Irradiation Creep Introduction to Irradiation Creep Irradiation Creep in the Absence of Swelling Irradiation Creep After Swelling Begins Disappearing Creep and Its Consequences Deformation Behavior in Response to Creep, Swelling and Gradients in Important Variables 6.6.5.1 Interactions Between Creep and Swelling-Induced Stresses 6.6.5.2 Strategies for Reducing the Consequences of Swelling and Creep 6.7 Irradiation-Induced Changes in Mechanical Properties 6.7.1 Austenitic Alloys Prone to Void Swelling 6.7.2 Swelling-Resistant Alloys 6.8 Summary 6.9 References .
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
420 422 423 427 434 438 439 439 444 449 450 453 460 467 475 478 480 482 483 483 489 496 502 508 509 512 515 515 530 533 534
420
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
List of Symbols and Abbreviations A B Bo C D E Jc t Ms r S
creep modulus creep rate creep compliance void concentration creep-swelling parameter energy fracture toughness length martensite start temperature mean void radius instantaneous volumetric swelling rate per dpa
a a v , ad S/ 0.1)
4.0
3.0 -100
-75
-50
-25
0
25
50
75
100
125
150
Distance From Core Midplane (cm)
Figure 6-2. Displacement effectiveness factors for several MOTA experiments, conducted in the mixed oxidefueled FFTF reactor, showing the relation to the various MOTA irradiation canisters (after Garner and Greenwood, 1992).
6.2 Influence of Reactor Environment
r
0.14 -
425
\
0.12 \ 0.10 Cross Section
\
0.08
[barns) 0.06 /
0.04 0.02 1
i 2
i
i 4
l
II
I
y
>
6
8 10
\
Cr
V
Figure 6-3. Cross sections for (n, a) reactions as a function of neutron energy for common elements used in structural steels (after Mansur and Grossbeck, 1988).
20
Energy (MeV)
1981c; Mann, 1982). This is due primarily to the absence of thermal neutrons, although the softer nature of the high energy neutrons also contributes. It is generally accepted, however, that production of gaseous elements by transmutation, especially helium, plays some role in the radiation-induced microstructural evolution (McElroy and Farrar, 1972; Mansur and Grossbeck, 1988). Relatively low levels of hydrogen are also produced by (n,p) reactions. Hydrogen is highly mobile in steels compared to helium and thus does not build up to levels that significantly influence microstructural evolution in LMR materials. While some helium arises from (n,a) reactions at low neutron energies with the small amounts of boron found in most steels, the major contribution in the early stages of the irradiation comes primarily from threshold-type (n,a) reactions with the major alloy components. This type of reaction occurs only above higher neutron threshold energies ( > 6 MeV). Figure 6-3 shows that nickel is the major contributor to helium production by (n,a) reactions and thus the helium generation rate scales almost directly with nickel content for a large number of commercial steels, as
shown in Fig. 6-4. The helium contributions from other elements are not sufficiently large enough to compete with that of nickel, regardless of the composition of the steel. A secondary helium-generation process also involving nickel produces helium via the two-step 58 Ni(n,y) 59 Ni(n,a) 56 Fe reaction (De Raedt, 1982; Greenwood, 1983). This process operates very strongly in mixed-spectrum reactors. 59 Ni is not a naturally occurring isotope and thus this contribution involves a delay relative to that of single step threshold (n,a) reactions. Since both steps of the reaction involve cross sections that increase with decreasing energy and the second step exhibits a resonance at 203 eV, the generation rate per dpa in LMRs increases strongly near the core boundaries and out-of-core. In recent studies designed to examine the effect of helium, some simple model steels have been doped with 59 Ni prior to irradiation in FFTF to avoid the delay associated with 59 Ni buildup (Simons et al., 1986; Garner etal., 1993d). As shown in Fig. 6-5, the increase in helium production rate near and beyond the core boundary becomes obvious even at low dpa levels for doped
426
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
600
10
• 80A
Helium appm 6
625
HAST-X
-
901 \ .
115
^
. ^
X-750
\
M-813
-
70^^0-979 A-286X*330
4 — _
718
PE-16
455^X
HT-9
10
I
I
I
I
I
I
I
20
30
40
50
60
70
80
90
% Nickel Figure 6-4. Predicted helium concentrations arising from threshold (n, a) reactions in typical commercial structural steels irradiated at core center of EBR-II to 1 x 1022 n cm" 2 (E > 0.1 MeV) or « 5 dpa during the early stages of the U.S. LMR materials program (unpublished calculations of E. P. Lippincott, R. W. Powell and K A. Garner of Hanford Engineering Development Laboratory, 1974).
specimens. In non-doped situations typical of fast reactor irradiation, however, it requires only a few dpa for the secondary process to approach and eventually surpass the contribution of the primary (n,a) process. The secondary process thus makes the major contribution, both in and out of the core, especially in oxide fueled cores which have softer neutron spectra than metal fueled cores. Another source of helium arises from the implantation of helium into the inner layer of fuel cladding. This helium comes from two major sources, namely from ternary fission events (two heavy fission fragments plus an alpha particle) in the fuel and from helium recoiling from the pins' helium cover gas as a result of collisions with neutrons (Garner etal., 1982). These two sources of injected helium are
slowly reduced during irradiation, however, as heavy fission gases build up in the space between the fuel pellet and the cladding. These gases interact with the energetic helium atoms, reducing their energy sufficiently to preclude most of them from reaching the cladding. Some studies have cited this early source of helium as contributing to the embrittlement of fuelled cladding and its performance during transient heating tests (Hunter and Johnson, 1979), although more recent studies have linked the major mechanism to the fission products cesium and tellurium (Duncan etal., 1981; Hamilton etal., 1981). In terms of component behavior, however, one of the most important factors is the presence of gradients in important environmental variables that exist across or
6.3 Phase Stability During Irradiation
along a component. Spatial gradients in temperature, neutron spectrum, stress and displacement rate have all been found to be very important in determining the evolution of microstructure and macroscopic properties. These variables are often timedependent, and temporal gradients in these parameters sometimes exert a very strong influence on component behavior. Examples of the influence of both types of gradients will be shown in later sections.
(a) 20
Level BC
1
2
3
59
15
4
5
6
7
8
Ni Doped
appm "dpT10
Undoped
-120
-80
-40
0
40
80
120
427
6,3 Phase Stability During Irradiation Even in the absence of neutron irradiation, structural steels can be viewed as being in a state of metastable equilibrium. The various phases and their distribution are determined first by composition and thermomechanical treatment, and second by the details of their service history, especially that of temperature and stress state. Extensive precipitation, due to aging at elevated temperatures, of carbides, intermetallics and many other phases occurs while in service. Table 6-1 lists the wide variety of crystal structures and phases reported to form in austenitic steels. Phases of equal complexity and diversity form in ferritic steels. With only a few exceptions, most major LMR components have been constructed from austenitic steels. Table 6-2 presents the compositions of the steels discussed in this chapter. The details and kinetics of phase formation in any given steel are often rather
Height, cm
(b)
1.2
Fe - 15Cr - 25Ni (Undoped) 1.0
Average Helium Generation Rate 0.6 appm dpa 0.4
Figure 6-5. (a) Calculated helium (appm)-to-dpa ratios for both 59Nidoped and undoped Fe-15Cr-45Ni alloy as a function of height and canister level after irradiation in FFTF for two MOTA cycles (after Garner et al., 1993 d). (b) Measured increases in helium generation rate during irradiation in FFTF of Fe-15Cr-25Ni (Garner and Oliver, 1993). See Fig. 6-2 for canister and basket locations.
428
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Table 6-1. Crystal structure and composition of phases in austenitic steels (after Harries, 1981). Phase
Crystal structure
Composition
f.c.c. pseudo-hexagonal f.c.c. f.c.c. (diamond type)
TiC; NbC Cr 7 C 3 ;(Fe,Cr) 7 C 3 (Cr 16 Fe 5 Mo 2 )C 6 ;(FeCr) 23 C 6 (Cr,Co,Mo,Ni)6C; (Ti,Ni)6Q (Nb,Ni)6C
f.c.c. f.c.c.
Nb(C,N); Ti(C,N) Cr 2 N
f.c.c.
(Fe,Cr,Mo)23(C,B)6
f.c.c. b.c. tetragonal orthorhombic b.c.c.
Ni3(Ti,Al) Ni 3 Nb Ni 3 Nb; Ni3Ti Ni(Al,Ti)
b.c. tetragonal hexagonal b.c.c. rhombohedral f.c.c. hexagonal
(F^Ni^Cr.Mo), Fe2Mo; Fe2Ti; Fe 2 Nb (Fe,Ni) 36 Cr 18 Mo 4 ; Cr 6 Fe 18 Mo 5 (Co,Ti)7(CrW)6 (Ti,Zr, V,Nb, Ta, Mn)6(Ni, Co), 6 Si 7 Fe-Cr-Mo
h.c.p. b.c.c. f.c.c.
Ti 4 S 2 C 2
Carbides MC M7C3 M 23 C 6 M6C Nitrides M(C,N) M2N Borocarbides M23(C,B)6 Geometrically close-packed phases
Y Y' 5
P Topologically close-packed phases a Laves X G R Others Boro-sulfides Cr-rich ferrite (o^) a-manganese sulfide
complex, however, being strongly dependent on many variables related to the original thermomechanical history, the service history, and the amount and distribution of various minor elements such as carbon, phosphorus, boron, etc. In some cases the phase evolution involves early formation of some phases such as carbides, which then alter the matrix composition sufficiently to set the stage for other phases to form, such as intermetallics. A review of the physical metallurgy of Fe-Cr-Ni austenitic steels with respect their potential
MnS
application to nuclear service has been presented elsewhere (Harries, 1981). During irradiation, however, the phase evolution can be significantly altered, both in its kinetics and in the identity and balance of phases that form (Russell, 1984; Nolfi, 1983). Phases can be altered in their composition from that found in the absence of irradiation and new phases can form that are not found on the equilibrium phase diagram of a given class of steels. In AISI 316 type steels these new or altered phases have been classified into radiation-
Table 6-2. Typical compositions of alloys discussed in this chapter (wt.%). Alloy
Fe
Ni
Cr
Ferritic
EM10 EM12 HT9 9Cr-lMo
bal. bal. bal. bal.
0.18 0.12 0.47 0.09
Austenitic
AISI310 RA-330 Incoloy 800a AISI316 M316 D9 FV548 AISI304 AISI304L AISI321 AISI347 AISI216 PCA DIN 1.4970 DIN 1.4981 DIN 1.4864 AMCR 0033
bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal. bal.
A 286 Nimonic PE16 a Inconel 706 a Inconel X-750a Inconel 625a Hastelloy X a Inconel 600
Precipitation strengthened
C
Mo
Mn
Nb
Al
Ti
Si
Other
8.76 9.58 12.0 8.61
0.105 0.086 0.020 0.081
1.05 1.91 1.03 0.89
0.48 0.92 0.50 0.37
_ 0.41 0.07
_ -
_ -
0.37 0.37 0.41 0.11
N = 0.024 V = 0.28 W = 0.5, V = 0.32 V = 0.21
19.7 36.05 33.7 13.7 13.8 15.8 11.8 9.35 9.26 9.5 9.1 6.7 16.6 15.2 16.0 35.0
>ility Duri
Type
rr
Z3 CD
=7
iatio
I 3
Nickel base
a
Incoloy and Inconel are registered trademarks of the International Nickel Company, Nimonic is a registered trademark of Henry Wiggin & Co., U.K., and Hastelloy is a registered trademark of the Cabot Corporation.
CD
430
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
induced phases, radiation-modified phases, and radiation-enhanced phases (Lee et al., 1981; Yang et al., 1981; Williams, 1982; Yang, 1987). These classifications are equally applicable to phases formed in other steels. Little (1993) has presented a review of the phase stability of ferriticmartensitic steels during irradiation. Radiation-induced alterations occur because new driving forces arise that do not occur in purely thermal environments. First of these new driving forces is the presence of very large supersaturations of point defects, especially at relatively low irradiation temperatures (250-550 °C). Not only are vacancies present in unchar-. acteristically high levels, thereby accelerating normal vacancy-related diffusional processes, but interstitials are also abundant. Solutes which can bind with either type of point defect tend to flow down any microstructurally-induced gradient of that defect, providing a new mechanism of solute segregation referred to as solute drag (Okamoto and Rehn, 1979). This has been judged to be particularly important for binding of smaller solute atoms with interstitials. A second new driving force is the inverse Kirkendall effect (Marwick etal., 1979) whereby differences in elemental diffusiv-
KIRKENDALL EFFECT
SOLUTE FLUX
C v (x)
VACANCY FLUX
INVERSE KIRKENDALL VACANCY FLUX '
ity (DFe, D Ni , DCr) via vacancy exchange lead to segregation of the slowest diffusing species at the bottom of sink-induced vacancy gradients. This mechanism is illustrated in Fig. 6-6 and is particularly effective in segregating nickel in Fe-Cr-Ni alloys at all sinks which annihilate vacancies, leading to nickel-rich shells or atmospheres on grain boundaries and other microstructural sinks. This segregation arises because the elemental diffusivities are different, with DCr > DFe > DNi at all nickel levels (Rothman et al., 1980; Esmailzadeh and Kumar, 1985; Esmailzadeh et al., 1985; Garner and Kumar, 1987). In Fe-CrMn steels, it is iron that segregates via the inverse Kirkendall mechanism, manganese being the fastest diffusing element (McCarthy and Garner, 1988). A third new driving force results from the action of the other two driving forces when operating on microstructural sinks that are produced only in irradiation environments. These are Frank interstitial loops, helium bubbles, and voids that may have developed from helium bubbles. Precipitates are often observed to form and to co-evolve with such radiation-induced sinks. Thus these sinks have been implicated as participating in the evolutionary path taken by the precipitates and thereby
FAST Al VACANCY DIFFUSERS / / ^ G R A D I E N T NEAR SINK
SOLUTE FLUX
SLOW DIFFUSING ELEMENTS SEGREGATE BY DEFAULT AT BOTTOM OF VACANCY GRADIENTS
SLOW DIFFUSERS Figure 6-6. Schematic illustration of the inverse Kirkendall segregation mechanism.
6.3 Phase Stability During Irradiation
influencing the microchemical evolution of the matrix (Maziasz and Braski, 1984; Maziasz, 1984; Mansur e t a l , 1986). When the solute drag mechanism, operating between interstitials and smaller size silicon atoms, combines with nickel segregation via the inverse Kirkendall mechanism, phases that are rich in nickel and silicon often form (y' and G phase for example), although in many steels they cannot form thermally. Other phases that are normally stable in the absence of radiation can be forced during irradiation to become enriched in these two elements. Not only does the formation of nickeland silicon-rich phases affect the mechanical properties of irradiated steels but the removal of these elements (as well as a few others) from the matrix appears to be the
< M - 9 . 8 X 1 0 2 2 n/cm2
455°C
431
major determinant of when the onset of void swelling and accelerated irradiation creep occurs in typical commercial steels (Garner, 1981 a, 1984; Garner, and Wolfer, 1981, 1984; Porter and Wood, 1979; Porter, 1984). This is demonstrated in Fig. 6-7, where both void swelling and Y'(Ni3Si) formation in an AISI 316 fuel pin tube both increase from negligible amounts to large amounts in response to a relatively small difference in temperature (Brager and Garner, 1978). The removal of nickel and silicon from the matrix by precipitation exerts a large effect on the effective vacancy diffusivity, resulting in a strong effect on void nucleation (Garner and Wolfer, 1981; Garner and Kumar, 1987). On a per atom basis, phosphorus has been shown to exert an even larger
• t - 10.0 X 10 2 2 n/cm2 T ~ 480°C
Figure 6-7. Correlated development of voids (shown in bright field) and y' precipitates (shown in dark field) observed in 20% cold worked AISI 316 cladding of the PNL-11-9R fuel pin irradiated in EBR-II, The density change profile shown in the top panel shows that the swelling varies strongly as a function of position (after Brager and Garner, 1978).
432
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
effect on the effective vacancy diffusivity (Garner and Kumar, 1987). Some phases such as phosphides and TiC, especially when precipitated on a very fine scale, are thought to be beneficial in resisting the evolution of nickel silicide type phases (Lee and Mansur, 1986; Lee et al., 1984; Maziasz, 1984). An example of the influence of phosphide precipitation in delaying swelling is shown in Fig. 6-8. It has been shown, however, that the segregation process eventually overwhelms these phases, causing their dissolution and replacement with nickel- and silicon-rich phases (Lee et al., 1984; Thomas, 1982; Itoh et al., 1987). At this point void nucleation and void growth are almost always accelerated. Shibahara et al. (1993) demonstrated that the incubation period for swelling was prolonged by extending the duration of the neutron exposure required to cause phosphide dissolution. Titanium additions were shown to lead to the maintenance of more stable phosphides. Some researchers, including Shibahara, prefer to ascribe the action of phosphides
(a)
" " ^
to their trapping of helium in subcritical bubbles rather than focus on their role in delaying the microchemical evolution of the alloy matrix. In general, it appears that the five most important elements in the determination of swelling and creep behavior of typical LMR steels are Ni, Si, P, Ti, and C, all of which participate strongly in precipitate formation. The latter three elements appear to exert most, but not all, of their influence on the thermally stable phases. These phases are most important at higher temperatures, while the nickel silicide type phases are most important at lower temperatures. This difference in temperature regimes will be shown to exert a very large influence on void formation and the resultant swelling profiles observed over the length of reactor components. There is another category of late-term phase changes that arise from radiation-induced nickel segregation and chromium depletion at various radiation-induced microstructural sinks, especially voids, when they come to dominate the microstructure.
(b)
Figure 6-8. Effect of phosphorus concentration on precipitate formation and void growth in annealed AISI 316 stainless steel irradiated in EBR-II to 7 x 1022 n cm" 2 (E > 0.1 MeV) at 600°C. (a) 0.04 wt.% P, (b) 0.08 wt.% P (after Garner and Brager, 1985 a).
6.3 Phase Stability During Irradiation
When AISI 304L stainless steel with 9.3 % Ni was irradiated in EBR-II at ^500°C, nickel segregation at void surfaces removed so much nickel from the matrix that it was almost completely transformed to ferrite, leaving austenite only as shells on the voids, as shown in Fig. 6-9 (Porter, 1979). AISI 316 stainless steel with « 1 2 % ' Ni usually resists this process, but when larger than average levels of silicon were added to the steel, the combined formation of nickel silicides and nickel-enriched void shells during irradiation caused a similar y -> a transformation of the matrix (Brager and Garner, 1981; Williams et al., 1987). In some model alloys prone to such transformations during aging, irradiation considerably enhances the process (Bullough et al., 1987). Ferromagnetic phases form at much lower levels and sizes long before swelling and large scale segregation occurs, however. Magnetic measurements show that progressive formation of very small magnetic centers occurred during neutron irradiation in a variety of austenitic steels. (Baron etal., 1974; Stanley and Garr, 1975; Stanley and Hendrickson, 1979; Stanley, 1979). These centers were identified by electron microscopy in type 321 stainless steel to be ferrite. The precipitates appear to form via small scale segregation, nucleation and growth at microstructural sinks, and not as a result of a martensitic transformation. Ferrite formation was found to be sensitive to the temperature, dpa level and alloy composition, especially that of molybdenum, titanium and carbon. As will be discussed in Section 6.6.1, there are measurable macroscopic strains and density changes that arise from most phase transformations. These strains have a strong impact on the interpretation of swelling and creep data and also on the early behavior of some structural components.
433
Figure 6-9. Solution annealed 304 L irradiated in EBR-II at ^500°C, (a) bright field electron micrograph, (b) dark field micrograph taken using a 110-oc reflection, (c) dark field micrograph showing austenite shells around voids (after Porter, 1979).
434
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
6.4 Dislocation Evolution During Irradiation Dislocations, their mobility, and impediments to their mobility determine the major features of the mechanical properties of a metal. Not only are dislocations also important in thermal creep but they have been cited as the major determinant of both irradiation creep and void swelling (Brailsford and Bullough, 1973; Matthews and Finnis, 1984; Wolfer, 1984). The basic driving force for creep and swelling is thought to be related to a slight preference or "bias" of dislocations towards absorbing interstitial atoms when confronted with equal fluxes of interstitials and vacancies (Wolfer and Ashkin, 1975). The stress sensitivity of the bias term is
2x10 22 n/cm2 (a)
thought to be the origin of irradiation creep. Since dislocations also serve as sinks for point defects and thereby generate point defect gradients in their vicinity, they also participate in the radiation-induced segregation process (Garner, 1981 a). Thus they participate in precipitate formation and the microchemical evolution of the alloy matrix. An example of y' formation on Frank interstitial loops is shown in Fig. 6-10. Since the radiation-induced phases often form on dislocation loops, it is not surprising that phases such as y' exhibit a dependence of size and density on irradiation temperature that is comparable to that of Frank loops. Dislocation densities vary strongly from steel to steel with the major determinant being the initial cold work level, the solute
7x10 22 n/cm 2 (b)
Figure 6-10. Dark field micrographs showing formation of y' in silicon-modified AISI 316 irradiated in EBR-II at 482 °C; (a) starting on the edge of Frank loops and growing inward; (b) after consolidation to a more compact three dimensional shape (courtesy of E. H. Lee, of Oak Ridge National Laboratory).
6.4 Dislocation Evolution During Irradiation
strengthening and the precipitate type and density. In cold worked steels there is usually a pronounced texture to the dislocation distribution that is specific to the type of reduction process involved in the cold working. In larger structural components there are usually spatial gradients in the effective cold work level as well as in the texture. Associated with the cold work and texture is a storage of energy in internal stresses (Challenger and Lauritzen, 1975; Bates etal., 1981b). This energy is available to participate in both thermallydriven and irradiation-affected processes, including thermal and irradiation creep, precipitation and phase changes. The release of this energy is facilitated by the application of external stress and elevated temperature (Garner et al., 1993 c). Typical annealed austenitic steels exhibit dislocation densities on the order of 10" 8 cm" 2 while cold worked materials are in the 10 11 - 10 12 cm" 2 range. The starting dislocation density is not maintained during irradiation, however. There is a tendency for the dislocation microstructure to evolve toward a saturation or quasi-steady state level that is independent of the starting state. This tendency has been observed in many simple f.c.c. metals (Al, Cu, Ni), model Fe-Cr-Ni austenitic alloys and commercial stainless steels (Garner, 1993). The process involves irradiation-assisted climb and glide of network components, annihilation of dislocations of opposite sign, and the nucleation, growth and unfaulting of Frank interstitial loops to form new network line length (Garner and Wolfer, 1982 a, b; Stoller, 1990). This process is also strongly affected by the application of stress. The most significant influence of applied stress is the development of a dislocation and loop structure that is highly anisotropic in its distribution of Burgers vectors (Garner
435
et al., 1979; Gelles et al., 1981; Garner and Gelles, 1988; Gelles, 1993). More discussion on the effect of stress on microstructural evolution, swelling and particularly irradiation creep will be presented later. For stainless steels the neutron-induced saturation density of network dislocations has been measured to be 6 ± 3 x l O 1 0 cm" 2 , relatively independent of starting state, temperature, displacement rate, He/dpa ratio and most other important variables. This saturation process involves an order of magnitude reduction in the dislocation density of cold worked steels and a comparable or larger increase in the density of annealed steels (Brager et al., 1977; Azam et al., 1979) as shown in Figs. 6-11 and 6-12. The bulk-averaged dislocation densities derived from the data presented in Fig. 6-12 exhibit much less scatter than do measurements acquired by microscopy. It is therefore easier to observe the dependence of dislocation density on other variables. Typical densities determined by microscopy are shown in Figs. 6-13 to 6-15. The scatter represents not only the local inhomogeneity of dislocation structure in cold worked metals but also reflects the difficulty of accurately measuring the dislocation density by microscopy. There is a general but incorrect perception that the neutron-induced saturation density of network dislocations is moderately or even strongly dependent on temperature (Stoller, 1990; Lucas, 1993). This perception arose primarily from the tendency of early U.K. papers to ambiguously report the total dislocation density, without specifying it as such. These data included both network and loop line length. When the first subsets of the data shown in Fig. 6-15 were published (Bramman et al., 1977), they were designated only as the "line dislocation density." Another contribution to the perception of a strong tern-
436
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors 1012 20% COLD-WORKED
NETWORK DISLOCATION DENSITY 1010 "SOLUTION ANNEALED
108
10x1022
2 4 6 8 NEUTRON FLUENCE (E>0.1 MeV)
Figure 6-11. Saturation of network dislocation density in both 20% cold worked and annealed AISI 316 after irradiation in EBRII at 500 °C (after Brager et al., 1977).
I 4>% COLD WORKED GED AT 600° 401 • S2S°C
66d°C 500°C
2
A(20)x1O~ (WIDTH A T MID-HEIGHT)
Figure 6-12. Independence of dislocation density in AISI 316 L on starting condition and temperature after irradiation in the RAPSODIE reactor, as measured by X-ray line broadening of the (311) austenite reflection (after Azam etal., 1979).
36
- ANNEALED
30
-AGED AT600°C
20
40
80
60
DISPLACEMENTS, dpa
II
II
1011 -
II
1
1
0
-
r NETWORK DISLOCATION DENSITY cm-2
•
-
D
cP
n
a
O
• • •
o
"
'_ c on -
•
O-
-
1010 I
u
FUEL PIN IDENTITY • PNL-9-30 Q PNL-11-9R o WSA-4-25 1 | 1 1 1 1 1 1 400 500 IRRADIATION TEMPERATURE, °C
_ -
600
Figure 6-13. Network dislocation densities measured by microscopy in cladding from three 20% cold worked AISI 316 fuel pins at doses ranging from 20 to 50 dpa (NRT) after irradiation in EBR-II (after Brager et al., 1977).
6.4 Dislocation Evolution During Irradiation 20% COLD-WORKED
ANNEALED
10.12
437
I
DOSE IN dpa (H/2) 10,11 -
30
29
28
_r
Pd
27
T
n-2
Figure 6-14. Network dislocation densities measured by microscopy in M 316 cladding from two fuel pins irradiated in DFR (after Brown and Linekar, 1974).
10.10 -
i
109 420
500
460
600
500
550
600
640
IRRADIATION TEMPERATURE. °C
perature dependence of line length arose from the early publication of relatively low fluence data for solution annealed steel (Barton et al., 1977). The annealed steel had not yet approached the saturation state, especially at higher temperatures. The data of Barton et al. and Bramman et al. dominate most data compilations of other researchers. i
I
i
The loop density and its associated line length are strongly dependent on irradiation temperature, as shown in Fig. 6-16, and therefore it is the loop line length that accounts for the apparent strong temperature dependence usually attributed to the network dislocation component. The data in Figs. 6-15 and 6-16 demonstrate that Frank loops also approach a saturation or 1
• Q
•
#
±
• D •
#
TOTAL DISLOCATION DENSITY (cm-2) 1 x 10 10
9
1 x 10
-
D
V1294 I 8.5% MAXIMUM r BURN-UP, 38 dpa (NRT) V1256
•
V1257
-
18.5% MAXIMUM BURN-UP, 67 dpa (NRT)
i
I
I
I
300
400
500
600
TEMPERATURE, °C
Figure 6-15. Total of network and loop dislocation density observed in M316 cladding from three fuel pins irradiated in DFR (after Brown and Fulton, 1979; Brown etal, 1983).
438
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
ioie;La.1
1
' >
FRANK LOOP DENSITY cm-3
1015
-
O.IMeV) \ V
O5.3
92\\
I
s, N ^ P N L - 9 - 3 0 PNL-11-9Fr \ >v »-»D\
: :
^5.0
4.00 0.1 MeV). Out-of-pile annealing of thermal control specimens was also performed. They found a wide distribution of length change (A///) measurements about the
mean, but mean length changes were identical among various nominally similar steels. They also found that there was a slight temperature dependence of the density change (AQ/Q0) in both the thermal and irradiation environments. As shown in Fig. 6-20, both the thermal control and irradiated specimens at ^1000h developed density changes of the same magnitude. It also appeared that the densification was anisotropic, leading to a length decrease somewhat larger than that predicted by VIAQ/QQ (AI/10X1.5AQ/3Q0). This anisotropy is thought to be a consequence of carbide precipitation on a dislocation network known to have a pronounced texture. Finally, the den-
442
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
A. Ap/p 0.24 ~ # Ap/p
Thermal Control, 6000 hrs
•
Thermal Control, 1000 hrs
3A///
I 3A$^ L
Thermal Control, 1000 hrs Irradiated, 1150 hrs in EBR-II
•
0.16
Ap/p 0
/ I I
% 0.08
0.00
- ^ ^ As Received i
i
i
i
i
400
500
600
700
800
Figure 6-20. Thermal shrinkage of both irradiated and unirradiated 20% cold worked 316 stainless steel (after Straalsund et al., 1972).
Temperature, °C
sity change observed at 6000 h was not much larger than that at 1000 h, implying a saturation of the densification process. Bates (1975) demonstrated that minor carbon variations have the greatest impact on net densification of annealed 316 stainless steel, with other elements exhibiting a smaller impact. Based on these various data sets, Garner and co-workers (1978) developed a correlation for densificationinduced strains in 20% cold worked 316 stainless steel and tested it with good results on dimensional change measurements made on tubes irradiated to relatively low neutron fluences in EBR-II. The correlation reflected the major features of the various experimental observations, namely a dependence on cold work level, carbon content and irradiation temperature, as well as an anisotropy in the distribution of strain. Bates and co-workers (1981 b) later performed a similar but more detailed study at much higher fluences on both annealed and cold worked 316 stainless steel cladding. They demonstrated that anisotropic growth indeed occurs both in the presence and absence of stress, and
that it is more pronounced for cold worked material. They also demonstrated conclusively that the growth arises as a consequence of the pronounced texture of the cladding. As mentioned earlier, however, the carbide densification sequence precedes a more sluggish evolution of intermetallic phases, namely a, % and Laves. The formation of these phases produces a net dilation of the austenitic lattice, causing a volume change of 2 - 3 % . Puigh et al. (1984) showed that not only were these strains distributed very anisotropically in cold worked AISI 316 but that their magnitude was such as to give the false appearance of the onset of tertiary creep during thermal aging of pressurized tubes. When the precipitation process was completed, however, the strain rate fell back toward that typical of the second stage of thermal creep. The formation of the intermetallic phases is known to be sensitive to stress, minor impurity elements, cold work and radiation, as outlined by Puigh et al. (1984). The composition and rate of formation of these phases can also be altered by radiation, as discussed earlier.
6.6 Dimensional Stability of Irradiated Steels
The sensitivity of intermetallic formation to minor differences in solute concentrations and environmental conditions can lead to significant and unexpected differences in component behavior. While investigating the origins of failures observed in fuel pin cladding, Hales (1978) showed that one heat of 20% cold worked 316 steel designated FFTF Core 1 developed extensive amounts of intermetallic phases around the entire circumference of fuel pins. Nominally identical companion pins constructed of a steel designated T lot formed significantly less intermetallic phases and the amount formed was larger on the inside of the cladding and on the side of the pin facing core center, consistent with the higher temperatures of those areas. Out-of-reactor aging studies conducted by Hales on these steels showed that the differences in the onset of a formation (Fig. 6-21) arose only from small 15'
19
az HI
151 in 141 EC
UJ
a.
§12
Figure 6-21. Second phase precipitation kinetics observed in two nominally identical AISI 316 steels during thermal aging (after Hales, 1978).
443
variations in the levels of "tramp" elements. Karnesky (1980) later described the transformation kinetics of a formation in typical 20% cold worked 316 fuel pin cladding, demonstrating that grain boundaries participate as short circuit paths for chromium diffusion, leading to a phase formation at grain boundary triple points. These precipitates were identified by Hales (1978) as crack initiation sites for cladding failures. The total amount of a and other intermetallics eventually formed in nominally similar heats is not strongly varying, however. It is considered significant that the relative onset of precipitation appears to correlate with the relative onset of swelling. Garner (1981b) has shown that the application of stress to these two heats of steel also accelerates the formation of intermetallic phases. As will be shown later, stress is another variable that affects the onset of swelling, due at least in part to its effect on precipitate evolution. The phase-related strains associated with the formation of the nickel silicide phases are harder to observe since they usually occur concurrent with the onset of void swelling. The presence of such strains can be inferred, however, from comparison of dimensional change data with density change data, as will be demonstrated in a later section. While the various strains discussed in this section may all appear to be small, the integrated strains over the length of a fuel pin or duct can be quite substantial. This can lead to unintended loads on some long components early in their lifetime. It can also complicate the analysis of creep and swelling data derived from dimensional change measurements, as will be demonstrated several times later in this chapter. One type of precipitate-related strain that will not be covered in this chapter is
444
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
that arising from irradiation of a precipitation-strengthened alloy in an annealed condition, with most solutes still in solution, y'-strengthened alloys such as PE16 or Inconel 706 can densify several percent as a result of precipitation occurring in reactor (Garner and Gelles, 1990). This phenomenon has been observed in experimental studies but is not relevant to actual reactor components since precipitation is usually accomplished prior to fabrication and use. Even then, however, commercial steels with high nickel levels can suffer appreciable reductions in lattice parameter, arising from thermally-induced order-disorder transformations (Marucco and Nath, 1983). 6.6.2 Strains Arising from Void Swelling: An Overview
The discovery by Cawthorne and Fulton (1967) of void swelling due to extensive cavity formation in 316 fuel pin cladding had a very large impact on the LMR community. The void volume was determined to be much too large to be a direct consequence of helium bubble accumulation only; thus the cavities were designated "voids". It quickly became obvious that if swelling did not exhibit saturation at some relatively low level, it would become a major factor limiting the lifetime of LMR components. An early hope that swelling was inherently self-limiting was slowly and progressively disappointed as larger instantaneous swelling rates and larger swelling levels were reached in various neutron or charged particle irradiation studies conducted to increasingly higher dpa levels (see Fig. 6-22). Early observations that swelling indeed saturated in some simple f.c.c. metals (see review by Garner, 1993) helped to maintain for some time the hope that the high swelling rates
observed in some austenitic alloys could eventually be tamed. Whereas the early void swelling levels were masked somewhat by densificationinduced strains, and could only be observed as relatively small distortions of pin diameter, swelling became much more obvious at higher voidage levels, dominating the microstructure (Fig. 6-23), with the macroscopic consequences becoming easily visible to the unaided eye (Fig. 6-24). The currently reported record for maximum swelling in AISI 316 during neutron irradiation is ^ 8 8 % at 510 °C in EBR-II (Garner and Gelles, 1990). The second highest level was also observed in EBR-II, where Hastelloy X reached 80% at 540 °C (Gelles, 1984). The largest reported swelling of AISI 316 (or any metal) is ^260%, reached during 140 keV proton irradiation at 625 °C (Kumar and Garner, 1983). The latter experiment also provides the only clear example of saturation of charged particle-induced swelling that did not arise from some artifact of the simulation procedure. Of course, swelling magnitudes of this level are only of scientific interest, being impossible to accommodate into an engineering design. Due to the importance of the void swelling phenomenon to the successful operation of LMR and later for fusion reactors, it became without a doubt the most intensively studied of all radiation damage phenomena, with several thousand scientific and engineering papers written on the subject. Many of the early studies involved the use of charged particle simulation techniques operating at displacement rates up to four orders of magnitude greater than that of the simulated reactor environment. Simulation studies were also used to study solute segregation when it became obvious that this process was either directly involved or at least coincident with void swelling.
6.6 Dimensional Stability of Irradiated Steels 6
20% Cold-Worked (U.S.A.)
5
o (0 DC O>
Annealed
/~1%/dpa
/ Current Model \ 1 \\ 1 1
/ 1
1 |
4 ~ 1 1 1 1 3
i
/
A
77
\
1 1 1 1 1 1 1
1L
/ / /
1
Current Model \ •-v SA316 / \ 1977, U.S.A. I 1 /
1
~W\A
2
^pr 400
U
1
1
500
600
V \ \
A
\^
445
\l 1/ L
rx
\\ U \
I1
1/
r
i
1
700 800 400 500 Temperature, °C
SAM316 850°C Age 1977, U.K.
1
600
XI
700
1
Figure 6-22. Chronological evolution of swelling predictions for AISI 316 in the U.S. LMR materials program, reflecting the tendency of predictions to increase as data became available at progressively higher swelling levels (Garner, previously unpublished). The swelling rate is in units of %/10 22 n cm" 2 (E > 0.1 MeV).
800
Figure 6-23. Temperature dependence of void swelling observed in FFTF first coreheatCN-13of20% cold worked AISI 316 at ^1.4xl023ncm~2 (£>0.1 MeV) or ^70dpa (courtesy of W. J. S. Yang of Westinghouse Hanford Company).
446
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
2©% CW 31ti
UNIRRAC CQNTRGI
FLUENCES BEYGWD FFTF GOAL
Figure 6-24. Easily observed swelling (^10% linear, « 3 3 % volumetric) in unfueled 20% cold worked AISI 316 cladding tube at 1.5 x 1023 n cm" 2 (E > 0.1 MeV) or «75 dpa at 510°C in EBR-II (after Straalsund et al, 1982). Note that, in the absence of physical restraints, all relative proportions are preserved during swelling.
Although the use of charged particle simulation studies was very useful in identifying many of the variables participating in microstructural evolution, it also led to or supported a number of early misperceptions concerning the parametric sensitivity of swelling. Some of these misperceptions were quickly discarded, such as the observed tendency for swelling to saturate during electron irradiation (Laidler and Mastel, 1972). When bulk-representative thickness criteria were established (Garner
and Thomas, 1973) and void-surface interactions were clarified (Laidler et al., 1976) swelling thereafter did not saturate, although most electron irradiation experiments were terminated before swelling levels exceeded 20-30%. The major problem with charged particle simulation experiments was that the accelerated displacement rates were unavoidably purchased at the expense of some very important trade-offs in other parameters later found to have a strong influence
6.6 Dimensional Stability of Irradiated Steels
on swelling. These trade-offs involved parameters that were designated "atypical variables" (Garner and Laidler, 1976; Garner et al., 1977) and the influence of these parameters became a separate field of study in itself. The most persistent of the early misperceptions are listed below: (1) Swelling could be defined in terms of microstructural interactions with point defects using quasi-steady state "rate theory" approaches that described the flow and absorption of point defects to the various microstructural sinks. These theories were largely physical in nature rather than chemical (Harkness and Li, 1971; Wiedersich, 1972; Brailsford and Bullough, 1972; Mansur, 1978; Mansur and Yoo, 1979). Differences later attributed to microchemical and nucleation-dominated processes were thought at that time to be an expression of a strongly variable steady-state swelling rate. (2) The steady state swelling rate was thought to exhibit a behavior that was strongly dependent on the initial dislocation density (cold work level), temperature and displacement rate in particular, but also on stress and other variables. (3) The swelling rate could be manipulated by compositional changes or thermomechanical processing to lower the maximum swelling rate and alter its temperature dependence. (4) The most important determinant of void nucleation and growth was the accumulation of helium or the segregation of residual gases such as oxygen and nitrogen. Each of these perceptions later proved to be largely or at least partially incorrect for stainless steels. A more correct perception arose as the result of many reactor studies conducted under better characterized irradiation and material conditions,
447
but did not involve fuel pins, which are subject to the time-dependent action of many variables. The new perceptions, first reviewed by Garner in 1984, presented a very different picture. The majority of the parametric sensitivities of swelling were found to express themselves only in the duration of the transient regime before and during the void nucleation phase. The steady state swelling regime proved to be remarkably insensitive to most of these variables, with the swelling rate of most Fe-Cr-Ni (Garner, 1984; Garner and Gelles, 1990) and Fe-Cr-Mn (Garner et al., 1987 a; Garner and McCarthy, 1990) austenitic alloys being « 1 % per dpa over a wide range of irradiation temperature. Although the transient regime of austenitic alloys was found to be very susceptible to metallurgical manipulation, eventually the very high swelling rate of 1 %/dpa would assert itself over most of the temperature range of LMR interest. While some significant progress has been made in understanding the radiationinduced swelling behavior of simple Fe-CrNi model austenitic alloys (Garner and Wolfer, 1984; Garner and Brager, 1985 a; Garner and Kumar, 1987; Coghlan and Garner, 1987; Muroga et al., 1991, 1992; Hoyt and Garner, 1991 a, b), there are still a number of areas in which the response of more complex structural alloys is not completely understood. For Fe-Cr-Ni ternary alloys it now appears to be relatively easy to understand the compositional and temperature sensitivity of their relatively short swelling incubation period, as well as its relatively abrupt termination, and the insensitivity to many variables of the steady state swelling rate of « 1 %/dpa. In more complex alloys, however, the transition between the incubation and steady state swelling regimes can be quite protracted. Whereas many steels quickly reach the
448
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
1 %/dpa swelling rate, others approach it very slowly with 1 %/dpa reached only at very high swelling levels. Simple rate-theory models based on void and dislocation microstructure alone are not capable of describing such prolonged transition behavior in typical structural steels. An excellent example of this problem was shown by Brown etal. (1983), who measured the swelling of cold worked M316 fuel cladding as a function of temperature and displacement damage. As shown in Fig. 6-25, this alloy did not reach the 1 %/dpa swelling rate during this experiment and displayed a protracted transient regime that was very sensitive to irradiation temperature and neutron flux variations that exist along the cladding. Brown and co-workers also measured the microstructural densities at several damage levels (see Fig. 6-15 for examples of these data). They found that the dislocation density saturated relatively early in the irradiation at a temperature-dependent level that did not change significantly with further irradiation. Only the void characteristics changed significantly between the different dpa levels. When the measured swelling rates were matched against those
10
20
30
40
50
60
70
predicted by Cawthorne (1979) using rate theory and the measured sink strengths (Fig. 6-26), it was apparent that there were significant differences in both magnitude and temperature dependence. In particular, Fig. 6-26 does not show the proportionality between swelling rate and the sink strength term (av ad)/(av + a d ) 2 required by simple rate theory. The dislocation sink strength is the total dislocation line length, and the void sink strength is 4 7i r C, where r is the mean void radius and C the void concentration. It is considered to be particularly significant that the microstructurally-based prediction exhibits the relative temperature independence of swelling rate that has been observed in experimental studies, but that the actual swelling rates do not yet exhibit this independence. This implies that these is another transient process superimposed on that involving the microstructural sinks alone. This process, unidentified at that time, apparently has its own temperature dependence and contributes strongly to controlling the instantaneous rate of void swelling. It is obvious, therefore, that some other nonmicrostructural variable must be in-
400
500 TEMPERATURE, °C
600
Figure6-25. Swelling behavior observed in 20% cold worked M316 in the DFR reactor (after Brown etal., 1983).
6.6 Dimensional Stability of Irradiated Steels
449
Figure 6-26. Comparison of measured swelling rates (solid lines) with those predicted by Cawthorne (dotted lines) based on rate theory and measured microstructural parameters. The swelling rate continues to increase in a temperaturedependent manner long after the microstructurally-based prediction loses its temperature dependence. The parameters av and ad are the sink strengths for voids and dislocations, respectively. The endof-life irradiation temperatures are shown above the figure (courtesy of C. Cawthorne). Distance Along Pin
volved. As discussed earlier, one of the major candidate mechanisms invoked to explain the inability of simple microstructurally-based models to fully explain the kinetics of void swelling has been the sluggish compositional evolution of the matrix, as various elements are removed from the matrix via precipitation or segregation (Garner, 1981a, 1984; Brager and Garner 1978, 1979; Porter, 1984). Such models concentrate on segregation and especially precipitate-related changes in matrix concentration of elements known to be very active in suppressing swelling, e.g., nickel, silicon and phosphorus. There is another element, however, whose role in swelling is usually thought not to be as strong as these other elements, but which is known to control much of the behavior of such steels in nonradiation environments. This element is carbon (Bates et al., 1981a), and its chemical activity in thermal environments is known to be influenced not only by its concentration but also by its distribution and interaction with other solutes such as phosphorus (Banerjee et al., 1968; Kegg et al., 1974; Hosoi and Wade, 1984). The distribution of carbon and other minor elements may also be influenced somewhat by both the mode of deformation and the resultant dislocation density,
as well as the temperature history prior to irradiation. Although cold working to produce high densities of dislocations is known to decrease swelling in solute-modified steels, not much is known about its interactions during irradiation with elements such as carbon. As will be demonstrated later, both the carbon concentration and the cold work level are known to strongly influence the rate of microchemical evolution in both radiation and nonradiation environments. As will be shown in a later section, the major features of the parametric dependencies of swelling result from the action of carbon on thermally-induced precipitation, the radiation-related factors which influence carbon's role, and the competition between thermally-induced phases and radiation-influenced phases. Figure 627 illustrates the range of swelling variability that a fuel pin constructed from any one heat of 316 stainless steel can exhibit in response to variations in thermomechanical starting state alone. 6.6.3 Variables Which Influence the Swelling of Steels
The onset of swelling in austenitic or ferritic steels based on iron, nickel and
450
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors Double-aged to precipitate all carbon from solution \
AD D'
Distance Along Fuel Pit
Figure 6-27. Schematic illustration of swelling behavior observed in EBR-II for one heat of AISI 316 stainless steel irradiated in various starting conditions (after Garner, 1981, adapted from Hofman et al, 1977).
chromium has been found to be very sensitive to an incredibly large number of variables, some of which are subtle enough in their action to have escaped detection initially. Thus, some of these variables were often uncontrolled in early experiments. Because of this complexity, it is often difficult to identify single variable experiments to separate and study the action of any one variable on swelling. Wherever possible, however, the data chosen for presentation in this section reflect the action of one or sometimes two variables. 6.6.3.1 Crystal Structure As mentioned earlier, all austenitic alloys based on the Fe-Cr-Ni and Fe-Cr-Mn systems exhibit a maximum steady state swelling rate of ^1%/dpa, even though some alloys such as Nimonic PE16 and Inconel 706 appear to resist entering the steady state regime to very high dpa levels (Yang and Makenas, 1985; LeNaour et al., 1987; Brown and Linekar, 1987; Garner and Gelles 1990). Pure nickel also exhibits this same initial swelling rate, but eventually saturates in swelling due to a collapse in dislocation density arising from its high
stacking fault energy (Stubbins and Garner, 1992; Garner, 1993). It initially appeared that 1%/dpa might be an f.c.c. "crystal constant", but pure copper was found later to swell at £s 0.5 %/dpa without saturation to at least 150 dpa, indicating that crystal structure alone does not define the steady state swelling rate (Garner, 1992b). Ferritic and ferritic/martensitic steels do not appear to swell at such large rates, however (Powell et al., 1981; Vitek and Klueh, 1983; Gelles and Thomas, 1983; Wassiliew et al., 1983; Bagley et al., 1987; Little, 1987). Even simple model Fe-Cr alloys swell at rates of 0.1 MeV), showing « 1 3 % and 0% swelling respectively. (Courtesy D. S. Gelles, of Pacific Northwest Laboratory.)
1991; Garner etal., 1992a; Nakajima et al., 1992; Matsui et al., 1992) while pure vanadium and other vanadium binaries do not swell at such high rates. This suggests that the bias must also depend on microchemical properties as well. Some crystal structures are more anisotropic in their properties (e.g., h.c.p. structures) and exhibit considerable anisotropy in their strains (see Chap. 7 on zirconium alloys). Given that anisotropies have been observed in austenitic stainless steels as a result of texture, dislocation evolution under stress, and precipitation-related
Figure 6-29. Visual comparison of postirradiation diameters of small rods of 20% cold worked 316 and HT9 irradiated at various temperatures to high exposures in EBR-II (after Garner and Gelles, 1990).
HT-9
316
425°C
2.0 x 10 23 n/cm 2
510°C
2.4 x 10 23 n/cm 2
540°C
2.3 x 10 23 n/cm 2
590°C
2.5 x 10 23 n/cm 2
650°C
2.5 x 10?3 n/cm 2
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
452
25.4
1( 0.9 0.8
-
2 0.7
0 6
I" 5
V
0.4 -
I °-3 -
1
?
CW316 \ _
/
/
- 6.3
- "**
• I
8
HT-9
•
\
•
u
I
!
I
i
i
i
i
i
i
i
10
12
14
16
18
20
22
24
26
28
!
-0.1
O
>
0.1 I
Q) CO
y
0.2
I
-- 12.7
CWD9
/
0
19.1
30
Peak Neutron Fluence (10 2 2 n/cm 2 , E>0.1 MeV)
Figure 6-30. Changes in length of various FFTF ducts subjected to the same irradiation conditions but varying only in the type of steel (after Hecht and Trenchard, 1990). D9 is a titanium-modified variant of 316 stainless steel.
cannot be ignored as they occur under a variety of conditions, especially in materials whose stress state was changed during irradiation (Garner and Gelles, 1988). An example is shown in Fig. 6-32. Figure 6-31 also demonstrates that the stress-enhanced portion of swelling is
strains, it is prudent to ascertain whether the distribution of swelling strains is truly isotropic. Both Fig. 6-18, presented earlier, and Fig. 6-31 show that swelling appears to be isotropic when concurrent phase-related density changes or growth processes are taken into account. Growth processes
To-525 °C
To-585 °C
12
15, 9.7, -161 °C -—£m/ * / 10
AV/Vo,
~
.. fully
/ / 60,9.7, -134 °C // /
8
1^7,9.7, OX
t t
%
6
t
m§\- isothermal 2 ~~ /JT J at ~ 475 °C t
0 //
/ /
/
I
I
I
I
/
/ jmSn, 10.0,-139 °C / ' /" 29, 9.9, -91 °C / A 0,10.0, -79 °C
/
/% 89.5, 8.7, -66 °C /A 0,9.0,-64 °C
4 _
/
I 5
/
/
/\
0 1 AUL, %
(MPa, 0t/1O22, Tf -To) I
I
I
I 6
Figure 6-31. Isotropic swelling observed in pressurized tubes of 20% cold worked 316 stainless steel irradiated in an experiment where the temperature fell slowly during irradiation. The offset from the origin is a consequence of strains related to y' formation, which is promoted by declining temperature (previously unpublished data of F. A. Garner and E. R. Gilbert of Pacific Northwest Laboratory).
6.6 Dimensional Stability of Irradiated Steels
453
isotropically distributed. With one exception, there is no previously published evidence to support the contention that this strain contribution was isotropic in nature. The single exception arose when Harbottle and Silvent (1979) demonstrated that lateterm stress effects on swelling in pure nickel irradiated in SILOE involved an isotropic partition of both the stress-free and stress-enhanced portions of swelling strain.
# Diamete O Length
Percent Change
6.6.3.2 Base Composition (AV/Vi + 1)1/3-1
Figure 6-32. Comparison of incremental changes in dimension with changes predicted from density change measurements for previously irradiated 304 L stainless steel fuel pin segments subjected to «10 dpa additional irradiation without stress at 560 °C (after Flinn and Kenfield, 1976; Brager et al., 1977). Diameter changes are consistently larger than length changes, reflecting either continued growth due to the stress state of the original fuel pin or possibly lateterm precipitation.
120
As shown in Figs. 6-33 and 6-34, it was charged particle simulation experiments by Johnston et al. (1974) that first demonstrated very clear trends of swelling of solute-free Fe-Cr-Ni ternary alloys with base composition and ion bombardment temperature. At a given temperature and displacement level, swelling first decreased strongly with increasing nickel content,
Fe-Cr-Ni ALLOYS 675 °C 140 dpa
100
15
20
30 35
85
100
WT.
Figure 6-33. Swelling measured by the step-height technique on annealed Fe-Cr-Ni ternary alloys after irradiation with 5 MeV Ni + ions to 140 dpa at 675 °C (after Johnston et al., 1974).
454
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
120
Fe - 30Cr - YNi
Fe-15Cr-XNi 15 Ni
80
35 Ni
575
625
675
725
775
575
625
675
725
775
Temperature, °C
Figure 6-34. Temperature and nickel dependence of swelling of Fe-Cr-Ni alloys at 140 dpa for two chromium levels after 5 MeV Ni + ion irradiation (after Johnston et al., 1974). Previously unpublished data at 775 °C supplied by T. Lauritzen of General Electric Company.
then reached a minimum value at some intermediate nickel content. Thereafter, swelling increased relatively slowly at higher nickel levels. Chromium's influence 60
50
o COMMERICAL ALLOYS (Cr: 18 ± 4%) 304 A BASE ALLOYS Fe - 15Cr - XNi
40
SWELLING %
30
20
IN706 PE 16 IN702 HAST-X IN600/ / IN625 \ /
10
20
40
60
80
NICKEL. w t %
Figure 6-35. Comparison of ion-induced swelling at 625°C for commercial alloys (14-22% Cr) and Fe15Cr-Ni ternary alloys (after Johnston et al., 1976).
was found to be monotonic, with increasing chromium causing higher swelling at all nickel levels. Ion-induced swelling was shown to peak strongly with temperature at all nickel and chromium levels. It was also shown by Johnston et al. (1976) that commercial alloys swelled less than comparable pure ternary alloys, but also exhibited the same strong dependence on nickel content, as shown in Fig. 6-35. These trends were observed in many different types of charged particle studies, some examples of which are shown in Figs. 6-36 through 6-38. Despite the uniformity of the simulation data, they were very misleading in several important respects. In particular, they strengthened the impression that the steady state swelling rate was a strong function of both nickel content and irradiation temperature, and therefore was susceptible to metallurgical alteration. As the first neutron data on the nickel dependence of swelling became available at rela-
6.6 Dimensional Stability of Irradiated Steels 25 18Cr ± 4%
20 -
200 KeV C+
Jp304
41 dpa, 625 °C 15
-
Swelling, %
10 --
-
6316
O 321
\ \
Ni
Inc 800
I 20
•
—
-
—
_
Inc 600 4_ _r^y
JO
40 Nickel, wt%
Figure 6-36. Dependence of swelling of commercial alloys (18 + 4% Cr) on nickel content at 625 °C after 200 keV C + irradiation to 41 dpa (after Terasawa, 1985). ^ ^ 20 ~¥*^^ / 16
^ \ \ \ 12 \ Swelling 8 -
\
4
o Ar" + A Ni
-\
08X20H35 05X5H35 05X15H35M3 05X15H36M2>
7
/
\
A
08X20H45 y / 03X20H45M454 / / jT*~
Figure 6-37. Dependence of swelling of Soviet commercial alloys on nickel content at 550 °C after Ar + or Ni + ion irradiation to 50 dpa (after Parshin, 1980).
08X20H80
^ / / _ _ ^ 08X20H55 ^ ^ ^ ^ I
0
20
40
Ions
Pure Nickel 80 dpa
/1'
1
lively low exposure levels, this impression was not dispelled, as can be seen in Figs. 6-39 through 6-41. As swelling data became available on high nickel alloys irradiated at higher fluence levels, it became obvious that the ion bombardment data were misleading, however. Nikolaev et al. (1985) and Vasina et al. (1985) both noted that the swelling of alloys based on Fe-37.3Ni-14.6Cr exhibited a high degree of swelling that was "incompatible" with Johnston's conclusion that steels near 35% Ni have a tendency toward low swelling, as shown in Figs. 6- 42 and 6-43.
50 dpa, 550°C
08X8H9 09X16H11M3
r^
455
60
I
80
100
Nickel, wt % 1.0
i
1
8
1
600°C 0.8-
-
Swelling
0.6 SWELLING RATE %/dpa e 0.4
\ SA316
\
»OFV548
7
DFR Reactor
6 -
30 dpa, 600 C
\
5 ~ 347\
Fe-22Cr-XNi + SOLUTES 3 -
O316L
O \ oM316 321
2 0.2-
INC825S INC 706 "0
20
INC 625 /
40 60 % NICKEL
80
PE16(3) ° A B * * 6
1 100
Figure 6-38. Effect of nickel on swelling of various alloys irradiated at 600 °C with 1 MeV electrons (after Levy et al., 1977).
n
I
10
I
lfE16(1)O
|
20 30 40 Nickel Equivalent of Matrix, wt.%
Figure 6-39. Dependence of swelling on nickel equivalent (Ni + Co + 0.5 Mn + 30C + 0.3 Cu + 25N) at 30dpa after irradiation in DFR at 600 °C (after Watkin, 1976).
456
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors 510 °C
Figure 6-40. Swelling of Fe-Cr-Ni ternary alloys in EBR-II at 13 dpa and 510 °C (after Bates and Johnston, 1977). 15 20 25 30 35
75
45 60 WT. % Ni —
85
Garner (1984) was the first to show convincingly that the strong dependence of neutron-induced swelling on alloy composition, starting condition and irradiation environment lay almost exclusively in the duration of the transient regime that precedes the onset of steady state swelling, which proceeds at ^1%/dpa thereafter.
400
4
-
2
-
20
30
40
50
At relatively low irradiation temperatures, the composition dependence of swelling in Fe-Cr-Ni ternaries disappeared altogether, as shown in Fig. 6-44. Above some temperature-dependent compositional threshold, the temperature and composition dependence of swelling were expressed in only the duration of the transient regime (Gar-
60
NICKEL (WEIGHT PERCENT) Figure 6-41. Relative swelling behavior of eight annealed austenitic alloys over a limited range of temperature (400-650°C) and neutron exposures corresponding to 16 to 27 dpa (after Bates and Powell, 1981).
457
6.6 Dimensional Stability of Irradiated Steels
ner and Brager, 1985 b, c) as shown in Figs. 6-45 and 6-46. Although the strong nickel dependence of swelling observed by Johnston et al. was also observed at higher neutron irradiation temperatures by Garner (Fig. 6-47), it was found not to be representative of neutron irradiation at lower and middle swelling regime temperatures. The strong dependence of swelling on nickel content observed in the charged particle simulation studies was misleading for several reasons. First, the ion irradiations must be conducted at higher temperatures to compensate for the upward shift in the temperature regime of swelling that occurs at the higher displacement rates used in charged particle irradiations (Westmoreland et al., 1975; Garner and Guthrie, 1975). Second, the limited range of most charged particles leads to a strong temperature-dependent influence of the specimen surfaces. The surface influence increasingly depresses the swelling rate at higher temperatures (Garner and Laidler, 1976; Bullough and
10
24 HOURS AT 750°C
BASE ALLOY SWELLING %
700 IRRADIATION TEMPERATURE, °C
Figure 6-43. Swelling of annealed Fe-37.3 Ni-14.6 Cr1.2Mn-0.28Si-0.05C in the BOR-60 reactor at « 52 dpa. Solute additions reduce swelling strongly, but aging of the solute-modified alloy at 750 °C leads to even larger swelling, reflecting thermal formation of nickel-rich phases which remove nickel and silicon from the matrix (after Nikolaev et al., 1985).
BOR-60 REACTOR
6 SWELLING 4
NEUTRON FLUENCE GIVEN IN UNITS OF 10 22 ncm- 2 tE>0.1 MeV)
400
10,000 HOURS AT 750°C
500
600
TEMPERATURE. °C
Figure 6-42. Swelling of annealed Fe-37.3 Ni-14.6 Cr2.7Mo-1.8Mn-0.28Si-0.05C-0.1P in the BOR-60 reactor at 53-59 dpa (after Vasina et al., 1985).
Haynes, 1977). Third, the temperature dependence of ion-induced swelling is also strongly depressed at both low and high irradiation temperatures by the injected interstitial effect (Mansur and Yoo, 1979; Garner, 1983; Plumton and Wolfer, 1984; Kumar and Garner, 1985). Fourth, it appears that the bias of dislocations toward interstitial absorption is a strong function of displacement rate (Rauh and Bullough, 1985; Rauh et al., 1992; Tenbrink et al., 1988). All of these factors strongly distort the temperature dependence of the swelling rate relative to that observed in neutron irradiations. Another consequence of these factors is that the maximum ion-induced
458
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors i
i
i
i
i
i
i o
i
i
i
i
r
26
22
18
SWELLING %
14
10
• 12.1% • 15.7% o 19.4% n 24.4%
1 2
I I I I I 3 4 5 6 7
I I I I I 8 9 10 11 x 10 22 n c m 2
Figure 6-44. Swelling of Fe-15 CrNi ternary alloys in EBR-II at temperatures between 400 and 510 °C for nickel levels between 12.1 and 24.4 wt.% (after Garner, 1984).
NEUTRON FLUENCE (E>0.1 MeV)
T i l l 593°C
Ti
650°C
J 10
20
30 0 10 20 30 0 " 10" NEUTRON FLUENCE (n cm-2, E>0.1 MeV)
20
L 30X10 2 2
Figure 6-45. The influence of temperature and nickel content on swelling of ternary Fe-15 Cr-Ni alloys in EBR-II (after Garner, 1984).
6.6 Dimensional Stability of Irradiated Steels
459
60
Fe-Ni-Cr TERNARY ALLOYS PRIOR TO BREAK-AWAY
Fe-35.2Ni-20.0Cr
-I
zFe-34.5Ni-15.1Cr
Fe-35.5Ni-7.5Cr -*-••
i
i
i
8 12 16 20 0 4 8 12 16 NEUTRON FLUENCE (E>0.1 MeV) n cm-2
i
i
i
20x1022
Figure 6-46. The influence of temperature and chromium level on the swelling of selected Fe-Cr-Ni ternary alloys in EBR-II (after Garner, 1984). 70
Fe-15Cr-XNi \
20.4 x 10 22 50 -
r 40 SWELLING % 30 20
V
\ -^
22.0x1022 9 _
^ - 1 ^ 2 1 . 9 xiO 2 2 10
V 0 10
i
20
30
40
50
-—-i— 60
WEIGHT PERCENT NICKEL
*T ^ 70 80
Figure 6-47. Swelling of Fe-15 Cr-Ni alloys at relatively high temperatures and neutron fluences (E > 0.1 MeV) in EBR-II (after Garner and Brager, 1985 b,c).
460
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
425°C
AISI310/
540°C
INC 800 40 Fe-Ni-Cr Ternary Alloys at Nickel Levels O.IMeV)
Figure 6-49. Swelling of annealed RA 330 in EBR-II, demonstrating that, at high nickel levels, the duration of the transient regime is quite protracted and very dependent on temperature (after Garner and Gelles, 1990).
A wide range of minor and not so minor solutes are found in the various steels employed in LMRs. The role of these elements on void swelling has been investigated in many studies using both charged particle and neutron irradiation. Examples of both types of irradiation studies were presented by Bates and Johnston (1977). The solutes can be ranked either in terms of the usual concentration ranges encountered or by their impact on swelling. There appears to be no direct correlation between the results obtained using the two types of rankings, however.
461
6.6 Dimensional Stability of Irradiated Steels
In the early stages of the U.S. LMR program, it was suspected that "trace" or "tramp" elements at very low and usually unspecified levels might exert a large influence on swelling. In the only reported neutron irradiation study to specifically address this issue, the swelling behavior in EBR-II of 316 alloys containing small and typical amounts (0.01-0.08 wt%) of Al, As, Co, Cu, N, Nb, Sn, Ta, V, W, B and P were compared with the response of high purity alloys (Garner and Brager, 1988). Of these elements it was shown that only phosphorus had a strong and reproducible influence on swelling (Fig. 6-50). Boron was not found to exert a perceptible or consistent influence on swelling, even though it is one of the major early sources of helium. Boron, nitrogen and sulphur were studied in another experiment on AISI 316 and also were found to show no significant effects (Bates, 1975).
The trace element study was a small part of a much larger experiment designated MV-III, in which 150 compositional variations of AISI 316 in various thermomechanical starting conditions were irradiated side-by-side (Garner and Brager, 1988). The major variations were in Ti, Si, Zr, Cr, Mo, C, and P, with each element having at least two and as many as five concentration levels. The major conclusion of the study was that each of the elements studied exerted their influence only on the duration of the transient regime. No combination of elements or starting conditions studied suppressed the tendency of this class of steels to eventually swell at ^1%/dpa. No other published neutron studies disagree with this conclusion whenever there are sufficient data as a function of neutron fluence to establish the steady state swelling rates.
0.01 Zr 0.10 Zr HIGH PURITY TRACE IMPURITIES
• •
30A
A O
ANNEALED A I S I 316 IN EBR-II
20% COLD-WORKED AISI 316 SWELLING %
HIGH PURITY SWELLING
o.oo
0.02
0.04
0.06
PHOSPHORUS (wt%)
0.08
0
Q.02
0.04
0.06
0.06
PHOSPHORUS (wt %)
Figure 6-50. The relative influence of phosphorus, zirconium and trace impurities on the swelling of AISI 316: (left) 20% cold worked condition at 425 °C and 55 dpa, (right) annealed condition at 540 °C and 51 or 76 dpa (after Garner and Brager, 1988).
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
462
Some minor elements in the MV-III experiment, such as zirconium (Fig. 6-50), were found to exert very little influence on swelling, although some synergisms were observed with titanium and carbon. Chromium additions were observed to increase swelling in a manner consistent with the results of Fe-Cr-Ni ternary studies discussed in the previous section. Molybdenum additions could either increase or decrease swelling, depending on the base alloy composition and especially on the irradiation temperature (Garner and Brager, 1988; Garner e t a l , 1993b; Herschbach et al, 1993), as shown in Fig. 6-51. Some elements such as P, Si, and Mo were initially thought to be monotonic with concentration in their influence on swelling, but it was later shown that relatively small amounts of each element actually increased swelling initially, with larger additions leading to a peak swelling and a decrease thereafter (Figs. 6-52 through 6-54). Some elements such as manganese exhibit the opposite behavior, first decreasing and then increasing the swelling of AISI 316 (Bates et al., 1981 c).
20
425°C p
15
£ !
0.04 0.04 Mn 2.0 Si 0.75
c
It is important to note, however, that some elements (Ti, C, Mo) may strongly suppress swelling in one temperature range but increase it in another. Titanium and carbon are particularly important in this respect. Some examples of the action of titanium at different temperatures are shown in Figs. 6-55 and 6-56. As indicated earlier, carbon plays a major role in both thermally activated and radiation-influenced precipitation. Figure 6-57 shows a typical example of carbon's influence in decreasing swelling at low temperatures in AISI 316, while Fig. 6-58 demonstrates its enhancement effect at higher temperatures. Figure 6-59 demonstrates that the net result of increasing carbon in AISI 316 is to shift the peak in swelling to higher temperatures, a process which occurs both in annealed and cold worked steels. As will be shown in a later section, the two regimes of precipitation, one based on carbon and one based on nickel silicides, induce a double peak swelling distribution whenever a sufficiently large enough range of temperature is studied. The two peak
540°C Cr 16.3 Ni 13.8 Ti 0.20
10 22
10
Figure 6-51. Influence of molybdenum on neutroninduced swelling in an annealed titanium-modified 316 stainless steel irradiated in EBR-II to 23 and 56dpaat425°Cand32, 50 and 75 dpa at 540 °C (after Garner etal,, 1993b).
,10.0 x 1 0 2 2
CO
7.6X
1.0
3.0
1.0
2.0
Molybdenum, Weight %
4.0
6.6 Dimensional Stability of Irradiated Steels
A o • V • O
SWELLING %
463
399°C, 8.2 dpa 510°C, 13.2 dpa 454°C, 9.5 dpa 482°C, 11.9 dpa 538°C, 13.6 dpa 593°C, 14.3 dpa 649°C, 14.3 dpa
Swelling, %
454°C, 9.5 dpa 482°C, 11.9 dpa
1.0
0.2
074
0.6
0.8~~Br 1.0
1.2
0
.02 .04 .06 .08 .10 .12 Weight Percent Phosphorus
WEIGHT PERCENT SILICON
Figure 6-52. Influence of (a) silicon and (b) phosphorus on the swelling in EBR-II of annealed Fe-25Ni-15Cr at various combinations of temperature and fluence (after Garner and Kumar, 1987).
distributions for both low carbon steels and titanium-modified moderate-carbon steels appear to exhibit similar behavior, however, as shown schematically in Fig. 6-60. It appears that the role of titanium is directly linked to that of carbon, since titanium carbides form easily at elevated temperatures and remove carbon from solution. In many studies, therefore, an important variable studied was the Ti/C ratio, which determines both the level of free carbon and the degree of TiC precipitation (e.g., see Tateishi, 1989) By using carefully selected and balanced solute additions, along with the appropriate thermomechanical treatment, it is possible to substantially prolong the incubation period of swelling. This has been done
Swelling, %
0.04
0.08 0
0.04
0.08
Phosphorus, Weight %
Figure 6-53. Non-monotonic swelling behavior vs. phosphorus content observed in Fe-16.2Cr-13.8Ni2.5Mo-2.0Mn-1.5Si-0.2Ti-0.04C alloys at 31 and 51 dpa in the annealed condition and 75 dpa in the 20% cold worked condition (after Garner and Mitchell, 1992).
464
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Fe-15Cr-20Ni-XMo
Fe-18Cr-8Ni-1Mo
Fe-18Cr-8Ni 0.5% Mo
Density Change %
10 1U
8 6 4 2
Fe-17Cr -16.7Ni-2.5Mo"
400
I
I
500
600
400
600
500
Temperature °C Fe-15Cr-20Ni-XMo
30 Fe-18Cr-8Ni Fe-18Cr-8Ni-1Mo
20 Density Change
0% Mo
10 Fe-17Cr-16.7Ni-2.
400
500
400
600
500
Figure 6-54. Effect of composition and temperature on neutron-induced swelling of several model Fe-Cr-NiMo alloys irradiated in EBR-II to 14 dpa (top) and FFTF to 31.836.6 dpa (bottom) (after Garner et al, 1993 b).
600
Temperature, °C
AXIAL LOCATION IN INCHES FROM BOTTOM OF FUEL -1.0
0.0
1.0
2.0
3.0
4.0
5.0
6.0
7.0
8.0
9.0 10.0 11.0 12.0 13.0 14.0 15.0 i
i
i
r
316 T-LOT CLADDING
50.8 101.6 152.4 203.2 254.0 304.8 AXIAL LOCATION IN MILLIMETERS FROM BOTTOM OF FUEL
355.6
Figure 6-55. Diameter change measurements of two 20% cold worked fuel pins irradiated side-by-side in EBR-II to a peak exposure of 10.5 x 1022 n cm" 2 (E > 0.1 MeV) or ^52 dpa (after Makenas et al., 1980).
465
6.6 Dimensional Stability of Irradiated Steels I
I I 20%CW316 , (3 CYCLES)
1 620C
I
i
i
i
r
oo
12.2 x 10 22 n/cm 2
20 -
0.047 N
0.007 N
COLD-WORKED AND AGED 20% COLD WORKED -20
-15
15
-10 -5 0 5 10 INCHES ABOVE CORE MIDPLANE
J
20
1
~~
1
1
1
> 1
0.01 % C
1
i
1
i
1
i
1
:
L
—
^ 0.06%C
f
10
I
"
—
1 0
I
^-0.035-0.039% C
J J1
2
J
1
425 C>C
/a
5S WELLING % 4 3
1
ANNEALED
L
Figure 6-58. Dependence of swelling in AISI 316 on carbon content and starting condition, indicating some possible synergism with nitrogen (after Bates et al., 1981 a).
Figure 6-56. EBR-II based swelling profile predictions for two nominally identical ducts to be irradiated in FFTF, illustrating the tendency of the D9C1 steel, a titanium-modified 316 variant, to swell more at lower temperature than unmodified 316 steel.
_
I
.02 .04 .06 .08 .10 .12 .14 .16 .18 .20 WEIGHT PERCENT CARBON
i 1 i 1 1 I i 20 30 40 50 60 70 DISPLACEMENTS PER ATOM
I
80
Figure 6-57. Effect of carbon content on swelling of three U.K. stainless steels at 425 °C in DFR (after Bramman et al., 1977).
466
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors ANNEALED
20% COLD-WORKED
(10.4)
10
° 0.021%) CARBON ° 0.033%[WEIGHT . 0.053%) PERCENT
10
-
\
8 -
8
(9.6)/ /
SWELLING
6
6
Al\< - > \
7 7
4 •
/
/
/
2
\\ \\
/
/ /
\\
/
/
/
/
Is i)l' /
V% \ (3.4) D
(6.1)i; _
I
400
i
450
500
550
400
450
500
550
TEMPERATURE, °C
Figure 6-59. Effect of carbon content, temperature and starting condition on swelling of AISI 316 after irradiation in RAPSODIE (after Dupouy et al., 1978). Neutron fluence is shown in parentheses in units of 1022 n cm" 2 (E > 0.1 MeV).
in all of the major national LMR programs, each by a slightly different route. An example from the Japanese LMR program (Itaki et al., 1987) is shown in Fig. 6-61. A somewhat similar approach employed in the U.S. LMR program was preSOLUTION-ANNEALED AISI 316
N
in
LOW \ ^/CARBON \ OR Ti-MODIFIED \ \ ^ / HIGH
TEMPERATURE
Figure 6-60. Schematic illustration of the relative influence of carbon and titanium on swelling of austenitic steels over the temperature range 350650 °C.
sented by Hamilton et al. (1989). The latter study explored the optimization of creep and mechanical properties as well. The optimization routes employed in the French LMR program are described by Seran et al. (1992) and in the Japanese program by Fujiwara et al. (1987) and Tateishi (1989). On a per atom basis, phosphorus and then silicon are the most effective suppressors of void growth over the entire temperature regime of swelling. The action of phosphorus is demonstrated in Fig. 6-62. In precipitation strengthened alloys which form ordered y' and/or y" precipitates at high densities, elements such as silicon, aluminum and titanium have sometimes been found to play a large role in swelling suppression. In this latter case the surfaces of the ordered precipitates probably serve as sinks for point defects and thereby reduce the supersaturations necessary to nucleate voids. This indirect role of solutes is
6.6 Dimensional Stability of Irradiated Steels
467
12 WITHOUT PHOSPHORUS ADDITIONS TWO HEATS OF 10% COLD-WORKED 316
TWO HEATS OF 20% COLD-WORKED 316
10% COLD-WORK TWO HEATS WITHOUT PHOSPHORUS ADDITIONS SWELLING %
TWO HEATS OF PHOSPHORUS-MODIFIED 316
20% COLD-WORK
TWO HEATS WITH ADDITIONS OF PHOSPHORUS AND BORON ~1%/dpa
' /
' /
ADDITIONS OF PHOSPHORUS, BORON, TITANIUM, AND NIOBIUM
TWO HEATS WITHOUT ADDITIONS
,7 0
5
10
~
20 x 1 0 2 2 n c m ~ 2
15
NEUTRON FLUENCE (E>0.1 MeV) Figure 6-61. Variation of swelling produced in AISI 316 at 500 + 50 °C as a function of cold work level and solute addition (after Itaki et al, 1987).
in addition to the direct action of the solute fraction still left in solution. The dual role of solutes while in solution or precipitates can sometimes lead to larger swelling after aging to produce precipitation, as shown earlier in Fig. 6-43.
6.6.3.4 Thermomechanical Treatment
Since void swelling has been shown to be related to the dislocation density and its distribution, it is not surprising that swelling was found in early studies to be
468
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
40
0.04 C 0.20 Ti 0.80 Si 2.5 Mo 2.0 Mn
30 SWELLING % 20
Figure 6-62. Influence of phosphorus level on swelling in EBR-II for a titanium-modified 316 steel at 425 and 540 °C (after Garner and Brager, 1985 a).
10
DISPLACEMENTS PER ATOM
strongly dependent on the cold work level. Initially, it appeared that cold work always decreased swelling (e.g., Fig. 6-63) and that it occurred by lowering the swelling rate. It later became obvious that the suppression of swelling by cold work was only temporary and manifested itself primarily as an extension of the transient regime or incubation period of swelling, as shown in Figs. 6-64 and 6-65. The degree of swelling suppression for a given level of cold work was found to vary strongly as a function of steel identity, irradiation temperature and other variables. In some cases, a moderately low level of cold work actually increased swelling, especially at higher irradiation temperatures. Each of these effects of cold work on swelling is demonstrated in Fig. 6-66. In some cases the swelling was initially suppressed by cold work, but then accelerated at higher dpa levels (Garner et al., 1993 c; Krasnoselov et al., 1983), as demonstrated in Fig. 6-67.
Another facet of thermomechanical processing is postproduction aging. Aging at temperatures well below those used for annealing encourages precipitate formation and stress relief, and is therefore sometimes used to adjust the starting properties of a steel. It became obvious early in the various LMR materials programs, however, that aging at intermediate (650800 °C) temperatures prior to irradiation often caused a large increase in swelling, as demonstrated in Figs. 6-68 and 6-69. The very early data presented in Fig. 6-69 were derived from long term sodium-exposed specimens which were later irradiated in EBR-II. These data were initially rejected by the radiation damage community, however. These very high swelling rates, approaching 1 %/dpa or ^ 5 % / 1 0 2 2 n cm" 2 (is > 0.1 MeV) were felt to represent some artifact of sodium leaching of solutes prior to irradiation, a process that would not occur sequentially during typical reactor operation. It was not the sodium exposure,
469
6.6 Dimensional Stability of Irradiated Steels
however, but the time at temperature which was found to be important. As more data became available, however, it was obvious that preirradiation aging was the most efficient way to develop the characteristic 1 %/dpa swelling rate, as shown in Fig. 6-70. The effect of aging is not always monotonic, and under some irradiation conditions it can sometimes delay swelling a little, as shown in Fig. 6-71. One of the most surprising features of some aging studies was the tendency of the swelling of annealed material to sometimes also be accelerated by aging, as shown in Figs. 6-68 and 6-72. This behavior is difficult to explain in terms of the cold work dependence of dislocation density alone. If one investigates the response of dislocation behavior to cold work and aging (e.g., see Fig. 6-73) it is found that aging often produces dislocation densities similar to that of 5-10% cold worked material, and yet the swelling behavior of the aged and cold worked materials is often very different. Brager and Garner (1979) and Garner and Porter (1982) later showed that both cold working and aging exerted much of their influence on the distribution and precipitation of solutes, and thereby on the associated microchemical evolution of the alloy matrix. Aging at intermediate temperatures prior to irradiation significantly accelerated the formation of carbide and intermetallic phases. At higher irradiation temperatures, however, the starting dislo-
Figure 6-63. Effect of cold work level on the swelling of AISI 316 for (a) U.S. steel at 650 °C in EBR-II for exposures of 33 and 50 dpa (after Brager and Garner, 1979); (b) French steel at various temperatures in RAPSODIE for exposures of 20-61 dpaF (Dupouy et al., 1978); and (c) Japanese steel in RAPSODIE at 550 °C and 48 dpa (Uematsu et al., 1979).
(a) 10 650 °C
1.0 x 10 23 n/cm2 (E > 0.1 MeV)
0.1 MeV)
10
400
20 Cold Work Level, %
450
500
550
TEMPERATURE, (c)
0
10
20 30
40
% COLD WORK
50
470
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors SOLUTION ANNEALED
20
1
I
1
I
I
4$
550°C
ANNEALED^/ T_ 15 -
10%CW\// 20% CW.
SWELLING 10 -%
5
f
A
-
/
/
s{ 30% CW
1
5
10
15 x 1022
NEUTRON FLUENCE
20% COLD-WORK
Figure 6-64. Effect of cold work level on swelling of AISI 316 at 550 °C (after Appleby et al., 1977). Note that cold work also affects the precipitate evolution.
450°C
SWELLING
1 2
3"
4
5
6
7
8
NEUTRON FLUENCE, n / c m
9 2
10 x 1 0 2 2
Figure 6-65. Effect of cold work level on swelling of AISI 304 at 450 °C in EBR-II (after Busboom et al., 1975). Note the diminishing effect of higher cold work levels compared to that shown in Fig. 6-64.
6.6 Dimensional Stability of Irradiated Steels
471
20
15
Swelling 10
500 Temperature, °C
400
600
20 Cold Work Level, %
Figure 6-66. Complex influence of cold work level on swelling of M316 stainless steel in DFR at 50 dpa (H/2) (after Brown et al., 1977).
Figure 6-67. Nonmonotonic influence of cold work vs. dpa on swelling of OKH16N15M3B (Fe-16Cr15Ni-3Mo-B) in BOR-60 at 447- 547 °C (after Krasnoselov et al., 1983).
16
14
12
10
£
-
-
68.4
r /
-
8
A
^ S A H h Aged
r
\
CW 25% + Aged
*63.4
' 1J '
/
CW25% \
46.9cr^
\
23.30
400
1
450
500
550 400
450
Temperature, °C
500
550
Figure 6-68. Effect of aging at 750 °C for 100 h on the swelling of (left) annealed low carbon 316 stainless steel and (right) 0.05 wt.% carbon cold worked 316 stainless steel, both after irradiation in RAPSODIE (after Dupouy et al., 1978). Exposures are in units of dpaF.
472
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
wards (1993) have investigated the role of cold work and aging in simple solute-free metals and alloys which do not undergo microchemical evolution. It was found that cold working decreases swelling at relatively low temperatures but actually increases swelling at higher temperatures and nickel levels. Even more important, it was shown that aging of such simple metals after cold working produced intermediate dislocation densities that were more stable during irradiation and thereby reached the saturation levels sooner. This induces an earlier onset of swelling at higher temperatures, as shown in Fig. 6-74. This phenomenon probably accounts for the non-monotonic influence of cold work on swelling of FV548 at 600 °C shown in Fig. 6-66. Note that a similar behavior did not occur in the range 400- 550 °C. One of the most surprising aspects of swelling in some early studies was the very strong influence of the final annealing temperature on swelling of some steels, examples of which are shown in Figs. 6-75 and 6-76. With hindsight, this is clearly an ex-
Aged at 704°C
0
2 4 6 8 10x10 22 2 Neutron Fluence, n cm" (E>0.1 MeV)
Figure 6-69. Comparison of neutron-induced density changes at 480 °C of three steels in EBR-II following prior long term exposure in flowing sodium at 704 °C (after Busboom et al., 1975).
cation density also plays a more direct role, arising from the slowness with which annealed materials approach saturation dislocation densities in this temperature regime. Garner (1993) and Garner and Ed40
1
I
650°C AGE _ 100 hr
30
I
1
10% CW 800°C, 1 hr
i
.
i
i
10% STRAIN AT 650°C, 100 hr AT 650°C i
-
SWELLING %
1%/dpa
/ /
20 - -
/
:
/ / 10
j
-
^1
20
1
4 0 6 0 0
d° 8 ° 2 0 4 0 6 0 0
DISPLACEMENTS PER A T O M
Figure 6-70.
I
i
Early development of 1 %/ dpa swelling rates in aged OKH16N15M3B during irradiation in BOR-60 at 400-550°C (after Krasnoselov et al., 1983).
I
2 0 4 0 6 0
80
6.6 Dimensional Stability of Irradiated Steels
473
40
540°C
540°C
Corel
Core 4
30
^-1%/dpa
• ^1%/dpa
f
•E 20 -
^ - 2 0 % Cold Work 1150 C Anneal
10 -
7 Jlr 10
— 1150°C Anneal + 650°C Age -1300°C Anneal I 20
' / r /
r- 20% Cold Work -1150°C Anneal -1150°C Anneal + 650°C Age 1300°C Anneal
/ FT
0
10
I 20
30X10 26
Neutron Fluence (n rrr2, E>0.1 MeV)
Figure 6-71. Variations in swelling observed at 540 °C in EBR-II for two nominally similar heats (Core 1 and Core 4 steels used to construct FFTF) of AISI 316 as a function of starting condition, showing influence of annealing and aging (after Garner et al., 1992 c).
15
316 T i
COLD-WORKED ANNEALED.
10
ANNEALED AND AGED COLD-WORKED AND AGED /
A /
SWELLING
15.15 Ti ANNEALED 10
Figure 6-72. Effect of starting state on swelling at 500 °C in PHENIX for two European Community steels (after EhrlichetaL, 1986). 50
100
DISPLACEMENTS PER ATOM
474
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
SOLUTION TREATED
20%C.W. • 100 HR ANNEAL AT I2OO°F
10% COLD WORKED
20% COLD WORKED
Figure 6-73. Dislocation arrays observed in AISI 316 as a function of cold work level and aging (courtesy of H. R. Brager of Westinghouse Hanford Company).
30 600°C 32 dpa
425°C 31 dpa
20 Swelling Cold-worked and Aged
Cold-worked and Aged
Figure 6-74. Temperature dependent influence of cold work and aging on the swelling in EBR-II of simple Fe-15Cr-Ni ternary alloys (after Garner, 1993).
10
10
20
30
40
* 10 20 Nickel Content, wt%
ample of the influence of temperature on the distribution of elements such as phosphorus and carbon, and the resultant impact on microchemical evolution. This type of temperature sensitivity exerts an equally strong influence in a much more subtle form that was not discovered until
30
40
50
much later. Garner et al. (1992 c) showed that swelling was sensitive to the details of annealing procedure used prior to cold working of the steel, a variable which varied strongly from one study to another and which was almost never considered or reported to be a significant factor.
6.6 Dimensional Stability of Irradiated Steels 20
15 1000 °C
>
1075 °C Swelling 10 1200 °C 5
-
1300°C^7i
400
X
500 600 Temperature, °C
Figure 6-75. Effect of annealing temperature on swelling of FV548 in DFR at 40 dpa (H/2) (after Brown et al, 1977).
ANNEALED AT 900°C
I5
475
accomplished by moving the very long tube through a moderately long furnace. Thus the feed rate, gas atmosphere and temperature profile of the furnace are not only important variables, but the ranges of variation in these variables are very atypical of the longer anneals and quenching methods employed in typical laboratory practice. Figures 6-77 and 6-78 demonstrate the sensitivity of swelling to the intermediate annealing temperature employed for one heat of 20 % cold worked AISI 316 irradiated in EBR-II. The feed rate was not found to be a very sensitive variable in this particular experiment but may well be important for other steels. The behavior of the steel in this study was explained in terms of the (1) availability of phosphorus, arising from the temperature dependence of phosphide dissolution and (2) the "effective" amount of cold work remaining from the previous pass through the furnace. The influence of both of these on carbide nucleation was also invoked by Garner et al. (1992 c) to explain the observed behavior. 6.6.3.5 Displacement Rate
DISTANCE ALONG PIN
-*-
Figure 6-76. Influence of annealing temperature on swelling of AISI 316 fuel pins in PHENIX (after Boutard et al., 1977).
In general, the method used to produce cold work in the laboratory varies strongly from that used in industrial production of cladding tubes or hexagonal ducts for LMRs. One major difference lies in the annealing treatment experienced by the steels prior to cold working. Since the production of a cold worked component often involves multiple stages of reduction, this in turn requires multiple intermediate anneals. These "temper" anneals are usually
The early charged particle simulation studies were most often interpreted in terms of rate theory considerations rather than incubation effects. One well known consequence of the accelerated displacement rates employed in such studies was an upward shift in the temperature regime of swelling by 40-60 °C per order of magnitude increase in displacement rate (Bullough and Perrin, 1972; Garner and Guthrie, 1975). Since the swelling rate was thought at that time to be strongly dependent on temperature, such "temperature shifts" would produce large increases or decreases in swelling, depending on the temperature regime in which the irradiation was being conducted.
476
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
(470 -> 459°C)
500°C
32
567°C
24
3.5x10 22 x 1 0 22
1020
1060
1100
1020
1060
1140
1100
Annealing Temperature, C Figure 6-77. Impact of temper annealing temperature prior to last cold reduction on the swelling of 20% cold worked AISI 316 in EBR-IL Open symbols denote a furnace feed rate of 20.3 mm/s and closed symbols 15.2 mm/s. This experiment fell slowly in temperature during irradiation and did not achieve the target temperatures indicated in the upper left hand corner (after Garner et al., 1992 c).
1121OC 30
0
4
8
12
16x1022
2
Neutron Fluence, n/cm (E>0.1 MeV) Figure 6-78. Influence of temper annealing temperature in determining the duration of the transient regime of swelling in EBR-II for 20% cold worked AISI 316 (after Garner et al., 1992 c).
400
500 600 Temperature, °C
700
Figure 6-79. Temperature shift of swelling observed in pure nickel during 2.8 MeV self-ion irradiation at displacement rates of 7 x 10 ~4 and 7 x 10 ~2 dpa/s (after Westmoreland et al., 1975).
10
i
6.6 Dimensional Stability of Irradiated Steels
477
20
dpa/sec
~ 4> = 18.5 x 10-7 1 dpa/sec 1
10
V : M :
_
4> = 13.0 U » r-|l
SWELLING % 5
0) = 7.0 — I • 1 1 ' I ^
1
1 1
A .D
T
10
l'D
•D i
D o
50 100 DISPLACEMENTS, dpa F
20 40 80 DISPLACEMENTS, dpa F
100
Figure 6-80. Effect of displacement rate on swelling of (left) annealed AISI 316 in RAPSODIE fuel pin cladding at 562°C and (right) 20% cold worked AISI 316 in PHENIX fuel cladding at 590-610°C (after Seran and Dupouy, 1982, 1983).
Early charged particle studies by Westmoreland et al. (1975) indeed demonstrated changes in the temperature dependence of swelling of pure nickel associated solely with increases in displacement rate, as shown in Fig. 6-79. Similar results were obtained in other charged particle studies, for example in copper (Glowinski et al., 1976) and AISI 316 stainless steel (Menzinger and Sacchetti, 1975; Sacchetti, 1977). Porter and Hudman (1980) later showed decreases in neutron-induced swelling of annealed AISI 304L at 385 °C and in 5 % cold worked AISI 316 at 400 °C with increasing displacement rate in the range 385-400°C. When it became apparent that the steady state swelling rate was not a strong function of temperature, it was realized that differences in displacement rate at most temperatures could only exert their influence on the duration of the transient regime of swelling. Porter and Garner
(1985) demonstrated that swelling and creep profiles developed along the length of 152 cm long creep tubes irradiated in EBR-II could be explained in terms of the influence of the displacement rate profile on the duration of the transient regime of swelling. Figures 6-80 and 6-81 demonstrate that relatively small differences in displacement rate have indeed been found to induce large changes in the duration of the transient regime at relatively high temperatures, but not in the posttransient behavior. Although defect supersaturations and thereby void nucleation are predicted to be enhanced at higher displacement rates and therefore should produce shorter incubation periods, increases in displacement rate frequently produce longer incubation periods. This is thought to be a consequence of several rate-dependent processes. First, the dependence of various precipitation sequences on displacement rate (nickel sili-
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
478
500°C
450°C I
I
1
I
10 _
_
1
550°C i
1
-
?- - ftJ //
I f f ff
SWELLING _
I
_
i
_
/
(j
_
/
i
i 1
OUTER / i ROWS 1
X o
cp
1
1
I
60
i
80
40
i
60
?|
| _ |
1°
CORE 'CENTER
- 1° r D
40
i| OO 1
• '
r-JL L-JJ (V— #O
rj
i
O°
i
80
20
40
60
DISPLACEMENTS, dpa p
Figure 6-81. Temperature dependence of effects of dpa rate and time on swelling of AISI 316 in RAPSODIE fuel pin cladding (after Seran and Dupouy, 1982).
cides) or on time at temperature (thermally stable phases) is very important in the rate of microchemical evolution of the alloy matrix. For example, Boulanger et al. (1983) explored the dose rate sensitivity of cold worked 316 and showed that Laves phase formation decreased progressively with increasing displacement rate. Swelling also decreased concurrently. Some small thermally stable precipitates such as phosphides also tend to dissolve at higher displacement rates, as cascade-induced resolution of such precipitates overwhelms thermal transport processes contributing to their formation (Garner et al., 1993 e). Second, the dislocation and loop evolution is also dependent on displacement rate. LeNaour et al. (1982) have demonstrated this dependence in annealed AISI 316. 6.6.3.6 Temperature
The determination of the temperature dependence of void swelling has been the
subject of extensive research over the last two decades, but was hindered by the strong concurrent influence of displacement rate and the inherent coupling of displacement rate and temperature in the various LMR reactors. The large influence of many other variables on swelling has also complicated the separation of temperature's individual role. It now appears that one can define for a given temperature two regimes of swelling (transient and steady state). There are three separate temperature regimes of steady state swelling at a given displacement rate, however, designated here as I, II and III, as shown in Fig. 6-82. Temperature regime I of steady state swelling covers the range from the lowest temperature where swelling occurs to the temperature where the 1 %/dpa plateau (regime II) begins. Finally, regime III covers the range in which the swelling rate declines from 1 %/dpa to zero. Regime I has been traditionally characterized as being controlled by the increas-
6.6 Dimensional Stability of Irradiated Steels
ing tendency toward direct recombination of vacancies and interstitials as the temperature decreases. More recently, however, the recognition has grown that the temperature dependence of recombination in traditional rate theory is not large enough to produce the observed temperature dependence of swelling. Nor is rate theory capable of producing a swelling rate of 1 %/dpa in Regime II using measured defect survival rates (Rehn, 1990) and traditional bias factors for dislocations (Skinner and Woo, 1984). A new, more strongly temperature-dependent process, designated the production bias concept, has been proposed by Woo and Singh (1990, 1992). This model incorporates into the rate theory the different formation and dissolution rates of both interstitial clusters and vacancy clusters. The applicability of the production bias concept to explain many features of swelling experiments has been demonstrated by Woo et al. (1992). Similar application of this concept to various creep-
1.2
1.0
x
f
I
0.8
i
n
I l I l l 1l
0.6
0.4 1 0.2 -
300
\1
'M
~1x10-6dpa/sec
\\ \\
/ /
J400
11
M
.
.
500 600 700 Irradiation Temperature, C
l i
800
900
Figure 6-82. Schematic representation of the three regimes of steady state swelling rate in AISI 316. The full range of regime III is still undetermined.
479
swelling experiments has been demonstrated by Woo and Garner (1992) and Wooetal. (1993). Garner and Wolfer (1984) demonstrated that rate theory descriptions of the steady state swelling rate depend very strongly on the value chosen for the vacancy migration energy. Whereas many earlier papers used 1.4 eV, they showed that, according to the literature, 1.1 eV was a more correct value, and that this value produced a relatively temperature-independent description of the swelling rate that was characteristic of regime II. Most of the earlier studies using 1.4 eV were attempting to reproduce the strongly peaked temperature profiles of swelling observed in ion-bombardment experiments and also observed in early swelling profiles taken from EBR-II control and safety rod thimbles. These strongly peaked profiles were later shown to be distorted in their apparent temperature dependence by the injected interstitial effect and surface proximity during ion bombardment (Garner, 1983), and the large axial gradients of displacement rate that exist across the small EBR-II core (Garner and Porter, 1984). Garner and Wolfer (1984) also demonstrated that the strong temperature dependences of void and dislocation microstructures were not reflected in the temperature dependence of the steady state swelling rate as long as the two temperature dependences were similar. Such similarity has been observed many times. The description of regime III swelling behavior is similar in almost all swelling theories and predicts that at higher temperatures, vacancy emission from voids begins to overwhelm the in-flow of excess vacancies. The temperature above which swelling does not occur has not been conclusively demonstrated, however, since as higher fluences are reached, swelling ap-
480
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
pears to extend in temperature range, due to the strong effect of temperature on the incubation period of swelling at higher temperatures. The most important determinant of the apparent temperature dependence of swelling for most LMR applications, however, is the duration of the transient regime of swelling. As shown in preceding sections, this parameter is sensitive not only to irradiation temperature but also to a very large number of other variables, all of which affect the microchemical evolution of the alloy matrix. This evolution is dominated by two separate regimes of precipitation, a factor which caused a persistent misperception to arise that the steady state swelling rate was not only strongly dependent on temperature, but exhibited two peaks in temperature. Garner (1985) showed that the apparent "double bump" in swelling rates measured for various 316 variants in DFR (Fig. 683) was partially an artifact of the axial
0.7
displacement rate profile and partially due to the tendency to assume that steady state swelling had been reached at the highest fluence level attained at each temperature. The curvature of each peak actually represented only the sluggish microchemicallydominated rate of approach to a 1 %/dpa swelling rate in each of the two precipitation regimes. Once swelling began, there was a different rate of approach to 1 %/ dpa for each bump, as shown in Fig. 6-84. The posttransient swelling within each of the two bumps was also found to proceed at similar rates, as demonstrated in Fig. 6-85 for annealed 316 fuel pins irradiated in RAPSODIE. 6.6.3.7 Temperature History
Changes in temperature during irradiation can sometimes have a pronounced effect on microchemical evolution and thereby on both void swelling and irradiation creep (Chin and Straalsund, 1978;
LINEAR APPROACH
0.6
• 0.5
Fe-17.2Cr-13.7Ni-2.4Mo -1.85Mn -0.63Si -0.035C
• ANNEALING TEMPERATURE o o1000°c] o a1100°C I f 1 h * *1200°C f o 01250°CJ
SWELLING 0.4 RATE %/dpa (H/2)
• age 24h at 850 °C
0.3
0.2
0.1 DFR IRRADIATION 0.0
350
400
450 500 550 600 TEMPERATURE °C
650
Figure 6-83. "Double bump" swelling rate profiles observed in annealed and aged M316 after irradiation in the DFR, as observed by Gittus, Watkin and Standring, and compiled by Garner (1985). The "apparent" swelling rates shown above are actually underestimates of varying degrees arising from the premature assumption of swelling having reached the regime where it is linearly proportional to displacement dose.
6.6 Dimensional Stability of Irradiated Steels (a)
481
0.7 4.2% per 10 2 2 n cm"2 (E > 0.1) 0.6 M316. ANNEALED AND AGED o 1000°C A 1100°C 1 hr + 8S0°C v 1200°C FOR 24 hrs D 12S0°C
0.5
SWELLING RATE %/dpa (H/2)
0 4
M316, ANNEALED O 1050°C, 1/2 hr 0.3
316, ANNEALED O 1070°C, AC
0.2
316L, ANNEALED • 1050°C, 1/2 hr
0.1
DATA ABOVE 475 °C ONLY
0.0
5
10
15
20
MAXIMUM SWELLING, % (b)
u./
\
s
2
4.2% PER 10 2 2 ncm" (E> 0.1) 0.6
0.5
/
0.2
0.1
/
DATA BELOW 475 °C y/ \
SWELLING 0.4 RATE %/dpa (H/2) 0.3
s
p
\
a
/
Y±o
s
/ /
ss
y
TREND LINE FOR
M316DATAAT>475°C-
y
• A / \ '7// / // / /
/
/
S
1t 5
10
15
Figure 6-84. "Apparent" swelling rate vs. maximum swelling plotted for a variety of U.K. steels for temperatures (a) greater than 475 °C and (b) less than 475 °C (Garner, 1985). In each case the swelling rate increases as the swelling increases, slowly approaching 1 %/dpa NRT (0.7%/dpa H/2 = 0.82%/dpa NRT).
20
MAXIMUM SWELLING, %
Bates and Gelles, 1978; Foster and Boltax, 1980; Garner etal., 1981b; Bates, 1981; Yang and Garner, 1982; Porter etal., 1990). A recent review of temperature history effects has also shown that results of many microstructurally oriented experiments were dominated by unintentional or unrecognized variations in temperature (Garner et al., 1993 f). Both gradual and abrupt decreases in temperature have been shown to sometimes significantly enhance the onset of swelling of 316 stainless steel, especially if the change occurs during the transient
regime of swelling. Nickel segregation at higher temperatures, followed by subsequent silicon segregation at lower temperatures, has been found to be a very effective way to nucleate and grow the y'(Ni3Si) phase (Garner etal., 1981b). Formation of this otherwise sluggishly growing phase was identified earlier as a precursor to rapid swelling of typical stainless steels. The impact on swelling of both deliberate or unintentional changes in irradiation temperature depends not only on the steel being studied, as well as the magnitude and
482 (a)
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors PEAK DOSE = 100 dpa F 5.67%
= 15 x 10 2 2 n/cm 2 (E>0.1 MeV) 4.64%
BOTTOM OF THE FUEL
TOP OF THE FUEL
already attained swelling rate if the temperature remained in regime II. Porter, Wood and Garner (1990) demonstrated, however, that posttransient changes from regime II to regime I drastically lowered the steady state swelling rate, reflecting the temperature-dependent operation of the production bias concept. In each of these two cases, the microstructures attained prior to the temperature change had little effect on the swelling after the change, indicating that details of microstructure play a lesser role in void swelling than originally anticipated.
80
100
DISPLACMENTS PER ATOM (dpap)
Figure 6-85. (a) Double-bump swelling profiles observed in FORTISSIMO fuel elements clad with annealed AISI 316 after irradiation in RAPSODIE (after Weisz, 1978). (b) Swelling behavior at each of the two peaks (after Leclere and Marbach, 1979).
direction of the change, but also on the accumulated exposure when the change was initiated. Sometimes the effect is small and sometimes it is very large. Yang and Garner (1982) showed that posttransient increases or decreases in temperature in regime II had no effect on changing the
6.6.3.8 Stress Misperceptions concerning the influence of stress on swelling have persisted longer than those related to most other variables. As reviewed by Garner et al. (1981 a), the possible role of stress on swelling was debated for some time as various contradictory data sets were compiled. A more detailed and recent review of data and theoretical descriptions of stress- enhanced swelling was presented by Hassan etal. (1992). The first conclusive evidence for stressaffected swelling in austenitic steels was reported by Bates and Gilbert (1975,1978) and has since been confirmed by many
6.6 Dimensional Stability of Irradiated Steels
studies. Using pressurized tubes, they demonstrated that the swelling of AISI 316 was linear with hoop stress up to the proportional elastic limit, decreasing thereafter at higher stress levels. Compared to the annealed condition, the linear range therefore persisted to higher stress levels in the cold worked condition. Similar results were obtained by Garner et al. (1981a), Herschbach et al. (1990, 1993), and Shamardin et al. (1990), although the swelling enhancement was not always linear with stress level. Ferritic-martensitic steels have also been shown to exhibit stress-enhanced swelling (Toloczko et al., 1994). Reflecting the then prevailing misperception that the swelling rate was strongly dependent on temperature and other variables, Bates and Gilbert (1978) modeled the stress effect as an enhancement of the swelling rate by the hydrostatic stress. This of course implied that compressive stresses would decrease the swelling. Garner etal. (1981a) noted that the rapidly expanding data field on AISI 316 and other alloys was more consistent with tensile stresses acting to shorten the transient regime of swelling. This suggestion was proven to be valid by Porter et al. (1983) and confirmed by other studies (Harbottle, 1977; Khera etal., 1980; Porter and Garner 1985; Seran et al., 1990). For some time, however, it was still assumed that the hydrostatic component of the stress state was the important variable, implying that compressive stress states would delay the onset of swelling. In one experimental series covering a wide range of stress states, however, Hall (1978) of Argonne National Laboratory studied the role of applied stresses on swelling. The preliminary but unpublished results suggested that any stress state - tensile, compressive, shear, or mixed - would
483
accelerate the onset of swelling, but this experiment was never completed. This proposal was later proven to be true (Sahu and Jung, 1985; Lauritzen etal., 1987). Since purely hydrostatic tests were never performed, this indicated that the deviatoric component of the stress state was most likely the important variable. This finding foreshadowed the strong linkage between swelling and irradiation creep, since the latter also responds to the deviatoric component of the stress state. The effect of stress on swelling was initially thought to be negligible at lower irradiation temperatures and to increase steeply at temperatures above 550 °C (Garner et al., 1981 a). It appears, however, that stress enhances swelling at all temperatures, but is more sluggish in exerting its influence at lower irradiation temperatures. A typical example is shown in Fig. 6-86. This relatively recent insight had a large impact on the analysis of irradiation creep experiments by Garner and co-workers in the early 1990s. 6.6.4 Strains Arising from Irradiation Creep 6.6.4.1 Introduction to Irradiation Creep
While the potentially deleterious impact of thermal creep at higher LMR service temperatures has long been known to be of engineering concern, the orders of magnitude increase in creep during relatively low temperature irradiation was as unexpected as was the discovery of void swelling. Figures 6-87 and 6-88 demonstrate that irradiation creep rates at temperatures below the thermal creep regime can be comparable to those of typical thermal creep conditions. The degree of enhancement of creep is greatest for lower irradiation temperatures, decreasing as the thermal creep regime is approached. While this enhance-
484
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors 20
450-460 c >Q
16 ~-~
•
12
-
8
-
4
-
L
o OMPa D 90 MPa 170 MPa
316Ti+p
_J/jb I 20
40 60 80 Displacements Per Atom
100
ment arises partially from the radiationinduced alteration of dislocation and precipitate microstructure, it is due primarily to the presence of both interstitial atoms and vacancies at very large supersaturation levels. Figure 6-89 shows that the measured activation energies associated with irradiation creep are typical of inter-
Irradiation Creepy-
800
-
/
400
/
stress 550°C
MPa
/
310°C 290°
355° 500°, A • • 500° 380° •440°
/
y
200
mn
/
•
o
33°
/
I
c/mermal Creep >^OOX
I I Mil
1
I I
i mil
i
10"4 10"3 Creep Rate, % per hr
i|
10"2
Figure 6-87. Influence of neutron irradiation on the creep rate of annealed KM6N15 steel (after Logunstsev etal., 1984).
Figure 6-86. Effect of stress on swelling of two modified 316 stainless steels irradiated in the form of pressurized tubes in PHENIX (after Dubuisson et al., 1992). 120
stitial-related processes rather than vacancy-related processes observed in the thermal creep regime. While there are a very large number of creep mechanisms that have been proposed, most fall in several broad categories with respect to the microstructural components and the defects involved. Matthews and Finnis (1988) have presented the most extensive review of these mechanisms and ranked them in terms of their relative plausibility. Garner and Gelles (1988) demonstrated that, based on microstructural evidence obtained from specimens with an absence of swelling, those mechanisms associated with the SIPA concept appear to be the most likely processes to control the rate of irradiation creep at doses, displacement rates and temperatures typical of LMR interest. Other proposed processes are most likely contributing as well, but do not appear to be rate-controlling. There is often a tendency to treat irradiation creep and thermal creep as separate and directly additive processes, but the situation is more complex, primarily because each of the various stages of thermal creep can be altered by the radiation-induced microchemical and microstructural evolutions. While Lovell et al. (1981) and
6.6 Dimensional Stability of Irradiated Steels
485
Figure 6-88. Effect of irradiation (solid lines) on thermal creep (broken lines) of annealed O3Kh20N45M4BCh, conducted in SM-2 at 575 and 600 °C, and at 650 °C in RBT-6 (after Grabova et al, 1986).
10'° 200
300
400
500
Applied Stress, MPa
Temperature, K 985 923 823
723
573
O DIN 1.4970 A DIN 1.4981 • Other Austenitics
A H = 1.16eV
Normalized Creep Rate
A H = 0.13eV
100
-
A H h = 1.63eV AHj = 0.09 eV 1
10"
1.0
J
I 1.2
L_J 1.4
I
I
L
1.6x10- 3
Inverse Temperature, K"1
Figure 6-89. Analysis of in-reactor creep data by Wassilew et al. (1987) showing that the rate controlling defect varies with temperature, with activation energies for volume migration of 1.63 eV for vacancies and 0.09 eV for interstitials.
Gilbert et al. (1987) report that irradiation appears to delay the onset of tertiary thermal creep in 20% cold worked AISI 316, other researchers found that irradiation generally accelerated thermal creep (Wassilew et al., 1987; Shibahara et al., 1993). A clear example of accelerated thermal creep is shown in Fig. 6-90 and an example of its consequences on creep rupture life is provided in Fig. 6-91, but it appears possible that thermal creep can be delayed or accelerated, depending on the steel, its thermomechanical starting state and the irradiation conditions. Garner et al. (1993 c) showed that the stored energy due to cold working and the microstructural resistance to the release of stored energy were particularly important in the high temperature creep response of austenitic steels. Since LMR design philosophy required that the temperature-stress regime of thermal creep be avoided, the remainder of this section will concentrate on irradiation creep. Due to the importance of this phenomenon to LMR and other reactor concepts, there have been a number of conferences that focussed primarily on irradia-
486
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors 10 100 MPa hoop stress
8
AD
6
%
4 2
0 600
700
900 800 Temperature, K
1000
Figure 6-90. Diametral strains observed in 20% cold worked PNC 316 pressurized tubes irradiated in FFTF (after Shibahara et al, 1993). Also shown are thermal creep strains at a comparable time and the same stress level.
tion creep (see J. Nucl. Mater., Vols. 65,90, and 159) and many other symposia where irradiation creep was one of the major subjects. In addition, periodic review articles have been published on irradiation creep measurements (Gilbert, 1971; Harries,
1977; Ehrlich, 1981; Coghlan, 1986; Simonen, 1990). Review of the many papers on this subject shows that our understanding of irradiation creep has been continuously evolving, with significant modifications being made even as this chapter is written. Compared to our understanding of void swelling, however, the emerging perception of creep has lagged considerably for a variety of reasons: (1) Compared to other types of irradiation experiments, creep experiments are more expensive to perform, and consume significant reactor space. Experimental trade-offs required by these limitations sometimes mask important features of the creep - swelling relationship. (2) Clark et al. (1984) demonstrated that unless creep and swelling predictions were generated from within the same experiment, the strong sensitivities of swelling to environmental and material history variables would obscure the creep-swelling relationship and produce incompatible
10°
700°C Annealed
Thermal Aged
Cold-Worked
f Irradiated in BR-2
10
1.3
1.4
1.5
1.6
1.7
1.8
T(13.5 + Logt R ),K
Figure 6-91. Effect of starting condition and irradiation on stress rupture behavior of DIN 1.4970 at 700 °C (Wassilew et al., 1987).
6.6 Dimensional Stability of Irradiated Steels
equations. Early experiments tended to study creep and swelling separately. (3) The early misperceptions of swelling discussed in a previous section were also reflected in the concurrent perceptions of creep. Additional complications arose from the difficulty of separating noncreep strains such as precipitate-related volume changes and stress-enhanced swelling. (4) The strong coupling of neutron flux and temperature profiles in LMR reactors often obscures the independent action of these two strongly related variables. (5) One very recent realization is that a fundamental and very reasonable set of assumptions underlying the most extensive creep studies were subtly defective. It was assumed that (a) thin-walled gas-pressurized tubes were excellent simulations of fuel pin cladding, (b) creep strain rates are linear with stress, and (c) increasing the stress levels above those typical of fuel pins (in order to better measure the creep strain) would not significantly perturb the simulation. Separately, as well as in aggregate, these assumptions become progressively more defective as swelling increases (Garner etal., 1994). One significant consequence is that fuel pin strains in general cannot be predicted using pressurized tube strains (Seran et al., 1990; Garner etal., 1987). Similar difficulties have occurred in predicting the behavior of hexagonal ducts at relatively low stress levels (Bump et al., 1973; Shields, 1981; Foster etal., 1993). There are two major ways to conduct in-reactor creep tests using gas pressurized tubes, and an examination of the different results from these two types of tests led to the most recent insights concerning the swelling-creep relationship (Garner et al., 1994). In the French LMR approach,
487
many short tubes were welded together to form long pins for insertion into subassemblies. Diameter changes were determined when the pins were withdrawn from reactor, and then the tubes were cut apart, sectioned and measured to determine density changes (Seran et al., 1990). Different stress levels and different dpa levels required the use of other tubes on different pins. This approach allows the determination of the actual stress-enhanced swelling level, but not the path by which it was reached, since only the average swelling rate over the entire irradiation could be determined. This represents a substantial underestimate of the final instantaneous swelling rate for steels with long swelling incubation periods. This approach also incorporates a relatively large influence of variations in displacement rate and temperature, introducing substantial scatter into the data fields. Initially, U.S. LMR tests were conducted in EBR-II in a similar manner, but without welding (Gilbert and Chin, 1981). The small size of EBR-II subcapsules required that nominally identical tubes at different stress levels be irradiated in different locations, sometimes varying from one irradiation segment to the next, thus introducing significant scatter. The tubes were not destroyed after a given irradiation segment but were reinserted into the reactor for further irradiation after diameter measurements were made. A similar approach involving the use of tubes as long as one meter in EBR-II reduced the scatter significantly and also allowed a study of flux effects along the tube length (Walters et al., 1977). Neither the short or the long tubes in EBR-II, however, involved the use of well-defined or well- controlled temperatures. In the short tube experiments conducted in the larger FFTF core, data scatter was
488
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
reduced significantly by the side-by-side irradiation at one well-defined and actively controlled temperature of all tubes at different stress levels (Puigh and Schenter, 1985). Since the tube diameters were also measured and returned for further irradiation, data scatter was reduced. This made it easier to determine the instantaneous diametral change rates and better illuminate the parametric dependences of creep. A weakness of this and the other EBR-II approaches, however, was that the stress-affected swelling rate was not determined. Based on earlier work by Garner et al. (1981a), it was assumed that at temperatures below ^55O°C stress enhancement of swelling was small and that the stressenhanced increment could be included as a creep contribution without significant consequences. Unfortunately, this assumption camouflaged the strong late-term feedback of swelling into the creep phenomenon. Much of the irradiation creep data from earlier studies have been developed using methods other than pressurized tubes, however. Other techniques involved tensile specimens, bent beams, helical springs, stress relaxation tests and measurements on actual reactor components. A summary of the relative advantages and disadvantages of each technique was provided by Harries (1977) and more recently by Coghlan (1986). Until recently, the data produced by these various techniques produced a picture of irradiation creep that was relatively simple. Irradiation creep strains at temperatures of LMR interest could be described as consisting of several minor contributions (precipitation-related dimensional changes, and primary creep resulting from nucleation of dislocation loops and/or by the relaxation and rearrangement of cold work-induced dislocations) and two major contributions. The major contributions
are described by the creep compliance, Bo, a quantity unrelated to void swelling, and a swelling-driven creep component. Although swelling is very sensitive to a variety of material and environmental variables, the instantaneous creep rate appears to be proportional only to the applied stress and the instantaneous swelling rate. The major components of the instantaneous creep rate B can be written in the form B = e/d = Bo + DS, where ija is the effective strain rate per unit stress and dpa, a is the effective stress, or (^3/2) <Jhoop for a pressurized tube, Bo is the creep compliance, D is the creep-swelling coupling coefficient, and S is the instantaneous volumetric swelling rate per dpa (Ehrlich, 1981; Garner, 1984). This expression is thought to be equally valid for austenitic and ferritic steels, but data on ferritic alloys are rather limited (see Puigh and Gelles, 1989). The French LMR program used a modified notation for the creep-swelling relationship, such that A = E/d = Bo + a S. As actually derived, however, the parameters a and D are not equal. This is because the swelling of French pressurized tubes or fuel pins was only measured once at the end of the irradiation. Thus a is an average value of D and was determined in conjunction with the average swelling rate Sj(4> i) over the entire irradiation. This is in contrast to the U.S. approach where the tube diameters were measured at a number of fluences, making it easier to determine instantaneous diametral change rates. If D is not a function of the swelling level or swelling rate, a and D should be identical. Regardless of the notation used, this type of relationship assumes that creep continues as long as swelling continues, and that D is invariant with respect to flux, stress and total swelling. As will be shown later, neither of these assumptions remain
6.6 Dimensional Stability of Irradiated Steels
As mentioned in Sec. 6.4, the primary and secondary creep regimes involve the development of a quasi-steady state or saturation dislocation microstructure, characterized by a stress-induced anisotropic distribution of Burgers' vectors (Garner and Gelles, 1988). The duration of the approach to saturation depends not only on the stress state and irradiation conditions, but also on the orientation of the stress state with respect to the pre-existing microstructure, which often exhibits a substantial texture. Therefore transient regimes of different magnitude can occur in the same steel for nominally similar irradiation conditions when conducted in different types of creep tests, as shown in Fig. 6-93. The secondary creep rates at a given stress level are quite similar, however. The magnitude of the primary and secondary stages of creep are often masked somewhat by phase-related dimensional changes that vary as a function of steel composition, starting state and irradiation temperature, as illustrated in Fig. 6-94. Note in this figure that if only the last data point was available on each curve, one would reach the conclusion that Bo was a
valid as voids become the dominant feature of the microstructure. It now appears that four separate regimes of irradiation creep exist, described as transient (primary) creep, steady state (secondary) creep in the absence of swelling, swelling-enhanced creep, and feedback-dominated or "disappearing" creep. No creep experiments on ferritic alloys involving significant levels of swelling have yet been conducted. 6.6.4.2 Irradiation Creep in the Absence of Swelling
Primary irradiation creep is characterized by an initially high creep rate that declines with continued irradiation. The creep rate eventually approaches a "steady state" rate that persists as long as void swelling does not occur. This stage is often called secondary creep and is characterized by a rate that is much larger than the thermal creep rate that would develop at the same temperature, as shown in Fig. 6-92. Note that the thermal creep rate is initially negative, reflecting the carbide-related densification that occurs in this steel. 15x10"4
489
138 MPa 454°C
Irradiation Creep
Thermally-Induced Densification and Creep
500
1000 Time, Hours
i
i
1500
2000
Figure 6-92. Comparison of creep rates observed in 20% cold worked 316 stainless steel in uniaxial tests during thermal aging or neutron irradiation in EBR-II (after Gilbert et al., 1972).
490
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors Figure 6-93. Comparison of creep in cantilever beams and in pressurized tubes for 20% cold worked 316 stainless steel, both irradiated in EBR-II (after McSherry et al., 1978; Puigh et al., 1982).
NICE-LOT EBR-II BEAMS BENDING 425°C*
PRESSURIZED TUBES 5 10 DISPLACEMENTS PER ATOM
15
Figure 6-94. Length changes observed in HFR during uniaxial creep tests of (a) three different cold worked steels at 370 °C; (b) 20% cold worked AMCR 0033 at different irradiation temperatures; (c) 20% cold worked AMCR 0033 in different starting conditions (after Hausen et al., 1988).
0.06
370 °C, 130 MPa
316
0.04
0.02
0.00
20% Cold Worked AMCR 0033
• 200 MPa 140 O 100 • 60 o 30 A
2.0
1.7
0.8 x 10-2 MPa-1
Stress-Free Swelling Used
>
StressEnhanced Swelling Used
Average value based on entire data set (1A-1G)
0.6
Djn
1.3 1.0
0.4
Stress-Free Swelling Used
0.8 x 10-2 MPa-1
0.2
_L
0.6
0
0.1
0.2
0.3% /dpa
AS/Adpa
Figure 6-126. Average creep-swelling coupling coefficients vs. average swelling rate for the experiment shown in Fig. 6-125, calculated for the first seven irradiation segments, and also for each of the last three segments designated IE, IF, and 1G (after Toloczko and Garner, 1994).
0.4
0.2
Stress-Enhanced Swelling Used
I 50
100
150
200
250
Hoop Stress, MPa
Figure 6-125. Average creep coefficients vs. hoop stress derived from the first seven irradiation segments of the PCA irradiation experiment shown in Fig. 6-114 for both the stress-insensitive swelling and stress-enhanced swelling cases (after Toloczko and Garner, 1994).
obvious even in the ignorance of the degree of stress enhancement (Garner and Toloczko, 1992; Garner etal., 1994). Thus, the very reasonable assumption that the creep-swelling coupling coefficient is independent of swelling is actually incorrect, and the limiting value of 0.6 x 10 ~2 MPa" 1 is representative primarily for cases where swelling is relatively low and the stress is relatively high. The total amount of creep for other situations depends on the time-dependent path taken by the swelling and stress. For situations where swelling is significant, the average creep coefficient will approach the asymptotic value.
6.6.5 Deformation Behavior in Response to Creep, Swelling and Gradients in Important Variables
At this point it is obvious that swelling and irradiation creep are very sensitive to environmental variables, temporal and spatial gradients of which exist throughout the LMR core. An example of the various interacting gradients is shown in Fig. 6127. The gradients in dpa rate, for instance, along the length of a component sometimes act to shape its profile in a variety of ways, accentuating or suppressing the tendency of some steels to exhibit complex shapes as a function of temperature. One consequence of such interactions is that peaks in component deformation often do not occur at the peak displacement level, as shown in Figs. 6-128 and 6-129. Gradients across a subassembly frequently lead to large variations in behavior of individual pins in a subassembly, as demonstrated in Fig. 6-130.
6.6 Dimensional Stability of Irradiated Steels
509
-^-Active Core—**
CW316Ti 6 _ Fast Flux, 10 1 5 n/(cm 2 s ) 4 2
jf—N.
0.5
600 Duct Mean Temp, °C 500
Fuel Column
Figure 6-129. Single-bump deformation profiles observed in fuel pin cladding irradiated in PHENIX (after Lippens et al., 1987). The peak deformation occurs well below the core centerline where the peak exposure level was 93 dpa.
400 -
10 Across-Flat / 8 Fast Flux / Gradient 6 / 10 1 4 n/(cm 2 S ) 4
\ \
\
80 Across-Flat 60 Gradient Temp, °C 40 20
Figure 6-127. Axial variations in environmental factors controlling the behavior of hexagonal ducts (after Boltax, 1992).
CWM316
The integrated swelling along the length of a pin also responds to such gradients as shown in Fig. 6-131. In this case, the length strains are due only to swelling and not irradiation creep, as defined by the Soderberg relationship (Soderberg, 1941; Norton and Soderberg, 1942). Note that a few pins from a different but nominally similar lot of tubing rise above their neighbors, reflecting the influence on swelling of small differences in composition and fabrication history. A similar difference in elongation of individual pins in a single fuel subassembly has also been shown by Reshetnikov (1991) to arise from compositional variations, primarily in silicon content.
AD
6.6.5,1 Interactions Between Creep and Swelling-Induced Stresses
0 Hot
100
200
300 Cold
Distance from Pin Base, mm
Figure 6-128. Double-bump deformation profiles observed in fuel pin cladding irradiated in DFR (after Lippens et al., 1987).
Gradients in swelling and creep can lead to interaction of components that can affect the reactivity of the core, the movement of control rods, the flow of coolant and the ability to withdraw and replace components as needed. Since creep and swelling are strongly coupled, stresses arising from such interactions never reach values that in themselves cause failure, but
510
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors AXIAL LOCATION FROM BOTTOM OF FUEL (mm) 38
0.249
i
190 I
I
343 I
I
495
648
800
I
6.32
ROW-6 FAST FLUENCE = 15 x 10 2 2 n/cm 2 316 SS
6.20
0.244
0.229
960
7.5 13.5 19.5 25.5 31.5 AXIAL LOCATION FROM BOTTOM OF FUEL (In.)
i-J5.82 37.5
Figure 6-130. Diameter profiles from a large number of fuel pins located in an outer row assembly with large flux gradients (after Makenas et al., 1990 a).
(a)
Figure 6-131. (a) Top of a bundle of D9 fuel pins irradiated to a peak fluence of 2.1 x 1023 n cm" 2 (E > 0.1 MeV), showing varying length of pins in response to gradients across the bundle in flux and temperature and also to small variations in pin fabrication history #nd composition, (b) An undistorted fuel pin assembly with nonswelling HT9 cladding at 1.9 x 1023 n cm" 2 (E > 0.1 MeV) (after Makenas et al., 1990 a).
they can cause distortions that may precipitate failure by other means, especially if they block coolant flow or lead to large increases in forces necessary to withdraw components. An early example of creepswelling interaction occurred in EBR-II hexagonal thimbles at high neutron exposures. As the ducts increased in diameter the stand-off buttons or "dimples" on each flat that separated adjacent ducts were depressed below the duct surface, forming a moat, as shown in Fig. 6-132. Differential swelling can arise from interaction of components constructed from the same alloy but exposed to different environmental factors. It can also arise between components made from different alloys. The combined action of differential swelling and irradiation creep can lead to distortions such as those shown in Figs. 6-133 and 6-134. Restraining forces applied to components that are swelling will tend to limit their movement, but irradiation creep will then distribute the restrained volume in-
6.6 Dimensional Stability of Irradiated Steels
Before Irradiation
Moat
After Irradiation
Figure 6-132. Deformation of spacer dimples on hexagonal guide thimbles in EBR-II (after Harkness et aL, 1970). The dimple height was reduced from 0.017 to 0.002 inches, with the moat being 0.006 inches wide and 0.008 inches deep.
511
crease toward unconstrained directions. Note in Fig. 6-135 that a distinct "waviness" developed in a highly irradiated pin bundle composed of 20% cold worked 316. The waviness corresponds to a permanent spiral deformation of the pins in response to restraining forces applied by the spiral wire wrap that separates each pin from its neighbors and also by the restraint and interaction with neighboring pins. This interaction also introduces ovality distortions into the cross section of the pin (Fig. 6-135 b). Interestingly, the periodic spiral deformation does not significantly affect the operation of the fuel assembly.
60
SA 316 Cladding and Wire Wrap 59 E
CO
5
Figure 6-133. Interaction between the spiral wire wrap separating individual fuel pins and the surrounding hex can in an EBR-II driver fuel assembly (unpublished data, courtesy of D. L. Porter).
58 peak exposure: ~1.6x10 2 3 n/cm2 (E>0.1 MeV) 57
i
I
1200
1000
1600
1400 Elevation, mm
n.
1 mm
STAPE16 ^Wrapper in PFR
/ dilation profile A ^ of wrapper
szzz V7V7
STM316
grids
T777
^supporting ^ CW316 fuel pins
f A+\ C ^
localized bulges of ~ 0.4 mm
^^ assumed non-bulged profile with maximum at ~ 9.4 mm —t-—
Figure 6-134. Grid-wrapper interaction due to different swelling behavior of two steels in a PFR subassembly (after Higginson and Lilley, 1990).
512
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Figure 6-135 a. Waviness induced by spiral deformation of wire-wrapped fuel pins clad with 20% cold worked 316 irradiated to 1.65 x 1023 n cm" 2 (E > 0.1 MeV) in FFTF (after Makenas et al, 1990a).
The hexagonal ducts which enclose the fuel pin assemblies also interact with each other as they deform in response to local variations ki temperature, displacement rate and stress level. Figure 6-136 demonstrates the types of problems that arise and which can limit the lifetime of individual ducts. 6.6.5.2 Strategies for Reducing the Consequences of Swelling and Creep
It is obvious that the reduction of swelling in itself resolves most of the prob-
lems discussed in earlier sections. While most approaches have concentrated on extending the incubation period of swelling in austenitic steels, the most successful strategy has been the replacement of cladding, wire and ducts constructed from austenitic steels with components constructed from ferritic-martensitic steels. The Core Demonstration Experiment conducted in the FFTF core center region using a partial core load of ten HT9 fuel assemblies and six blanket assemblies demonstrated the feasibility of this strat-
6.6 Dimensional Stability of Irradiated Steels
513
AXIAL LOCATION FROM BOTTOM OF FUEL (mm) 38
190
343
495
648
800
960
316 SS 0.250 " FAST FLUENCE = 15 x
6.35
0.245
6.22 E 6.10
5.97 5.84
0.230
1.5
4.5 7.5 10.5 13.5 16.5 19.5 22.5 25.5 28.5 31.5 34.5 37.5 AXIAL LOCATION FROM BOTTOM OF FUEL (In.)
Figure 6-135 b. Diameter measurements from a single FFTF fuel pin. Each curve is a diameter profile taken at a different circumferential orientation, showing ovality induced by swelling-creep interactions with wire wrap and neighboring pins (after Makenas et al, 1990 a).
egy (Baker et al., 1993; Ethridge and Waltar, 1989; Bridges et al., 1991). While design allowances can be made to absorb the dimensional changes and distortions caused by swelling, there are also other approaches to mitigate its impact. One of these involves the incorporation of TOP LOAD PAD YOKE
WITHDRAWAL FORCE ,
CORE * BARREK0ABOVECORE LOAD PAD YOKEv
CORE SUPPORT STRUCTURE
Figure 6-136. Schematic representation of the major factors which limit the lifetime of FFTF ducts (after Leggett and Walters, 1993).
restraint systems that reduce or redirect some fraction of component bowing. The core restraint system in FFTF, for instance, performed very well in limiting the overall distortion of the core while maintaining the ability to remove fuel assemblies easily after reactor shutdown (Sutherland, 1984; Hecht and Trenchard, 1990). Another approach involves the periodic rotation of subassemblies that lie across large flux gradients. This has been found to be effective in relaxing the effects of swelling-induced bowing, but to be ineffective in ameliorating the bowing due to relaxation of stresses induced by thermal gradients (Shields, 1981). The lifetime of radial reflector subassemblies in FFTF was found to be extended significantly by rotation, with the relative benefit directly correlatable to the accumulated fluence level when rotation occurred (Hecht and Trenchard, 1990).
514
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
600
400
-
/ / /
s^—^ \ Fe-15Cr-45Ni Neutron Irradiated at 450 °C to 12.5 dpa
Fe-15Cr-45Ni •
Unirradiated\ ^ ^ ^ \ ^
I
i
I
i
20
10
Figure 6-137. Changes observed in room temperature tensile properties of Fe-15Cr-25Ni and Fe-15Cr-45Ni model alloys after irradiation in EBRII to 12.5 dpa at 450 °C (after Garner and Toloczko, 1993).
\
Fe-15Cr-25Ni \
200
1
40
30
Strain, %
INTERMEDIATE 10.7xl0 2 2 n/cm 2 a MECHANISM I
2.8xlO 22 n/cm 2
UNIRRADIATED PLASTIC DIMPLING
CHANNEL FRACTURE o UNIFORM ELONGATION OPROP. ELASTIC LIMIT D YIELD STRENGTH
100
SYMBOLS OPEN - 5C3 CRT CROSSED - 5 A 3 CRT CLOSED-3A1SRT
60
S
EBR-II 304 SS IRRAD. AT 700 °F TESTED AT 700 °F
40
2 § o
ii
20 ELONGATION
I
3
o% 4
5
6
7
8
9
10
11
FLUENCE, x 10 2 2 n/cm 2 ( E > 0.1 MeV)
Figure 6-138. Increase in strength, loss of ductility and change in failure mode observed in annealed 304 after irradiation at 370 °C in EBR-II as control and safety rods thimbles (after Fish et al., 1973).
6.7 Irradiation-Induced Changes in Mechanical Properties
515
cal starting state, both at temperatures where voids do not develop (Fig. 6-140) and at temperatures where voids form relatively easily (Fig. 6-142). The saturation strength level of 316 stainless steel is very dependent on irradiation temperature, however, as shown in Figs. 6-141 to 6-143, reflecting the strong temperature dependences of the densities of voids, dislocation loops and precipitates. Similar temperature-dependent saturation of strength was observed in annealed 1.4988 stainless steel (Ehrlich, 1985). The saturation density of network dislocations was shown earlier to be relatively independent of temperature, however, and to be lower than the level found in typical cold-worked steels before irradiation. Thus, cold-worked steels usually soften at higher irradiation temperatures as the network dislocation density quickly falls to the saturation level during irradiation, especially since the densities of other microstructural components develop at relatively low levels at high temperatures.
6.7 Irradiation-Induced Changes in Mechanical Properties 6.7.1 Austenitic Alloys Prone to Void Swelling
Irradiation-induced embrittlement of stainless steels has long been recognized as one of the major potential limitations of LMR performance (Bloom, 1972; Bagley et al, 1972; Harries, 1979). The irradiation at LMR-relevant temperatures of even the simplest austenitic alloys in the annealed condition usually leads to an increase in strength and a reduction of ductility, as demonstrated in Figs. 6-137 and 6-138. Another characteristic of irradiated austenitic alloys is a loss of plasticity that occurs when the yield strength approaches the ultimate strength, especially at lower irradiation temperatures, as shown in Fig. 6-139. Consistent with the saturation-ofmicrostructure concept introduced in Sect. 6.4, the saturation yield strength is usually independent of the alloy's thermomechani-
1200
_
370 °C
1000
Ultimai e Strength
#
800 Strength MPa
600
y y ^ ^ Yield Strength 400
o 200 -
i
i
i
I
l
I
i
i
I
10x10 2 2 Neutron fluence, E > 0.1 MeV, n cm" 2
Figure 6-139. Saturation of irradiation-induced hardening of annealed AISI 304 L irradiated at 370 °C in EBR-II and tested at 370 °C, showing the plasticity loss that results when the yield strength approaches the ultimate strength early in the irradiation (after Holmes and Straalsund, 1977).
516
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
COLD-WORKED 800 YIELD STRENGTH, MPa SOLUTION-ANNEALED
400
i
Figure 6-140. Evolution of yield strength in M316 stainless steel irradiated at 300 °C in DFR (after Bagley et al., 1972). Voids do not form easily at this temperature.
DFR IRRADIATION 20
30
40
50
DISPLACEMENTS, DPA (N/2)
1100
-
1
i
i
1
1
!
1
T" '
1
1000
900
q-ft o p •
O
^ ***
800
0
-
.427°C
— — ~*?~
700 YIELD STRENGTH MPa
.483°C
0
.—-—
-
600 500
-
a
400
3_538°C u_593°C
0
300
704° c
200
>C
100 ^ - « _ 8 1 6 C >C 0
i
i
-
i
1
1
1
1
2 3 4 5 6 7 8 NEUTRON FLUENCE, n/cm2 (E>0.1 MeV)
10x1(F
Figure 6-141. Evolution of yield strength in AISI 316 stainless steel irradiated in EBR-II at temperatures within the void swelling regime (after Garner et al., 1981 c).
The saturation strength level is also sensitive to the displacement rate, as shown in Fig. 6-144, reflecting the dependence of the various microstructural components on this important variable. Similar behavior was also observed in comparisons of strength changes developed in irradiations conducted in the center and peripheral regions of PHENIX (Dupouy et al., 1983). In addition, the transient regime is usually,
but not always, sensitive to the displacement rate, as demonstrated in Fig. 6-145. There is one important exception to the saturation concept as defined earlier, however. If the irradiation temperature is low enough, such that dislocation mobility and the associated self-annihilation of dislocations are strongly limited by the combined action of defect recombination and high densities of small defect clusters induced
6.7 Irradiation-Induced Changes in Mechanical Properties 600
I
I
I
I
400 200
I
I
I
650°C
ANNEALED
0
I
I
I
1
538°C _
600 YIELD STRENGTH MPa
I
-20% COLD-WORKED
517
20% COLD-WORKED
400 200
ANNEALED
0
i—r
i
i
427°C 800 600 400
ANNEALED
200
I
I
I 3
I 4
I
I
I
I 8
I 9 10
Figure 6-142. Influence of temperature and neutron exposure on evolution of yield strength in 20% coldworked AISI 316 irradiated in EBR-II (after Garner et al., 1981 c).
NEUTRON FLUENCE,
Yield Stress MPa
630
-
560
-
490 420
-
770 -
UTS MPa
700
-
630 -
560 -
490
390
450
480
510
540
Irradiation and Test Temperature, C Figure 6-143. Irradiation-induced changes in the tensile properties of PHENIX hexagonal wrappers constructed from cold worked 316Ti stainless steel, showing the tendency of cold worked steels to harden at relatively low temperatures and soften at higher temperatures (after Lippens et al., 1987).
by radiation, then annealed and coldworked microstructures may not approach the same saturation strength level. An example is shown in Fig. 6-146. Also demonstrated in Fig. 6-146 is the almost total loss of plasticity that can occur at relatively low dpa levels ( 0.1 MeV). The break occurred at «32% swelling (after Makenas et al., 1990 b).
_
527
perature in two sibling fuel pins clad with 20% cold-worked D9 stainless steel. Each pin was subjected to a different physical insult, but both failed at the same place. Due to the nature of the flux and temperature profiles of these pins, the temperaturedependent critical swelling level was exceeded at ^460°C and ^ 3 2 % swelling before it was exceeded at other temperatures. Note that the pins did not fail at the peak swelling level of ^ 3 8 % , indicating that it is not only the void volume but its distribution which is the important variable. As stated earlier, it is thought that the void surface area is most important. Based on these failures, Baker et al. (1993) have proposed that exposure limits be placed on D9-clad pins such that 10% diametral strain arising from swelling is the maximum allowed. Steels with very low ( < 12%) nickel levels are even more prone to this problem due to their earlier onset of swelling and their enhanced tendency toward martensitic transformation. Very early in the U.S.
3 x1023
si
Figure 6-162. Conditions under which severe room temperature embrittlement was observed in two sibling D9 clad fuel pins after irradiation in FFTF (after Garner et al., 1993 a). ABOF and EOL denote above-bottom-of-fuel and end-of-life, respectively.
Break on Pin 312
350
0
10
20
30
Location ABOF, in.
40
400
500
EOL Mid-Wall Temperature
600
528
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
LMR program, for example, an AISI 304 stainless steel control rod thimble (with ^ 9 % Ni) was observed to fail in a very brittle manner after irradiation in EBR-II to swelling levels of 7 to 9%. The damage occurred while the thimble was inadvertently but lightly loaded in a bending mode during room temperature handling (Flinn et al, 1973). At that time, the microscopic origin of the failure process was not yet understood, however, and failure was ascribed only to the flow localization phenomenon. While the fracture toughness of various unirradiated stainless steels can be quite different, it appears that all steels studied undergo the same general evolution in toughness during irradiation. Mills (1982, 1987) has shown that three regimes of evolution occur. The first regime involves a low-dose threshold exposure range
(< 1 dpa) where there is no loss of toughness, and the second involves an intermediate exposure range (1-10 dpa) where toughness decreases very 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. As shown in Fig. 6-163, however, 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 Fig. 6-164. At high test temperatures, the fracture mode changes from transgranular to intergranular in nature, reflecting
11070 kJ/mz
Irrad. Temp. = 400-427°C Test Temp. = 427°C 500 • • V • • +
400
316SS "I 304 SS Unnealed Inconel 600 Incone! 800j CF8 SS as-cast 308 SS Weld
300
200
Base Metal
100 -
Figure 6-163. Irradiation-induced evolution of J c fracture toughness in various austenitic steels and welds (after Mills, 1982). 5
10 15 Neutron Exposure, dpa
20
25
6.7 Irradiation-Induced Changes in Mechanical Properties
529
Transgranular Fracture 11.0 to 11.3x1022 n/cm2 (E > 0.1 MeV) 100
i
A • ••
80 _377 A 400 °C
If
> 388 °C
Test Temp. = 538 °C
Fatigue Precracked Zone
60 I
Intergranular Fracture irradiation Temperature
o
I
40 --
20 — ( A r—*** Fatigue Precracked Zone
382 °C * i
200
300
i
i
400 500 Test Temperature, C
i
600
700
Test Temp. = 649 °C
Figure 6-164. Dependence on test temperature of fracture toughness and fracture mode of highly irradiated 20% cold worked 316 (after Huang and Wire, 1983).
the effect of test temperature on both matrix strength and helium embrittlement at grain boundaries. The level of helium needed to promote high temperature embrittlement is not very high, however, and can easily be reached after moderate neutron exposure in alloys with the lowest nickel level. In general, it appears that the influence of helium on mechanical properties is not as pronounced as was originally feared in the early 1970s. As reviewed by Garner and Greenwood (1993), it seems that many of the early experiments on the effects of helium were strongly affected by the influence of uncontrolled variables, primarily temperature, temperature history and displacement rate. In better-controlled experiments, it was shown that the helium generation rate does indeed influence the details of microstructural evolution some-
what, but in most cases the alteration is insufficient to significantly affect the macroscopic mechanical properties. There is still one unresolved area where helium's influence may be significant in stainless steels. As discussed by Mansur and Grossbeck (1988), large helium levels may influence the creep rupture life, especially during very slow strain rate tests. This possibility is supported by data that show that higher failure strains observed in coldworked AISI 316 pressurized tubes at lower stress levels and lower strain rates result from a transition from triple-point wedgecracking to grain boundary cavitation (Gilbert et al., 1987). The triple point cracking to cavitation transition is shifted to lower stresses by irradiation and is expected to be enhanced by higher helium levels. It is much easier to obtain tensile data in irradiated steels than to obtain fracture
530
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
toughness data. Most toughness data are concentrated at irradiation temperatures in the 400-500 °C range and were derived from specimens cut from hexagonal ducts. Since the microstructural origins of the two mechanical properties are presumably identical, attempts have been made to develop tensile-toughness correlations and thereby extend toughness predictions to higher fluence levels and higher irradiation temperatures. Several of these efforts appear to have been reasonably successful (Huang and Wire, 1983; Hamilton etal., 1984). 6.7.2 Swelling-Resistant Alloys The early observations that void swelling could be reduced by the use of higher nickel levels led to the initiation of many studies involving high nickel alloys, especially in the range 30-45 wt.% nickel. It soon became obvious, however, that near-total suppression of swelling was often obtained at the total expense of ductility. For example, in a series of alloys strengthened with y' or y'/y" precipitates, it was shown that severe embrittlement developed after low exposure irradiation 0 1 0 dpa) in HFIR (Yang and Hamilton, 1982; Hamilton and Huang, 1986). The embrittlement increased at higher temperatures, usually reaching zero ductility at ^600°C. It was shown that the embrittlement arose primarily from radiation-induced formation of y' phase on the grain boundaries, with the relatively large levels of helium formed in HFIR playing only a second-order role. This conclusion is supported by the observation that similar microstructural developments and concurrent embrittlement occurred in other highnickel alloys irradiated in LMRs at much lower helium generation rates (Yang et al., 1985).
One of the most swelling-resistant highnickel alloys studied was Inconel 706, but it was found in both the French and U.S. LMR materials programs to become progressively embrittled, quickly reaching zero ductility (Huang and Fish, 1985; Bajaj et al., 1981; Vaidyanathan et al., 1982; Cauvin et al., 1987). The embrittlement appeared to be specific to high test temperatures ( > 500 °C) rather than to any specific range of irradiation temperature. Microscopy examination showed that swelling was not involved in the embrittlement of this alloy, but radiation-induced precipitation was clearly demonstrated to be responsible. Not only did irradiation significantly enhance precipitation on y' and y" within the grains, but it caused extensive segregation of nickel and other solutes on the grain boundaries. This, in turn, leads to precipitation of y' and T]phase on the grain boundaries (Thomas, 1980; Bell and Lauritzen, 1982; Le Naour etal., 1987; Yang and Makenas, 1985). While the interior of the grains was hardened, the grain boundaries were weakened, yielding a nil ductility trough in the vicinity of 500° to 600 °C. A typical consequence of such embrittlement is shown in Fig. 6-165. Another high-nickel alloy that swells only a moderate amount ( < 5 % ) up to very high dpa levels is Nimonic PE16. This alloy is also subject to radiation-induced segregation, as shown earlier in Fig. 6-17, but not quite as prone to segregation-induced embrittlement at a given temperature. Yang (1982) showed that in unirradiated PE16, a nil ductility trough occurred in the vicinity of 700 °C and corresponded to a similar trough observed in pure Ni3(Al,Ti) single crystals. In irradiated PE16, the trough gradually shifts down toward lower temperatures. The ductility degradation in PE16 was attributed by
6.7 Irradiation-Induced Changes in Mechanical Properties
531
Figure 6-165. Failure of a fuel pin constructed from annealed Inconel 706 irradiated in EBRII (after Yang and Makenas, 1985). The failure occurred between two positions operating at447 o C(5.5xl0 2 2 n/cm 2 ) and526 o C(6.1xl0 2 2 n/cm 2 (E > 0.1 MeV)).
Yang to brittle cleavage failure of the relatively thick y' layer formed on the grain boundaries during irradiation. While not used in the U.S. program due to its lack of irradiation and fabrication experience, PE16 was better known in the U.K. and was selected for service in PFR, where « 70 % of the core loading consisted of fuel pins clad in PE16 (Lippens et al. 1987; Walker and Lightowlers, 1990). This alloy maintained considerable strength in PFR even at the highest dpa levels, but its ductility decreased somewhat with increasing irradiation temperature. The ductility loss was not sufficient, however, to significantly impact its overall performance in PFR (Brown et al., 1991). A loss of toughness, accompanied by channel fracture, has also been observed in PE16. In this case, however, the loss does not arise from void swelling, but rather from segregation and precipitation (Little, 1987). Ferritic steels have been found to be very resistant to void formation, with swelling much less than 1 % at exposures in excess of 100 dpa (Gilbon et al., 1989; Gelles, 1990 a, b; Seran et al., 1992; Little, 1993). As shown earlier in Fig. 6-158, the toughness and tearing modulus of HT9 is maintained at higher fluence, presumably due to its resistance to swelling. The major materials problem in ferritic steels is the radiation-induced elevation in ductile-to-
brittle transition temperature (DBTT) and a concurrent drop in the upper shelf energy (Hu and Gelles, 1983, 1987; Gilbon et al., 1989; Seran et al., 1992). As demonstrated in Figs. 6-166 and 6-167, this shift is very dependent on irradiation temperature. To date, however, HT9 has served very well in FFTF without significant problems. This is largely a consequence of the fact that the shift in DBTT tends to quickly saturate in HT9, as demonstrated in Fig. 6-167, keeping the transition temperature well below the fuel handling temperature, the lowest in-core temperature experienced by LMR core components. Similar radiation-induced shifts in DBTT have been observed in other ferritic steels, such as 12Cr-lMoVW (Klueh and Alexander, 1992) and 9Cr-lMo (Hu and Gelles, 1987). The shift in DBTT is most pronounced at lower irradiation temperatures and reflects the strong temperature dependence of irradiation hardening, as demonstrated in Figs. 6-168 and 6-169. Powell et al. (1986) also showed that the ADBTT of HT9 saturated, with very little change at higher irradiation temperatures, as shown in Fig. 6-170. A similar dependence of room temperature hardness on irradiation temperature was observed in 10-12% Cr ferritic-martensitic steels by Little and Stoter (1982). There is still no general agreement on the origin of the embrittlement, although precipitation is
532
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
50
1
'
HT-9 HAZ Tj = 390°C—«• — 500°C
35
200
0 50 100 TEST TEMPERATURE (°C)
*
_
a / 30 _ 25 #
/ 20 _ /
15 _
/
Figure 6-166. Total fracture energy observed in HT9 base metal and weldments after irradiation in EBR-II to 13 dpa at various irradiation temperatures (after Hu and Gelles, 1983).
rQTXM 10 / 5 i 100
I
0 100
0
-50
50
i 150
i 200
TEST TEMPERATURE (°C)
FFTF Refueling Temperature-
200
Irradiated at 380°-400°C
160
w
DBTT
o 120
•
80
f
40 n
/
405°-430°C [
1
i
I
2
I
3
I
4
I
i
I
5
•
D
6
7
8
Neutron Fluence, 1 0 2 2 n/cm 2 (E>0.1 meV)
i 9
Figure 6-167. Saturation observed in the DBTT shift of HT9 after neutron irradiation (unpublished data courtesy of W. L. Hu, Westinghouse Hanford Company).
533
6.8 Summary 1
1
150
i
iu 200 DC D DC LU
Q- 160
-
u
•EMPERA1
Q
O
t 120 Z DC IUJ _J
80
-
-il
^RANGE OF ^ ^ ^ ^ ALL DATA
-
L
t
DC 00
0
t
10
L
O 40 "
O D O 300
1• I
400
500
600
I
1
1
700
Figure 6-168. Dependence of DBTT on irradiation temperature for HT9 irradiated in FFTF (unpublished data courtesy of W. L. Hu, Westinghouse Hanford Company).
2.0 HT9
Test Temperature 205-232°C " Fluences: 1.7 to 16.5x10 22 n/cm 2 _
.2
(ACO 1 Duct)
1.6
-
O 1.4
-
0 1.2
(f) £ 1.0
O
O°
,8 0.8"
30
Figure 6-170. Shift in ductile to brittle transition temperature in HT9 as a function of neutron fluence and irradiation temperature (after Powell et al., 1986).
IRRADIATION TEMPERATURE, °C
1.8"
20 Fluence (dpa)
o\ O ®° -
0.6 300 400 500 600 700 Irradiation Temperature (°C)
Figure 6-169. Dependence of yield strength on irradiation temperature for HT9 irradiated in FFTF (after Huang, 1992 b).
probably most important. The relative role of helium in embrittlement of these steels is still being debated. The radiation-induced evolution of the tensile and impact properties of various ferritic-martensitic steels has been examined in terms of their microstructural and microchemical changes. The results show many similarities with the behavior of austenitic steels (Agueev et al., 1989; Gilbon etal., 1989; Maziasz and Klueh, 1992; Dvoriashin et al., 1992).
6.8 Summary The operating environment of LMRs subjects its structural materials to unprecedented levels of phase alteration, dimensional change and modification of mechanical properties. While many of the operating phenomena have often been studied as separate entities, it is apparent in hindsight that all of the various phenomena participate in an intimately linked co-evolution driven by the production of large supersaturations of vacancies and in-
534
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
terstitials. While each of these various processes can in themselves limit the performance and lifetime of core components, it is the void swelling phenomenon that eventually becomes the dominant determinant of phase stability, irradiation creep and mechanical property changes. The major strategy, therefore, for extending component lifetimes and insuring safe operating conditions requires that swelling be either reduced in rate or delayed until some other nonstructural consideration such as fuel burn-up determines the lifetime. The various national fast reactor programs have all succeeded in this goal, either by the use of ferritic or high nickel austenitic alloys, or by delaying the onset of void swelling in conventional stainless steels via compositional and thermomechanical modifications. Those few remaining problems that arise from gradients in environmental variables can be ameliorated either by innovative designs that reduce the impact of swelling and irradiation creep, or by operational procedures such as duct rotation.
6.9 References Agueev, V. S., Bykov, V. N., Dvoryashin, A. M., Golovanov, V. N., Medvedeva, E. A., Romaneev, V. V., Shamardin, V. K., Vorobiev, A. N. (1989), in: 14th Int. Symp. on Effects of Radiation on Materials, Vol. 1, STP1046. Philadelphia, PA: ASTM, pp. 98-113. Anantatmula, R. P. (1984), J. Nucl. Mater. 125, 170181. Anderson, K. R., Garner, F. A., Stubbins, J. F. (1992), Effects of Radiation on Materials: 15 th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 835-845. Appleby, W. K., Hilbert, R. R, Bailey, R. W. (1972), Proc. Conf. Irradiation Embrittlement and Creep in Fuel Cladding and Core Components. London: British Nuclear Energy Society, pp. 209-216. Appleby, W. K., Bloom, E. E., Flinn, J. E., Garner, F. A. (1977), Proc. Int. Conf Radiation Effects in Breeder Reactor Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 509-527.
Azam, N., DeLaplace, X, LeNaour, L. (1979), in: Proc. Conf. Irradiation Behavior of Metallic Materials for Fast Reactor Core Components, Ajaccio, Corsica: Poirier, X, Dupouy, X M. (Eds.). Paris: Commissariat a l'Energie Atomique, pp. 177-182. Bagley, K. Q., Barnaby, X W, Fraser, A. S. (1972), Proc. Conf. Irradiation Embrittlement and Creep of Fuel Cladding and Core Components. London: British Nuclear Energy Society, pp. 143-153. Bagley, K. Q., Little, E. A., Levy, V., Alamo, A., Ehrlich, K., Anderko, K., Calza Bini, A. (1987), Proc. Int. Conf Materials for Nuclear Reactor Core Applications, Vol. 2, Bristol. London: British Nuclear Energy Society, pp. 37-46. Bajaj, R., Shogan, R. P., DeFlitch, C , Fish, R. L., Paxton, M. M., Bleiberg, M. L. (1981), Effects of Radiation on Materials: 10th Inter. Symp., STP 725. Philadelphia: ASTM, pp. 326-351. Baker, R. B., Bard, F. E., Leggett, R. D., Pitner, A. L. (1993), /. Nucl. Mater. 204, 109-118. Banerjee, B. R., Dulis, E. X, Hauser, X X (1968), Trans. ASM 61, 103. Barmore, W., Ruotola, A., Raymond, E., Mukherjee, A. (1983), J. Nucl. Mater. 117, 258-263. Baron, X L., Cadalbert, R., Delaplace, X (1974), J. Nucl. Mater. 51, 266-268. Barton, P. X, Eyre, B. L., Stow, D. A. (1977), J. Nucl. Mater. 67, 181-197. Bates, X F. (1975), Properties of Reactor Structural Alloys after Neutron or Charged Particle Irradiation, STP 570. Philadelphia: ASTM, pp. 369-387. Bates, X F. (1981), J. Nucl. Mater. 98, 71-85. Bates, X F , Gelles, D. S. (1978), J. Nucl. Mater. 71, 365-367. Bates, X F , Gilbert, E. R. (1975), /. Nucl. Mater. 59, 96. Bates, J. F , Gilbert, E. R. (1978), /. Nucl Mater. 71, 286. Bates, X F, Johnston, W. G. (1977), Proc. Int. Conf Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 625-644. Bates, X F, Powell, R. W. (1981), J. Nucl. Mater. 102, 200-213. Bates, X F , Paxton, M. M., Straalsund, X L. (1973), Nucl. Technol. 20, 134. Bates, X F , Powell, R. W, Gilbert, E. R. (1981a), Effects of Radiation on Materials: 10th Int. Symp., STP 725. Philadelphia: ASTM, pp. 713-734. Bates, X F , Cummings, W. V., Gilbert, E. R. (1981 b), /. Nucl. Mater. 99, 75-84. Bates, X F , Garner, F A., Mann, F. M. (1981c), /. Nucl. Mater. 103-104, 999-1004. Bell, W. L., Lauritzen, T. (1982), Effects of Radiation on Materials: Eleventh Int. Symp., STP 782. Philadelphia: ASTM, pp. 139-151. Bloom, E. E. (1972), Proc. Conf Irradiation Embrittlement and Creep in Fuel Cladding and Core Components. London: British Nuclear Energy Society, pp. 93-102.
6.9 References
Boltax, A. (1992), Fusion Technology 21, 1921-1926. Boulanger, L., Le Naour, L., Levy, V. (1983), Proc. Conf. Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys, Vol. 1, Brighton. London: British Nuclear Energy Society, pp. 1-4. Boutard, J. L., Brun, G., Lehmann, X, Seran, J. L., Dupouy, J. M. (1977), in: Proc. Conf. Irradiation Behavior of Metallic Materials for Fast Reactor Core Components, Ajaccio, Corsica: Poirier, X, Dupouy, X M. (Eds.). Paris: Commissariat a l'Energie Atomique, pp. 137-144. Boutard, X L., Carteret, Y, Cauvin, R., Guerin, Y, Maillard, A. (1983), Proc. Conf. Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 109-112. Brager, H. R., Garner, F. A. (1978), J. Nucl. Mater. 73, 9-19. Brager, H. R., Garner, F. A. (1979), Effects of Radiation on Structural Materials: 9th Inter. Symp., STP 683. Philadelphia: ASTM, pp. 207-232. Brager, H. R., Garner, F. A. (1981), Proc. Symp. Phase Stability During Irradiation, Pittsburgh. Warrendale: The Metallurgical Society of AIME, pp. 219-235. Brager, H. R., Garner, F. A., Gilbert, E. R., Flinn, X E., Wolfer, W. G. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials. New York: The Metallurgical Society of AIME, pp. 727-755. Brager, H. R., Blackburn, L. D., Greenslade, D. L. (1984) /. Nucl. Mater. 122-123, 332-337. Brailsford, A. D., Bullough, R. (1972), J Nucl. Mater. 44, 121. Brailsford, A. D., Bullough, R. (1973), Philos. Mag. 27, 49. Bramman, X L, Brown, C , Watkin, X S., Cawthorne, C , Fulton, E. X, Barton, P. X, Little, E. A. (1977), Proc. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 479-507. Bridges, A. E., Waltar, A. E., Leggett, R. D., Baker, R. B., Gneiting, B. C. (1991), Proc. Int. Conf. Fast Reactor and Related Fuel Cycles, Kyoto, Japan, Vol. III. Atomic Energy Society of Japan, pp. PI.21-1 Pl.21-10. Brown, C , Fulton, E. X (1979), U.S./U.K. Seminar on Cladding and Duct Materials, CONF-7910113. Richland, WA: U.S. DOE, pp. 11-29. Brown, C , Linekar, G. A. B. (1974), Dislocation Densities in ST and 20% Cold Worked Type M316 Pin Cladding Irradiated to 30 dpa in DFR, UKAEA Report TRG-M-6571. London: U.K. Atomic Energy Authority. Brown, C , Linekar, G. A. B. (1987), Proc. Int. Conf. Materials for Nuclear Reactor Core Applications, Vol. 2, Bristol. London: British Nuclear Engineering Society, pp. 219-222. Brown, C , Fulton, E. X, Watkin, X S., Mazey, D. X, Hudson, X A. (1977), Report on U.S./U.K. Clad-
535
ding/Duct Materials Specialists Group Meeting, Harwell, U.K. Rep. TC-1028, Part II. Richland, WA: Hanford Engineering Development Laboratory. Brown, C , Sharpe, R. M., Fulton, E. X, Cawthorne, C. (1983), Proc. Conf. Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 63-67. Brown, C , Levy, V., Seran, X-L., Ehrlich, K., Roger, R. X C , Bergman, H. (1991), Proc. Int. Conf. Fast Reactor and Related Fuel Cycles, Kyoto, Japan, Vol. I, pp. 7.5.1-7.5.10. Bullough, R., Haynes, M. R. (1977), J. Nucl. Mater. 68, 286. Bullough, R., Perrin, R. C. (1972), Proc. Int. Conf. on Radiation-Induced Voids in Metals, AEC Symp. Series 26, Albany, pp. 769-797. Bullough, R., Wood, M. H., Little, E. A. (1981), Effects of Radiation on Materials: Tenth Conf, STP 725. Philadelphia: ASTM, pp. 593-604. Bullough, C. K., Jenkins, X K., Williams, T. M. (1987), Proc. Int. Conf. Materials for Nuclear Reactor Core Applications, Bristol. London: British Nuclear Energy Society, pp. 285-291. Bump, T. R., Huebotter, P. R., Phipps, R. D., Strain, R. V. (1973), Argonne National Laboratory Report ANL-7986. Argonne, IL. Busboom, H. X, McClellan, G. C , Bell, W L., Appleby, W K. (1975), "Swelling of Types 304 and 316 Stainless Steel Irradiated to 8 x 1022 n/cm 2 ", General Electric Company Report GEAP-14062. Sunnyvale, CA: General Electric Company. Cauvin, R., Schauff, R., Robouille, O. (1987), Proc. Conf. Materials for Nuclear Core Applications. London: British Nuclear Energy Society, pp. 187190. Cawthorne, C. (1979), Proc. U.K./U.S. Fast Reactor Exchange Meeting on Cladding and Duct Materials. Richland, WA: U.S. DOE. Cawthorne, C , Fulton, X E. (1967), Nature 216, 515517. Challenger, K. D., Lauritzen, T. (1975), Properties of Reactor Structural Materials after Neutron or Particle Irradiation, STP 570. Philadelphia: ASTM, pp. 388-403. Chin, B. A., Straalsund, X L. (1978), J. Nucl. Mater. 74, 260-266. Clark, R. W, Kumar, A. S., Garner, F. A. (1984), J. Nucl. Mater. 155-157, 845-849. Coghlan, W A. (1986), Int. Metals Rev. 31, 245-257. Coghlan, W A., Garner, F. A. (1987), Radiation-Induced Changes in Microstructure, Part 1: 13th Int. Symp., STP 955. Philadelphia: ASTM, pp. 289314. De Raedt, Ch. (1982), Proc. Conf. Fast, Thermal and Fusion Reactor Experiments, Salt Lake City. La Grange Park, IL: ANS, p. 226. Dubuisson, P., Maillard, A., Delalande, C , Gilbon, D., Seran, X L. (1992), Effects of Radiation on Ma-
536
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
terials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 995-1014. Duncan, D. R., Panayotou, N. R, Wood, W. L. (1981), ANS Trans. 38, 265. Dupouy, J.-M., Erler, X, Huillery, R. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 83-93. Dupouy, J. M., Lehmann, X, Boutard, X L. (1978), Proc. Conf. Reactor Materials Science, Vol. 5. Alushta USSR. Moscow: U.S.S.R. Government, pp. 280-296. Dupouy, X M., Sagot, X P., Boutard, X L. (1983), in: Proc. Conf. on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 157 — 160. Dvoriashin, A. M., Dmitriev, V. D., Khabarov, V. S. (1992), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia, PA: ASTM, pp. 1180-1189. Ehrlich, K. (1981), /. Nucl. Mater. 100, 149-166. Ehrlich, K. (1985), /. Nucl. Mater. 133-134, 119126. Ehrlich, K., Calza Bini, A., Levy, V., Brown, C , Lippens, M. (1986), Proc. Int. Conf. Reliable Fuels for Liquid Metal Reactors, Tuscon. LaGrange Park, IL: ANS, pp. 4.1-4.16. Elen, X D., Fenici, P. (1992), J. Nucl. Mater. 191-194, 766-770. Esmailzadeh, B., Kumar, A. S. (1985), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 468-480. Esmailzadeh, B., Kumar, A. S., Garner, F. A. (1985),
/. Nucl. Mater. 133-134, 590-594. Ethridge, X L., Waltar, A. E. (1989), Trans. ANS 60, 286-287. Fish, R. L., Straalsund, X L., Hunter, C. W., Holmes, X X (1973), Proc. Symp. Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys, STP 529. Philadelphia: ASTM, pp. 149-164. Fissolo, A., Cauvin, R., Hugot, X P., Levy, V. (1990), Effects of Radiation on Materials: 14th Int. Symp., Vol. II, STP 1046. Philadelphia: ASTM, pp. 700713. Flinn, X E., Kenfield, T. A. (1976), Proc. Workshop on Correlation of Neutron and Charged Particle Damage, CONF-760673. Oak Ridge, TN: U.S. Dept. of Energy, pp. 253-290. Flinn, X E., Krajcinovic, D., Phipps, R. D., Franklin, D. G., Miller, S. C. (1973), Evaluation of Ex-Reactor Loading Event of High-Fluence EBR-II ControlRod Thimble 5E3, ANL/EBR-068. Argonne, IL: Argonne National Laboratory. Foster, X P., Boltax, A. (1980), Nucl. Technol 47,181. Foster, X P., Wolfer, W. G., Biancheria, A., Boltax, A. (1972), Proc. Conf. Irradiation Embrittlement and Creep in Fuel Cladding and Core Components. London: British Nuclear Energy Society, pp. 273-281.
Foster, X P., Houghtaling, T. K., Boltax, A. (1993), J. Nucl. Mater. 202, 275-285. Fujiwara, M., Uchida, H., Ohta, S., Yuhara, S., Tani, S., Sato, Y. (1987), Radiation Induced Changes in Microstructure: 13th Int. Symp., STP 955. Philadelphia: ASTM, pp. 127-145. Garafalo, F., Wriedt, H. A. (1962), Ada Metall. 10, 1007. Garner, F. A. (1981a), Proc. Symp. Phase Stability During Irradiation, Pittsburgh. Warrendale: The Metallurgical Society of AIME, pp. 165-189. Garner, F. A. (1981b), Damage Analysis and Fundamental Studies Quarterly Progress Report DOE/ ER-0046/5. Richland, WA: U.S. DOE, pp. 198218. Garner, F. A. (1983), /. Nucl. Mater. 117, 177-197. Garner, F. A. (1984), /. Nucl. Mater. 122-123, 459471. Garner, F. A. (1985), Proc. Symp. Optimizing Materials for Nuclear Applications. Warrendale: The Metallurgical Society of AIME, pp. 111-139. Garner, F. A. (1993), J. Nucl. Mater. 205, 98-117. Garner, F. A., Brager, H. R. (1985a), /. Nucl. Mater.
133-134,511-514. Garner, F. A., Brager, H. R. (1985 b), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 187-201. Garner, F. A., Brager, H. R. (1985 c), Proc. Symp. Optimizing Materials for Nuclear Applications, Los Angeles. Warrendale: The Metallurgical Society of AIME, pp. 87-109. Garner, F. A., Brager, H. R. (1988), /. Nucl. Mater. 155-157, 833-837. Garner, F. A., Edwards, D. X (1993), Fusion Reactor Materials Semiannual Progress Report DOE/ER0313/14. Richland, WA: U.S. DOE, pp. 125-126. Garner, F. A., Gelles, D. S. (1988), /. Nucl. Mater. 159, 286-309. Garner, F. A., Gelles, D. S. (1990), Effects of Radiation on Materials: 14th Int. Symp., Vol. II, STP 1046. Philadelphia: ASTM, pp. 673-683. Garner, F. A., Greenwood, L. R. (1992), Fusion Reactor Materials Semiannual Progress Report DOE/ ER-0313/12. Oak Ridge, TN: U.S. DOE, pp. 5458. Garner, F. A., Greenwood, L. R. (1993), Mater. Trans., JIM 34, 985-998. Garner, F A., Grossbeck, M. L. (1994), Fusion Materials Semiannual Progress Report DOE/ER-0313/ 16. Oak Ridge, TN: U.S. DOE, in press. Garner, F. A., Guthrie, G. L. (1975), Proc. Conf. on Radiation Effects and Tritium Technology for Fusion Reactors, CONF-750989, Vol. 1. Gatlinburg, TN: U.S. DOE, pp. 491-518. Garner, F. A., Kumar, A. S. (1987), Radiation-Induced Changes in Microstructure: 13th Int. Symp., Part 1, STP 955. Philadelphia: ASTM, pp. 289314. Garner, F A., Laidler, X X (1976), Proc. Workshop on Correlation of Neutron and Charged Particle Dam-
6.9 References
age, CONF-760673. Oak Ridge, TN: Oak Ridge National Laboratory, pp. 177-240. Garner, F. A., McCarthy, I M. (1990), Proc. Symp. on Reduced Activation Materials for Fusion Reactors, STP 1047. Philadelphia: ASTM, pp. 19-29. Garner, F. A., Mitchell, M. A. (1992), /. Nucl. Mater. 187, 223-229. Garner, F. A., Oliver, B. M. (1993), Fusion Reactor Materials Semiannual Progress Report DOE/ER0313/14. Oak Ridge, TN: U.S. DOE, pp. 47-50. Garner, F. A., Porter, D. L. (1982), Effects of Radiation on Materials: 11th Conf, STP 782. Philadelphia: ASTM, pp. 295-309. Garner, F. A., Porter, D. L. (1984), Proc. Conf. on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 41 -44. Garner, F. A., Porter, D. L. (1988), /. Nucl. Mater. 155-157, 1006-1013. Garner, F. A., Puigh, R. J. (1991), J. Nucl. Mater. 179-181, 577-580. Garner, F. A., Thomas, L. E. (1973), Effects of Radiation on Substructure and Mechanical Properties of Metals and Alloys, STP 529. Philadelphia: ASTM, pp. 303-325. Garner, F. A., Toloczko, M. B. (1992), Fusion Reactor Materials Semiannual Progress Report DOE/ER0313/12. Oak Ridge, TN: U.S. DOE, pp. 145-147. Garner, F. A., Toloczko, M. B. (1993), J. Nucl. Mater. 206, 230-248. Garner, F. A., Wolfer, W. G. (1981), J. Nucl. Mater. 102, 143-144. Garner, F. A., Wolfer, W. G. (1982 a), Effects of Radiation on Materials: 11th Conf, STP 782. Philadelphia: ASTM, pp. 1073-1087. Garner, F. A., Wolfer, W. G. (1982b), University of Wisconsin Fusion Materials Report UWFDM-484. Garner, F. A., Wolfer, W. G. (1984), J. Nucl. Mater.
122-123,201-206. Garner, F. A., Powell, R. W, Diamond, S., Lauritzen, T, Rowcliffe, A. F , Sprague, J. A., Keefer, D. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 543-569. Garner, F. A., Cummings, W. V., Bates, J. R, Gilbert, E. R. (1978), Densification-Induced Strains in 20% Cold Worked 316 Stainless Steel During Neutron Irradiation, HanfordEngineering Development Laboratory Report HEDL-TME-78-9. U.S. DOE. Garner, F. A., Wolfer, W. G., Brager, H. R. (1979), Effects of Radiation on Structural Materials, STP 683. Philadelphia: ASTM, pp. 160-183. Garner, F. A., Gilbert, E. R., Porter, D. L. (1981a), Effects of Radiation on Materials: 10th Conf, STP 725. Philadelphia: ASTM, pp. 680-697. Garner, F. A., Gilbert, E. R., Gelles, D. S., Foster, J. P. (1981b), Effects of Radiation on Materials: 10th Conf, STP 725. Philadelphia: ASTM, pp. 698712.
537
Garner, F. A., Hamilton, M. L., Panayotou, N. R, Johnson, G. D. (1981c), J. Nucl. Mater. 103-104, 803-808. Garner, F. A., Hunter, C. W, Johnson, G. D., Lippincott, E. P., Schiffgens, J. O. (1982), Nucl. Technol. 58, 203-217. Garner, F. A., Brager, H. R., Gelles, D. S., McCarthy, J. M. (1987a), /. Nucl. Mater. 148, 294301. Garner, F. A., Porter, D. L., Makenas, B. J. (1987 b), J. Nucl. Mater. 148, 279-287. Garner, R A., Heinisch, H. L., Simons, R. L., Mann, R M. (1990), Radiation Effects and Defects in Solids 113, 229-255. Garner, R A., Gelles, D. S., Takahashi, H., Ohnuki, S., Kinoshita, H., Loomis, B. A. (1992a), J. Nucl. Mater. 191-194, 948-951. Garner, R A., Hamilton, M. L., Shikama, T, Edwards, D. J., Newkirk, J. W. (1992b), J. Nucl. Mater. 191-194, 386-390. Garner, F. A., Bates, J. R, Mitchell, M. A. (1992c), /. Nucl. Mater. 189, 201-209. Garner, F. A., Toloczko, M. B., Puigh, R. J. (1992d), Fusion Reactor Materials Semiannual Progress Report, DOE/ER-0313/12. Oak Ridge, TN: U.S. DOE, pp. 148-162. Garner, F. A., Chastain, S. A., Makenas, B. J. (1993 a), Fusion Reactor Materials Semiannual Progress Report DOE/ER-0313/15. Oak Ridge, TN: U.S. DOE, in press. Garner, R A., Lauritzen, T, Mitchell, M. A. (1993 b), Effects of Radiation on Materials, 16th Int. Symp., STP 1175. Philadelphia: ASTM, pp. 803-815. Garner, F. A., Hamilton, M. L., Eiholzer, C. R., Toloczko, M. B., Kumar, ,A. S. (1993 c), Effects of Radiation on Materials: 16th Int. Symp., STP 1175. Philadelphia: ASTM, pp. 696-713. Garner, F. A., Hamilton, M. L., Greenwood, L. R., Stubbins, J. P., Oliver, B. M. (1993 d), Effects of Radiation on Materials, 16th Int. Symp., STP 1175. Philadelphia: ASTM, pp. 921-939. Garner, F. A., Miyahara, K., Newkirk, J. W, Kinoshita, H. (1993 e), / Nucl. Mater. 199, 132142. Garner, F. A., Sekimura, N., Grossbeck, M. L., Ermi, A. M., Newkirk, J. W, Watanabe, H., Kiritani, M. (1993f), / Nucl. Mater. 205, 206-218. Garner, F. A., Toloczko, M. B., Woo, C. H. (1994), Fusion Materials Semiannual Progress Report DOE/ER-0313/16. Oak Ridge, TN: U.S. DOE, in press. Gelles, D. S. (1984), /. Nucl. Mater. 122-123, 207210. Gelles, D. S. (1990a), Metall. Trans. 21A, 10651071. Gelles, D. S. (1990 b), Effects of Radiation on Materials: 14th Int. Symp., Vol. 1, STP 1046. Philadelphia: ASTM, pp. 73-97. Gelles, D. S. (1993), /. Nucl. Mater. 205, 146-161.
538
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Gelles, D. S., Thomas, L. E. (1983), Proc. Topical Conf. on Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird. Warrendale: The Metallurgical Society of AIME, pp. 559-568. Gelles, D. S., Garner, F. A., Brager, H. R. (1981), Effects of Radiation on Materials: 10th Int. Symp., STP 725. Philadelphia: ASTM, pp. 735-753. Gilbert, E. R. (1971), Reactor Technology 14, 258285. Gilbert, E. R., Blackburn, L. D. (1970), Proc. 2nd Int. Conf Strength of Metals and Alloys, Pacific Grove, CA. Metals Park, OH: ASM, pp. 773-777. Gilbert, E. R., Chin, B. A. (1978), Trans. ANS 28, 141-142. Gilbert, E. R., Chin, B. A. (1981), Effects of Radiation on Materials: 10th Int. Symp., STP 725. Philadelphia: ASTM, pp. 665-679. Gilbert, E. R., Kaulitz, D. C , Holmes, I X, Claudson, T. T. (1972), Proc. Conf. Irradiation Embrittlement and Creep in Fuel Cladding and Core Components. London: British Nuclear Energy Society, pp. 239-251. Gilbert, E. R., Chin, B. A., Duncan, D. R. (1987), Metal. Trans. ISA, 79-84. Gilbon, D., Seran, X L., Cauvin, R., Fissolo, A., Alamo, A., LeNaour, R, Levy, V. (1989), in: Effects of Radiation on Materials: 14th Int. Symp., STP 1046. Philadelphia: ASTM, pp. 5-34. Gittus, J. H. (1972), Phil Mag. 25, 345. Glowinski, L. D., Lanore, J. M., Fiche, C , Adda, Y. (1976), J. Nucl. Mater. 61, 41-52. Grabova, R. B., Goncharenko, Yu. D., Gorbatov, V. K., Kosenkov, V. M., Losev, N. P., Rogozyanov, A. Ya., Samsonov, B. V. (1986), Fiz. Metal. Metalloved. 61, 1164-1169. Greenwood, L. R. (1983), J. Nucl. Mater. 115, 137142. Greenwood, L. R., Garner, F. A., Heinisch, H. L. (1993), J. Nucl. Mater. 191-194, 1096-1100. Grossbeck, M. L., Horak, J. A. (1988), J. Nucl. Mater. 155-157, 1001-1005. Grossbeck, M. L., Mansur, L. K., Tanaka, M. P. (1989), Effects of Radiation on Materials, 14th Int. Symp., STP 1046, VolII. Philadelphia: ASTM, pp. 537-550. Hales, X W. (1978), Trans. ANS 28, 153-155. Hall, M. M. (1978), Trans. ANS 28, 146-147. Hamilton, M. L., Huang, F. H. (1986), Proc. Symp. Use of Small-Scale Specimens for Testing Irradiated Material, STP 888. Philadelphia: ASTM, pp. 5 16. Hamilton, M. L., Johnson, G. D., Hunter, C. W., Duncan, D. R. (1981), Proc. Conf. Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys, Vol. 1. London: British Nuclear Energy Society, pp. 211-214. Hamilton, M. L., Garner, F. A., Wolfer, W. G. (1984), J. Nucl. Mater. 122-123, 106-110. Hamilton, M. L., Huang, F. H., Yang, W. X S., Garner, F. A. (1987), Influence of Radiation on Material
Properties: 13th Int. Symp., Part II, STP 956. Philadelphia: ASTM, pp. 245-270. Hamilton, M. L., Johnson, G. D., Puigh, R. X, Garner, F. A., Maziasz, P. X, Yang, W. X S., Abraham, N. (1989), Proc. Symp. Residual and Unspecified Elements in Steel, STP 1042. Philadelphia: ASTM, pp. 124-149. Harbottle, X E. (1977), /. Nucl. Mater. 66, 258-265. Harbottle, X E., Silvent, A. (1979), Proc. Conf. Irradiation Behavior of Metallic Materials for Fast Reactor Core Components, Ajaccio, Corsica. Paris: Commissariat a l'Energie Atomique, pp. 391 — 397. Harkness, S. D., Li, C.-Y. (1971), Metall. Trans. 2, 1457. Harkness, S. D., Kestel, B. X, Franklin, D. G. (1970), Reactor Development Program Progress Report ANL-775. Argonne, IL: Argonne National Laboratory, pp. 5-6. Harries, D. R. (1977), /. Nucl. Mater. 65, 157-193. Harries, D. R. (1979), /. Nucl. Mater. 82, 2-21. Harries, D. R. (1981), Proc. Conf Mechanical Behavior and Nuclear Applications of Stainless Steel at Elevated Temperatures, Varese, Italy. London: The Metals Society, pp. 1-14. Hassan, M. H., Blanchard, X P., Kulcinski, G. L. (1992), Stress-Enhanced Swelling: Mechanisms and Implications for Fusion Reactors, University of Wisconsin Fusion Technology Report UWFDM-901. Madison, WI: University of Wisconsin. Hausen, H., Schule, W, Cundy, M. R. (1988), Fusion Technology 88, 905-909. Hecht, S. L., Trenchard, R. G. (1990), Proc. Int. Conf Fast Reactor Core and Fuel Structural Behavior, Inverness. London: British Nuclear Energy Society, pp. 275-282. Herschbach, K., Bergmann, H.-X (1990), /. Nucl. Mater. 175, 143-146. Herschbach, K., Schneider, W., Bergmann, H.-X (1990), Effects of Radiation on Materials: 14th Int. Symp., Vol. II, STP 1046. Philadelphia: ASTM, pp. 570-587. Herschbach, K., Schneider, W, Ehrlich, K. (1993), /. Nucl. Mater. 203, 233-248. Higginson, P. R., Lilley, R. X (1990), Proc. Conf. Fast Reactor Core and Fuel Structural Behavior. London: British Nuclear Energy Society, pp. 307-314. Hofman, G. L., Truffert, X, DuPouy, X-M. (1977), /. Nucl. Mater. 65, 200-203. Holmes, X X, Straalsund, X L. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 53-63. Hosoi, Y, Wade, N. (1984), Proc. Japan-France Seminar on Fundamental Aspects of Mechanical Properties and Microstructure Evolution of Stainless Steels at High Temperature, Tokyo. University of Tokyo, pp. 55-62. Hoyt, X X, Garner, F. A. (1991a), / Nucl. Mater. 179-181, 1096-1099.
6.9 References
Hoyt, J. I , Garner, F. A. (1991b), Fusion Reactor Materials Semiannual Progress Report DOE/ER0313/9. Oak Ridge, TN: U.S. DOE, pp. 81-86. Hu, W. L., Gelles, D. S. (1983), Proc. Topical Conf. on Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird. Warrendale: The Metallurgical Society of AIME, pp. 631-645. Hu, W. L., Gelles, D. S. (1987), Influence of Radiation on Mechanical Properties: 13th Int. Symp., Part II, STP 956. Philadelphia: ASTM, pp. 83-97. Huang, R H. (1992 a), Engineering Fracture Mechanics 43, 733-748. Huang, F. H. (1992 b), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 1267-1286. Huang, F. H., Fish, R. L. (1985), Effects of Radiation on Materials: Twelfth Int. Symp., STP 870. Philadelphia: ASTM, pp. 720-731. Huang, F. H., Wire, G. L. (1983), Proc. Conf Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 135-138. Hunter, C. W, Johnson, G. D. (1979), Proc. Int. Conf. Fast Breeder Reactor Fuel Performance, Monterey. ANS/AIME, pp. 478-488. Itaki, T, Yuhara, S., Shibahara, I., Kubota, H., Itoh, M., Nomura, S. (1987), Proc. of Int. Symp. Materials for Nuclear Core Applications, Bristol. London: British Nuclear Energy Society, pp. 203-210. Itoh, M., Onose, S., Yuhara, S. (1987), Proc. 13th Int. Symp. Radiation-Induced Changes in Microstructure, Part 1, STP 955. Philadelphia: ASTM, pp. 114-126. Johnston, W. G., Rosolowski, J. H., Turkalo, A. M., Lauritzen, T. (1974), J. Nucl. Mater. 54, 24-40. Johnston, W. G., Lauritzen, T, Rosolowski, J. H., Turkalo, A. M. (1976), Proc. Conf. Radiation Damage in Metals, Cleveland. Metals Park, OH: American Society of Metals, pp. 227-266. Karnesky, R. A. (1980), The Precipitation Kinetics of the Sigma Phase in 20% Cold Worked Type 316 Stainless Steel. Hanford Engineering Development Laboratory Report HEDL-6788. U.S. DOE. Kegg, G. R., Silcock, J. M., West, D. R. F. (1974), Metal Science 8, 337. Khera, S. K., Schwaiger, C , Ullmaier, H. (1980), / Nucl. Mater. 92, 299-305. Klueh, R. L., Alexander, D. J. (1992), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 1256-1266. Krasnoselov, V. A., Prokhorov, V. I., Kolesnikov, A. N., Ostrovskii, Z. A. (1983), Atomnaya Energiya 54(2), 111-114. Krasnoselov, V. A., Kolesnikov, A. N., Prokhorov, V. I (1987), Atomaya Energiya 63, 240-242. Kruglov, A. S., Bul'Kanov, M. E., Bykov, V. N., Pevchikh, Yu. M. (1980), Atomnaya Energiya 48, 258-259. Kumar, A. S., Garner, F. A. (1983), /. Nucl. Mater. 117, 234-238.
539
Kumar, A. S.? Garner, F A. (1985), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 493-506. Lai, J. K. (1983), Materials Science and Engineering 61, 101-109. Lai, J. K., Chastell, D. I, Flewitt, P. E. J. (1981), Proc. Int. Conf on Solid-Solid Phase Transformations. New York: The Metallurgical Society of AIME, pp. 781-785. Laidler, J. J., Mastel, B. (1972), Proc. 13th Annual Meeting of the Electron Microscopy Society of America. Baton Rouge, LA: Claitors, pp. 680-681. Laidler, J. I, Garner, F. A., Thomas, L. E. (1976), Proc. Conf. Radiation Damage in Metals. Metals Park, OH: American Society for Metals, pp. 194-226. Latanison, R. M., Ruff, A. W, Jr. (1971), Metall. Trans. 2, 505-509. Lauritzen, T, Vaidyanathan, S., Bell, W. L., Yang, W J. S. (1987), Radiation-Induced Changes in Microstructure: 13th Int. Symp., Part I, STP 955. Philadelphia: ASTM, pp. 101-113. Le Naour, L., Vouillon, N., Levy, V. (1982), Effects of Radiation on Materials: 11th Int. Symp., STP 782. Philadelphia: ASTM, pp. 310-324. Le Naour, L., Hugon, M. P., Grosjean, P., Maillard, A., Seran, J.-L. (1987), Proc. Int. Conf. Materials for Nuclear Reactor Core Applications, Vol. 2, Bristol. London: British Nuclear Engineering Society, pp. 211-217. Leclere, I , Marbach, G. (1979), Proc. DOE/CEA/ BMFT Exchange Meeting, June 5-14, 1979, Saclay, France. Hanford Engineering Development Laboratory Report TC-1577, Vol. 1. Richland, WA: U.S. DOE. Lee, E. H., Mansur, L. K. (1986), /. Nucl. Mater. 141-143, 695-702. Lee, E. H., Maziasz, P. X, Rowcliffe, A. F. (1981), Proc. Symp. on Phase Stability During Irradiation, Pittsburgh. Warrendale: The Metallurgical Society of AIME, pp. 191-218. Lee, E. H., Mansur, L. K., Rowcliffe, A. F. (1984), /. Nucl. Mater. 122-123, 299-304. Leggett, R. D., Walters, L. C. (1993), J. Nucl. Mater. 204, 23-32. Levy, V., Azam, N., Le Naour, L., Didout, G., DeLaplace, J. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: American Institute of Mining, Metallurgical and Petroleum Engineers, pp. 709-725. Lewthwaite, G. W, Mosedale, D. (1980), J. Nucl. Mater. 90, 205-215. Lippens, M., Ehrlich, K., Levy, V., Brown, C , Calza Bini, A. (1987), Proc. Conf. on Materials for Reactor Core Applications, Bristol. London: British Nuclear Energy Society, pp. 177-182. Little, E. A. (1983), Proc. Conf. on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 139-142.
540
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Little, E. A. (1987), Proc. Int. Conf. Materials for Nuclear Reactor Core Applications, Vol. 2, Bristol. London: British Nuclear Energy Society, pp. 4 7 55. Little, E. A. (1993), 1 Nucl. Mater. 206, 324-334. Little, E. A., Stoter, L. P. (1982), Effects of Radiation on Materials: Eleventh Conf, STP 782. Philadelphia: ASTM, pp. 207-233. Little, E. A., Bullough, R., Wood, M. H. (1980), Proc. Roy. Soc. (London) A372, 565. Logunstsev, Ye. N., Safonov, V. A., Tyumentsev, S. N., Kozlov, A. V., Nalesnik, V. M. (1984), Fiz. Metal. Metalloved. 57, 802-807. Loomis, B. A., Smith, D. L., Garner, F. A. (1991), /. Nucl. Mater. 179-181, 111-111. Lovell, A. I , Chin, B. A., Gilbert, E. R. (1981), /. Mater. Sci. 16, 870. Lucas, G. E. (1993), "The Evolution of Mechanical Property Change in Irradiated Austenitic Stainless Steels", J. Nucl. Mater. 206, 287-305. Maillard, A., Touron, H., Seran, J. L., Chalony, A. (1993), Effects of Radiation on Materials: 16th Int. Symp., STP 1175. Philadelphia: ASTM, in press. Makenas, B. X, Jost, J. W, Hales, J. W. (1980), Trans. ANS 34, 256-257. Makenas, B. X, Chastain, S. A., Gneiting, B. C. (1990 a), Proc. LMR: A Decade of LMR Progress and Promise. La Grange Park, IL: ANS, pp. 176183. Makenas, B. X, Chastain, S. A., Gneiting, B. C. (1990 b), Dimensional Changes in FFTF Austenitic Cladding and Ducts, Westinghouse Hanford Company Report WHC-SA-0933VA. Richland, WA: Westinghouse. Mann, F. M. (1982), Transmutation of Alloys in MFE Facilities as Calculated by REAC, HEDL-TME81-37. Richland, WA: Hanford Engineering Development Laboratory. Mansur, L. K. (1978), Nuclear Technology 40, 5. Mansur, L. K., Grossbeck, M. L. (1988), /. Nucl. Mater. 155-157, 130-147. Mansur, L. K., Yoo, M. H. (1979), /. Nucl. Mater. 85-86, 523-532. Mansur, L. K., Lee, E. H., Maziasz, P. X, Rowcliffe, A. F. (1986), /. Nucl. Mater. 141-143, 633-646. Marlowe, M., Appleby, W. K. (1973), Trans. ANS 16, 95-96. Marucco, A., Nath, B. (1983), Proc. Conf Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys. London: British Nuclear Energy Society, pp. 59-62. Marwick, A. D., Piller, R. C , Sivell, P. M. (1979), /. Nucl. Mater. 83, 35-41. Matsui, H., Gelles, D. S., Kohno, Y. (1992), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 928-941. Matthews, X R., Finnis, M. W. (1988), J. Nucl. Mater. 159, 257-285. Maziasz, P. X (1984), /. Nucl. Mater. 122-123, 472486.
Maziasz, P. X, Braski, D. N. (1984), J. Nucl. Mater. 122-123, 311-316. Maziasz, P. X, Klueh, R. L. (1992), in: Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia, PA: ASTM, pp. 1135-1156. McCarthy, X M., Garner, F A. (1988), /. Nucl. Mater. 155-157, 877-882. McElroy, W N., Farrar IV, H. (1972), Proc. Conf. on Radiation-Induced Voids in Metals, Albany, CONF710601. U.S. Atomic Energy Commission, pp. 187-229. McSherry, A. X, Patel, M. R., Marshall, X, Gilbert, E. R. (1978), Trans. ANS 28, 146. Menzinger, K, Sacchetti, F. (1975), /. Nucl. Mater. 57, 193-197. Mills, W. X (1982), Irradiation Effects on the Fracture Toughness of Austenitic Fe-Cr-Ni Alloys, Hanford Engineering Development Laboratory Report HEDL-TME-82-17. Richland, WA: Hanford Engineering Development Laboratory. U.S. DOE. Mills, W X (1987), Nucl. Technol. 82, 290-303. Mosedale, D., Harries, D. R., Hudson, X A., Lewthwaite, G. W, McElroy, R. X (1977), Proc. Int. Conf Radiation Effects in Breeder Reactor Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 209-218. Muroga, T, Garner, F. A., Ohnuki, S. (1991), /. Nucl. Mater. 179-181, 546-549. Muroga, T, Garner, F. A., McCarthy, J.M., Yoshida, N. (1992), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 1015-1033. Nakajima, H., Yoshida, S., Kohno, Y, Matsui, H. (1992), J. Nucl. Mater. 191-194, 952-955. Nelson, R. S., Etherington, E. W, Smith, M. I. (1972), UKAEA Report TRG-2152(D). London: U.K. Atomic Energy Authority. Neustrov, V. S., Shamardin, V. K. (1991), Atomnaya Energiya 71, 345-348. Neustrov, V. S., Shamardin, V. K. (1993), Effects of Radiation on Materials: 16th Int. Symp., STP 1175. Philadelphia: ASTM, pp. 816-823. Neustrov, V. S., Golovanov, V. N., Shamardin, V. K. (1990), Atomnaya Energiya 69, 223-226. Nikolaev, V. A., Kursevich, I. P., Zhukov, O. N., Lapin, A. N. (1985), Atomnaya Energiya 59(3), 200-204. Nolfi, F. W, Jr. (1983), Phase Transformations During Irradiation. London: Applied Science. Norgett, X X, Robinson, J. T, Torrens, I. M. (1975), Nucl. Eng. Design 33, 50. Norton, S. H., Soderberg, C. R. (1942), Trans. ASME 64, 769. Odette, G. R. (1988), /. Nucl. Mater. 155-157, 921927. Okamoto, P. R., Rehn, L. E. (1979), J. Nucl. Mater. 83, 2-23. Parshin, A. M. (1980), Voprosy Atomnoj Nauki i Tekhniki. Ser. Fizika Radiats. Provrezhdenii i Radiats. Materialovedenie, No. 3, pp. 20-29.
6.9 References
Plumton, D. L., Wolfer, W. G. (1984), /. Nucl. Mater. 120, 245-253. Porollo, S. I., Vorobyev, A. N., Dmitriev, V. D., Bibilashvili, Yu. K., Golovnin, I. S., Kalashnik, G. V., Romaneyev, V. V. (1990), Proc. Conf. Fast Reactor Core and Fuel Structural Behavior, Inverness. London: British Nuclear Energy Society, pp. 237-241. Porter, D. L. (1979), /. Nucl Mater. 79, 406-411. Porter, D. L. (1984), Effects of Environmental Variables on Void Swelling in Austenitic Stainless Steel, ANL-84-42. Idaho Falls, ID: Argonne National Laboratory. Porter, D. L., Garner, F. A. (1985), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 212-220. Porter, D. L., Garner, F. A. (1987), in: Influence of Radiation on Material Properties: 13th Int. Symp., Part II, STP 956. Philadelphia, PA: ASTM, pp. 11-21. Porter, D. L., Garner, F. A. (1988), J. Nucl. Mater. 159, 114-121. Porter, D. L., Hudman, G. L. (1980), Trans. ANS34, 230-231. Porter, D. L., Wood, E. L. (1979), /. Nucl. Mater. 83, 90-97. Porter, D. L., Takata, M. L., Wood, E. L. (1983), /. Nucl. Mater. 116, 272-276. Porter, D. L., Wood, E. L., Garner F. A. (1990), Effects of Radiation on Materials: 14th Int. Symp., Vol. II, STP 1046. Philadelphia: ASTM, pp. 551559. Porter, D. L., Hudman, G. D., Garner, F. A. (1991), /. Nucl. Mater. 179-181, 581-584. Powell, R. W, Peterson, D. T, Zimmerschied, M. K., Bates, J. F. (1981), J. Nucl. Mater. 103-104, 969974. Powell, R. W, Johnson, G. D., Hamilton, M. L., Garner, F. A. (1986), Int. Conf Reliable Fuels for Liquid Metal Reactors, Tuscon. La Grange Park, IL: ANS, pp. 4-17-4-29. Puigh, R. X, Garner, F. A. (1990), Effects of Radiation on Materials: 14th Int. Symp., Vol. II, STP 1046. Philadelphia: ASTM, pp. 527-536. Puigh, R. J., Gelles, D. S. (1989), Fusion Reactor Materials Semiannual Progress Report DOE/ER0313/6. Oak Ridge, TN: U.S. DOE, pp. 201-224. Puigh, R. J., Schenter, R. E. (1985), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 795-802. Puigh, R. X, Gilbert, E. R., Chin, B. A. (1982), Effects of Radiation on Materials: 11th Int. Symp., STP 782. Philadelphia: ASTM, pp. 108-121. Puigh, R. X, Lovell, A. X, Garner, F. A. (1984), /. Nucl. Mater. 122-123, 242-245. Rauh, H., Bullough, R. (1985), Philos. Mag. 52, 333356. Rauh, H., Bullough, R., Matthews, X R. (1992), Philos. Mag. 65, 53-70. Rehn, L. E. (1990), /. Nucl. Mater. 174, 144-150.
541
Reshetnikov, F. G. (1991), Proc. Int. Conf Fast Reactors and Related Fuel Cycles, Kyoto, Japan, Vol. I, Atomic Energy Society of Japan, pp. 7.6-1-7.6-7. Rothman, S. X, Nowicki, L. X, Murch, G. E. (1980), J. Phys. F. Met. Phys. 10, 383-398. Russell, K. C. (1984), Prog. Mater. Sci. 28, 229-434. Sacchetti, F. (1977), /. Nucl. Mater. 64, 115-120. Sahu, H. K., Jung. P. (1985), /. Nucl. Mater. 136, 154-158. Sasmal, B. (1981), Proc. Int. Conf Solid-Solid Phase Transformations. New York: The Metallurgical Society of AIME, pp. 775-779. Schule, W, Hausen, H. (1994), Proc. of ICFRM-6, Stresa, Italy. /. Nucl. Mater., in press. Seran, X L., Dupouy, X M. (1982), Effects of Radiation on Materials: 11th Int. Symp., STP 782. Philadelphia: ASTM, pp. 5-16. Seran, X L., Dupouy, X M. (1983), Proc. Conf Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys, Vol. 1., Brighton. London: British Nuclear Energy Society, pp. 2 2 28. Seran, X L., Touron, H., Maillard, A., Dubuisson, P., Hugot, X P., Le Boulbin, E., Planchard, P., Pelletier, M. (1990), Effects of Radiation on Materials: 14th Int. Symp., STP 1046. Philadelphia: ASTM, pp. 739-752. Seran, X L., Levy, V, Dubuisson, P., Gilbon, D., Maillard, A., Fissolo, A., Touron, H., Cauvin, R., Chalony, A., Le Boulbin, E. (1992), Effects of Radiation on Materials: 15th Int. Symp., STP 1125. Philadelphia: ASTM, pp. 1209-1233. Shamardin, V. K., Golovanov, V. N., Povstyanko, A. V, Neutroev, V. S., Bibilashvili, Y. K., Golovin, I. S., Kalashnik, G. V, Romaneev, V. V. (1990), Effects of Radiation on Materials: 14th Int. Symp., STP 1046. Philadelphia: ASTM, pp. 753-765. Shcherbak, V, Dmitriyev, V. D. (1987), Fiz. Metal. Metalloved. 64, 592-595. Shibahara, I., Ukai, S., Onose, S., Shikakura, S. (1993), J. Nucl. Mater. 204, 131-140. Shields, Jr., X A. (1981), Nucl. Technol. 52, 214-227. Simonen, E. P. (1990), Metall. Trans. 21 A, 10531063. Simons, R. L., Brager, H. R., Matsumoto, W. Y. (1986), /. Nucl. Mater. 141-143, 1057-1060. Skinner, B. C , Woo, C. H. (1984), Phys. Rev. B30, 3084. Sniegowski, X X, Wolfer, W. G. (1984), Proc. Topical Conf Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird. Warrendale: The Metallurgical Society of AIME, pp. 579-586. Soderberg, C. R. (1941), Trans. ASME 63, 731. Spitznagel, X, Stickler, R. (1971), Correlation Between Precipitation Reactions and Bulk Density Changes in Type 18-8 Austenitic Stainless Steel, Westinghouse R&D Report 71-1D4-CLADS-P1. Pittsburgh, PA: Westinghouse. Stanley, X T. (1979), J. Nucl. Mater. 85-86, 787-791.
542
6 Irradiation Performance of Cladding and Structural Steels in Liquid Metal Reactors
Stanley, J. T., Garr, K. R. (1975), Metall. Trans. 6A, 531-535. Stanley, J. T., Hendrickson, L. E. (1979), / NucL Mater. 80, 69-78. Stoller, R. E. (1990), Metall. Trans. 21A, 1829-1837; also: (1987) Oak Ridge National Laboratory Report ORNL-6430. U.S. DOE. Straalsund, J. L. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: American Institute of Mining, Metallurgical and Petroleum Engineers, pp. 191-207. Straalsund, J. L., Day, C. K. (1973), NucL Technol. 20, 27. Straalsund, J. L., Paxton, M. M. (1972), NucL Technol. 13, 99. Straalsund, J. L., Guthrie, G. L., Larson, T. J. (1972), Length Changes in FTR Prototypic Cladding Irradiated in EBR-II, Hanford Engineering Development Laboratory Report HEDL-TME-72-95. Richland, WA: Hanford Engineering Development Laboratory. Straalsund, J. S., Powell, R. W., Chin, B. A. (1982), /. NucL Mater. 108-109, 299-305. Stubbins, J. E, Garner, F. A. (1992), /. NucL Mater. 191-194, 1295-1299. Sutherland, W. H. (1984), Proc. Specialists Meeting on Predictions and Experience of Core Distortion Behavior, Report IWGFR-54, Winslow. England: Inter. Atomic Energy Agency, p. 1/7-2. Tateishi, Y. (1989), J. NucL Sci. Technol. 26,132-136. Tenbrink, I , Wahi, R. P., Wollenberger, H. (1988), /. NucL Mater. 155-157, 850-855. Terasawa, M. (1985), Proc. 1st Sino-Japanese Symp. Metal Physics and Physical Metallurgy, Beijing, China. Tokyo: Science University of Tokyo, pp. 43-52. Thomas, L. E. (1980), Proc. Symp. Phase Stability During Irradiation. Warrendale: The Metallurgical Society of AIME, pp. 237-255. Thomas, L. E. (1982), Proc. 40th Annu. Meeting Electron Microscopy Society of America. Washington D. C : Electron Microscopy Society of America, p. 597. Toloczko, M. B., Garner, F. A. (1994), Proc. ICFRM6, Stresa, Italy. J. NucL Mater., in press. Toloczko, M. B., Garner, F. A., Eiholzer, C. R. (1992), /. NucL Mater. 191-194, 803-807. Toloczko, M. B., Garner, F. A., Eiholzer, C. R. (1994), Proc. ICFRM-6, Stresa, Italy. J. NucL Mater., in press. Uematsu, K., Kodama, T, Ishida, Y, Suzuki, K., Koyama, M. (1979), Proc. Inter. Conf. on Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: The Metallurgical Society of AIME, pp. 571-589. Vaidyanathan, S., Lauritzen, T, Bell, W. L. (1982), Effects of Radiation on Materials: Twelfth Int. Symp., STP 870. Philadelphia: ASTM, pp. 127138.
Vasina, N. K., Kursevich, I. P., Kozhevnikov, O. A., Shamardin, V. K., Golovanov, V. N. (1985), Atomnay a Energiya 59(4), 265-267. Vitek, J. M., Klueh, R. L. (1983), Proc. Topical Conf Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird. Warrendale: The Metallurgical Society of AIME, pp. 551-558. Walker, B. I , Lightowlers, M. A. (1990), Proc. Conf Fast Reactor Core and Fuel Structural Behavior, Inverness. London: British Nuclear Energy Society, pp. 63-67. Waltar, A. E., Reynolds, A. B. (1981), Fast Breeder Reactors. New York: Pergamon Press. Walters, L. C , McVay, G. L., Hudman, G. D. (1977), Proc. Int. Conf. Radiation Effects in Breeder Reactor Structural Materials, Scottsdale. New York: American Institute of Mining, Metallurgical and Petroleum Engineers, pp. 277-294. Wassiliew, C , Herschbach, K., Materna-Morris, E., Ehrlich, K. (1983), Proc. Topical Conf. on Ferritic Alloys for Use in Nuclear Energy Technologies, Snowbird. Warrendale: The Metallurgical Society of AIME, pp. 607-614. Wassiliew, C , Ehrlich, K., Bergmann, H.-J. (1987), Influence of Radiation on Materials Properties: 13th Int. Symp., Part II, ASTM STP 956. Philadelphia: ASTM, pp. 30-53. Watkin, J. S. (1976), Irradiation Effects on the Microstructure and Properties of Metals, STP 611. Philadelphia: ASTM, pp. 270-283. Watkin, J. S., Standring, J. (1974), Anisotropic Swelling in M316 and PE16 PFR Fuel and Structural Materials, UKAEA Report TRG-M-6436. London: U.K. Atomic Energy Authority. Weiss, B., Stickler, R. (1970), Phase Instabilities During High Temperature Exposure of 316 Austenitic Stainless Steel, Westinghouse R&D Report 701D4-STABL-P1. Pittsburgh, PA: Westinghouse. Weisz, M. (1978), Proc. Alushta Conf. Radiation Damage in Materials, Vol. 5, Moscow, U.S.S.R. Government, pp. 148-176. Westmoreland, J. E., Sprague, I A., Smidt Jr., F. A., Malmberg, P. R. (1975), Radiation Effects 26, 1-16. Wiedersich, H. (1972), Radiation Effects 12, 111. Williams, T. M. (1982), Effects of Radiation on Materials: 11th Int. Symp., STP 782. Philadelphia: ASTM, pp. 166-185. Williams, T. M., Boothby, R. M., Titchmarsh, J. M. (1987), Proc. Int. Conf Materials for Nuclear Reactor Core Applications, Bristol. London: British Nuclear Energy Society, pp. 293-299. Wolfer, W. G. (1984), J. NucL Mater. 122-123, 367378. Wolfer, W G., Ashkin, M. (1975), /. Appl. Phys. 46, 547. Wolfer, W G., Garner, F. A. (1984), Damage Analysis and Fundamental Studies Quarterly Progress Report DOE/ER-0046/17. Richland, WA: U.S. DOE, pp. 58-69.
6.9 References
Woo, C. H., Garner, F. A. (1992), /. Nucl. Mater. 191-194, 1309-1312. Woo, C. H., Singh, B. N. (1990), Phys. Status Solidi B159, 609. Woo, C. H., Singh, B. N. (1992), Phil. Mag. A65, 889. Woo, C. H., Singh, B. N., Garner, F. A. (1992), J. Nucl. Mater. 191-194, 1224-1228. Woo, C. H., Garner, F. A., Holt, R. A. (1993), Effects of Radiation on Materials: 16th Int. Symp., STP 1175. Philadelphia: ASTM, in press. Yang, W J. S. (1982), J. Nucl. Mater. 108-109, 339346. Yang, W J. S. (1987), Radiation-Induced Changes in Micro structure: 13 th Int. Symp., STP 955. Philadelphia: ASTM, pp. 628-646. Yang, W. J. S., Garner, F. A. (1982), Effects of Radiation on Materials: 11th Int. Symp., STP 782. Philadelphia: ASTM, pp. 186-206. Yang, W. J. S., Hamilton, M. L. (1982), J. Nucl Ma-
ter. 108-109, 339-346. Yang, W. J. S., Makenas, J. B. (1985), Effects of Radiation on Materials: 12th Int. Symp., STP 870. Philadelphia: ASTM, pp. 127-138. Yang, W. J. S., Brager, H. R., Garner, F. A. (1981), Proc. Symp. Phase Stability During Irradiation, Pittsburgh. Warrendale: The Metallurgical Society of AIME, pp. 257-269. Yang, W. J. S., Gelles, D. S., Straalsund, J. L., Bajaj, R. (1985), J. Nucl. Mater. 132, 249-265.
543
General Reading Nolfi, F. W, Jr. (1983), Phase Transformations During Irradiation. London: Applied Science. Russell, K. C. (1984), "Phase Stability Under Irradiation", Prog. Mater. Sci. 28, 229-434. Waltar, A. E., Reynolds, A. B. (1981), East Breeder Reactors. New York: Pergamon Press. Series of Proc. Int. Symp. on Radiation Effects in Materials: STP 683 (1979); STP 725 (1980); STP 780 (1982); STP 870 (1985); STP 955 and 956 (1987); STP 1046 (1989); STP 1125 (1992); STP 1175 (1993). Philadelphia: ASTM. Proc. Int. Conf on Radiation Effects in Fast Breeder Reactor Structural Materials (1977). New York: The Metallurgical Society of AIME. Proc. Conf on Phase Stability During Irradiation (1981), Pittsburgh. Warrendale, PA: The Metallurgical Society of AIME. Proc. Conf. on Dimensional Stability and Mechanical Behavior of Irradiated Metals and Alloys (1983), Brighton. London: British Nuclear Energy Society. Proc. Int. Conf. on Materials for Nuclear Reactor Core Applications (1987), Bristol. London: British Nuclear Energy Society. Proc. LMR: A Decade ofLMR Progress and Promise (1990). La Grange Park, IL: ANS.
7 Zirconium Alloys in Nuclear Applications Clement Lemaignan CEA/Centre d'Etudes Nucleaires de Grenoble/DTP/SECC, Grenoble, France Arthur T. Motta Nuclear Engineering Department, The Pennsylvania State University, University Park, PA, U.S.A.
List of 7.1 7.1.1 7.1.2 7.2 7.2.1 7.2.2 7.2.2.1 7.2.2.2 7.2.3 7.2.3.1 7.2.3.2 7.3 7.3.1 7.3.1.1 7.3.1.2 7.3.1.3 7.3.1.4 7.3.1.5 7.3.1.6 7.3.1.7 7.3.2 7.3.2.1 7.3.2.2 7.3.2.3 7.3.2.4 7.3.3 7.4 7.5 7.6
Symbols and Abbreviations History High Temperature Water Reactors Current Use Fabrication and Products Processing Microstructure Alloys and Alloying Elements Heat Treatments and Resultant Microstructure Properties Mechanical Properties Diffusion Data In-Reactor Behavior Irradiation Damage and Irradiation Effects Displacement Calculations Irradiation Effects in the Zr Matrix Irradiation Effects on Second Phases Irradiation Growth Irradiation Creep Changes in Mechanical Behavior Charged-Particle Irradiation Corrosion Behavior General Corrosion Behavior .. Oxidation of the Precipitates Water Radiolysis Hydrogen Pickup Pellet-Cladding Interaction Challenges Acknowledgements References
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
2 4 4 4 5 5 7 10 17 18 18 22 24 24 24 25 28 29 32 33 34 36 36 37 39 40 41 46 47 47
2
7 Zirconium Alloys in Nuclear Applications
List of Symbols and Abbreviations a, c a 0 , c0 A , d D D|l ,D ± E Ed Es Fx hD1 Klc ^iscc m, n Q SL R t T
directions 1010 and 0001, respectively lattice parameters constant unit cell vectors parallel and perpendicular to the basal plane of Zr oxide layer thickness diffusion coefficient diffusion coefficient parallel, perpendicular to the c-axis energy displacement energy position of the most stable interstitial resolved fraction of basal planes in the direction x diffusion enthalpy fracture toughness SCC stress intensity factor exponents activation energy ratio of thickness reduction to diameter reduction gas constant time absolute temperature
£ e ex v (E) a 0
strain strain rate strain rate in direction x number of displaced atoms in a cascade stress neutron flux
AECL ASTM b.c.c. b.c.t. BU BWR CANDU CAP CRNL DAD DHC dpa f.c.c. h.c.p. HVEM IAEA
Atomic Energy of Canada Ltd. American Society for Testing Materials body-centered cubic body-centered tetragonal burn-up boiling water reactor Canadian deuterium uranium reactor cumulative annealing parameter (£A) Chalk River National Laboratory diffusion anisotropy difference delayed hydride cracking displacements per atom face-centered cubic hexagonal close-packed high voltage electron microscope International Atomic Energy Agency
List of Symbols and Abbreviations
IGSCC I-SCC LHGR LWR MIBK PCI PKA ppm PWR RBMK RX R&D SCC SEM SIPA STEM SOCAP SR TBS TEM Trex UTS VVER YS
iodine intergranular stress corrosion cracking iodine stress corrosion cracking linear heat generation rate light water reactor methyl-isobutyl-ketone (process) pellet-cladding interaction primary knock-on atom part per million pressurized water reactor Russian graphite-moderated boiling water reactor fully recrystallized research and development stress corrosion cracking scanning electron microscope stress induced preferential absorption scanning transmission electron microscope second order cumulative annealing parameter stress relieved to be specified transmission electron microscopy tube-reduced extrusion ultimate tensile strength Voda-Voda energy reactor, Russian type PWR yield strength
7 Zirconium Alloys in Nuclear Applications
7.1 History 7.1.1 High Temperature Water Reactors
Soon after the observation of the fission of uranium 235, L. Szilard and F. JoliotCurie recognized the possibility of using the chain reaction phenomenon as a source of energy. Initially, test reactors were designed with no constraints on thermal efficiency. The aim was then to understand neutron physics and to study the behavior of materials under irradiation. Low temperature, pool type reactors were constructed in which the structural material used was exposed to a comparatively mild environment. Aluminum and beryllium alloys were used for core components, due to their low thermal neutron capture cross section and acceptable corrosion rate in water below 100 °C. Once nuclear power reactors for submarine propulsion and production of electricity were designed, thermal efficiency became mandatory, and materials had to be found that could withstand the high temperature of the coolant, usually water. Zirconium (Zr), with its very low thermal neutron capture cross section, was a potential candidate, but had poor ductility and corrosion resistance. The first pressurized water reactors were loaded with fuel claddings and other structural elements (guide tubes and grids) made of stainless steel. An improvement in neutron efficiency was a driving force for the development of industrial type Zr-based alloys. At the end of World War II, the nuclear submarine program undertook a large effort in that field. Systematic testing and research and development (R & D) by the U.S. Navy resulted in the development of an efficient hafnium (Hf) separation process and industrial scale ingot production procedures. During the test of a series of binary and
ternary alloys, an accidental contamination of a Zr-2.5% Sn (Zircaloy-1) melt by stainless steel caused the serendipitous discovery of an alloy of good corrosion behavior. Composition variations around this alloy led to Zircaloy-2. Zircaloy-3, a very low tin variant, was soon abandoned in favor of the better Zircaloy-4, a Ni-free variant designed to decrease hydrogen pickup. Similar tinkering with compositional variations to improve corrosion resistance and strength led to the development by the Soviet Union of another family of alloys using the Zr-Nb binary system, later used by Canada as well. The Zr-Nb system allowed the possibility of obtaining a fine two-phase structure that leads to higher strength. The subject of Zr metallurgy has merited books (Lustman and Kerze, 1955) and reviews (Douglass, 1971; Cheadle, 1975) in the past. Detailed aspects of the metallurgy of the IV-A series (Ti, Hf, and Zr) can be found in Vol. 8, Chap. 8, of this Series. It is the purpose of this work to review the use of Zr for nuclear applications and to present some of the more recent developments in the field. 7.1.2 Current Use
In today's nuclear power reactors, Zr alloys are commonly used for structural components and fuel cladding. For light water reactors (LWR), the common choices are Zircaloy-4 in pressurized water reactors (PWR) and Zircaloy-2 in boiling water reactors (BWR). The heavy watermoderated natural uranium CANDU reactor (Canadian deuterium uranium), as well as the Russian RBMK reactor, use Zr-Nb alloys. In fuel assemblies and bundles, claddings are made out of Zircaloy-2 or Zircaloy-4. Those components are exposed to the fission products at the inner surface
7.2 Fabrication and Products
at temperatures close to 400 °C. At the outer surface they are in contact with light or heavy water at coolant temperatures (from 280 to 350 °C). Typical heat fluxes across the cladding are in the range of 30-50 W • cm" 2 . Those tubes have different geometries, depending on reactor design (Fig. 7-1). In PWR's fuel rods claddings are 4 to 5 meters long and have a diameter of 9 to 12 mm for a thickness of 0.6 to 0.8 mm. BWR fuel rods are usually slightly larger. In CANDU, fuel bundles are short - 0.5 m - to allow on-line refuelling. The cladding is very thin - 0.4 mm - and is designed to collapse around the UO 2 pellets early during irradiation. In the Russian VVER's the fuel rod geometry is similar to PWR's but the usual cladding alloy is Zr-1 % Nb. Structural components of the fuel assemblies are guide tubes and grids that compose the skeleton. They have to withstand mechanical stresses during normal or accidental operation as well as the oxidizing hot water. In BWR's each assembly is surrounded by a Zircaloy-2 channel box that avoids cross-flow instabilities of the two-phase coolant. Geometrical stability of those components is a critical aspect of core design as it affects fuel loading capability, cooling efficiency and neutron physics behavior of the core. In the case of CANDU's and RBMK's, the coolant is separated from the moderator and flows around the fuel bundles in pressure tubes, usually made of Zr-Nb alloys. Those large components (10 m x 20 cm x 5 mm) are considered as a structural part of the reactor with a design life of tens of years. They are thus exposed to a high irradiation fluence (up to 3 x 1026 n m" 2 ) in contact with the coolant on the inner surface. Mechanical stability of those large components affects the overall geometry of the reactor.
7.2 Fabrication and Products 7.2.1 Processing Zirconium is commonly found in nature associated with its lower row counterpart in Mendeleev's table, hafnium. Most of the common Zr ores contain between 1.5 and 2.5% Hf. Due to its high thermal neutron capture cross section, Hf needs to be removed from Zr for nuclear applications. The most frequently used ore is zircon (ZrSiO4) with a worldwide production of about one million metric tons per year. Most of the zircon is used in its original form or in the form of zirconia (ZrO2) as foundry die sands, abrasive materials or high temperature ceramics. Only 5% is processed into Zr metal and alloys. The processing of Zr alloy industrial components is rather complex due to the reactivity of the metal with oxygen. The general scheme is presented in Fig. 7-2: Ore processing, zirconium/hafnium separation, reduction to metal, alloy melting, hot and cold deformation processing. The first step is to convert the zircon into ZrCl 4 , though a carbo-chlorination process performed in a fluidized bed furnace at 1200 °C. The reaction scheme is the following: ZrO 2 ( + SiO2 + HfO2) + 2 C + 2 Cl2 -> -> ZrCl 4 (+ SiCl4 + HfCl4) + 2 CO After this step, Zr and Hf are separated using one of the two following processes: (i) Wet chemical: after reaction with ammonium thiocyanate (SCNNH 4 ) a solution of hafnyl-zirconyl-thiocyanate (Zr/Hf)O(SCN)2 is obtained. A liquidliquid extraction is performed with methyl-isobutyl-ketone (MIBK, name of the process). Hf-free ZrO 2 is obtained after several other chemical steps: hydrochlonation, sulfation, neutralization with NH 3 ,
7 Zirconium Alloys in Nuclear Applications
(c)
(a)
Figure 7-1. Fuel cladding and other components, made of Zr alloys, used in different reactor types: (a) PWR fuel assembly (courtesy FRAGEMA), (b) BWR fuel assembly and channel (courtesy GEc), (c) CANDU fuel assembly and surrounding pressure tube.
7.2 Fabrication and Products
and calcination. ZrCl 4 is the final result of a second carbo-chlorination process (Stephen, 1984). (ii) Direct separation process: this is an extractive distillation within a mixture of KCI-AICI3 as solvent at 350 °C. The vapor phase, generated by a boiler at the lower part of the distillation column, is enriched in Hf, while the liquid phase traps the Zr (Moulin et al., 1984b). In either case, Zr metal is obtained by a reduction of ZrCl 4 in gaseous form by liquid magnesium, at about 850 °C in an oxygen-free environment. Residual quantities of Mg and MgCl4 are removed from the "sponge cake" by distillation at 1000 °C. After mechanical fracturing, the pieces of sponge are sorted, giving the basic product for alloy ingot preparation. High purity Zr can be obtained by the Van Arkel process. This consists of the reaction of Zr with iodine at moderate temperature, gaseous phase transport as Zrl 4 and decomposition of the iodide at high temperature on an electrically heated filament, the iodine released being used for the low temperature reaction in a closed loop transport process, according to the following scheme: Zr + 2I 2 - ZrI 4 (g)... Zrl 4 -> Zr + 2I 2 (g) t I 250-300 °C
1300-1400°C
For industrial alloys, a compact of sponge containing the alloying elements O (in the form of ZrO 2 ), Sn, Fe, Cr, Ni, and Nb - in the desired composition, is melted in a consumable electrode vacuum furnace, usually three times. These vacuum meltings reduce the gas content and increase the homogeneity of the ingot. Typical ingot diameters range between 50 and 80 cm, for a mass of 3 to 8 metric tons. Industrial use of Zr alloys requires either tube- or plate-shaped material. The first
step of mechanical processing is forging or hot rolling in the (3 phase, at a temperature close to 1050 °C. Hot extrusion is used to obtain tube shells or Trex (tube-reduced extrusion), while hot rolling is used for flat products. For Zircaloys, at that stage a (3 quench is performed to increase the corrosion resistance of the final product. This treatment controls the distribution of second phase particles, if no further processing is performed above 800 °C (Schemel, 1977). Further reduction in size is obtained by cold rolling either on standard or pilger-rolling mills. Low temperature recrystallization is performed between the various size reduction steps. 7.2.2 Microstructure Pure zirconium crystallizes at ambient temperature as an hexagonal close-packed metal, with a c/a ratio of 1.593 (i.e., a slight compression in the c-direction compared to the ideal ratio of 1.633). Lattice parameters are a o = 0.323nm and co = 0.515nm (Douglass, 1971). The thermal expansion* coefficients have been measured by Lloyd (1963) on single crystals. The difference in thermal expansion coefficients between the a and odirections (see Table 7-1) implies that the c/a ratio tends towards the ideal ratio at higher temperatures - i.e., towards a more isotropic behavior. At 865 °C, Zr undergoes an allotropic transformation from the low temperature h.c.p. a phase to body centered cubic P phase. On cooling, the transformation is either martensitic or bainitic, depending on the cooling rate, with a strong epitaxy of the a platelets on the old (3 grains according to the scheme proposed by Burgers (1934) : (0001)a||{110}p
and
lMeV). Loops exhibiting inside contrast with a 1212 diffraction vector are vacancy in nature (beam direction close to [1213]). Close to the grain boundary, vacancy loops form preferentially, in contrast to the bulk, where some interstitial loops also form (micrograph courtesy of M. Griffiths, AECL, Chalk River).
dislocations of the type 1/2 [0001] and 1/2 [1123] are found, which supports the general idea of development of -type dislocations under irradiation. Voids
Contrary to the behavior of stainless steel, Zircaloy does not exhibit significant void formation under neutron irradiation (Farrell, 1980). Under electron irradiation swelling is observed in Zr which was previously injected with helium (He) (Faulkner and Woo, 1980). Accordingly, TEM studies have not found many cavities and voids in neutron or charged-particle irradiated Zr alloys. This was confirmed by Baig et al. (1989) who could not find any voids in irradiated Zr by small angle neutron scattering. The reason is thought to be that the gases that could stabilize voids and cavities (O,H,N) are very soluble in the Zr matrix, where their equilibrium concentration can be quite high (Figs. 7-4 and 7-5). Some cavities have nevertheless been found in crystal bar Zr after neutron irradiation, in the interstitial denuded zone near grain boundaries (Griffiths et al., 1988). This is probably due to the high
(a)
(b)
„. ...
l
Figure 7-18. (a) Formation of -component dislocations in Zircaloy-4 irradiated in a PWR at 580 K to a fluence of 8.5 x 1025 n • m~ 2 . Compare with Fig. 7-16 b taken under same diffraction conditions, (b) Association of -component dislocations with amorphous Zr(Cr, Fe)2 precipitate, in course of dissolution. Arrows indicate steps similar to those in Fig. 7-19 Band C.
28
7 Zirconium Alloys in Nuclear Applications
vacancy concentration in those regions. For reasons unknown, voids are also found near second phase particles. Their total number density is however quite low and they are not thought to have much effect in the irradiation behavior of zirconium alloys. 7.3.1.3 Irradiation Effects on Second Phases Crystalline to Amorphous Transformation (Amorphization) of Precipitates One of the most striking effects of irradiation on Zr alloys is the crystalline to amorphous transformation (amorphization) observed in the intermetallic precipitates Zr(Cr, Fe) 2 and Zr 2 (Ni, Fe), commonly found in Zircaloys. Those precipitates, described in Sec. 7.2.2, undergo amorphization under neutron irradiation as reported by Gilbert et al. (1985) and confirmed by Yang et al. (1986). The transformation has been observed in Zircaloys irradiated at 550 to 620 K (from both LWR fuel cladding and structural material) and at 350 K (calandria tubes from CANDU reactors). At 350 K, both types of precipitates are completely amorphous after very low fluences (0.5 to 1 dpa). At the higher temperature range, the Zr 2 (Ni, Fe) precipitates are completely crystalline, while the Zr(Cr, Fe) 2 precipitates are partially amorphous having developed a "duplex" structure, consisting of an amorphous layer that starts at the precipitatematrix interface, and gradually moves into the precipitate until the precipitate is completely amorphous. This is shown in Fig. 7-19 A. A crystalline core is present, as evidenced by the stacking fault contrast, while an amorphous layer has been formed that will eventually envelop the whole precipitate. Amorphization is associated with a depletion of iron from the amorphous
layer into the Zr matrix, while the Cr concentration in the precipitate remains constant. It is thought amorphization occurs because the irradiated crystalline structure is destabilized with respect to the amorphous phase due to the accumulation of irradiation damage. A review of the experimental evidence and theoretical models for amorphization of those precipitates under neutron and charged particle irradiation was given by Motta et al. (1991) and Motta and Lemaignan (1992). Precipitate Dissolution and Reprecipitation After amorphization, precipitate dissolution is accelerated, as illustrated in Figs. 7-19 B, C, and D. A serrated interface is formed as the precipitates dissolve, indicating either a dissolution of the precipitate along preferential directions (Yang, 1989), or a minimization of interfacial energy by faceting. After precipitate dissolution, and post irradiation annealing, Fe and Cr reprecipitate in the matrix, forming Cr-rich precipitates close to the original particle, and iron rich precipitates further away. This illustrates the difference in mobility of those two elements in Zr, mentioned in Sec. 7.2.3.2. It is possible that some of the Fe and Cr remains in solid solution. Although precipitate dissolution may be accelerated by the amorphous transformation, it also happens under neutron irradiation at higher temperatures, where amorphization does not occur, so amorphization is not a precondition for dissolution. There have been reports of irradiation induced precipitation and dissolution of other types of precipitates in zirconium alloys. ZrSnFe precipitates have been
7.3 In-Reactor Behavior
29
Figure 7-19. Precipitate amorphization and dissolution under neutron irradiation (courtesy of W. I S. Yang, GE). (A) The duplex structure in a Zr(Cr, FE)2 precipitate in Zircaloy-4 irradiated to 5 x 1025 n • m~ 2 at 561 K, showing a crystalline core surrounded by an amorphous layer. Iron is depleted in the amorphous layer. (B), (C), and (D): Zr(Cr,Fe) 2 precipitate dissolution along preferential crystallographic directions after irradiation to 1 4 . 7 x l 0 2 5 n m ~ 2 . The interface faceting is arrowed in (B).
found in Excel alloys (Zr-3.5% Sn-1.0% Nb-1.0% Mo) following neutron irradiation to 1.5 x 10 26 n • m~2 at 690 K (Griffiths, 1988). This was paralleled by the observation by Woo and Carpenter (1987) of Zr 5 Sn 3 precipitates in Zircaloy-2 following irradiation to 7.4 x 10 24 n * m~ 2 ( > 1 MeV) at 875 K (total dose about 3 dpa). These last precipitates were later found to be associated with Fe. Also in Zr-2.5% Nb, the Fe-rich (3 phase loses its Fe to the a phase, and extensive precipitation of fine-sized Nb rich precipitates in the a phase has been reported as a result (Coleman etal., 1981). This is shown in Fig. 7-20. Outside irradiation, by contrast, the p phase decomposes under thermal annealing to a mixture of Zr-86% Nb and the intermetallic (Zr, Nb) 3 (Fe, Cr). Second phase redissolution and redistribution of alloying elements, can have important consequences for irradiation growth and corrosion resistance, as explained below and in Sec. 7.3.2.
7.3.1.4 Irradiation Growth
Irradiation growth refers to the dimensional changes at constant volume of an unstressed material under irradiation (Fidleris, 1988). For Zr single crystals, irradiation growth consists of an expansion along the ^-direction, and corresponding contraction along the c-axis. In polycrystalline materials the situation is more complex, since grain boundaries can act as biased sinks for point defects, so that grain shape and orientation play a large role. Usually, however, growth behavior of polycrystalline materials also consists, of expansion along the a-direction and contraction along the c-axis. As noted in Sec. 7.2.3.3, the fabrication process of Zr alloy components induces a texture. For Zircaloy cladding tubes, prism planes are preferentially aligned perpendicular to the axial (longitudinal) direction, which means that irradiation growth causes the axial length to increase and the cladding diameter and thickness to dimin-
30
7 Zirconium Alloys in Nuclear Applications
Figure 7-20. p-Nb precipitation in the a phase of Excel alloy (micrograph courtesy of R. W. Gilbert, AECL, Chalk River).
ish. Besides the possibility of failure by bowing in restricted rods (Franklin and Adamson, 1988), this has an impact on design and safety, the growth distortions setting operational burnup limits. A considerable amount of experimental data has been gathered on this phenomenon (Rogerson, 1988), which will be briefly summarized here. Irradiation growth is influenced by microstructural variables such as amount of cold work, residual stresses and alloying additions, as well as by irradiation variables such as flux and temperature. Figure 7-21 shows the influence of cold work and temperature on irradiation growth (Adamson, 1977). The growth strain in cold-worked material follows in general a constant linear slope with fluence, which is higher at higher temperatures. Annealed material has considerably lower strain rates, at least in the initial part of irradiation exposure. However, after a
0.4
0.3
CW, 550 K
I 0.2 CO
„ RX, 550 K
I 0.1
- RX, 350-K
2 4 6 8 Fast Fluence (n/m2) (E>1 MeV)
10x1025
Figure 7-21. Growth strain versus fluence for different amounts of cold work, at several irradiation temperatures from Rogerson (1988). The dashed lines show the growth strain for cold-worked (CW) material: the strain is nearly linear with fluence and increases with temperature. The solid lines show the growth strain for annealed material: the growth rates are noticeably smaller than in cold-worked material, and also increase with temperature. After the "breakaway" fluence, growth rates in annealed material are comparable to those of cold-worked material. "Breakaway" growth in annealed materials has been linked to the development at that fluence of -component dislocations (see Fig. 7-18).
7.3 In-Reactor Behavior
fluence of about 3 x 10 25 n • m 2, there occurs the "breakaway phenomenon", the strain rates jumping to approximately the values observed in cold worked materials (Fidleris, 1988). The observation of breakaway growth has been linked to the development of -component dislocations (Holt and Gilbert, 1983, 1986). In general, when a high dislocation density exists, it dominates the growth behavior, which is then linear and steady state. According to rate theory (Wiedersich, 1972), this corresponds to a sink-dominated regime, where most of the defects created are lost to sinks, in this case dislocations. Transient behavior is observed in annealed materials or at the start of irradiation. When breakaway occurs, the observed growth rates are characteristic of a sink-dominated regime, as seen in Fig. 7-21. Alloying elements can also have an effect on the growth rates. It has been suggested (Griffiths, 1988) that impurity elements, including Fe, stabilize embryo component loops, enabling the development of -component dislocations referred to above. As noted in Sec. 7.2.3.2, iron enhances Zr self diffusion, which is another possible mechanism of enhancing growth, by enhancing point defect motion to dislocations. The influence of Sn has been measured by Zee et al. (1984): it was found that Zr-0.1% Sn grows considerably more than Zr-1.5% Sn. It is possible that this is a consequence of iron atom trapping by Sn atoms, thereby precluding Fe from enhancing growth. Recently, a new group of alloys with low Sn content, notably Zirlo (Boman et al., 1991) and the Russian alloy Zr 1% Nb 1% Sn 0.4% Fe (Nikulina et al., 1993) have shown noticeably reduced irradiation growth. Irradiation growth is due to the partitioning of interstitials and vacancies to dif-
31
ferent sinks that are anisotropically distributed in the material. Those sinks can be cold-work or irradiation-induced dislocations of or or character, and grain boundaries of different orientations. One example is the formation of either interstitial loops on -planes or vacancy loops on and -planes, or both, as initially proposed by Buckley (1961). Buckley's original model gave SXK1-3FX
(7-1)
where ex is the strain rate due to growth in direction x and Fx is the resolved fraction of basal planes in the direction x. Since basal loops were initially not observed in irradiated material, alternative models were proposed that included other vacancy sinks such as grain boundaries (Carpenter and Northwood, 1975), or considered the effect of anisotropic defect migration (Woo, 1988). Holt (1988) has presented a rate theory calculation that includes anisotropic point defect migration, texture effects, and all the point defect sinks proposed in previous models. It is found that growth in cold worked samples can be explained by a partitioning of vacancies to -network dislocations and interstitials to -type dislocations, while the accelerated growth observed in annealed material is due to the partitioning of vacancies to newly formed -type loops and interstitials to -type loops. In order to rationalize the linear growth rate at 350 K, a vacancy migration energy less than 0.7 eV is necessary, as initially shown by Buckley et al. (1980). This supports the Fe-enhanced vacancy migration mechanism mentioned in Sec. 7.2.3.2. Holt was also able to explain the simultaneous contraction along the longitudinal and transverse directions observed in some specimens as a grain size effect, operational when grain size falls below 1 Jim in
32
7 Zirconium Alloys in Nuclear Applications
highly oriented grains. The parameters adjusted to the behavior of Zircaloy were then used to successfully predict the growth of Zr-2.5% Nb (Holt and Fleck, 1991). 7.3.1.5 Irradiation Creep
Irradiation creep refers to the slow deformation under an external stress, experienced by a material under irradiation. In anisotropic materials, such as Zr, creep has always to be separated from growth in experimental situations. A convenient way to do this is to assume that creep and growth are linearly additive and define the irradiation creep strain as the additional strain that results when the deformation process takes place under an external stress. The creep strain is then the total strain minus the growth strain. As shown in Fig. 7-22, the creep deformation rate of Zr alloys is increased under neutron flux. Due to the objective of using Zr alloys in reactor, this phenomenon has been subject to a large amount of experimental work, reviewed by Fidleris (1988),
s — A. ' y*P t) ' o ' e
/
0.6/
At=2x 10 1 4 0.4/ 0.2-
0-
n-cnr 2-S-1
/ ^t=4xi( D13 n-cnrr2-s-1
/ . . . 4>t=o
/ / £••••1000
in order to improve the knowledge of the parameters that control creep. Some of this work has been performed in materials test reactors, and some during detailed examinations of the behavior of structural material in power reactors. The experimental procedures used for those analyses, are based upon pressurized tube expansion, stress relaxation measurements of thin plates, tensile testing or shear testing (i.e., springs loaded in tension). The last method is a specific experiment designed to analyze the pure creep behavior. For example, Causey et al. (1984) were able to derive a contribution of about 30% slip on the 1/3 {1011} system in addition to the classical -type basal slip by testing Zr-2.5 Nb in pure shear using compressed helical springs or twisted tubes at 573 K. They also showed a weak dependence of creep rate on dislocation density. Irradiation creep in Zircaloy-2 is anisotropic and dependent on texture as shown by Harbottle (1978). For practical purposes and for the limited range of operating parameters, the irradiation creep strain rate may be described accurately with a simple equation of the form:
2000 3000 time (hours)
4000
Figure 7-22. Diametral creep of RX Zircaloy under internal pressure (330 °C, aH = 150 MPa). The irradiation affects both the primary and the secondary creep, and the creep rate is increased in proportion to the dose rate.
(/-2)
where the effect of flux, stress and temperature can be separated. The usual values of those exponents are m = 0.6 to 1, n close to 1. The fact that n is close to 1 means that large creep strains can be sustained without failure. The activation energy, Q, is low, in the range of 5 to 15 kJ • mol" *, i.e., 0.05 to 0.16 eV • at" 1 (Franklin, 1982). The task of finding a mechanistic model that explains irradiation creep in Zr alloys is a daunting one, in view of the complexity of the alloy, its inherent single-crystal and texture anisotropy, and the effect of the irradiation field. For example, since void
7.3 In-Reactor Behavior
swelling is practically nonexistent in Zr alloys (Sec. 7.3.1.2), the standard rate theory model applicable to cubic metals coupling irradiation creep and void swelling (Brailsford and Bullough, 1972), cannot be used for Zr alloys. In addition, many of the parameters needed for mechanistic models, such as defect-defect interactions, are not known for Zr, further complicating the endeavor. Matthews and Finnis (1988) recently reviewed the literature on mechanisms of irradiation creep. It is thought that deformation during irradiation creep occurs by a combination of dislocation climb and glide, the climb being controlled by the stress-modified absorption of point defects at dislocations. According to the so-called SIPA (stress induced preferential absorption) mechanism (Bullough and Willis, 1975), dislocations that have their Burgers vectors parallel to the applied uniaxial stress, preferentially annihilate interstitials than vacancies, leading to dimensional changes due to the dislocation climb itself and to the subsequent dislocation glide (Woo, 1979). Matthews and Finnis (1988), analyzing the possible origins for SIPA conclude that the elastodiffusion model proposed by Woo (1984) is the strongest candidate, since it is a first-order effect, involving the migration anisotropy of interstitials in an applied stress field. The climb-assisted glide (I-creep) and SIPA mechanisms have recently been employed by Woo et al. (1990) to derive analytic expressions of their contribution to the single crystal stress compliance tensor and used a self-consistent model that treats each grain as an inclusion embedded in an anisotropic medium, for deriving expressions that relate the polycrystalline creep compliance tensor with that of a single crystal. Using a self consistent model, and a numerical technique, and using the ana-
33
lytic expressions derived by Woo et al. (1990), Christodoulou et al.(1992) derived the relative contributions of basal, pyramidal and prismatic slip systems to the total strain in Zircaloy-2 and Zr-2.5% Nb. It was concluded that -type dislocations account for over 90 % of the total strain. The question of the influence of the microstructure is far from being resolved. Microstructural features that appear to influence creep are the grain shape (in the case of two-phase material the distribution of the P phase), and the dislocation structure. For example, when the grain size is relatively fine, the stress-induced anisotropic diffusion to grain boundaries can contribute significantly to the creep rate. Until the contribution of different mechanisms as a function of the microstructural features of the material is assessed, irradiation creep cannot be analyzed on the basis of a single creep mechanism. 7.3.1.6 Changes in Mechanical Behavior Due to the high concentration of defects produced by neutron irradiation (point defects, dislocations), dislocation slip is inhibited and thus the yield strength increases after irradiation. This effect is rapidly saturated at a fluence which varies with irradiation temperature. For power reactors, the increase of the yield strength (YS) or UTS saturates above about 5 x 10 24 to 10 25 n m" 2 . This saturation value is the same for both SR and RX Zircaloys. This is due to the fact that the initial microstructure no longer controls the deformation mechanisms. The increased density of dislocation loops controls the critical shear stress for dislocation glide. Since the increase in loop concentration under neutron irradiation saturates rapidly due to overlap of absorption volumes for interstitials and vacancies, the increase in yield strength saturates as well.
7 Zirconium Ailoys in Nuclear Applications
34
This increase in yield strength is associated with a reduction in ductility, affecting both the uniform elongation (through a reduction of the strain hardening exponent), and the total elongation, decreasing from about 20% to 2 - 4 % (Morize, 1984). This effect is clearly illustrated in Fig. 7-23 from Price and Richinson (1978). 7.3.1.7 Charged-Particle Irradiation
For the sake of completeness, the use of charged-particles (electrons and ions) for experimental and theoretical studies of ir-
600-
400-
C\J
d 200 A—
10 18 10 19 10 20 10 21 10 22 (a) Integrated neutron flux (£>1 MeV) ( n crrr 2 )
-.0(b)
10 18
10 19
10 20
10 21
10 22
Integrated neutron flux (£>1 MeV) ( n crrr 2 )
Figure 7-23. Effect of irradiation on the mechanical properties: (a) Due to defect build-up, the yield strength (YS) and fracture strength increase with dose, but the effect saturates at about 1024 n • m~ 2 (dashed lines: 10% cold worked, solid lines: annealed yield strength), (b) The same effect occurs on ductility. In the case of cold worked material, the initial high dislocation density masks the effect of irradiation.
radiation effects in Zr alloys is discussed in this section. The motivation for studying irradiation effects with charged-particles is several fold: firstly, because their displacement rates are orders of magnitude higher than under neutron irradiation, equivalent doses in dpa are achieved correspondingly faster; secondly, the effect of experimental variables such as temperature, particle type and dose rate can be studied with relative ease, and, finally, the irradiated samples are usually not radioactive, making them much easier to handle. In general, neither the results obtained in such experiments have a one-to-one correspondence with the results from neutron irradiation nor should such close correspondence be expected in principle, as the irradiations are quite different. Electron irradiation differs from neutron irradiation both in the absence of collision cascades, and in that the electron beam is focused, causing a much higher electron flux and displacement rate. For ion irradiation, the principal difference is that the rate of ion energy loss per unit target thickness is much higher than that for neutron irradiation, which causes only a relatively thin layer close to the surface to be irradiated, albeit at a much higher dose rate. Since irradiation effects are dependent on the balance between irradiation damage and thermal annealing, any shift in the displacement rate can affect the observed effects. Also, bulk effects of neutron irradiation, such as irradiation-induced creep and growth, and irradiation hardening are difficult to study with near-surface irradiation. If, however, the characteristics of each type of irradiation are taken into account, then charged-particle irradiation can provide valuable information on the mechanisms of irradiation damage and microstructural evolution, which can be applied to neutron irradiation.
7.3 In-Reactor Behavior
Most of the effects reviewed in Sec. 7.3.1 have also been studied with charged-particle irradiation, often with different results. Extensive studies of microchemical evolution in Zircaloy-2 and Zircaloy-4 under proton irradiation have been performed by Kai et al. (1990, 1992). It was found that proton irradiation may induce variations in the chemical composition of intermetallic precipitates and increase the solute content in the matrix. These changes appear to increase the nodular corrosion resistance of the alloy. Other investigations of the effect of irradiation on uniform oxidation in Zircaloy-2 and 4 have been conducted by Pecheur et al. (1992) and are reported in Sec.7.3.2. Irradiation-induced segregation and precipitation has been observed in at least two different circumstances. Cann et al. (1992) have observed that 3.6 MeV proton irradiation of annealed Zr-2.5 Nb at 720 K 770 K to 0.94 dpa causes a fine dispersion of Nb-rich platelike precipitates to appear within the oc grains that is absent upon equivalent heat treatment of a control sample, in agreement with the neutron irradiation results of Coleman et al. (1981), reported in Sec. 7.3.1.3. Segregation of tin and precipitation as (3-Sn at the surface of thin Zircaloy-2 foils has also been reported after 5.5 MeV proton irradiation to 1 dpa at 350 K (Motta et al., 1992), a result that has not been observed after neutron irradiation of calandria tubes at equivalent temperatures. Amorphization of intermetallic precipitates in Zircaloy-2 and 4 under charged particle irradiation has been extensively studied, both experimentally (Motta et al., 1989; Lefebvre and Lemaignan, 1989), and theoretically (Motta and Olander, 1990). The critical temperature for amorphization of Zr(Cr,Fe)2 precipitates under electron irradiation was found to be 300 K
35
lower than under neutron irradiation. For Ar ion irradiation, although the critical temperature was similar to that under neutron irradiation, the transformation morphology was quite different, amorphization no longer starting at the precipitate matrix interface. As mentioned above, bulk effects such as irradiation creep and growth and hardening, are more difficult to simulate with charged particle irradiation, but some attempts have been made, especially in the area of irradiation creep and growth. Parsons and Hoelke (1989) developed a "cantilever" beam method, which relates the irradiation-induced creep and growth of a cantilevered Zr beam to its deflection when exposed on one side to 100 keV Ne ion irradiation. The results of the experiment differed from those of neutron irradiation, but no detailed modeling effort has been undertaken as yet to rationalize the differences. When both bulk and near-surface irradiation can be understood in terms of the atomic level mechanisms, then neutron irradiation can be simulated with charged particle irradiation. For example, formation of voids under electron irradiation was observed in a Zr sample that had been previously charged with He (Faulkner and Woo, 1980), and dislocation loop formation can be studied with electron irradiation (Woo et al., 1992). This last study incidentally, supported the idea that microstructural evolution in Zr alloys is affected by the large diffusion anisotropy, as mentioned in Sec. 7.2.3: irradiation-induced voids changed their shape under electron irradiation, preferentially shrinking in the -direction, the expected direction of arrival of the interstitial flux.
36
7 Zirconium Alloys in Nuclear Applications
7.3.2 Corrosion Behavior 7.3.2.1 General Corrosion Behavior
Zr alloys are highly resistant to corrosion in common media and are used for that reason in the chemical industry (Tricot, 1989). Those alloys are however not immune to oxidation and in the high temperature water environment found in a power reactor (280-340 °C at 10-15 MPa), corrosion and hydriding controls the design life of fuel rods and other components. Several international meetings have been organized to discuss this important design parameter, the latest ones having been reviewed by Franklin and Lang (1991). In the early stages, a thin compact black oxide film develops that is protective and inhibits further oxidation. This dense layer of zirconia is rich in the tetragonal allotropic form, a phase normally stable at high pressure and temperature. As described in Fig. 7-24, the growth of the oxide layer thickness d, follows a power law, usually described by a quadratic relationship (7-3)
docjt
The activation energy of 130 kJ • mol x (1.35 eV • at" 1 ) for corrosion in the dense oxide regime (Billot et al., 1989) is equivalent to activation of the diffusion of oxygen in zirconia (Smith, 1969). Oxygen is considered to diffuse from the free surface of the oxide as O 2 ", by a vacancy mechanism through the zirconia layer, and to react with the zirconium at the matrix-oxide interface, but recent studies support the fact that those ions diffuse mostly through grain boundaries (Godlewski et al., 1991). For very thin oxide layers, the zirconia grows in epitaxy on the metallic zirconium (Ploc, 1983). The Pilling-Bedworth ratio, or ratio of the oxide volume to the parent metal volume, is equal to 1.56 for zirconia. Thus as the oxide grows, the stress buildup due to the volume expansion associated with oxidation induces a preferential oxide orientation that reduces the compressive stresses in the plane of the surface (David et al., 1971). This gives rise to various fibrous textures. Those compressive stresses are one of the factors that explain the stabilization of the tetragonal form of zirconia in this layer. A chemical effect of the alloying elements could also be invoked to explain that stabilization.
200 180
o Alpha annealed (RXA) _ O Stress relieved annealed (SRA)
SRA! —
Least square fitted line
£ or Y ^ a t a spfefld around average - (Summary of 1600 data points)
-~ 120
Figure 7-24. Uniform corrosion behavior of Zircaloy-4 in water at 633 K. The oxide thickness follows a power law up to a transition (1.5 to 2 mm thickness) and then remains linear. The metallurgical state of the material affects also the corrosion rate (Clayton and Fisher, 1985).
§>100 D) 80 ^
60 40 20 0 200 300 Time (days)
400
500
7.3 In-Reactor Behavior
As the oxidation proceeds, the compressive stresses in the oxide layer cannot be counterbalanced by the tensile stresses in the metallic substrate and plastic yield in the metal limits the compression in the oxide. The tetragonal phase becomes unstable and the oxide transforms to a monoclinic form (Godlewski et al., 1991). This martensitic-type transformation is associated with the development of a very fine interconnected porosity that allows the oxidizing water to access closer to the corrosion interface (Cox, 1969). The size of those pores has been measured to be very small: Cox (1968) has used a modified mercury porosimeter to determine pore sizes smaller than 2 nm. His work has been confirmed by nitrogen absorption kinetics to pore sizes smaller than 1 nm, corresponding to a volume fraction of the order of 1% (Ramasubramanian, 1991). Once this transition has occurred, only a portion of the oxide layer remains protective. The corrosion kinetics are therefore controlled by diffusion of oxygen only through the dense protective oxide layer next to the metal substrate. Since the thickness of this layer remains constant, in the range of 1 |im (Garzarolli et al., 1991), the corrosion rate is constant after this transition (Fig. 7-24). In this dense oxide layer the structure of the zirconia, which controls the posttransition corrosion kinetics, is complex and still under discussion. Starting from the metaloxide interface, a very thin layer of amorphous oxide has been reported under particular conditions, about 10 nm in thickness (Warr etal., 1991). The existence of epitaxy shows that this layer, if present, should not be continuous. It is followed by a zone of very small zirconia crystallites, 10-20 nm, that become larger in diameter and columnar in shape further into the oxide layer (Bradley and Perkins, 1989).
37
For thick oxide layers, in excess of 50 |im, the oxide may spall, leaving zirconia particles free to flow in the cooling water, and giving rise to a much thinner oxide film. On the one hand, this process could be beneficial since it reduces the metal-oxide interface temperature, which is the parameter controlling oxidation kinetics. However, design considerations on remaining cladding thickness after oxidation does not allow for massive spalling in commercial reactors. Other undesirable consequences of spalling are the buildup of activity in the coolant and safety considerations, since those particles can interfere with the functioning of valves. In BWRs, as shown in Fig. 7-25, nodular corrosion is the limiting design consideration. This behavior, specific to the boiling water reactor, can be reproduced by testing alloys in steam at 500 °C (Schemel, 1987). Several mechanisms have been proposed for the nucleation of those nodules, leading to various possible sites for nodule nucleation: metallic matrix grain boundaries, local rupture of the continuous dense oxide at an early stage of growth, local variations in composition and precipitate densities or crystallographic orientation of clusters of grains (Charquet et al., 1989b; Ramasubramanian, 1989). As the corrosion progresses, the nodules grow in size and thickness and their number density increases leading to a complete coverage of the metal. The P quench structure described in Sec. 7.2.2.2 significantly improves the resistance of Zircaloy-4 to nodular corrosion, however as higher burnups are reached uniform corrosion could then become a problem. 7.3.2.2 Oxidation of the Precipitates Since most of the metallic alloying elements present in Zircaloys are added for
38
7 Zirconium Alloys in Nuclear Applications
(a)
(b)
(c)
Figure 7-25. (a) Typical aspect of nodular corrosion of Zircaloy-4 obtained in a 500 °C steam test, (b) Higher magnification of a cracked nodule in Zircaloy-4, obtained during a 10.3 MPa steam test at 500 °C for 25 h, showing the extensive oxide cracking associated with the process. In addition to circumferential cracks in the flanks of the nodules, a vertical crack can also be seen running between the nodules [Courtesy of NFIR (Nuclear Fuel Industry Research Group)], (c) Greater detail of the nodule, showing that the upper portion of the nodule consists of a series of parallel, cracked, oxide layers (Courtesy of NFIR).
improving corrosion resistance, the mechanism of their interaction with the oxidation front is of great importance. Due to the fine structure of the precipitates no specific observation of this process was available until recently, when a few studies have been performed using advanced STEM. It was shown that the precipitates are incorporated in metallic form into the oxide layer, and oxidize only afterwards, deeper into the oxide layer. In particular, iron has been shown to remain unoxidized in the dense oxide layer (Garzarolli et al., 1991; Pecheur etal., 1992). The precipitates slowly release part of their iron in the oxidized matrix so oxide chemistry changes as the oxidation proceeds. It is found that the oxidation of iron coincides with the oxide transition from tetragonal to monoclinic. Also, some of the crystalline intermetallic precipitates are found to be amorphous after incorporation into the oxide layer (Pecheur et al., 1992). This transformation is not caused by irradiation since those are also observed in nonirradiated oxidized material. The crystal-to-amorphous transformation in the oxide layer could be caused by changes in chemistry, notably hydrogen intake (the nearest neighbor distance in the amorphous phase is bigger than in the corresponding precipitates amorphized by irradiation in the Zr matrix). The effect of the release of iron in the oxide is still under consideration. A correlation has been found between the oxidation of the precipitates and the tetragonalto-monoclinic transformation of the zirconia, but without clear knowledge of its origin (Godlewski et al., 1991). A chemical effect on the stabilization of the tetragonal phase of the oxide may be present. An additional point is the role played by the precipitates as possible short circuits
7.3 In-Reactor Behavior
for the electron current in the oxide film. In order to balance the charges transported by the O 2 ~ ions across the film, electrons have to flow from the metal to the water. As zirconia is a good insulator, metallic precipitates could enhance this flow and be sites of preferential corrosion, a process considered for nodule nucleation (Cox, 1989). A full description of the role played by the precipitates in the oxidation process is still not available. To date only partial, and often unexplained, observations have been made. For instance, a surprising observation is the opposite impact of precipitate size on corrosion rate in PWR's and BWR's as reported by Garzarolli and Stehle (1986). As seen in Fig. 7-26, the best microstructure for resisting localized corrosion in a BWR consists of fine precipitates (diameter below 0.15 jim ), whilst in PWR's, precipitates smaller than 0.1 jxm increase the uniform corrosion rate.
CO
39
1
0.02 "-S 30-
*r
out-of-pile o
10 •
/500°C/16h ' \
1
*
^
° ^ ^o3£aoLo 350°C/
1a -
0.02 0.1 0.8 Average diameter of precipitates (|iim) Figure 7-26. Influence of precipitate size on localized and uniform corrosion: The optimal microstructure for corrosion resistance is strongly dependent on corrosion conditions (Garzarolli and Stehle, 1986).
7.3.2.3 Water Radiolysis The interaction of irradiation particles (neutron, gamma, beta) with the coolant water leads to the formation of radiolytic oxidizing species, that are assumed to be responsible for the large increase in corrosion rate of Zr alloys in reactor, compared to similar out-of-reactor conditions. Water radiolysis in the reactor core has been analyzed in great detail with the objective of knowing the steady state concentration of the radiolytic species during irradiation and their possible effect upon the zirconium corrosion rate (e.g., Burns and Moore, 1976). The main process of radiolysis is an instantaneous decomposition of the water molecules by interaction with the electrons in spurs (small volume of high interaction along the path of the electron), giving birth to metastable species that re-
combine in a large variety of possible ways. The complexity of the recombination reactions can be illustrated by the large number (35 to 40) of reactions to be considered, each with its own rate. The concentrations of the intermediate and final products depend strongly on irradiation rate and initial conditions and therefore, local changes in irradiation intensity lead to drastic changes in the concentrations of radiolytic species. Some species appear for very short times as intermediate steps to long-lived species (Lukas, 1988). Computer programs are thus used for the computation of the evolution of those species versus initial chemical and irradiation conditions (Buxton and Elliott, 1991). One way to decrease the build-up of oxidizing radiolysis products is to add hydro-
40
7 Zirconium Alloys in Nuclear Applications
gen to the cooling water. Then a general reduction of stable radiolytic species occurs and a reduction in corrosion rate of thin oxide ensues. However this effect is not as efficient in boiling conditions compared to pressurized ones, due to the segregation of H 2 in the steam phase (Ishigure et al., 1987). Radiolytic enhancement of corrosion can also occur in the case of PWR's, in post transition thick oxide films, or during corrosion testing of coupon specimens in a reactor corrosion loop. There, although no steam is expected to be present, the water chemistry in the pores of the thick oxide should be much like that present in BWRs, i.e., a two-phase regime (Johnson, 1989). Some particular cases of localized corrosion enhancement have recently been linked to an increase in metastable oxidizing species due to P- radiolysis. The irradiation enhancement of corrosion is well observed for thick oxide films, where a threeto four-fold enhancement is commonly seen (Marlowe et al., 1985). The reported occurrences are characterized by the presence of dissimilar materials in the vicinity of the corroding material. As recently reviewed, enhanced corrosion has been observed in front of stainless steel, copper cruds, platinum inserts and in the case of gadolinia bearing poison rods (Lemaignan, 1992). In all the reported cases, strong p emitters are present and the local energy deposition rates of those particles within the coolant are much larger than the bulk radiolysis contribution due to the neutrons and gammas. Thus, in reactor, the local intensity of the radiolysis has to be considered, as it will change the local chemistry for alloy corrosion. In addition to the radiolysis described above, corresponding to chemical evolution by reaction in the bulk, specific considerations should be given to the fact that
the pores have a large surface to volume ratio, leading to additional reactions at the surface of the pores, instead of a recombination between them as in the bulk. For instance, the H 2 O 2 molecules can be adsorbed on the surface of the pores leading to the reaction:
instead of recombining in two steps with 2 H as the standard back reaction to water after radiolysis. This leads to a higher oxygen potential at the surface of the ZrO 2 , giving a higher corrosion rate.
7.3.2.4 Hydrogen Pickup During oxidation of Zr alloy components in reactors or in autoclaves, the reduction of water by the Zr alloy follows the general reaction scheme:
As described above, the reaction proceeds in several steps and the oxidation progresses at the metal-oxide interface by diffusion of O 2 " ions through the oxide. The reduction of the water molecules at the coolant-oxide interface releases hydrogen as radicals H + . They are chemically adsorbed at the tips of the oxide pores and their evolution controls the behavior of the hydrogen. Most of them recombine, creating hydrogen molecules that escape through the pore and dissolve into the coolant. A limited amount can ingress in the oxide and diffuse through to the metallic matrix and then react with Zr for the formation of hydrides, when the terminal solid solubility is exceeded. Corrosion experiments performed using tritium-doped water have shown that the H dissolved in the coolant is not trapped by the matrix, but that radicals obtained dur-
7.3 In-Reactor Behavior
ing the reduction of water are necessary for hydrogen pick-up (Cox and Roy, 1965). The diffusion coefficient of H in ZrO 2 has been measured with difficulty because of its very low value and because of the contribution of grain boundary diffusion and surface diffusion to diffusion in the pores of the oxide. Recent measurements by Khatamian and Manchester (1989) have confirmed earlier results in the range of 1(T 17 to 10~ 1 5 m 2 - s " 1 at 400 °C, depending on the condition of measurements and chemical composition of the alloy on which the oxide is grown. With such low values, the penetration of H in the oxide is limited, and the zirconia layer is indeed protective with respect to H ingress. Care should be taken not to introduce catalyzers of the hydrogen molecule dissociation reaction that could enhance H pick-up (or hydrogen uptake) - i.e., the fraction of the hydrogen produced by reduction of water that is trapped into the Zr alloy. In that regard, Ni additions are to be avoided. This is the main reason for suppression of that element in Zircaloy-4. By reducing Ni, hydrogen pick-ups are generally in the range of 15% or less for standard PWR fuel cladding. BWR values tend to be around 20%. In the case of CANDU pressure tubes, another mechanism of H pick-up has to be considered: the H in solution within the coolant diffuses through the stainless steel end fittings either directly into the pressure tube, or into the annulus gap. In the latter case it is then in contact with the outer part of the pressure tube. In order to avoid interaction with the pressure tube, a remaining ZrO 2 layer has to be maintained. The size of the protective oxide is controlled by O dissolution into the metal (Cheadle and Price, 1989), and therefore oxidizing agents are added to the gas annulus to replenish oxygen in the oxide layer and pre-
41
vent its degradation and enhanced H pickup (Urbanic et al, 1989) . One of the consequences of hydrogen ingress into Zr can be delayed hydride cracking (DHC) (Cheadle et al., 1987). The mechanism involves the ingress of hydrogen into a Zr-alloy component, its migration up a stress or thermal gradient and its concentration in the regions of low temperature or higher tensile stress. When the local concentration exceeds the terminal solid solubility (which happens first in colder regions), the hydride phase precipitates. As can be seen in Fig. 7-9, the terminal solid solubility of H in Zr is quite low, e.g., 80-100 ppm at 300°C. At low temperature, the hydrides crack when the stresses are high enough, the mechanism then being repeated until failure. The concentration of hydrogen in a tube stressed in bending can be seen in Fig. 7-27 a, where the variation of hydride orientation along the tube thickness (and therefore stress state) is shown. The inner part of the tube is under tensile hoop stresses and the hydrides precipitate along radial planes, while in the outer part of the tube hydrides tend to precipitate along the habit plane, i.e., a more tangential direction. Figure 7-27 b shows hydride formation at the tip of a crack due to local stress buildup. Pressure tubes in the CANDU Pickering 3 and 4 reactors failed by this mechanism. With respect to this behavior, specifications may be required on cladding or pressure tube texture, usually measured by the morphology of hydride precipitation upon intentional addition of hydrogen. 7.3.3 Pellet-Cladding Interaction Pellet-cladding interaction (PCI) associated with iodine intergranular stress corrosion cracking (IGSCC) is a mode of fuel rod failure that has been observed after
42
7 Zirconium Alloys in Nuclear Applications
(a)
(b)
Figure 7-27. (a) Hydriding/D2 pickup, showing stress dependence of hydride formation. The inner part of the tube is under tensile hoop stresses and the hydrides precipitate along radial planes, while in the outer part of the tube hydrides tend to precipitate along their habit plane, i.e., in a more tangential direction (micrograph courtesy of C. E. Coleman, AECL/ Chalk River), (b) Hydride formation at a crack tip in Zr-2.5 Nb due to the local buildup of stresses (photo courtesy of D. Rogers, AECL/Chalk River).
fast variations in the linear heat generation rate (LHGR) (i.e., the thermal power released per unit length of fuel rod, typically 15-25 kW • m" 1 ), in fuel that has undergone significant burn-up (BU). The first reported observation of PCI failure on BWRs was rapidly diagnosed as a mechanical interaction between the cladding and the UO 2 pellet, associated with chemical interaction of some fission products with the Zircaloy cladding (Rosenbaum, 1966). This led to a spreading interest to identify other types of reactors where this problem may have been of concern. This subject has been reviewed recently by Cox (1990), see also Chap. 3. The main mechanisms invoked to explain the failure are described schematically in Fig. 7-28. A combined contribution of stresses induced by fuel pellet expansion due to LHGR increase and the presence of an active corrosion agent, the iodine, created in the fuel rod as fission product, induces failure by stress corrosion cracking (SCC). The occurrence of the problem led the international fuel community to set up large R & D programs. There were two main objectives:
7.3 In-Reactor Behavior
(1) For practical purposes, design parameters were analyzed on integral tests on fuel rods: Fuel rods of various design and irradiation histories were tested in irradiation devices of test reactors. After a short irradiation at the LHGR close to the one used in the power reactor, the irradiation power was increased at a given rate and the behavior of the rod was analyzed with respect to its capacity to support the change in LHGR without failure. This type of experiment gives information on the maximum power allowable for a fuel rod, usually expressed as a function of the BU. For fresh fuel rods, an open gap exists in operation and large power changes are acceptable: Maximum power levels
Figure 7-28. Cladding strain induced by pellet thermal expansion upon increase of linear heat generation rate (after Levy, 1974).
43
above 50 kW m * are common. This limit decreases slowly with BU up to 20 GW • d • t" 1 , when it stabilizes in the range of 40-45 kW • m" 1 . Various examples of the testing procedures used and of the results obtained can be found in the proceedings of a series of IAEA specialist meetings, e.g., IAEA (1982). (2) For understanding mechanisms and for remedial aspects, analytical work was carried out in laboratories on the basic mechanisms involved: The analytical work was focused on the SCC behavior of fuel cladding materials. This aspect of the problem is discussed in more detail here. The stresses are induced by the thermal expansion of the pellet during power transients. During steady state operation, the pressure of the coolant causes the cladding to creep down to the fuel pellet. In addition, fuel swelling at low temperature, due to fission products, contributes to gap closure. In PWRs this phase may take a year or two, while BWRs keep their gap open longer. Once the gap is closed, any change in pellet dimension is transferred to the cladding. For a standard power change, typically from 20 to 40 kW • m" 1 , the change in centerline temperature induces a thermal expansion of about 0.25% for each 10 kW • m~ *. This strain is high enough to induce stresses close to or above the yield strength (Porrot et al, 1991). During the time at which the fuel is maintained at high power, a stress level of one half to two thirds of the yield strength and the presence of a corrosive agent are sufficient to initiate and propagate a crack that leads to fuel rod failure. In addition to iodine (I), fission products such as cesium (Cs) and cadmium (Cd) were suspected early on as candidate corroding agents. Those elements are
44
7 Zirconium Alloys in Nuclear Applications
known to induce intergranular fracture of Zr alloy by a liquid metal embrittlement mechanism. The type of fracture surface morphologies was a main reason for rejection of those species. On the other hand, the higher fission yield of Cs compared to I, introduced a discussion on the availability of this agent for interaction with the stressed cladding, due to the low vapor pressure of iodine in equilibrium with solid Csl at fuel temperature. The intense radiolysis due to the fission recoils in the interior of the fuel rod, has been shown to be able to increase the effective pressure of iodine in equilibrium with Csl to a level where SCC has been shown to occur (Konashi, 1984). In order to understand the mechanisms involved in iodine-induced SCC of zirconium alloys, tests were performed on various types of materials in different conditions. The most frequent tests were either constant load tests, where the time to failure is expressed versus the applied stress in the iodine environment, or constant strain rate tests, where the reduction in ductility due to iodine is analyzed (Brunisholtz and Lemaignan, 1987). A reduction in time or strain to failure is observed and several parameters have been shown to be of importance: internal surface condition, metallurgical state and texture of the material. In an early experiment performed by Peehs et al. (1979), the effect of basal plane orientation with respect to stress, was rec-
basal pole orientation
ognized to be critical: machining a tube out of a thick Zircaloy plate, they were able to obtain an angular variation of the texture around this test sample. After testing under iodine, the density of cracks was found to be strongly dependent upon the relative orientation of the basal planes. The densities of cracks were higher when the oplanes were in the direction of the crack growth (Fig. 7-29). To explain those results, the characteristic aspect of SCC fracture surface has to be analyzed in order to explain the mechanism of rupture. Figure 7-30 shows a SEM view of a recrystallized Zircaloy tested in tension in iodine environment. The fracture surface is transgranular and consists mostly of large transgranular pseudocleavage areas where fracture occurs by SCC interconnected by fluted walls in which plastic deformation is evident. Crystallographic analysis of the pseudo-cleavage areas have shown that they consist of basal planes, while the fluted walls are located on prismatic planes. The propagation of the SCC fracture surface on the basal plane is enhanced by the strong decrease in surface free energy of this plane when iodine is adsorbed on it (Hwang and Han, 1989). Fluted walls are obtained by plastic deformation on the primary slip system to connect the different pseudocleavage planes in a ductile type rupture mode. Due to the perpendicular directions of the slip on the prismatic plane and of the
angular distribution of crack density
crack density, relative
Figure 7-29. Density of crack developed during an iodine SCC test of a tube machined out of a plate. Along the circumference, the orientation of the -planes changes from radial to tangential. The cracks tend to develop when they can grow following the c-planes (Peehs et al., 1979).
7.3 In-Reactor Behavior
Figure 7-30. Fracture surface of RX Zr by SCC cracking in iodine. The crack grows by pseudo-cleavage on the basal plane and those brittle areas are connected by fluted prismatic planes on which ductile rupture occurred.
pseudo-cleavage on the basal plane, no plastic deformation can contribute to the reduction of tensile stresses on this plane. Thus the relative orientation of the basal plane with respect to the applied tensile stresses is a critical parameter, and the effect of the texture is remarkable. Constant stress and fracture mechanics tests have indeed shown that when the otype planes tend to be aligned with the macroscopic crack surface, the susceptibility to SCC increases (Knorr and Pelloux, 1982). For cladding tubes, where the tensile stresses are the hoop stresses due to the pellet expansion, the best texture corresponds to a maximum intensity of the odirection in the radial direction. Figure 7-31 shows the susceptibilities to SCC in iodine environment of a series of claddings with different textures but processed from the same ingot, plotted as a function of the angle between the maximum intensity of the cpoles and the radial direction (Schuster and Lemaignan, 1992). It is clear that the more the basal planes are aligned with the radial direction, the less susceptible is the cladding to SCC. The stress intensity fac-
45
tor (Klscc) for SCC crack growth was found to be dependent upon the same parameters. Measurements on different orientation of cracks in Zircaloy plates or crack propagation through the wall of different claddings gave ^ ISCC in the range of 3.5 to 6 MPa • m 1/2 . The lowest Klscc corresponds to the highest orientation of the oplanes in the propagation direction. Due to the fast SCC crack growth rates measured, in the range of 1 to 10 jam • s" 1 , the nucleation step is the longest part of the SCC lifetime (Brunisholtz and Lemaignan, 1987); therefore it controls the overall SCC behavior. The intermetallic precipitates have been suggested to be places of crack initiation as preferential sites for iodine corrosion (Kubo et al., 1985). Another action of the precipitates related to crack initiation would be, in the case of intergranular precipitates, the inhibition of grain boundary sliding thereby causing a local stress buildup which enhances crack formation. The last mechanism to be considered is the large strain incompatibilities, from one grain to the next one during plastic deformation, found in Zr alloys, due to the limited number of slip systems available. Due to the fact that plastic deformation on the prismatic system does not relax
100 80
o ^ 60 0
of* S
20 15
20 25 30 35 40 45 50 Angle of the c-direction with the radial direction of the tube (degree)
Figure 7-31. Effect of texture on SCC susceptibility in iodine. The same alloy was used to process tubes of different textures, and the susceptibility to I-SCC is found to decrease for a more radial texture.
46
7 Zirconium Alloys in Nuclear Applications
stresses for crack formation on the basal plane, those incompatibilities are efficient mechanisms for crack initiation (Kubo etal., 1985). An early solution to the PCI problem was to reduce the power change rates in order to reduce the hoop stresses in the cladding by relaxation during the loading time. Specific procedures have been implemented with success, but the drawback is a loss of power availability, driving R & D work to a solution obtained by fuel design. For practical purposes, remedies have been found and tested to avoid PCI type failures. In the reactors most prone to this phenomenon, advanced claddings have been developed with success. For the CANDU reactors, where the fuel bundle maneuvers during on power refuelling induce large local changes in LHGRs, the "CANLUB" design consists of an internal coating of graphite in internal surface of the cladding (Wood and Kelm, 1980). "Barrier" fuel rods for BWRs contain an internal layer of pure recrystallized Zr that is coextruded with the Trex and cold rolled with the base Zircaloy metal to obtain a perfect metallurgical bond (Rosenbaum et a l , 1987).
7.4 Challenges At the end of this review of the current knowledge of the behavior of Zr alloys for nuclear application, it may be of interest to list some open questions for a better understanding of the physical mechanism involved in specific properties. Indeed although the use of Zr alloys has proven a good engineering solution, several properties are still subject to a poor scientific understanding of their origins. Among them, the following topics, related to deformation, point defects and corrosion, can be
proposed for further basic work: • Due to the high anisotropy of the Zr matrix, and the limited number of deformation systems available, thermomechanical processing gives rise to the textures described in Sec. 7.2.3.3. Although the texture obtained after a given thermo-mechanical processing route is well known, its accurate prediction is impossible today for a new alloy or a major change in processing. Improvement in the understanding of the deformation of Zr alloys could be focused on the different shear and twin systems active as a function of temperature, stress distribution, alloy composition and structure, irradiation fluence. The knowledge of the critical shear stresses for their activation is extremely limited. If well measured, a consequence would be a better evaluation of the local strain during plastic deformation or thermal treatments (e.g., strain incompatibility, local crystal rotation, texture development). In addition the interaction of the cold-work dislocations with the irradiation defects is probably responsible for the reduction of strain anisotropy after irradiation, by activation of new slip systems. This has to be analyzed. In connection with in-reactor deformation, the growth phenomenon is one of the limiting design considerations of the structural components. The reason for the occurrence of the -type dislocations after a fluence o f 3 x l 0 2 5 n m ~ 2 and their impact on growth, is still an open question, the contribution of minor alloying elements in stabilizing those dislocations being suspected. Detailed understanding of this process is highly desirable. • In the field of point defects, one should consider the alloying elements in solid
7.6 References
solution and irradiation-induced point defects. The complex diffusion behavior of the foreign atoms, either in substitution or as interstitials, leads one to suspect many cases of solute-defect interaction, giving rise to irradiation induced segregation or accelerated diffusion. With respect to that behavior, the current effort to compute the formation and migration energies of different defect configurations involving one or several atoms, using molecular dynamics and to obtain experimental data on point defect energies and mobilities should be continued. In particular, the anisotropy of diffusion is an important parameter for irradiation growth and creep behavior. • For corrosion, although a general scheme of the phase transformation in the zirconia responsible for the transition is available, the details of the mechanisms are still to be described accurately. In particular, the effect of alloying elements and their distribution in the Zr matrix is of importance. The localized corrosion rate enhancement under irradiation can be correlated to local radiolysis due to p flux, but a quantitative model is still unavailable. In the same area, in order to reduce the hydrogen pick-up during corrosion, the process should be studied in more detail, in particular the transport mechanism through the zirconia film. Zirconium alloys are paralleled only by steels in the large amount of irradiation data available. However, their irradiation behavior is more complex than that of stainless steel, due to the anisotropy of the h.c.p. Zr and to the behavior of the alloying elements. The increasing amount of basic data available leads to the hope that the prediction of the in-reactor behavior will
47
soon come from mechanistic models based on a more fundamental understanding of the processes involved, rather than from empirical correlations.
7.5 Acknowledgements This work was completed while one of us (ATM) was at AECL/Chalk River Laboratory, and this review was enriched by the comments and criticism from several scientists there: Vince Urbanic and Malcolm Griffiths carefully reviewed the manuscript; and Rick Holt, Gavin Hood, Nick Christodoulou and Robert Ploc made special contributions to various specific sections. Daniel Charquet, of Cezus, and John Harbottle of the Stoller Corporation also provided comments on the manuscript. The authors sincerely thank all of them.
7.6 References Abriata, I P., Bolcich, I C. (1982), Bull. Alloy Phase Diagrams 3(1), 1710-1712. Abriata, J. P., Garces, X, Versaci, R. (1986), Bull Alloy Phase Diagrams 7(2), 116-124. Adamson, R. B. (1977), ASTM-STP633. Philadelphia, PA: Am. Soc. Test. Mat., pp. 326-343. Arias, D., Abriata, I P. (1986), Bull. Alloy Phase Diagrams 7(3), 237-243. Arias, D., Abriata, J. P. (1988), Bull. Alloy Phase Diagrams 9(5), 597-604. Armand, M., Givord, J. P., Tortil, P., Triollet, G. (1965), Mem. Sclent. Rev. Metall. 62, 275-283. Asada, T., Komoto, H., Chiba, N., Kasai, Y. (1991), 9th Int. Conf. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat, pp. 99-118. ASTM (1990), Standard B 353.89. Philadelphia, PA: Am. Soe. Test Mat. Bacon, R. (1988), J. Nucl. Mater. 159, 176-189. Baig, M. R., Messoloras, S., Stewart, R. J. (1989), Rad. Effects 108, 355-358. Bangaru, N. Y. (1985), J. Nucl. Mater. 131, 280-290. Bhanumurthy, K., Kale, G. B., Khera, S. K. (1991), J. Nucl. Mater. 185, 208-213.
48
7 Zirconium Alloys in Nuclear Applications
Billot, P., Beslu, P., Giordano, A., Thomazet, J. (1989), 8th Int. Conf. Zr Nucl. Ind., June 19-23, 1988, San Diego, CA: ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat., pp. 165-184. Boman, L. H., Kilp, G. R., Balfour, M. G., Schoenberger, G. (1991), 9th Int. Conf. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132, Philadelphia, PA: Am. Soc. Test. Mat., poster session. Bradbrook, J. S., Lorimer, G. W., Ridley, N. (1972), /. Nucl. Mater. 42, 142-160. Bradley, E. R., Perkins, R. A. (1989), Proc. Tech. Com. Meet. IAEA: Fundamental Aspects of Corrosion in Zr-Base Alloys in Water Reactor Environments, Sept. 11-15, 1989, Portland, OR, pp. 101106. Brailsford, A. D., Bullough, R. (1972), J. Nucl. Mater. 44, 121-135. Brunisholz, L., Lemaignan, C. (1987), 7th Int. Conf Zr Nucl. Ind.: ASTM-STP 939. Philadelphia, PA: Am. Soc. Test. Mat., pp. 700-716. Buckley, S. N. (1961), Properties of Reactor Materials and the Effects of Irradiation Damage: Litter, D. J. (Ed.). London: Butterworths, p. 443. Buckley, S. N., Bullough, R., Hayns, M. R. (1980), /. Nucl. Mater. 89, 283-295. Bullough, R., Willis, J. R. (1975), Phil. Mag. 31, 855861. Burgers, W. G. (1934), Physica 1, 534-546. Burns, W. G., Moore, P. B. (1976), Rod. Effects 30, 233-242. Buxton, G. V., Elliott, A. J. (1991), Proc. JAIF Int. Conf Water Chem. Nucl. Power Plants, Apr. 22-25, 1991, Fukui City, Japan. Carpenter, G. J. C , Northwood, D. O. (1975), /. Nucl. Mater. 56, 260-266. Cann, C. D., So, C. B., Styles, R. C , Coleman, C. E. (1992), Precipitation in Zr-2.5 Nb Enhanced by Proton Irradiation, presented at: Conf. Microstructural Evolution Metals During Irradiation, Muskoka, Canada, Sept.-Oct. 1992. To be published in J. Nucl. Mater. (1993). Causey, A. R., Holt, R. A., McEwen, S. R. (1984), 6th Int. Conf. Zr Nucl. Ind., June 28-July 1, 1982, Vancouver, Canada: ASTM-STP 824. Philadelphia, PA: Am. Soc. Test. Mat., pp. 269-288. Charquet, D., Alheritiere, E. (1985), Proc. Workshop: Second Phase Particles in Zircaloys, Erlangen, F.R.G. Kerntechnische Gesellschaft, pp. 5-11. Charquet, D., Hahn, R., Ortlieb, E., Gros, I P . , Wadier, I F. (1989 a), 8th Int. Conf. Zr Nucl. Ind., June 19-23, 1988, San Diego, CA: ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat., pp. 405-422. Charquet, D., Tricot, R., Wadier, J. F. (1989b), 8th Int. Conf. Zr Nucl. Ind., June 19-23, 1988, San Diego, CA: ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat., pp. 374-391. Cheadle, B. A. (1975), The Physical Metallurgy ofZr Alloys, CRNL Report, 1208.
Cheadle, B. A., Alridge, S. A. (1973), /. Nucl. Mater. 47, 255-258. Cheadle, B. A., Price, E. G. (1989), IAEA Tech. Com. Meet. Exch. Operational Safety Exp. Heavy Water Reactors, Feb. 20-24, 1989, Vienna, Austria, AECL 9939. Cheadle, B. A., Coleman, C. E., Ambler, J. F. (1987), 7th Int. Conf. Zr Nucl. Ind., ASTM-STP 939. Philadelphia, PA: Am. Soc. Test. Mat., pp. 224240. Christodoulou, N., Causey, A. R., Woo, C. H., Klassen, R. J. (1992), 10th Conf. Zr Nucl. Ind., Philadelphia, PA, 1993: ASTM-STP 1175. Philadelphia, PA: Am. Soc. Test. Mat., submitted. Clayton, J. C , Fischer, R. L. (1985), Proc. ANS Topical Meet. Light Water Reactor Fuel Performance, Orlando, FL, April 21-24 1985. La Grange Park, IL: ANS, pp. 3-1-3-5. Coleman, C. E., Gilbert, R. W, Carpenter, G. J. C , Weatherly, G. C. (1981), in: Phase Stability Under Irradiation, AIME Symp. Proc: Holland, J. R. et al. (Eds.). Metals Park, OH: Am. Inst. Mech. Eng. Cox, B. (1968), J. Nucl. Mater. 27, 1-47. Cox, B. (1969), J Nucl. Mater. 29, 50-66. Cox, B. (1989), Proc. Tech. Com. Meet. IAEA: Fundamental Aspects of Corrosion in Zr Base Alloys in Water Reactor Environments, Sept. 11-15, 1989, Portland, OR, pp. 167-173. Cox, B. (1990), J. Nucl. Mater. 172, 249-292. Cox, B., Roy, C. (1965), The Use of Tritium as Tracer in Studies of Hydrogen Uptake by Zr Alloys, AECL-2159. David, G., Geschier, R., Roy, C. (1971), /. Nucl. Mater. 38, 329-339. Dawson, C. W, Sass, S. L. (1970), Met. Trans. 1, 2225-2233. Douglass, D. L. (1971), The Metallurgy of Zirconium, Atomic Energy Review, Vienna. Farrell, K. (1980), Rad. Eff. 53, 175-194. Faulkner, D., Woo, C. H. (1980), /. Nucl. Mater. 90, 307-316. Fidleris, V. (1988), /. Nucl. Mater. 159, 22-42. Franklin, D. G. (1982), 5th Int. Conf Zr Nucl. Ind., Aug. 4-7, 1980, Boston, MA: ASTM-STP 754. Philadelphia, PA: Am. Soc. Test. Mat., pp. 235267. Franklin, D. G., Adamson, R. B. (1988), /. Nucl. Mater. 159, 12-22. Franklin, D. G., Lang, P. (1991), 9th Int. Conf Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 3-32. Fuse, M. (1985), J. Nucl. Mater. 136, 250-277. Garzarolli, F , Stehle, H. (1986), IAEA Symp. Improvements Water Reactors Fuel Tech. Utilization, Stockholm, Sweden: IAEA SM 288124, pp. 387407.
7.6 References
Garzarolli, R, Seidel, H., Tricot, R., Gros, X P. (1991), 9th Int. Conf. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 395-415. Gilbert, R. W., Griffiths, M., Carpenter, G. I C. (1985), /. Nucl. Mater. 135, 265-268. Godlewski, I , Gros, X P., Lambertin, M., Wadier, X R, Weidinger, H. (1991), 9th Int. Conf. Zr Nucl Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 416436. Griffiths, M. (1988), /. Nucl. Mater. 159, 190-218. Griffiths, M. (1989), /. Nucl. Mater. 165, 315-317. Griffiths, M., Gilbert, R. W, Carpenter, G. X C. (1987), / Nucl. Mater. 150, 53-66. Griffiths, M., Gilbert, R. W., Coleman, C. E. (1988), J. Nucl. Mater. 159, 405-416. Griffiths, M., Gilbert, R. W., Pidleris, V. (1989), ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat., pp. 658-677. Gros, X P., Wadier, X R (1989), Proc. Tech. Com. Meet. IAEA: Fundamental Aspects of Corrosion in Zr Base Alloys in Water Reactors Environments, Sept. 11-15, Portland, OR, pp. 211-225. Guibert, E. R., Duran, S. A., Bement, A. L. (1969), ASTM-STP 456. Philadelphia, PA: Am. Soc. Test Mat, pp. 210-226. Harbottle, X (1978), Phil. Mag. 38, 49. Holt, R. A. (1988), J. Nucl. Mater. 159, 310-338. Holt, R. A., Causey, A. R. (1987), J. Nucl. Mater. 150, 306-318. Holt, R. A., Pleck, R. G. (1991), 9th Int. Conf. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 218-229. Holt, R. A., Gilbert, R. W. (1983), J. Nucl. Mater. 116, 127-130. Holt, R. A., Gilbert, R. W. (1986), / Nucl. Mater. 137, 185-189. Hood, G. M. (1988), J. Nucl. Mater. 159, 149-175. Hood, G. M., Schultz, R. X (1974), Acta Met. 22, 459-464. Hood, G. M., Zou, H., Schultz, R. X, Roy, X A., Jackman, X A. (1992), J Nucl. Mater. 189, 226230. Hwang, S. K., Han, H. T. (1989), /. Nucl. Mater. 161, 175-181. IAEA (1982), Specialist Meeting on Power Ramping and Cycling Behavior of Water Reactor Fuel, Sept. 8-9, Petten, Holland: IWGFPT/14. Ishigure, K., Takagi, X, Shiraishi, H. (1987), Radiat. Phys. Chem. 29, 195-199. Isobe, T, Matsuo, Y. (1991), 9th Int. Conf Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 346367. Johnson, A. B. (1989), IAEA Tech. Com. Meet.: Fundamental Aspects of Corrosion ofZb-base Alloys in Water Reactor Environments, Sept. 11-15, Portland, OR: IWGFPT/34, pp. 107-119.
49
Jostsons, A., Kelly, P. M., Blake, R. G., Parrell, K. (1979), Proc. 9th Symp. Effect Radiation on Materials: ASTM-STP 633. Philadelphia, PA: Am. Soc. Test. Mat., pp. 46-61. Kai, X X, Huang, W. I., Chou, H. Y. (1990), J. Nucl. Mater. 170, 193-209. Kai, X X, Tsai, C. H., Hsieh, W. F. (1992), 15th ASTM Int. Symp. Effects Irradiation Mater. ASTM-STP 1125. Philadelphia, PA: Am. Soc. Test. Mat., pp. 355-372. Khatamian, D., Manchester, R D. (1989), J. Nucl. Mater. 166, 300-306. Kearns, X X, Woods, C. R. (1966), J. Nucl. Mater. 20, 241-261. King, A. D., Hood, G. M., Holt, R. A. (1991), J. Nucl. Mater. 185, 174-181. Knorr, D., Pelloux, R. M. (1982), Met. Trans. 13 A, 73-83. Konashi, K. (1984), /. Nucl. Mater. 125, 244-247. Kubo, T, Wakashima, Y, Imahashi, H., Nagai (1985), J. Nucl. Mater. 132, 126-136. Lemaignan, C. (1992), J. Nucl. Mater. 187, 122-130. Levy, S. (1974), Nucl. Eng. & Design 29, 157-162. Lloyd, L. T. (1963), ANL 6591. Lucas, G., Pelloux, R. M. (1981), Nucl. Tech. 53, 4 6 57. Lukas, S. R. (1988), /. Nucl. Mater. 158, 240-252. Lustman, B., Kerze, R, (1955), The Metallurgy ofZr. New York: McGraw-Hill. MacEwen, S. R., Faber, X Jr., Turner, A. L. P. (1983), Acta Metall. 5, 657-676. MacEwen, S. R., Christodoulou, N., Tome, C , Jackman, X, Holden, T. M., Faber, X Jr., Hitterman, R. L. (1988), 8th Int. Conf. Text. Mat. (ICT0M8): Kallend, X S., Gottstein, G. (Eds.). Warrandale, PA: Met. Soc, pp. 825-836. Marlowe, M. O., Armijo, I S . , Cheng, B., Adamson, R. (1985), Proc. ANS Top. Meet. Light Water Reactor Performance, April 21-24, Orlando, FL. La Grange Park, IL: ANS, pp. 3-73, 3-90. Matsuo, Y. (1987), J. Nucl. Sci. Tech. 24, 111-119. Matsuo, Y (1989), 8th Int. Conf. Zr Nucl. Ind., June 19-23, 1988, San Diego, CA: ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat., pp. 678-691. Matthews, X R., Finnis, M. W (1988), J. Nucl. Mater. 159, 257-285. Mclnteer, W. A., Baty, D. L., Stein, K. O. (1989), 8th Int. Conf. Zr Nucl. Ind., June 19-23, 1988, San Diego, CA: ASTM-STP 1023. Philadelphia, PA: Am. Soc. Test. Mat, pp. 621-640. Miquet, A., Charquet, D., Michaut, C , Allibert, C. H. (1982), /. Nucl. Mater. 105, 142-148. Morize, P. (1984), Ann. Chim. Fr. 9, 411-421. Motta, A. T, Lemaignan, C. (1992), in: Ordering and Disordering in Alloys: Yavari, A. R. (Ed.). London: Elsevier Applied Science, pp. 255-275. Motta, A. T, Olander, D. R., Machiels, A. X (1989), 14th ASTM Int. Symp. Effects Irradiation Mater., ASTM STP1046. Philadelphia, PA: Am. Soc. Test. Mat., pp. 457-469.
50
7 Zirconium Alloys in Nuclear Applications
Motta, A. T., Olander, D. R. (1990), Ada MetalL Mater. 38(11), 2175-2185. Motta, A. T., Lefebvre, K, Lemaignan, C. (1991), 9th Int. Symp. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 718-739. Motta, A. T., Lemaignan, C , Olander, D. R. (1992), 15th ASTM Int. Symp. Effects Irradiation Mater., Nashville 1990, ASTM STP 1125. Philadelphia, PA: Am. Soc. Test. Mat, pp. 689-702. Moulin, L., Reschke, S., Tenekhoff, E. (1984a), 6th Int. Conf Zr Nucl. Ind., June 28- July 1,1982, Vancouver, Canada: ASTM-STP 824. Philadelphia, PA: Am. Soc. Test. Mat., pp. 225-243. Moulin, L., Thouvenin, JP., Brun, P. (1984b), 6th Int. Conf. Zr Nucl. Ind., June 28-July 1,1982, Vancouver, Canada: ASTM-STP 824. Philadelphia, PA: Am. Soc. Test. Mat., pp. 37-44. Murthy, K. L., Adams, B. L. (1985), Mater. Sci. Eng. 70, 169-180. Nash, P., Jayanth, C. S. (1984), Bull. Alloy Phase Diagrams 5(2), 1 111 -1779. Nikulina, A. V., Markelov, V. A., Peregud, M. M., Voevodin, V. N., Panchenko, V. L., Kobylyansky, G. P. (1993), 3rd Int. Conf. Evolution of Microstructure in Metals During Irradiation, Muskoka, Canada, Sept.-Oct. 1992, to be published in the /. Nucl. Mater. Northwood, D. O., Gilbert, R. W, Bahlen, L. E., Kelly, P. M., Blake, R. G., Jostsons, A., Madden, P. K., Faulkner, D., Bell, W, Adamson, R. B. (1979), J. Nucl. Mater. 79, 379-394. Olander, D. R. (1976), Fundamental Aspects of Nuclear Reactor Fuel Elements, TID-26711-P1. Springfield, VA: National Technical Information Service. Parsons, I R., Hoelke, C. W. (1989), 14th ASTM Int. Symp. Effect Irradiation Mater., ASTM STP 1046. Philadelphia, PA: Am. Soc. Test. Mat., pp. 689702. Pecheur, D., Lefebvre, F., Motta, A. T, Lemaignan, C, Wadier, J. F. (1992), J Nucl. Mater. 189, 318332. Peehs, M., Stehle, H., Steinberg, E. (1979), 4th Int. Conf Zr Nucl. Ind.: ASTM-STP 681. Philadelphia, PA: Am. Soc. Test. Mat., pp. 244-260. Ploc, R. A. (1983), /. Nucl. Mater. 113, 75-80. Porrot, E., Eminet, G., Baudusseau, C , Lemaignan, C. (1991), IAEA Tech. Com. Meet. Post Irrad. Eval Techn. Reactor Fuel, Sept. 11-14, 1990, Workington, England: IWGFPT/37, pp. 77-82. Price, E. G., Richinson, P. J. (1978), AECL 6345. Ramasubramanian, N. (1989), Proc. Tech. Com. Meet. IAEA: Fundamental Aspects of Corrosion in Zr Base Alloys in Water Reactors Environments, Sept. 11-15, Portland, OR: IWGFPT/34, pp. 36-44. Ramasubramanian, N. (1991), 9th Int. Conf. Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 613623.
Rogerson, A. (1988), / Nucl. Mater. 159, 43-61. Rosenbaum, H.s Lewis, J. E. (1977), J. Nucl Mater. 67, 237-282. Rosenbaum, H., Davis, J.H., Pon, I Q . (1966), GEAP 5100-5. Rosenbaum, H., Rand, R. A., Tucker, R. P., Cheng, B., Adamson, R. B., Davies, J. H., Armijo, I S . , Wiesner, S. B. (1987), 7th Int. Conf Zr Nucl. Ind.: ASTM-STP 939. Philadelphia, PA: Am. Soc. Test. Mat., pp. 675-699. Sayers, C. M. (1987), J Nucl. Mater. 144, 211-213. Schemel, J. H. (1977), ASTM Manual on Zirconium and Hafnium: ASTM-STP 639. Philadelphia, PA: Am. Soc. Test. Mat. Schemel, J. H. (1987), 7th Int. Conf Zr Nucl. Ind.: ASTM-STP 939. Philadelphia, PA: Am. Soc. Test. Mat., pp. 243-256. Schuster, I., Lemaignan, C. (1992), / Nucl Mater. 189, 157-166. Shaltiel, D., Jacob, I., Davidov, D. (1976), J. Less. Comm. Met. 53, 117-131. Simpson, L. A., Chow, C. K. (1987), 7th Int. Conf Zr Nucl. Ind.: ASTM-STP 939. Philadelphia, PA: Am. Soc. Test. Mat., pp. 579-596. Smith, T. (1969), J. Electrochem. Soc. 112, 560-567. Stephen, W. W. (1984), 6th Int. Conf. Zr Nucl Ind., June 28-July 1,1982, Vancouver, Canada: ASTMSTP 824. Philadelphia, PA: Am. Soc. Test. Mat., pp. 5-36. Tenekhoff, E. (1978), Met. Trans. 9A, 1401-1412. Tenekhoff, E. (1988), Deformation Mechanisms, Texture and Anisotropy in Zirconium and Zircaloys: ASTM-STP 966. Philadelphia, PA: Am. Soc. Test. Mat. Tricot, R. (1989), Materiaux et Techniques, AvrilMai, 1-32. Tricot, R. (1990), Rev. Gen. Nucl Jan.-Fev., 8-20. Urbanic, V. F., Warr, B. D., Manolescu, A., Chow, C. K., Shanahan, M. W. (1989), 8th Int. Conf Zr Nucl. Ind., San Diego, CA, June 1988, ASTM STP 1023, 20-34. Wadekar, S., Banerjee, S., Raman, V. V., Asundi, M. K. (1991), 9th Int. Conf Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 140-155. Warr, B. D., Elmoselhi, M. B., Brenenstuhl, A., Lichtenberger, P. C , Newcomb, S. B., Mclntyre, N. S. (1991), 9th Int. Conf Zr Nucl. Ind., Nov. 1990, Kobe, Japan: ASTM-STP 1132. Philadelphia, PA: Am. Soc. Test. Mat., pp. 740-757. Weatherly, G. C. (1981), Acta MetalL 29, 501-512. Wiedersich, H. (1972), Rad. Effects 12, 111-125. Williams, C. D., Gilbert, R. W. (1966), J. Nucl. Mater. 18, 161-166. Woo, C. H. (1979), ASTM-STP 683. Philadelphia, PA: Am. Soc. Test. Mat., pp. 640-655. Woo, C. H. (1984), / Nucl. Mater. 120, pp. 55 ff. Woo, C. H. (1988), J. Nucl. Mater. 159, 237-256. Woo, C. H., Singh, B. N. (1992), Phil. Mag. A 65, 889-912.
7.6 References
Woo, C. H., Causey, A. R., Holt, R. A. (1990), invited paper in: Diffusion and Defect Solid State Data, in press. Woo, C. H., Holt, R. A., Griffiths, M. (1992), in: Materials Modeling from theory to Technology. Bristol: Institute of Physics Publications, pp. 55-60. Woo, O. T., Carpenter, G. I C. (1987), / NucL Mater. 159, 397-404. Wood, J. C , Kelm, J. R. (1980), IAEA Spec. Meet on: Pellet Cladding Interaction in Water Reactors, Riso, Denmark: IWGFPT/8. Yang, W J. S. (1989), 14th Int. Symp. Effects Rad. Mater., 1988 Andover, MA: ASTM-STP1046. Philadelphia, PA: Am. Soc. Test. Mat., pp. 442456. Yang, W I S., Tucker, R. P., Cheng, B., Adamson, R. B. (1986), /. NucL Mater. 138, 185-195. Zee, R. H., Rogerson, A., Carpenter, G. J. C , Watters, J. (1984), J. NucL Mater. 120, 223-229. Zhou, H. (1992), Chalk River, unpublished results.
51
Zuzek, E., Abriata, J. P., San Martin, A., Manchester, F. D. (1990), Bull Alloy Phase Diagrams
11(4), 385-395.
General Reading ASTM-STP series of conferences on Zr in the nuclear industry nos. 1 to 9. Philadelphia, PA: Am. Soc. Test. Mat. Cheadle, B. A. (1975), The Physical Metallurgy ofZr Alloys, CNRL Report, 1208. Conferences on microstructural evolution under irradiation: Journal of Nuclear Materials 90 (1980) and 159 (1988). Douglass, D. L. (1971), The Metallurgy of Zirconium, Atomic Energy Review, Vienna. Tenckhoff, E. (1978), Met. Trans. 9A, 1401-1412.
8 Structural Materials Wolfgang Dietz
Siemens AG, Power Generation Group, KWU, Bergisch Gladbach, Federal Republic of Germany
List of 8.1 8.2 8.2.1 8.2.2 8.3 8.3.1 8.3.2 8.3.2.1 8.3.2.2 8.3.2.3 8.3.2.4 8.3.3 8.3.4 8.3.4.1 8.3.4.2 8.3.5 8.3.5.1 8.3.5.2 8.3.5.3 8.3.5.4 8.3.5.5 8.3.5.6 8.3.6 8.3.6.1 8.3.6.2 8.3.6.3 8.3.6.4 8.3.6.5 8.3.6.6 8.3.7 8.3.8 8.3.8.1
Symbols and Abbreviations Introduction Codes, Standards and Quality Assurance Codes and Standards Quality Assurance Within the Nuclear Industry Materials for Water-Cooled Reactors Survey of Light Water Reactor (LWR) Systems Water Chemistry of Light Water Reactors Primary Water Chemistry of Pressurized Water Reactors (PWRs) Secondary Water Chemistry Boiling Water Reactor (BWR) Water Chemistry Corrosion Phenomena in Water Requirements for Materials Selection Materials Technology of Light Water Reactor Components Fabrication Operating Experience Materials Properties of Ferritic Steels Tensile Properties Toughness Behavior Thermal Aging Effects Effects of Radiation on Reactor Pressure Vessel Steels Corrosion Resistance of Carbon and Low Alloy Ferritic Steels Erosion-Corrosion Properties of Austenitic Stainless Steels and Ni-AUoys Tensile Properties and Toughness Thermal Embrittlement (TE) of Cast Austenitic Stainless Steels Corrosion Cracking of Sensitized Type 304 and 316 Stainless Steel in the BWR Environment Effects of Irradiation on Tensile Properties Irradiation-Assisted Stress Corrosion Cracking (IASCC) Performance of Special High Alloyed Materials Activity Build-Up and Contamination of the Primary System by Corrosion Products Corrosion Effects in the Steam-Water System Steam Generator Tubing - General Aspects
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.
56 61 62 62 63 64 64 67 67 68 69 69 70 77 77 82 84 84 84 86 87 90 94 95 95 96 96 98 98 99 100 101 101
54
8 Structural Materials
8.3.8.2 Observations on Primary Water Stress Corrosion Cracking (PWSCC) of Alloy 600 104 8.3.8.3 Behavior of Alloy 800 (Modified) 106 8.3.8.4 Behavior of Alloy 690 107 8.3.8.5 Materials for Steam Turbines of Light Water Reactors 108 8.3.8.6 Power Plant Condensers 109 8.3.9 Friction and Wear 110 8.4 Pressurized Heavy Water Reactors (PHWRs) Ill 8.4.1 Overview of the Reactor Ill 8.4.2 Structural Materials in the Calandria Reactor Vessel Requirements and Selection .. 113 8.4.3 Behavior of Structural Materials in the PHWR Primary System 114 8.4.3.1 Oxidation and Hydriding 114 8.4.3.2 Mechanical Properties 115 8.4.3.3 Irradiation Creep and Growth of PHWR Zirconium Alloys 116 8.4.4 Materials in Candu Steam Generators 117 8.5 Structural Materials for Fast Breeder Reactors (FBRs) 118 8.5.1 General Background 118 8.5.2 Requirements for Materials Selection 121 8.5.3 Materials Selection for the Main Components of Fast Breeder Reactors .. 123 8.5.4 Materials Technology for Fast Breeder Reactor Components 127 8.5.4.1 Fabrication 127 8.5.4.2 Operating Experience 127 8.5.5 Materials Properties of the Austenitic Stainless Steels 128 8.5.5.1 Tensile Properties 128 8.5.5.2 Notch Impact Toughness Behavior 130 8.5.5.3 Creep and Creep Rupture Strength 130 8.5.5.4 Fatigue and Creep-Fatigue Interaction 133 8.5.5.5 Irradiation Behavior 136 8.5.6 Ferritic Steels 137 8.5.6.1 Tensile and Creep Properties 137 8.5.6.2 Aging Effects 139 8.5.6.3 Fatigue and Creep-Fatigue Interaction .. . 139 8.5.6.4 Corrosion Behavior (Waterside) 139 8.5.7 Fracture Mechanics Data 140 8.5.8 Corrosion Behavior and Compatibility Problems of Structure Materials in Sodium 141 8.5.8.1 General Considerations 141 8.5.8.2 Behavior of Austenitic and Ferritic Steels 142 8.5.8.3 Faulted Conditions of the Sodium System 143 8.5.9 Friction and Wear: Tribological Performance 144 8.6 Advanced Gas-Cooled Reactors (AGRs) 144 8.6.1 General Background 144 8.6.2 Requirements on Materials and Materials Selection 146 8.6.3 Oxidation of Structural Materials in the Reactor Coolant 147
8 Structural Materials
.6.3.1 L6.3.2 L6.3.3 1.6.3.4 1.7 1.7.1 1.7.2 1.7.2.1 1.7.2.2 !.7.2.3 L7.3 1.7.4 1.7.4.1 1.7.4.2 1.7.5 1.7.5.1 1.7.5.2 1.7.5.3 1.7.5.4 1.7.5.5 L7.6 1.7.7 5.8 1.8.1 >.8.2 L9 >.9.1 L9.2 >.9.3 L10 }.ll
Mild Carbon Steels 9 Cr 1 Mo Steel Austenitic Stainless Steel Irradiation Effects in the Reactor Pressure Vessel (RPV) Steel and Radiation Embrittlement High-Temperature Gas-Cooled Reactors (HTGRs) Background Technical Features of High-Temperature Gas-Cooled Reactors General Description Coolant Chemistry Water Chemistry of High-Temperature Gas-Cooled Reactors Design Requirements and Selection of Materials Materials Technology for High-Temperature Gas-Cooled Reactor Components Fabrication Survey of Operating Experience Behavior of Structural Materials Under Operating Conditions Tensile Properties Toughness Properties (Charpy V-Notch Test) Creep and Creep Rupture Fatigue and Creep Fatigue Fracture Mechanics Irradiation Behavior of High Temperature Gas-Cooled Reactor Materials Corrosion of Structural Materials in Helium Research Reactors (RRs) Embrittlement at L o w Temperature for a Ferritic Pressure Vessel Behavior of Aluminum Alloys Structural Materials for Components in the Fuel Cycle Outside the Reactor Structural Materials for Storage Facilities of Fuel Elements Transportation Casks and Long-Term Resistant Packages Structural Materials for Spent Fuel Reprocessing Plants Acknowledgements References
55
147 148 149 149 150 150 151 151 151 153 153 158 158 158 158 159 159 159 161 161 161 162 165 165 166 167 168 168 170 171 172
56
8 Structural Materials
List of Symbols and Abbreviations a A C Cx C(t\ C t , C*(t) Cv CVN E J Jc J IC JR K KQ
AK N Pm
Q R R
T t
k A (DBTT) e s Aet Asel, Asinel a % crt
crack length total elongation at fracture (tensile test) constant X (X = C, Cr, Nb, ...) content of steel crack-tip parameters for stress intensity for creep crack growth rates (see ASTM 1457-92) Charpy V-notch impact energy Charpy V-notch impact energy = C v energy J-integral, fracture resistance parameter, allowing for crack growth J-integral at crack initiation J-integral at crack initiation (ASTME 813-88) material property in terms of J to describe the crack growth resistance material-dependent parameter plain strain fracture toughness, not validated by all KIC requirements of ASTM E399-83 plain strain fracture toughness as defined by ASTM Standard E399-83 lower bound fracture toughness reference curve for ferritic alloys stress intensity factor range in FCG number of cycles number of cycles to failure primary stress activation energy gas constant i^-curve (fracture mechanics) ultimate tensile strength yield strength reference nil-ductility transition temperature allowable stress for short-, long-term loading, respectively absolute temperature time time to initiate cracking nil-ductility transition temperature shift of the 41 J transition temperature ductile-to-brittle transition temperature shift strain strain rate total strain range change of elastic, inelastic strain, respectively stress rupture strength hoop stress
List of Symbols and Abbreviations
57
600 mm) from steels can be significantly improved by the inside, offering lower costs and flexibilcold working or addition of nitrogen. Cold ity especially for local repairs. worked steels (about 20%) are used, for example, for FBR fuel cladding materi8.3.6 Properties of Austenitic Stainless als. In LWR fabrication technology, cold Steels and Ni-Alloys working is of interest for the construction Austenitic stainless steels have major of the piping system (bending of small diapplications, such as weld cladding of the ameter pipes, elbows). But owing to associ-
96
8 Structural Materials
ated corrosion risks the allowable cold work strain from the fabrication is limited (about 10%). 8.3.6.2 Thermal Embrittlement (TE) of Cast Austenitic Stainless Steels For more than a decade there has been concern about the susceptibility of cast stainless steels with a duplex ferriticaustenitic microstructure to thermal embrittlement by aging at reactor operating (or higher) temperatures. This type of degradation is manifested by an increase in DBTT, by a loss of Charpy impact energy and fracture toughness, or an increase in fatigue crack growth. Less extensive thermal aging effects have been observed for tensile properties. Recent reviews have been given by Chung (1992), Bogie et al. (1992), and during an international workshop by various authors (see Pumphrey e t a l , 1990). Mostly, the TE is quantified by the room-temperature Charpy V-notch impact energy after aging of the specimens at higher temperatures. Then these data are correlated to fracture toughness. Because a realistic aging of a component for 250 000 h or more (end-of-life conditions, plant life extension) at 280-340 °C plant operation temperatures cannot be produced, it is customary to simulate the metallurgical structure by accelerated aging tests at or near 400 °C. For the extrapolation to realistic conditions it is necessary to understand the kinetics of the aging mechanisms and to quantify the activation energy. In a first approach the aging mechanisms can be treated on the basis of a thermally activated process. In this case the kinetics are simply expressed by an Arrhenius-type equation, that is,
where t2 and tx are the times taken to reach an equivalent toughness at the absolute temperatures T2 and 7 \ , respectively, Q is the activation energy, and R is the gas constant (Chung, 1992; Jaske and Shah, 1987). Multiple correlation analyses have been made to correlate the activation energy with the chemical composition and the ferrite content. Activation energies in the range of 20-50 kcal/mol ( - 8 0 - 2 1 0 kJ/ mol) have been observed. TE increases with the ferrite content and depends on the alloy chemistry and process history. Saturation effects (CVN values of 20-30 J/cm2 at 25 °C) have been observed with significant cast-to-cast variations. Type 316 (CF8M) steel embrittles more than type 304 (CF3) steel. Several metallurgical processes have been identified in association with TE. TE is related to the formation of a Cr-rich a' phase and a Ni-rich and Si-rich G phase in the ferrite. At temperatures above 400 °C the precipitation of carbides on the austenite-ferrite phase boundary also plays a significant role. Reliable extrapolations to end-of-life (EOL) properties, therefore, need data from aging near operation temperatures. Life assessment studies need fracture mechanics considerations for materials aged in reactor components up to 60 years. Component experiences are available now up to 15 years (Ringhals NPP). For further verification of these procedures important programs are still going on in the U.S.A. (Chopra, 1992) and in France (Bonnet et al., 1990). 8.3.6.3 Corrosion Cracking of Sensitized Type 304 or 316 Stainless Steel in the BWR Environment The BWR recirculation piping circuit using standard type 304 or 316 stainless steel (carbon content 0.05-0.07 wt.%) de-
8.3 Materials for Water-Cooled Reactors
veloped heat-affected weld zone cracking problems and intergranular stress corrosion cracking (IGSCC) in the cold-bent piping. These problems led to the initiation of research programs to understand these corrosion effects. It was clearly demonstrated by SCC laboratory experiments that the presence of a BWR environment with a higher oxygen content than the PWR and trace impurities initiated the intergranular stress corrosion cracking. H 2 O 2 and H 2 SO 4 strongly enhance IGSCC, even more than oxygen (e.g. Lungberg et al., 1987). Besides the oxidizing water environment, significant stresses - residual stresses always exist in welded structures - and a sensitized metallurgical structure are also essential for producing IGSCC. For austenitic stainless steels of type 304 heating between 400 and 850 °C gives rise to the phenomenon of "sensitization", which results from the depletion of chromium at the grain boundaries caused by grain boundary precipitation of chromium carbides. Intergranular attack can occur if there is subsequent exposure to certain oxidizing media. The sensitization of the weld HAZ results from heating and cooling through the sensitization temperature range. Cooling rates during the welding process and the carbon content are correlated (see Fig. 8-10), but the nitrogen content, grain size, prior cold work, maximum annealing temperature, and strain during cooling are also relevant factors for sensitization (Solomon, 1984). The sensitivity of austenitic stainless steels is usually tested by standard procedures, for example ASTM 262, practice A or E, and they are classified to be sensitive or not. New electrochemical test methods (e.g. for quantitative modeling of sensitization) and other tests (Autass, Strauss) are used to prove sensitization
97
0.0001 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08 wt%C
Figure 8-10. Critical cooling rate versus carbon content to avoid sensitization of type 304 steel for two ASTM test procedures (Solomon, 1984). Tm: annealing temperature before cooling. H5, E3, E5, D4, Al: heats with variable C content. For cooling rates above curve A 262 A there were no precipitations (sensitizations) in these tests.
behavior of stainless steels. The sensitivity of cold-worked type 304 to IGSCC was mainly caused by the formation of martensitic structures. Welded cladding in the RPV with 5-ferrite of some percent is usually less sensitive to cracking. A low 5-ferrite content or local cold work may be the reason for some of the reported failures. The IGSCC problem in BWRs has been solved by the use of improved materials, a modification in water chemistry, and the development of residual stress improvement measures (see the proceedings of the series of symposia on environmental
98
8 Structural Materials
degradation of materials in nuclear power systems). In the U.S.A. and in Japan, type 316 NG stainless steel (SS) was developed and used as a replacement material. For BWRs with a modified type 347 SS (see Table 8-9, ISO 4550) no incidents were reported (Tenckhoff and Erve, 1992). However, under certain conditions stabilized SS such as type 321 may not be totally immune to SCC in high temperature water. IGSCC can occur in an overheated HAZ, coarse grained and sensitized. Methods for improving the residual stress state near welds, for example by induction heating, have been developed in the U.S.A. and in Japan. Hydrogen water chemistry (HWC) has been successfully used to reduce dissolved oxygen levels by additions of H 2 . Critical impurity species produced by radiolysis in the reactor core are now controlled by more sophisticated water-chemistry analysis. Summarizing, the problem of IGSCC is now considered as an operational problem, however, the cracks have cost utilities several billion dollars during the past decade (Jones and Nelson, 1990). Today it is mainly a question of economics as to which mitigating solution - HWC, stress relief treatment, weld overlay, replacement - is chosen. 8.3.6.4 Effects of Irradiation on Tensile Properties Displacement type of damage, structural changes due to radiation-enhanced diffusion and effects of transmutation products are the main sources of property changes and degradation in the radiation environment (see Chap. 9, by Schilling and Ullmaier, and Chap. 6, by Garner, in this Volume). In evaluating data of irradiated materials for design implications (mainly ductili-
ty exhaustion), results from experiments with relevant irradiation temperatures and damage rate have to be considered. Relevant data for core internal structures with maximum fluence levels between 10 21 and 1022 n/cm2 (E> 1 MeV) are available from a U.S. experimental test reactor (ETR) program (e.g. Lovell, 1968) and other programs (Jacobs, 1987; Leitz, 1991). The use of the large amount of data from the FBR program is limited (because of the difference in temperatures and dose rates to the LWR conditions). Due to radiation hardening an increase of yield strength is observed with a corresponding loss in ductility (see Fig. 8-11). There are some small differences in the behavior of the various austenitic steels used. Effects of microstructure such as cold work on property changes have also been observed. It is known from FBR irradiation at about 400 °C that, at fluences above 1022 n/cm2 (£>0.1 MeV), metallurgical variants of austenitic SS show the same scatter band for ductility. The fracture strain is low (2-4% at 280- 300 °C) but ductile behavior is observed (high reduction in cross section, "dislocation channeling effects"). Irradiation also has the general effect of lowering the work hardening rate and reducing the relaxation behavior. Properties of spring elements or bolting assemblies may be affected for fluences above 10 21 n/cm2 {E>\ MeV). 8.3.6.5 Irradiation-Assisted Stress Corrosion Cracking (IASCC) Nonsensitized austenitic steels and nickel alloys used for BWR and PWR core components show intergranular SCC after long-term exposure, and sensitization of the SS by irradiation has been observed. Irradiation-assisted stress corrosion cracking (IASCC) refers to all cases in which
8.3 Materials for Water-Cooled Reactors
99
o
I 2 LL1
1011
10 FLUENCE (n/cm2,E>1 MeV)
Figure 8-11. Effect of neutron fluence on total elongation of irradiated austenitic SSs [data from Jacobs (1988); data trend of FBR irradiation for a fluence &t >10 2 2 n/cm 2 ].
environmental cracking is accelerated by radiation, whether it acts singly or in combination to alter water chemistry, material microchemistry, hardness, creep, etc. (see review by Andresen and Ford, 1989). IASCC can occur at low stresses, and a tendency to increased effects in BWRs can be observed. Their susceptibility to IASCC, for example the severity of intergranular fracture, decreases markedly when the oxygen dissolved in the water is reduced. IASCC effects can be correlated with the percentage increase in yield strength and have a threshold of about 5xl0 2 O n/cm 2 (E> 0.1 MeV) for nonsensitized austenitic SS. It is the objective of ongoing research programs to explain how IASCC is correlated with fluence-dependent microstructural features, transmutations by neutrons, and radiolysis of water.
The use of high purity austenitic materials may be a way to reduce IASCC (Garzarolli et al., 1987). For BWRs H 2 water chemistry may reduce sensitivity (Jones and Nelson, 1990). The modeling of IASCC was included in the environmentally assisted cracking models by a slip dissolution/oxide film rupture mechanism (Andresen and Ford, 1992). By this relationships with primary water SCC (PWSCC) can be obtained. Up to now IASCC has not significantly affected plant availability, but IASCC may become more important in the future when the NPPs are aging and the life extension is evaluated for 60 years. 8.3.6.6 Performance of Special High Alloyed Materials High alloyed materials such as austenitic SS and their performance were
100
8 Structural Materials
considered in the previous sections. The performance of the group of high alloyed materials such as Alloy 600, Alloy 690, and Alloy 800 will be described in a special section about steam generator materials (see Sec. 8.3.8). It should be mentioned that especially Alloy 600 and suitable weld metals such as Alloy 82 and 182 were employed in the primary system for other components, too. SCC similar to PWSCC in the SG has been observed in components such as nozzles, especially in BWRs [see contributions in Nucl. Eng. Design (1990), 124]. Recent observations of cracking of French reactor vessel head penetrations have raised discussions about reactor safety and will entail a repair of components. The behavior of some other special alloys will be briefly discussed in the following. Nickel-based alloys such as Alloy X 750 and Alloy 718 are mainly used for the fabrication of relaxation-resistant parts such as bolts and springs or as core internals. The microstructure of these materials has been developed and optimized for aerospace applications. It was experienced that the alloys with these metallurgical properties need a special optimization for LWRs because of SCC. Efforts towards further optimization are still continuing. In many NPPs worldwide, bolts made of nickel Alloy X 750 cracked in the nearcore position. It was necessary to exchange the bolt materials (Grove and Petzold, 1985; and others). In German NPPs Alloy X 750 was replaced by austenitic steels, which was also the replacement material for Alloy 600 nozzles. Martensitic stainless steels of the type CrNi 12 3 or precipitation hardened alloys (17-4 PH) are used as materials for pumps, bolts, and valves. The hardness of these alloys is important for a good operation performance. Their hardness may be in-
creased by aging. If their hardness exceeds 350 HV these alloys become sensitive to SCC. In summary, all these cracking observations have increased the need for examination of components and mitigating actions (improved materials and designs, water chemistry, etc.). 8.3,7 Activity Build-Up and Contamination of the Primary System by Corrosion Products
When the first generation of commercial LWRs approached maturity from 1970 onward, high dose rates around the coolant circuits emerged as a widespread so-called occupational radiation exposure (ORE) or "man-rem" problem (Comley, 1985). The origin of the high out-of-pile radiation fields is linked to the chemical composition of the structural material in contact with water and the release of soluble or insoluble corrosion products in the coolant system. Details of activation processes are described in Chap. 5 of this Volume. The "man-rem" problem is mainly related to the long-lived 60 Co isotope. Co is contained in stainless steel and Ni alloys as an impurity on the order of 0.1 wt.% in improved grades of materials (see Table 8-8). Co is also included (in the range of 60 wt.% Co) in Stellite alloys, which are known to be some of the best wear-resistant alloys in LWRs. Stellite seems to contribute much more to activity build-up than Co impurities in the structural material (Friedrich, 1989). A reduction of Stellite in German KWU NPPs and Candu pressure tube stations resulted in a significant reduction of the primary circuit dose rates, too. Therefore, a first priority for new NPPs is that high cobalt alloys should be excluded wherever possible and replaced by alternative alloys (see Sec. 8.3.9).
8.3 Materials for Water-Cooled Reactors
Other methods to minimize the Co contamination of the primary circuit are pretreatment of the steel surfaces prior to operation or an optimization of the water chemistry in the existing LWRs (see Chap. 5 of this Volume). On the other hand, these possibilities are limited because modifications in water chemistry may introduce stress corrosion risks. In the last ten years a significant improvement has been made in the LWR chemical technology, which has led to substantially lower occupational radiation exposures (Amey and Johnson, 1990). 8.3.8 Corrosion Effects in the Steam-Water System
A good performance of the components depends on proper design, the selection of appropriate materials, and the water chemistry. Within the steam-water system the steam generator tubing bundle with very severe loading conditions is an excellent example showing that all factors have to be optimized. Other components in this system needed remedial measures in the past, too (Sees. 8.3.5.5 and 6). In the following we have to limit the contribution to some selected areas. 8.3.8.1 Steam Generator Tubing General Aspects
In a recent overview of stress corrosion experiences with Alloy 600 in U.S. PWRs (Paine, 1990) it is mentioned that the various types of corrosion are now a cause of great concern. In the U.S.A. repair or replacement measures are in progress at more than 50 NPPs. Primary water stress corrosion cracking (PWSCC) of Alloy 600 tubes is the main degradation effect found in SG units of French reactors (de Keroulas and Lunven, 1990), too. For many years utilities and
101
vendors have had important research programs to understand PWSCC and to find ways to prevent PWSCC of Alloy 600 or to repair cracked tubes. The main features of a PWR steam generator (SG) are shown schematically in Fig. 8-12. The heat exchanger tubing is present in the SG as a vertical U-bend through which the pressurized primary water coolant with a specified chemistry (see Sec. 8.3.2) circulates. The coolant enters the tube bundle with a temperature and pressure of about 316 °C and 16 MPa and exits with a temperature about AT = 35 K lower. The pressure of the secondary side is around 6 MPa and the temperature is about 270 °C. The chemistry of the secondary water side is also very specific. The water contains ammonia and hydrazine (controlled pH) as well as small amounts of soluble and insoluble impurities. During operation corrosion products - metal oxides ("sludge") or ions such as Na + , SO4", etc. - may accumulate at the SG tube sheet, in the annulus formed between the tubes and the tube sheet, and in the tubeto-tube support plate crevices. The local chemistry may be of acid or alkaline type. These are critical places where Alloy 600 tubes corroded in the secondary system. The various types of corrosion have been the reason for changes and modifications in SG materials over the last 25 years. At the beginning austenitic stainless steels were used. Failures mainly due to chlorideinduced stress corrosion cracking (SCC) occurred. For example, in 1967 in the U.S.A. the type 304 tubing of an SG cracked before reactor operation during the initial pressure test (Carlson and Kratzer, 1976). This event made replacement by more crack resistant alloys necessary. Austenitic steels, for which specific corrosion effects by faulted local chemistry have been reported by Mummert (1992),
102
8 Structural Materials
are still used in Russian PWRs but with another SG design. Alloy 600 (mill annealed) was the prime choice for the U.S., Japanese, and French PWRs. Alloy 800 (modified) is used in SGs designed by Siemens/KWU. When the SSC susceptibility of mill-annealed Alloy 600 in pure water became evident, new plants in France changed to Alloy 600 TT (thermally treated) and practically simultaneously a program to qualify Alloy 690 was started (1978/79). In 1984 the construction of the first French PWR designed with Alloy 690 TT began. For the first British PWR (Sizewell 'B') Alloy 690 TT will also be used (Airey et al., 1990). Replacement SGs use both Alloy 690 and Alloy 800 mod. (Stubbe et al., 1990). SGs with Alloy 690 have been in operation since 1989. The operating experience with SGs using Alloy 800 mod. has not indicated the same generic problems as for Alloy 600. About 235 000 tubes are working under normal operating conditions (Bouecke, 1990). The chemical composition of the three materials is given in Table 8-13. Figure 8-12 shows the critical damage points in steam generators and gives a short description of these points. A survey of the problems, their occurrence, their causes, and counteractive measures is summarized in Table 8-14.
Figure 8-12. Schematic features of a PWR steam generator with regard to corrosion effects (Bouecke et al., 1989).
Table 8-13. Chemical composition of steam generator tubing (in wt.%, see Table 8-8 for other elements).
Material Alloy 600a" Alloy 800 d' Alloy 690c' f a
C
Ni
Cr
Fe
Ti
Al
Co
0.010-0.050 70 32-35 >58
14-17 20-23 28-31
6-10 >39.5 7-11
10 ~6 s~ * uniform elongations greater than 10% for operating temperatures have been measured. The effect of sodium on tensile properties and irradiation effects will be discussed separately, also the reduced ductility during creep (strain rates 70 J, data up to 300 J) in the as-received condition. Generally, this type of test is used as an indicator of structural instability and degradation of properties by thermal aging, especially for weld metals. Prior creep or deformation also reduces the CVN values, but these effects are not of concern. The degradation in CVN is caused by various types of precipitations (carbides, a-phase, etc.), which differ for the austenitic grades. Generally, the degradation increases with temperature in the range of 500-800°C and with time and is more significant for weld metal with some 8 ferrite owing to the microstructural instability of the 5 ferrite at high temperatures. The correlation of CVN data with time and temperature is often represented in a Larson-Miller-type of plot (see Marshall, 1984). An extrapolation of CVN data for 550 °C to end-of-life properties of austenitic weld metals results in values of about 50 J at room temperature, which are acceptable. Within the existing scattering there is a trend for the 19 Cr 12 Ni 2 Mo weld metal to be more susceptible to embrittlement than the 16 Cr 8 Ni 2 Mo type (Smith and Farrar, 1993). Fracture toughness data are considered together with other fracture mechanics data in Sec. 8.5.7.
8.5.5.3 Creep and Creep Rupture Strength
At elevated temperatures mechanical properties become time dependent due to thermal activation, and thus strain-rate dependent. If components operate at higher temperatures only for a short period low temperature design procedures (less expensive) may be used within the limits of a temperature-time relation ("creep crossover" curve). To prevent failure modes due to creep effects, by design, the primary stresses in the components are limited to temperature- and time-dependent stress limits (st). The creep properties used to determine the allowable stress limits, for example the st values, may be found in Table 8-25, with further information in Fig. 8-28. Creep damage owing to secondary stress relaxation has to be limited, too. The strength properties to be used for design in the creep regime are given for the materials in highest safety classes in the nuclear codes. Non-nuclear codes may be used in the design of the other components. The usual scattering of the data is about ± 20 % for the average stress values, which may be reduced by stricter fabrication routes and control of chemical analysis and microstructure. Mean values of some materials of interest are given in Fig. 8-29. The effect of welding is investigated by creep and creep-rupture tests on weld metal and weldments (transverse to the welding direction). The behavior of weldments is usually described by "weldment-factors" - the ratio of the creep strength of the weldment or weld metal to the base metal - which are applied to the creep-rupture properties of the base metal. Such factors are within the range between 0.8 band 1.0. The stress-strain behavior of the different metallurgical structures of a weldment is
8.5 Structural Materials for Fast Breeder Reactors (FBRs)
131
Table 8-25. Definition of st values (allowable stresses) for austenitic steels (base metal)a. Allowable stress for long-term loading
Symbol
Definition
sr(T,t)/1.5
sr
Creep rupture strength at design temperature T and time t 1% strain limit Onset of tertiary creep; for definition see Fig. 8-28.
a10/o 0.1 MeV). For the fuel and breeder elements special design procedures have been developed (see Chap. 6 of this Volume). For both the reflector elements and the near core permanent structures either the structural design codes for unirradiated material have to be applied below a "threshold dose" or special design procedures have to be used. At present, the threshold dose is of the order of 1 dpa (displacements per atom) for low-temperature irradiation (T500°C (Laue and Bauermeister, 1989).
8.5 Structural Materials for Fast Breeder Reactors (FBRs)
A major concern is the effect of irradiation on the toughness and ductility of the core support structures (displacement-type damage, low-temperature embrittlement) and on creep/fatigue/ductility at high temperature (high-temperature helium embrittlement of grain boundaries). Up to now irradiation was not considered a prime selection criterion for the base metal of FBR components. The austenitic grades type 304, 316, and 321 show similar behavior in radiation hardening/ductility loss or hightemperature embrittlement. There are significant heat-to-heat variations. Reducing the boron content below 10 ppm, optimizing the microstructure (fine grain size, homogeneous boron distribution, etc.) and reducing the impurity level may give better performance for near-core permanent structures, especially at higher temperatures. For the weld metal low-temperature embrittlement is more critical. The variation of the toughness data for irradiation at about 400 °C is large. Marshall (1984) gives a summary of the data available from the open literature up to this date. There is still some work going on within the surveillance programs for operating plants and for the verification of design data for future projects. Some recent information about irradiation effects is reported by Tavassoli (1990) for the ongoing program. It is well established now that highstrength Ni alloys, which are of interest for structural materials for above-core structures, generally embrittle more than austenitic or ferritic steels due to irradiationinduced grain boundary weakening. Details may be found elsewhere (ASTM conferences). 8.5.6 Ferritic Steels The ferritic materials listed in Table 8-23 have been selected according to their phys-
137
ical and mechanical properties, their fabricationability, weldability, availability, and anticipated operation performance after preliminary laboratory tests. Strength properties can be found in the various codes. Environmental or aging effects and fracture mechanics are usually investigated in R&D programs for the verification of materials selection and safety assessments. In the following, some of the properties of general interest are introduced. The sodium corrosion aspects are covered in a separate section. The waterside corrosion with regard to ferritic steam generator materials is briefly discussed. The use of ferritic steels in nuclear technology was surveyed in the Snowbird Conference (1983). 8.5.6.1 Tensile and Creep Properties The main tensile and creep rupture properties of selected steel grades given in Table 8-23 are summarized in Table 8-26. An important difference between the various grades is the allowable stress for the design of the components. In Fig. 8-36 the allowable stresses of two ferritic and austenitic steels are compared. The modified 9 Cr curve approaches the austenitic steel curve. The improved strength of the modified 9 Cr steel in comparison to standard 9 Cr or 2 lA Cr 1 Mo results in a significant reduction in the size of FBR steam generators. The creep properties of the common ferritic steels are verified by experimental data in the range between 50 and 100 x 103 h. For the new modified 9 Cr steel, more information is necessary to validate the extrapolated properties in the ASME code, especially in the range of 525-550°C, including weld metal and weldment behavior. Recent failures in coal-fired power plants have identified the soft-fine-grained region in the heat affected zone of weldments made of low alloyed
138
8 Structural Materials
Table 8-26. Main tensile and creep properties of steels used for fast breeder steam generator tubesa> b. Creep at 105 h
Tensile strength 400 °C 500 °C
20 °C
(MPa) (MPa)
(%)
500 °C
550 °C
600 °C
(MPa) (MPa) (MPa) (MPa) (MPa) (MPa) (MPa) (MPa)
10 CD 9,10 2iCr-lMo
325 490-640 >20
216
196
88
110
49
58
29
31
10 CD Nb 9,10 2iCr-lMo-lNb
295 470-610 >18
201
177
98
110
44 44
47
18
20
Z 10 CD 9 9Cr-lMo
340 540-700
280
Z10CDNbV9,2 9Cr-2Mo-NbV
390 590-740 >15
284
265
Alloy 800
210
>30
162
155
Z6CNT18-12B 18-10 Ti (321)
195 490-690
>40
120
110
po.2
an
>520
62
134
d 0R are minimum values;
£
st
9 0 0 ° C , depending upon the welding process. Further R & D work will reduce uncertainties in knowledge of high-temperature weldment behavior, especially for transition welds (austenitic/ ferritic). 8.7.5.4 Fatigue and Creep Fatigue
Fatigue loading of HTGR components will result either from cycling stresses (e.g. vibrations, HCF) or from temperature changes inducing thermal stresses and strains in the structure (LCF). Some results from the experimental programs will be highlighted which may be of specific interest for component design. More specific and detailed information is given by various authors (see, e.g., Nucl. TechnoL, 1984; Rao et al., 1988). A different behavior is observed in the very high temperature regime compared with the lowtemperature regime and FBR operating temperatures: (i) with increasing temperature the fatigue strength is reduced compared to the range between 400 and 550 °C; (ii) the endurance limit observed at low temperatures diminishes owing to creep effects (damage) at the high temperature; (iii) with increasing temperature and decreasing strain rate the cycling hardening diminishes, that means recovering and softening occur in the microstructure in a complex manner; (iv) creep and stress controlled cyclic loading interact in a complex manner; (v) relaxation due to faster creep rates occurs more rapidly for fatigue with hold time conditions than at lower temperatures. It is beyond the scope of this section to go into details of the complex material per-
161
formance in the creep fatigue interaction regime and life-fraction rule considerations. We refer to ongoing R & D published in the recent literature. 8.7.5.5 Fracture Mechanics
For the behavior of flawed structures of HTGR components the same fracture mechanics parameters as for the other systems at high temperatures (e.g. FBRs) have to be considered: fracture toughness, fatigue crack growth (FCG), creep crack growth (CCG), and crack growth under creep fatigue conditions. For the fundamentals of high-temperature fracture mechanics under creep and creep fatigue conditions we refer to Riedel (1987). Data on creep and fatigue crack growth for temperatures up to about 1000 °C including the i?-ratio and frequency effects were reported very recently by Rodig et al. (1992) for Alloy 800 and Alloy 617. Creep-fatigue crack growth was evaluated by Rodig et al. (1991). The creep crack behavior was evaluated on the basis of the fracture mechanics parameters KY and C*. Fracture toughness data for Alloy 617 are shown in Fig. 8-46 (Huthmann et al., 1989). Comparing the data with the CVN data in Fig. 8-43 shows that the effect of thermal aging is much less significant. Procedures for the application of the Charpy V-notch data in the same way as for LWRs (see Sec. 8.3.5.4) have to be used with care for the design of HTGR components. 8.7.6 Irradiation Behavior of High Temperature Gas-Cooled Reactor Materials
Irradiation effects have to be considered in the design of the RPV and the internal structures. The design temperature of the RPV is in the range between 200 °C and 400 °C. The 200 °C for the HTR modules is
162
8 Structural Materials
1000
NiCr23Coi2Mo,25#C AGED 20000h/650'C
800
(a) J
0
1000
0
1 2 3 4 5 6 CRACK EXTENSION A a (mm)
7
NiCr23Co12Mo,650 C A6ED 20000 h/650 #C
1 2 3 4 5 6 CRACK EXTENSION A a (mm)
7
Figure 8-46. J-R curves of Ni Cr 23Co 12Mo (Alloy 617) at room temperature and 650 °C after aging (20000 h at 650 °C) (Huthmann et al., 1989).
about 100 K below the temperature at which most of the data of irradiated specimens for RPVs of LWRs were generated. It is known from experimental programs (Pachur, 1988) that a reduction of the irradiation temperature may further increase the DBTT compared to LWR conditions (see Fig. 8-47), but the effects of the anticipated fluence of the order of 10 18 n/cm2 (E > 1 MeV) are small. Findings on spectrum effects (Mansur and Farrell, 1990) at low irradiation temperatures (about 50 °C) have produced a need for a better understanding of radiation effects and their temperature depen-
dence under HTGR environmental conditions. The information available about the Japanese HTGR with a temperature of 400 °C for the RPV steel (2% Cr 1 Mo) indicates no significant irradiation effects (Nucl. TechnoL, 1984). The irradiation effects on graphite are treated in Chap. 5 of this Volume. Metallic materials, for example for the absorber rods, are exposed to a neutron fluence of the order 1022 n/cm2 (E>0.1 MeV) in the THTR design or in the PNP concept (Thiele et al., 1990). At operating temperatures of about 400 °C the radiation hardening scarcely affects the ductility. A significant degradation of the ductility is observed at emergency conditions for the absorber rods with temperatures around at 850 °C. High-temperature embrittlement, which is observed for all austenitic steels and Ni alloys, increases significantly at lower strain rates. An optimization of alloys and a modification of the HTR module design with reduced radiation exposure at these loading conditions were the engineering measures taken to combat these problems. 8.7.7 Corrosion of Structural Materials in Helium The corrosion of metallic materials in HTGR helium has been surveyed by various authors (Nucl. Technol., 1984; Quadakkers and Schuster, 1985; Graham, 1990). The observation of corrosion of reactor components in helium (e.g., FSV, Peach Bottom) is very rare and without an indication of problems in the operating temperature range. Most of the information is available from laboratory experiments simulating HTGR environmental conditions (see Table 8-29).
8.7 High-Temperature Gas-Cooled Reactors (HTGRs)
300-
163
oA533B,HSST03,LS (.21C/.012R.015S,.024AI,.11Cu/.62Ni)
ANiCrMo 9
(.15C;.010R.017S,.021AI/.15Cu,3.2Ni)
250
~
Neutron Ruence:6-7x10l9n/cm2[E»iMeVi Irradiation Time: 624h ( FRJ-2 ^ 200-
w
Ni-ions
\5MeV
|1O2
c o
>v \
£ 1
protons 5MeV
^
•aio \
210"* in
neutrons
8 to-6
5MeV
10"v-8
j
10"r10
102
\ |\ P
for e"
103 10* 105 Recoil energy T (eV)
Ni 106
1
107
cles with different mass as they slow down in a given target material. From the examples of Fig. 9-4, the following features are evident: light charged particles mostly transfer recoil energies which are little above the threshold energy for displacement, TA9 producing one or a few Frenkel defects per recoil event, and the distance between events is large compared to the dimensions of the defect zones. With increasing mass of the bombarding ions the number of recoils with T$>Td increases and their distance decreases. Therefore, going from He to Xe ions leads to an increase of the mean size and to smaller separation of the defect zones (later called displacement cascades) until these cascades finally
Figure 9-3. Differential cross section der/df for transferring a recoil energy T to nickel atoms as a function of T for different irradiation particles.
- • • • " "
60 keV He 3
\
"*
able energy by head on collisions and is given by T
max
=
(M + fn)2
.
.
.
.
-
.
,
—
*
"
•
•
*
"
and Tmax
M We2
2 £
(9-3)
for nonrelativistic particles and relativistic electrons, respectively. M and m are the masses of the target atom and the bombarding particle, respectively. Since the spatial arrangement and many other features of the radiation-induced defect structure (see Sec. 9.4.3) depend on the recoil energy T transferred to the PKA, it is important to characterize the recoil spectra produced by the different bombarding particles. A qualitative way to illustrate the damage patterns produced by different recoil energies is to follow the paths of parti-
500 keV C u * 3
J /*> 0
100
1 MeV Xe 1 3 0
200
300
x(nm) Figure 9-4. Results of TRIM calculations of the trajectories of ions with increasing mass and energy in Cu. The shaded areas indicate the locations (projected onto the drawing plane) of all the knock-ons with energies T > Td. The median transferred energies T1/2 are 0.2, 3, 20 and 50 keV, respectively, for the He3, N 14 , Cu 63 and Xe 130 ions shown.
190
9 Physics of Radiation Damage in Metals
overlap. Fast neutrons (not shown in Fig. 9-4) produce cascades of similarly large sizes as in the case of 1 MeV Xe ions. Their distance is, however, very large because of the small cross section da/dT of neutrons (see Fig. 9-3). For a quantitative characterization of recoil spectra it is very informative to plot the fraction W(T) of the damage energy associated with recoil events with T" > Td and V smaller than a given value T. An example is given in Fig. 9-5 for the same particles as in Fig. 9-3. Considering the logarithmic scale for T in Fig. 9-5 it is evident that the different irradiation particles transfer their energy to the lattice atoms in quite different energy regimes because of their different masses and/or interaction potentials. If we compare, e.g., the median recoil energies, T1/2, (that are the recoil energies up to which half of the displacements are produced), this energy is about 60 eV (^2T d ) for 5 MeV electrons but about 60keV (^2000T d ) for fast reactor neu-
trons, with the other projectiles lying between these extreme values. Besides T1/2, the primary recoil events are characterized by the steepness of W(T) indicating the energy range of the recoils which contribute significantly to displacement damage. For example, neutrons produce recoils in a rather narrow range of T, whereas for light ions practically all recoil energies, starting from Td up to Tmax, contribute about equally to displacement damage. 9.2.3 The Production of Foreign Atoms
In addition to the production of Frenkel defects, the properties of nuclear materials can be degraded by the creation of foreign elements by nuclear reactions (dE/dx\n in Sec. 9.2.1). Particularly important are reactions where gases are generated [such as (n,a), (n,n'a), (n,p), (n,d), (n,t), etc. reactions] because gaseous atoms, especially helium, are thought to degrade the longterm mechanical integrity of some reactor
107 Recoil energy T (eV)
Figure 9-5. Fraction W(T) of the damage energy producing recoils in nickel with T' smaller than a given value T. For the self ion, Ni, two curves are shown: The solid line (step at 5 MeV) applies if the bombarding Ni-ion is considered as a PKA starting at the sample surface. The dotted curve is valid if energy transfers to Ni atoms excluding the bombarding Ni ion are considered.
9.2 Primary Interactions Between Radiation and Solids
components, This has already been recognized in the midsixties during the development of alloys for core components of fast breeder reactors (Barnes, 1965; Harries, 1966). Since the cross sections for the above nuclear reactions are usually appreciable only above neutron energies of a few MeV (Fig. 9-6), this problem is even more severe in materials envisaged for the first wall and for the blanket components in future fusion reactors which will be exposed to 14 MeV neutrons. Three types of reactions are the main source of helium in a material M with atomic mass A and nuclear charge Z: Fast neutron induced (n, a) reactions, such as A-
3
3 6 9 12 Neutron energy (MeV)
191
15
Figure 9-6. Cross section for the production of helium by (n, a) processes as a function of the neutron energy for different materials (SS: austenitic stainless steel; Kulcinski et al, 1975).
(some MeV) (9-4 a) Although the cross sections vary considerably from isotope to isotope (see Fig. 9-7 for 14 MeV fusion neutrons), they are substantial for all nuclei, i.e., the production of helium by fast neutrons cannot be avoided by selecting alloys of special composition. This is in contrast to helium producing reactions involving thermal neutrons, nth. For steels there are important especially the two step reaction
+ Jnth !|Ni + Jn*-
(9-4 b) +y + fHe (4.76 MeV)
with cross sections of about 0.7 and 10 barn, respectively, and the reaction 10i
£He (1.47 MeV)
(9-4 c)
with a cross-section of 4010 barn. Equations (9-4) show that the a particles ejected from the reacting nuclei have energies of some MeV. This has two important consequences: (a) During slowing down in
the surrounding material, the nuclear stopping, d£/dx| d , of the oc particles increases with decreasing energy (Fig. 9-1), i.e., a high density of defects is created at the end of the range where the a particle comes to rest in the lattice. This means that the defect environment of each helium atom (which determines its further behavior) is mainly produced by the helium atom itself and only weakly dependent on the previous defect structure or on defects created by other displacement processes, (b) In most alloys the distribution of the different constituents is not uniform which should lead to spatially varying (n,oc) reaction rates. However, since the ranges of MeV a particles in medium-density solids are several micrometers, many of these nonuniformities are largely smeared out. Even in the worst case, i.e., complete segregation of an element (e.g., boron in grain boundaries), the "enhancement factor" is only moderate, e.g., about 2 for a grain size of 25 jim.
9 Physics of Radiation Damage in Metals
50
100 Atomic weight
150
Whereas our knowledge on production rates, atomistic properties and influence of helium on macroscopic properties is rather advanced (for a recent review see Donnelly and Evans, 1991), the situation is much less clear for the large variety of neutron-induced solid transmutation products. At first sight, they do not seem to be a problem since in most cases (a few examples are given in Table 9-2) the amounts produced during the life time of reactor components are far below their respective solubility limits. However, phase diagrams obtained under equilibrium conditions may not be applicable to the irradiation state. Coupling of defect and impurity fluxes can lead to strong impurity segregation at sinks. On the other hand, also precipitates initially present can dissolve or grow under irradiation. Systematic investigations of these complex processes have started only a few years ago. Although illuminating results are now available for several model alloys and some selected stainless steels (see e.g.,
200
Figure 9-7. Cross sections for 14 MeV neutron-induced nuclear reactions producing helium in different elements. The dotted lines indicate gross trends of the cross sections for (n, a) and (n, n' a) reactions, respectively, as a function of the target mass number (Qaim, 1981).
Nolfi, 1983; Maziasz, 1988), it will take many years of experimental studies and theoretical efforts to acquire guidelines which allow an assessment of the role of impurity atoms (initially present or radiation-induced) in the behavior of materials in an irradiation environment.
Table 9-2. Examples of chemical changes in materials exposed to a fusion environment with 1 M W/m2 neutron loading (Kulcinski, 1976). Material
Production rate (appm/yr)
Solubility atO.4Tm
Mg Si
400 40
4
AISI 316 stainl. steel
Mn V Ti
1200 200 50
60 20 3
V
Cr Ti
130 80
100 70
Al
Produced element
(at.%)
0.006
9.3 Production of Individual Frenkel Pairs
9.3 Production of Individual Frenkel Pairs Single Frenkel pairs constitute the basic elements of radiation damage. They are generated by low energy knock-ons, with PKA energies typical below 100 eV produced in electron irradiation experiments. In Sees. 9.3.2 and 9.3.3 we discuss the experimental information on the threshold energies for displacement and on the initial defect arrangement in individual Frenkel pairs. Direct insight into the displacement process is obtained mostly from computer simulation studies which are presented in Sec. 9.3.1. 9.3.1 The Displacement Process For a detailed understanding of the collective motion of the atoms following the energy transfer to a PKA and leading to the generation of Frenkel defects (FD) computer simulations were of eminent importance. These so-called molecular dynamics simulations (MDS) have been pioneered by Vineyard. For reviews see Beeler (1983) and Diaz de la Rubia and Guinan (1991). For an assembly of 103 to 105 atoms interacting via suitably chosen pair potentials - or many body potentials - in a crystal one solves in small time steps the
"Explosion"
RCS
Focuson
193
classical equations of motion and so follows the individual positions (and kinetic energies) of all atoms as a function of time. The potentials chosen should correctly reproduce static properties of the crystal such as its structure, cohesive energy, elastic constants, surface energy, vacany formation energy, etc. Visualization of the displacement event is greatly aided by a movie film. As a summary of a typical displacement event in Cu, Fig. 9-8 shows the tracks of the involved atoms starting from their ideal lattice positions and leading to their final positions in the damaged lattice. The typical features that can be recognized from Fig. 9-8 are the following: Initially a randomization of the recoil energy among the atoms sitting closest to the PKA takes place. From this region, which resembles a local explosion, a so-called replacement collision sequence (RCS) evolves. In this sequence each atom successively pushes out its neighboring atom, i.e., a kind of shock wave travels with supersonic velocity ( « mach 2) along an atomic row (Foreman et al., 1992 a). During the impact (see Fig. 9-10) each atom approaches its neighbor to about half of their original distance, and later on, after the momentum has been transferred, relaxes further towards the site of its collision partner. This replacement continues until so much energy has been
Figure 9-8. MD-simulation of a single displacement event in Cu. The figure shows the motions of the atoms in a (100) plane of the f.c.c. lattice, initiated from a PKA with T = 50 eV and recoil direction 10° off the [110] axis in the plane. The PKA (shaded in its initial position) comes to rest next to the generated vacancy. The interstitial is produced at the position of the other shaded atom where the RCS got stuck.
194
9 Physics of Radiation Damage in Metals
lost to the surrounding lattice that the last atom in the sequence cannot displace its neighbor sufficiently to finally take his place. This atom, therefore, has to fall back towards its starting point. This site, however, is occupied by the preceding atom in the RCS so that two atom now have to share one lattice site, i.e., an interstitial is formed there. Indeed, in most materials studied so far, the stable self interstitial atom (SIA) has the so-called dumbbell configuration in which two atoms share one lattice site. For f.c.c. and b.c.c. metals the dumbbell axis is along and directions, respectively (Fig. 9-9). The atoms ahead of the point where the SIA was deposited in the RCS all relax back to their original sites. Their collective motion forms a so-called focuson, in which only momentum and no mass is transported along the chain, in contrast to a RCS. For an easy spreading of focusons (and RCSs), the momentum of the recoiling atoms must be steered towards the chain direction. As visualized in Fig. 9-10 this is realized best for lattice rows with small values of d, i.e., along close packed lattice directions, e.g., along and in the f.c.c. and and in the b.c.c. lattice. b.c.c.
fxx.
P
Figure 9-9. Observed stable SIA configurations: Two lattice atoms (black) share one lattice site forming a so-called dumbbell with axis along in f.c.c. and along in b.c.c. metals.
Figure 9-10. Focusing condition in an isolated atom
row: (p2Td. The main difference to the single displacement events discussed in Sec. 9.3 is that the PKA now slows down over distances much larger than a lattice distance. Thereby it creates a so called displacement cascade (DC). This is a hierarchy of secondary, tertiary, etc. displacements from which, after some atomic rearrangements, a stable defect pattern evolves. This whole process occurs in a very short time in the order of picoseconds (ps). It, therefore, cannot be followed with present experimental techniques. However, in the last few years extensive computer simulation experiments have clarified many details of the defect production process. The insight gained from these studies will be reviewed in Sec. 9.4.1. In Sees. 9.4.2 to 9.4.4 the conclusions drawn from the computer simulations are then compared with the experimental findings on final defect numbers, defect arrangements, atomic mixing and phase transformations produced by high energetic displacement events. 9.4.1 Computer Simulation of Displacement Cascades
Different computer simulation techniques have been used to study the complex process of the slowing down of an energetic ion in a crystal. For reviews see Beeler (1983), Heinisch (1990) and Diaz de la Rubia and Guinan (1991). From these studies it is now evident that, conceptually, the defect production by high energetic recoils can be divided into two consecutive stages. First the collision cascade, then the spike (see also Table 9-1). In the collision cascade (CC) the PKA generates secondary, tertiary and higher
9.4 Displacement Cascades
generation recoils in the lattice. It lasts, depending on PKA energy, about 0 . 1 0.3 ps, i.e., less than a typical atomic vibration time. The basic elements of a CC are two body collisions between rather energetic recoils and lattice atoms at rest. For this reason this stage is also called the ballistic phase. At the end of this phase all recoils have slowed down to energies below Td, so that they cannot knock out further lattice atoms. The result of the CC are the recoils, i.e., atoms displaced by more than an atomic distance from their starting points, together with the vacancies they left behind. The spike which follows the cascade must be considered as a local region where the majority of the atoms are (temporarily) in violent motion (original definition by Seitz and Koehler, 1956). The intensity of this motion decreases as the spike cools down in times of the order of 10 ps. More specifically we can think of the nascent spike, present after 0.3 ps, to contain a central region in which the kinetic and potential energies of the CC have been randomized and thus converted into heat. Simultaneously, at the periphery a shockfront has built up. In the subsequent time interval, 0.3 to about 3 ps, from this shockfront SIAs are injected into the surrounding lattice e.g., by RCSs, loop gliding or other mechanisms. Simultaneously in the core of the spike a liquidlike droplet forms which, at about 3 ps, has cooled down below the melting temperature of the lattice. Subsequently it solidifies. This process can be viewed as an extremely rapid recrystallization of the undercooled liquid spike core which then cools down to ambient temperatures. In this recrystallization process a vacancy-rich core, called a depleted zone (DZ) is left behind. The different steps of the formation of a DC are summarized in Table 9-1.
199
Before going into further details a comment on the different notations found in the literature seems appropriate. Although the spike concept was introduced rather early in the field of radiation damage, it has been a long standing open question whether the energy from the CC dissipates slowly enough to cause significant rearrangement by thermal diffusion of atoms. Therefore, over many years either the cascade aspect, based on the binary collision approximation and ignoring the many body aspects, or the (thermal) spike concept, using classical heat flow and chemical rate equations and ignoring the collision aspects, have been overemphasized. In this period it has become customary to use the word "cascade" as a generic term, e.g., in characterizing the net effect of a high energy recoil as a "displacement cascade", "defect cascade", or "damage cascade" or in calling the atomic motions from the beginning of the CC to the end of the spike "cascade dynamics" or to call the cooling phase of a spike "cascade collapse". In the following we restrict the term "displacement cascade", DC, to describe the overall effect of a high energy recoil, comprising the initial CC and the subsequent spike phase. In our terminology the spike comprises both a hot core ("thermal spike") together with a shock front at its periphery ("plasticity spike"). 9.4.1.1 Collision Cascades
The most rigorous method for following the development of a displacement cascade is the MDS. This treats both the initial CC and the spike phase in a unified manner. Its two main drawbacks have been overcome recently: First, at least for metals, rather reliable interatomic potentials, taking into account also many body aspects, are now available from the embedded atom tech-
200
9 Physics of Radiation Damage in Metals
nique and used in MDSs (Diaz de la Rubia and Guinan, 1991; Foreman et al, 1992 b). Secondly, the increased power of modern supercomputers made it now possible to treat rather large crystals of up to 106 atoms and therefore to simulate displacement cascades, up to T = 25 keV in Cu (Diaz de la Rubia and Guinan, 1991). For the initial CC most of our information has, however, been obtained from computer simulations using the binary collision approximation. In these calculations one follows, one at a time, with the computer, the continually branching collision sequences between moving atoms and atoms at rest until the post-collision energy of all moving atoms has dropped below a given value, e.g., below Td. The result of such a calculation is shown in Fig. 9-14 in the form of a trajectory map of all atoms recoiled with energies larger than 8 eV. In its central and its peripheral part this map illustrates clearly the decreasing distances between collisions when the recoil atoms
8
Figure 9-14 Projection on the {100} plane of the knock-on-atom trajectories calculated for a PKA with T = 5 keV in b.c.c. iron (Beeler, 1983). The thick line represents the trajectory of the PKA, the dotted lines those of the secondary knock-ons. Higher order knock-ons are presented by the thinner alternately solid and dashed lines.
slow down. CCs calculated under the assumption given above are also called "linear" cascades. In amorphous materials such cascades can then be treated analytically by linearized Boltzmann equations or by Monte Carlo programs (TRIM program, Biersack and Haggmark, 1980). An important consequence of its linear character is that the number of recoils with energies larger than a given value T' should increase linearly with T/Tr. This property is used in Sec. 9.4.2 to estimate the total number Nd(T) of atoms displaced by a PKA of energy T. Up to a given value of T = Tsc the CCs are rather compact, i.e., the vacancies are contained in a volume whose linear dimensions are of the order of the PKA range. Larger CCs, however, tend to branch or to break up into individual subcascades. An example of such a cascade configurations is given in Fig. 9-15. The basic reason for the break up of a cascade into subcascades is the following: With increasing T the average distance the PKA travels before creating higher energetic secondary recoils increases and can become larger than the linear dimensions of the so created secondary cascades. This is, however, a statistical process and in this respect the occurrence of subcascades is a fluctuation phenomenon. An additional effect which would influence the cascade configurations may come from planar channeling. This influence can be judged from simulations of CCs in amorphous materials. For Cu such simulations showed that the cascades in crystalline material were only slightly larger in dimension and aspect ratio than in amorphous material (Heinisch, 1990). Because of their irregular shape it is very difficult to make meaningful estimates of the energy density, 6CC, and on the vacancy density, cv, in CCs. According to the linear character of the CC and with a narrow
201
9.4 Displacement Cascades
Figure 9-15. Three-dimensional view of a 100 keV collision cascade in Cu generated by the MARLOWE computer code (Heinisch, 1990). The heavy dots represent vacancies and the faint dots SIAs. All FDs with vacancy-SI A distances of less than 0.8 nm have been eliminated from this plot (spontaneous recombination). Three well separated subcascades (1 to 3) are recognized, subcascades 2 and 3 each have three lobes.
definition, counting only the volume enclosed by the collided atoms, these densities should be constant, e.g., 9CC « 2 eV/at. and cy « 3 at.% for Cu. The other extreme would be to measure the volume inside a regular parallelepiped just enclosing the whole cascade or to calculate something like a radius of gyration for the distribution of the recoiled atoms. With these volume definitions the average density would decrease as 1/T for higher T (Winterbon etal., 1970). Such definitions, however, seem to be somewhat artificial because they do not take into account the local concentration of the damage from one CC into separate subcascades or lobes. Recently Heinisch and Singh (1992) reported a systematic computer study of the structure of CCs for several metals and as
a function of the PKA energy. From a total of almost 104 CCs simulated with the MARLOWE code they were able to extract meaningful statistical averages of the local distributions of vacancies. Trying to avoid the ambiguity from visual inspections of CCs they defined "regular subcascades" by the condition that they must contain locally a higher than average vacancy density and that the separation between the subcascades must be larger than the subcascade diameter. The mutual distances between subcascades defined in this manner, turned out to be broadly distributed with average values around 50 atomic diameters independent of atomic number and PKA energy. Inspecting Fig. 9-15 more closely one can also recognize that the vacancy arrangement within the subcasqade is not spherical but that the subcascades themselves contain branches or lobes, i.e., could be considered as adjacent smaller cascade units, often called "subcascades", too. Fig. 9-16 shows that the so calculated average number, nsc, of regular subcascades per PKA increases as a function of damage energy, Tdam [defined by Eq. (9-2)].
100
200
300
400
500
600
Figure 9-16. Average number, nsc, of subcascades per PKA versus Tdam. Points connected by solid lines are from computer-simulations using the MARLOWE code (Heinisch and Singh, 1992).
202
9 Physics of Radiation Damage in Metals
The observed linear relation Tiun>Tu>
(9-5)
where Tsc is the average damage energy per subcascade, is a direct consquence of the linear character of the CC. This means that at high recoil energies the subcascades can also be considered as new independent damage units, a concept introduced by Merkle (1976) who also first observed subcascades in the TEM. On the average each of these units contains not only an equal amount, Tsc, of recoil damage energy but also equal number of defects, atomic volumes etc. Heinisch and Singh (1992) also showed that Tsc is about equal to that PKA energy above which CCs start to break up into subcascades and below which the CCs appear as one unit. Because the nuclear stopping power increases strongly with Z, we expect that the average distance travelled by the PKA between high energetic recoils decreases strongly with Z. This has the consequence that the breakup into subcascades occurs at higher 7^c values for heavier metals. From the data shown in Fig. 9-16 one obtains Tsc values of 25 and 170 keV for Cu and Au, respectively.
within the spike core by equating the average kinetic energy of the core atoms to (3/2) kB Tsp (kB is the Boltzmann constant). As seen from Fig. 9-17 b, for Cu the maximum temperature reaches 4 times the melting temperature Tm = 1356 K in the spike center. At such high values of the local temperature the existence of a (hot) molten zone in the spike has to be expected. Wolf et al. (1990) have shown that already at about 30% superheating above the thermodynamic melting temperature inside a crystal a mechanical shear (phonon) instability occurs, which causes a solid- to liquid-phase transition within one or a few lattice vibration times. 30 20 10
! "10
t= 0.25 ps t = 1.41 ps t = 3.47 ps
Vn-J
-20 -30
10
20
30
U0
Distance from center [%)
(b) 9.4.1.2 Spikes We now turn to the spike phase of the DC. Here binary collision models are inadequate and MDS is essential to describe the randomization of the PKA energy among the atoms in the spike and the collective nature of the resulting atomic motions. The simulations show that equipartition between kinetic and potential energy is achieved very quickly in the center of a nascent spike, typically within 0.25 ps, i.e., within 1 to 2 lattice vibrations. From then on a temperature, Tsp, can be defined
t= 0.25 ps t = 1.41 ps t= 3.47 ps
0 "0
•***,*, 10 20 30 Distance from center (X)
40
Figure 9-17. Radial profiles of atomic density (a) and temperature (b) at three instants of time within a 5 keV spike in Cu (MDS by Diaz de la Rubia et al., 1989).
9.4 Displacement Cascades
203
IttlSSfSfJ:
mssthr nut sis
i
iiiiiiiiiii ISSsii
iiliiiiliii H iiiiiiiiiii!
11! 5nm
The cross section through the center of a 10 keV MD cascade in gold shown in Fig. 9-18 exemplifies this disordered liquid-like region surrounded by the well preserved crystal lattice. The local atomic arrangement within the spike core can be characterized quantitatively by calculating a radial distribution function, g(r), which measures the probability to find, within one atomic volume, another atom at distance r. Figure 9-19 compares g(r) for the atoms within the spike core with g(r) of liquid copper (solid line). The striking sim-
Figure 9-18. {100} projection of instantaneous atomic configurations at 1 ps after the PKA event within a cross sectional slab of 3 atomic planes near the center of a 10 keV cascade in Au. MDS by Averback (private communication). Note that the shown computer print has slightly different scales in the x and y directions of the atomic positions. By an inclined viewing of the graph the shock front and the wake along the RCSs are clearly visible.
ilarity between the pair-correlation functions demonstrates clearly the liquidlike structure of the spike. The complete destruction of the crystallinity is well documented by the absence of the (200) peak at r = 3.6 A which appears in the MD-calculated g (r) at later times when the melt in the spike solidifies. MDS has also demonstrated that the local atomic density within the core of the nascent spike is reduced, e.g., by approximately 10 to 15% for Cu (see Fig. 9-17 a). To balance the atomic dilution within the
calculated for liquid Cu
Figure 9-19. Radial distribution function g (r) at two selected times within a 5 keV spike in Cu. MD-simulations by Diaz de la Rubia et al. (1989).
204
9 Physics of Radiation Damage in Metals
core a ridge of compressed materials builds up at the periphery of the spike as is also evident from Figs. 9-17 a and 9-18. This high density ridge is the manifestation of the shock-front (see Fig. 9-18) initiated during the radial propagation of the CC. The local mechanical stress in this shockfront temporarily reaches the order of the theoretical critical shear stress, /i/(2 n)9 with JLL being the shear modulus. After this description of the nascent spike we now discuss the first cooling phase which we call the spike relaxation phase (Table 9-1). During this phase, lasting about 3 ps, stable interstitial defects are deposited outside the molten core. The core itself cools down below Tm, remaining, however, in the liquid - now supercooled state. We start with a discussion of the interstitial formation process. Because of the abundance of free volume in the molten core interstitial defects cannot survive there. Only interstitials ejected from the shock-front far enough into the surrounding lattice have a chance to escape recombination. Therefore, this ejection process is the dominant process which determines
the formation of stable FDs. To conserve the number of atoms within a DC for each interstitial surviving at the periphery, a vacancy must finally be retained in the spike core. The mechanisms by which interstitials are produced by the shock-front are not yet fully understood. One process observed in MDS and shown in Fig. 9-20 is the ejection of RCs. This process seems to be similar to the single displacement event discussed in Sec. 9.3. However, in contrast to the single displacement events, where SIA deposition by RCSs is observed in Cu only along the close packed directions and , MDS suggest also ejection of RCSs along from cascades (see Fig. 9-20). The distances at which the interstitials are deposited seem to depend critically on the temperature of the surrounding lattice and on fine details of the interatomic potential. Typical average RCS lengths observed in MDS for Cu are 15 and 23 A corresponding to about 6 and 10 replacements along (Foreman et al, 1992 a; Diaz de la Rubia and Guinan, 1991).
Figure 9-20. Three-dimensional view of the lattice sites (indicated by the crosses) where atomic replacements have occurred in a 5 keV DC in Cu. The radius of the central molten core (not shown) is approximately 1.8 nm. The final positions of interstitials are indicated by the dots. MDS result by Diaz de la Rubia et al. (1989).
9.4 Displacement Cascades
Recently MDS has - apart from RCSs revealed another process by which interstitials are produced in the form of clusters at the periphery of the spike (Diaz de la Rubia and Guinan, 1991; Foreman et al, 1992 b). In this process, platelets of interstitials form in the shock front, either spontaneously or by rapid agglomeration of transient single SIAs. These plateles have the form of perfect {111}-dislocation loops and are therefore highly glissile along the direction of their Burgers vector 1/2 . By first gliding away from the molten spike core (see Fig. 9-21) they seem to escape recombination. In the later period of the spike relaxation, when the pressure wave of the shock-front relaxes back and compresses the expanded spike core again, these loops can return part way. However, in order to survive recombination, they must remain at a sufficient distance not to be incorporated into the shrinking molten core. The elastic interaction of the peripheral interstitial cluster with the vacancy core leads to a preferential nucleation and survival of loops with Burgers vectors tangential to the DC. In summary, MDS suggests that among the surviving SIAs more than 50% are present in form of small clusters (see Fig. 9-22) formed either by a cooperative process, or by the aggregation of single SIAs ejected from neighboring RCSs. From this description it also becomes clear that the number of surviving stable SIAs and their partitioning among single entities and clusters will depend not only on Tdam but also on the specific mechanical and structural properties of the material under consideration. Having discussed the interstitial ejection from the spike periphery we now return to the hot molten core. From this core heat is dissipated into the surrounding lattice mainly by thermal collisions, i.e., via phonons. Also electrons may be heated
205
Time (ps)
Figure 9-21. Position of the interstitial loop (shown in Fig. 9-22 in its final location) during the cooling of a 25 keV spike in Cu (MDS by Diaz de la Rubia and Guinan, 1991). The positions are given as distances from the spike center along the loop-glide direction. The solid line gives the positions the loop would have if it glided with the velocity of an elastic C44-shear wave.
Figure 9-22. Three-dimensional view of the final arrangement of vacancies (circles) and interstitials (dots) in a 25 keV DC in Cu (MDS by Diaz de la Rubia and Guinan, 1991). The arrow points to a loop of 17 SIAs on a {111} plane.
206
9 Physics of Radiation Damage in Metals
within the spike and for metals, because of the dominance of electrons in thermal conductivity, a large contribution to the spike cooling would be expected, at first sight. Following, however, the discussion given by Flynn and Averback (1988), the decisive parameter is the mean free path, X, of electrons in the melt. For liquid Cu, X = 45 A whereas for liquid Ni or Fe estimates of X are 8 and 5.5 A, respectively. These X values have to be compared with the dimension, Rsp, of the spike core. For Rsp < X electron-phonon coupling has no influence on the cooling of the spike. This situation is typical for Cu. However if i?sp > X, the electrons will make a kind of random walk within the spike and thereby acquire, at each collision, an energy /cB@D, where @D is the Debye temperature. In this way electrons can heat up to Tsp if Rsp/X > 5 to 10. If by this heating also the degeneracy of the electron system can be lifted (i.e., no longer Fermi energy > kB Tsp), then, especially for d band metals, the heat capacity of the electronic system would strongly increase and a large fraction of 7^am could be dissipated very rapidly by hot electrons leaving the atoms in the spike core relatively cool. Beyond these qualitative considerations, however, there is little quantitative information on the role of electron-phonon coupling in spike cooling. Although it might be extremely important at present no experimental evidence exist for or against it. The cooling of spikes has been followed in several MDS. For Cu, e.g., from Fig. 9-17 b it can be seen that after about 3.5 ps the central temperature in a 5 keV spike has dropped below the melting temperature. Recently Alurralde et al. (1991) have modeled the cooling of spikes in different metals and for different PKA energies. They used a binary collision code to determine the recoil energy distribution in a CC
and converted this into an initial temperature profile. The further cooling, i.e., volume and temperature of the molten zone as a function of time, was then calculated using the classical equations for heat conduction in an isotropic medium. By expressing the time steps used in their calculation in units of C/(a %), they avoided the necessity of knowing explicitly the variation of specific heat per atom, C, mean atomic distance, a, and thermal conductivity, x, with temperature. However, since the phase transition in the CC is far from thermodynamic melting, the definition of a proper melting criterion is difficult; for example one would have to quantify the minimum superheating necessary for the melt zone to spread rapidly enough, i.e., with a velocity of the order of the velocity of sound vs (Protasov and Chudinow, 1982). In addition part of the energy in the CC could go into electronic heating and thus spread much more rapidly than the temperature motion of the ions. Nevertheless, these calculations show that the melt zone defined as the volume where Tsp > Tm first expands only a little and then starts to shrink after typically 2 ps. Figure 9-23 shows, for Cu as a function of the PKA energy, the calculated maximum volume of the molten spike region, Fsp,m, and the time, t sp?m , it takes until Tsp has dropped below Tm everywhere in the spike. These results compare well with the results from MD simulations for Cu. If also for the other metals electron-phonon coupling is neglected, the data reported by Alurralde et al. (1991) for Cu, Ni, Ag, Fe and Pd suggest the following approximate relationship: sp, m
dam
14fe B T m
(9-6)
i.e., a linear increase of Vsvm with 7^am. Because of Eq. (9-5), this linearity is not surprising for Tdam > Tsc, where indepen-
9.4 Displacement Cascades
10
100
1000
100
1000
7"(keV)
10 T(keV)
Figure 9-23. Maximum volumes (a) and lifetimes (b) of the molten region(s), calculated for PKAs of different energies T in Cu by Alurralde et aL (1991). The points * in (a) are the results of a MDS.
dent subcascades are formed. For smaller compact cascades the linearity implies that the energy density, 0 CC , deposited in the CC is independent of % a result typical for linear cascades. Also the scaling of Vsp^m with the melting temperature Tm seems plausible. Equating the total (kinetic and potential) energy per atom to 3/cBTsp, Eq. (9-6) suggests an initial heating of the spike core up to Tsp « 4.7 Tm. This is, however, an upper limit which neglects that part of Tdam has been dissipated into the surrounding of the melt (20 to 30%) and also assumes that there is no heating of conduction electrons. The calculated "lifetime" Tsp,m, as shown in Fig. 9-23 b, is almost constant above a PKA energy of 20keV in Cu. This is a
207
consequence of the break up of the cascades into subcascades which then cool independently. The small increase observed in Fig. 9-23 b above T « 20 keV could reflect a proximity-effect in subcascade cooling. For T < 20 keV, i.e., for compact cascades, T spm is found to decrease approximately as r sp m oc T0'5. The increased spike cooling rate responsible for this decrease of T sP,m is a direct consequence of the increasing surface to volume ratio with decreasing cascade size. The precise meaning and the absolute magnitude of r sp m are difficult to assess, e.g., the values quoted in Fig. 9-23 b were obtained by adjusting the results given by Alurralde et al. (1991) in arbitrary units to the MDS results of Diaz de la Rubia and Guinan (1991). Also i sp?m is not the real time when all the melt has completely recrystallized. At quenching rates of approximately 1015 K/s and typical temperature gradients at the spike periphery of 100 K/A a high degree of undercooling is expected to occur in the melt region. This undercooling is, indeed, evident from the MD simulations: At spike cooling times for which the temperature has already dropped below Tm radial distribution functions typical for liquids are still observed (see e.g. Figs. 9-17b and 9-19). Nevertheless T spm can serve as an approximate number to characterize the average lifetime of the molten spike volume. Another question is: To what minimum DC size can the concept of a molten spike core be applied? One criterion necessary to define a molten zone is that the number of atoms within the melt should at least be comparable to the number of atoms within the transition region to the crystalline surrounding. Taking a thickness of 1 atom layer for the transition between liquid and solid one arrives at a minimum value of F s p m of about 100 Fa corresponding to a PKA energy T « 0.5 keV. A similar mini-
208
9 Physics of Radiation Damage in Metals
mum value can be obtained by requiring that the average melt lifetime r sp m cannot be smaller than the minimum time it takes for the liquid core to solidify with a velocity of the order of vs. We now turn to the final stage of the spike cooling, the solidification of the liquid melt core. All MDSs show that this process occurs rather rapidly due to the rapid heat dissipation into the surrounding lattice. After a typical time of 10 ps the temperature in the spike center has dropped practically to ambient. As discussed by Turnbull (1984) the velocity of the melt front during the solidification approaches its limiting value vs9 the sound velocity in the melt. The frequency of meltto-solid infacial rearrangement is then of the order of the lattice vibration frequency. The undercooling necessary to drive the solidification with such a high speed was estimated to be AT/Tm « 0.3 in agreement with MDS (Diaz de la Rubia and Guinan, 1991). Only for pure metals or random solid solutions it is possible for the interfacial atoms to epitaxially rebuild the original crystal structure within the short time available. For more complicated structures, e.g., ordered alloys, non equilibrium phases will be quenched in (see Sec. 9.4.5). What happens to the vacancies in the spike? At the beginning of the spike relaxation phase, the vancancies are distributed as "free volume" in the liquid. From the observed density reduction in the core (see Fig. 9-17 a) one can estimate that this "free volume" corresponds initially to a vacancy concentration of m.ore than 10%. During the relaxation phase most of this "free volume" is, however, filled up again by a back™ flow from the compressed lattice in the shock-rim (see Fig. 9-17 a). At the end of the spike cooling phase, only so many vacancies can survive as stable SIAs have been deposited outside the melt zone.
When the spike core solidifies these vacancies are then quenched and form the depleted zone (DZ). They are, however, not distributed at random. The infilling process of the spike from the surrounding shock-front and a kind of zone refining process at the crystallization front drive the vacancies towards the center of the spike (Protasov and Chudinov, 1982). Under favorable conditions they agglomerate there into immobile clusters or loops (see e.g., Fig. 9-22, 9-26 and the experimental results in Sec. 9.4.3). 9.4.2 Damage Function and Definition of dpa (Displacements per Atom)
The "true" damage function is defined as the average number, v(T), of Frenkel defects produced by a PKA of energy T starting in a random crystal direction. This number must not be confused with the "displacement" (damage) function defined by the number, JVd, of calculated "displacements" in a CC of energy T. As discussed in Sec. 9.4.1 an identification of Nd(T) with v(T) is questionable especially for larger recoil energies, % where spike effects are important. Historically the first estimate of Nd(T) was given by Kinchin and Pease (1955). This estimate used the idea that displacements in a CC occur as long as after each collision both atoms have an energy still somewhat larger than an (isotropic) threshold energy, Td, i.e., sufficient for both collided atoms to escape the vacancy created in the last collision. This means that, on the average, Nd « T/(2 Td). Taking into account electronic energy losses [by introducing Tdam(T% see Eq. (9-2) and Fig. 9-2] and the fact that the colliding atoms do not behave as hard spheres (by the factor 0.8) one arrives at the so called NRT approximation named after Norgett, Robinson
9.4 Displacement Cascades
209
and Torrens (Norgett et al., 1975). It is (9-7) T dam 2.5 Td T d < T < 2.5 Td T < Td This function is shown in Fig. 9-24 for Cu. At high PKA energies Nd increases linearly with Tdam. At low energies, JVd would also reproduce the average threshold behavior of single displacement events by taking T d =r d ? a v . Today Nd serves as a widely used standard to calculate the so-called number of displacements per atom, dpa, as a dose unit for a given irradiation condition. For a discussion see Garner et al. (1990). According to ASTM/E 521-89 (1989) the dose unit dpa is defined as
dpa = •Emax
= j
T" m ax
t
J
(9-8)
Figure 9-24. Damage function, v(T), and displacement function, Nd (T) for Cu. The dashed curve shows Nd from Eq. (9-7), using Td = 30 eV (Table 9-3). The "experimental" damage function v(T) (solid curve) was obtained by multiplying Nd(T) with £°(T) estimated from Fig. 9-25 and, for T < 100 eV, by evaluating the experimental threshold surface (King et al., 1982). The arrows denote Td min and Td.
]Nd(T)d<j(E,T)d has been introduced (Reiley and Jung, 1977), where is the average over the recoil spectrum centered at T1/2. From
210
9 Physics of Radiation Damage in Metals
1.2 1.0
o • • A 4>
0.8
electrons ions fast neutrons thermal neutrons MDS
0.6
0.4 0.2 0 101
102
103
10*
105
TO6
r V2 (ev) Figure 9-25. Damage efficiency £° (i.e., FDs produced per dpa) as a function of the median recoil energy T1/2 in copper, deduced from resistivity measurements on specimens irradiated with different particles at 4.2 K. Compilation of data by Averback et al. (1978) and Kinney et al. (1984), using Td = 30 eV and AQFD = 2.5 j^Qm. The symbols connected by the solid lines give results from MDSs by Foreman et al. (1992 b) and by Diaz de la Rubia and Guinan (1991), respectively. For the irradiations with heavy ions, e.g., 500 keV Cu, T1/2 refers to the (secondary) recoils giving partly overlapping DCs. If the implanted ion is counted as a PKA the data point for Cu must be plotted at T1/2 = Tdam = 240 keV.
Fig. 9-25 it can be clearly recognized that £° decreases with increasing T1/2 until, at T1/2 > 20 keV, £° becomes constant, £tn « 0.25. This constancy of £ is expected for subcascade formation. Values of £°n for other elements are collected in Table 9-3. In Fig. 9-25, £° values obtained from MDS for different PKA energies are also shown for comparison. Their steeper initial decrease, compared to the experimental £°(Tlf2) values, is consistent with the fact that the T spectra produced by the light ions are smeared out and contain large contributions from Tdf min < T < 2 Td. This is also the reason why £° > 1 for the e~irradiations, with T1/2 near Td = 30 eV. In this case ATd underestimates the real defect production as can be judged also from Fig. 9-24. TJ^fn is the average energy necessary to produce a stable FD by fast neutrons and characterizes the defect production at large
recoil energies. From Table 9-3 can be seen that for most metals this energy is 4 to 8 times larger than Td min . As explained in Sec. 9>.4.1.2, the average number of defects produced in a larger DC is solely determined by the number of SIAs ejected far enough from the shock-front of the spike to survive incorporation and recombination at the hot spike core during spike relaxation and cooling. Therefore TJ^n cannot be related directly to the minimum recoil energy, Td?min, to produce an isolated FD. Similarly, the increase in the damage efficiency below T1/2 ~ Tsc reflects the fact, that SIAs can escape more easily from small spikes. The concept of displacements leading directly to interstitial-type defects within the center of the DC becomes meaningless if applied to the hot liquidlike spike core. The extraction of an "experimental" damage function, v(T), from the measured
9.4 Displacement Cascades
damage rates for different irradiation particles requires a deconvolution with respect to the different recoil spectra. This is possible only with limited accuracy. Nevertheless the qualitative features of the v(T) curve shown in Fig. 9-24 should be correct: At low energies v(T) is determined by the detailed form of the threshold energy surface Td(Q): The steep increase of v(T) near Td min reflects the onset of defect production by recoils near the "easy" crystal directions which are and for Cu. The first levelling off of v(T) reflects the steep increase of Td (Q) between the angular windows for easy displacements (see Fig. 9-12). When T exceeds the threshold energy for the "hard" displacement directions, e.g., around , the defect production increases again and levels off as soon as all possibilities for single displacements are exhausted. The subsequent increase is then due to the onset of multiple displacements. At energies T&TSC& 25 keV the cascades split up into subcascades as explained in Sec. 9.4.1. At these higher energies nsc [see (Eq. (9-5)] and, therefore, also v(T) both increase linearly with Tdam. Very frequently simplified theoretical calculations of v(T), e.g., by the fast MARLOWE code, are used. These calculations are based on the binary collision approximation, so that the atomic rearrangement during the spike phase is not incorporated. This rearrangement can be approximated by eliminating later on all those vacancies and displaced atoms of the calculated CC which are separated by less than a chosen distance. By a proper choice of this recombination distance these models can give correct numbers of stable FDs and they also yield a vacancy rich DZ surrounded by a halo of interstitials. However, they are inadequate descriptions of the finer details in the defect arrangement such as the formation of vacancy or interstitial clusters
211
and the extensive atomic mixing occurring in the spike (see Sec. 9.4.4). 9.4.3 Defect Arrangement in Displacement Cascades
Three experimental techniques have mainly been applied to study the defect arrangements in DCs: Field ion microscopy (FIM), transmission electron microscopy (TEM) and diffuse X-ray scattering (DXS). The FIM technique is the only one capable of atomic resolution. It produces a point-projection image of the atoms on the surface of an approximately hemispherical tip with a typical radius of 5 to 30 nm. For stable FIM images, the tip is cooled to cryogenic temperatures and it must be made of a material with an evaporation field higher than the ionization field of the imaging gas. Most success has been obtained with refractory metals such as W, Pt and Ir. The interior of a specimen can be examined by increasing the applied voltage and thereby removing atoms from the surface through the field-evaporation effect. In this manner it is possible to dissect a volume of about 105 nm 3 atom by atom and to take three-dimensional maps of the vacancy and interstitial configurations. In the most extensive study of that kind Seidman et al. (1987) investigated DCs introduced by in situ ion implantation into W and Pt tips. Typical vacancy arrangements within the DZs created in W are shown in Fig. 9-26. Although the different ions produced about the same total numbers of vacancies, their arrangement is quite different. The W projectile produces a high density CC resulting in an average vacancy concentration = 27 at.% in the DZ. Because of their larger mean free path the lighter projectiles produce more widespread CCs
212
9 Physics of Radiation Damage in Metals
Figure 9-26. Visualizations of typical vacancy arrangements as determined by FIM in individual DCs (after Seidman et al., 1987). DCs created by implantation of different 30 keV ions into W at 18 K. Vacancies are represented by open circles. The rods connect vacancies at first nearest neighbor positions. The surrounding lattice is omitted for clarity. The diagonal dimensions represented by (a), (b) and (c) are 4, 6.5 and 17 nm, respectively. The total number of vacancies is about 155 in each case.
with, e.g., = 3.6 at.% in the DZ of the Cr + ion. The volume in which the vacancies are contained has a rather irregular shape. As indicated in Fig. 9-28 this DZ volume can be defined by the stacking of parallelepipeds drawn around the outermost vacancies within each slab. This volume was also used to define the average vacancy concentrations given above. With increasing size of the CC the tendency for cascade branching becomes apparent, e.g. as in Fig. 9-26 b. The vacancies within the DZ are not distributed at random. This is illustrated by Fig. 9-27 which shows radial distribution functions, gv (r), evaluated for the vacancy
arrangements produced by different implanted ions. gy (r) gives the fraction of lattice sites arround each vacancy occupied by another vacancy within a shell at distance r. For a random distribution, gy(r) should be constant (= 0.05^ B
>0.0lJ v
i
T
-4- V m
'
—>
vm
+s
->
im I,
+ s +S
v,
Annihilation at sinks
D
T/Tm >0.3
Vm + A n(AVJ + S Im
VTm >0.2 >
—» m'
-4- V
V
Vacancy emission
+ A
w(AIJ + S
a
t m
lm + i, r
11,V1 Im, Vm lg If, VjS A, An (AVJ, (AIJ
-y _
]
V
Interaction with foreign atoms
+ .-
T/Tm >0.05^
0
]
Radiation damage effect
s s s
T/Tm >0.2 T/Tm
^ (AVJ -> S+A n
m
E C,D C,D
>0.2 "V
^± (AIJ
-> S+ An
denote single interstitials, vacancies with m = 1,2,3 denote mobile single and multiple interstitials, vacancies denotes5 a glissile loop of g interstitials denote immobile clusters of i interstitials, vacancies denote*5 sink such as network dislocation, grain boundary, surface (internal, external) denote single foreign atoms and precipitates of n foreign atoms denote complexes of a foreign atom with Vm and Im; these complexes may be also mobile
The letters A to F in the last column refer to the first column in Table 9-5.
hit by another DC. (Also at lower temperatures the build-up of irradiation induced defect does not continue indefinitely. In this case the defect concentrations saturate as a result of spontaneous recombination processes from cascade overlaps. This leads to FD concentrations of the order of 10 ~3, much higher than the steady-state concentrations considered here.) The condition r\m = £ for the achievement of a quasistationary state is somewhat relaxed if the sinks absorb Els much more effectively than EVs. In this case the mobile vacancies
visit the recombination sites around a (possibly) stable interstitial cluster lt remaining from a DC more often than the mobile interstitials and the cluster will shrink even without thermal I± emission. The times necessary to establish stationary concentrations of the mobile interstitials and vacancies are typical of the order of their migration times the sinks S. For vacancy clusters, Yi9 this time corresponds to their thermal decay time plus the migration time of the liberated Vx to S. Since the interstitials move much faster than the vacancies
9.5 Defect Production and Reactions at Elevated Temperatures
233
Table 9-5. Macroscopic radiation damage effects. Effect
Process
Irradiation induced seg- A-atoms couple to I and regation and changes in V-fluxes to sinks; nonequilibrium phases due to miprecipitate structure crochemical changes
Temp, range
Important for . . .
T > 0.2 Tm
corrosion, weldability and indirect influence on all following effects B - F
Low temperature embrittlement, increase of ductile to brittle transition temperature
hardening by dislocation loops and irradiation-induced precipitates
Irradiation creep under mechanical load
stress-induced bias for preferential absorption of I and 0.2 71 < T < 0.4 71 V at favorably oriented dislocations
most nuclear materials under irradiation and mechanical stress
D Irradiation growth
anisotropic nucleation and 0.1 T < T < 0.3 T growth of dislocation loops
noncubic materials (Zr and alloys, U, graphites)
E Void swelling
bias for preferential absorption of I at dislocations and consequent surplus of vacancies flowing to voids
austenitic steels for core components in LMFBR and 1st wall comp. in FR
Helium high temperature embrittlement under creep and fatigue loads
nucleation and growth of He bubbles on GB leading to premature intergranular failure
the stationary concentrations of the EVs will always be much higher than the stationary concentrations of the Els. There are two limiting cases which can be distinguished for the defect situation in the quasisteady state: It can either be recombination dominated or sink dominated. In the sink dominated regime, basically, the stationary concentration of the EVs is much smaller than the concentration of the existing sinks and vice versa in the recombination dominated regime. After the quasisteady state situation has been established an equal total number of interstitials and vacancies has to arrive at the entirety of the sinks per unit time. This is a consequence of the overall balance between total I and V numbers in the simul-
0.1 71 < T < 0.3 71
0.3 71 < T < 0.5 Tm
T > 0.45 T
b.c.c.-steels and refractory alloys for pressure vessels
1st wall structures in FR core components LMFBR control rods HTR
taneously occurring production and recombination processes. Given only one type of sink, e.g., only dislocations or only voids, due to the I-V balance in the defect fluxes - on the average - no climb of the dislocation or growth of the voids will occur. This is the basic reason that any changes in the sink structure of a given sample will occur on a time scale much larger than the times necessary for the establishment of quasi-steady concentrations of the EVs and Els. Slow changes in the sink structure may be caused by a redistribution of atoms from one sink to another (or between members of the same sink type). Also the casual stabilization of some irradiation deposited Yt clusters may contribute to a slow change in the concen-
234
9 Physics of Radiation Damage in Metals
trations of sinks. The stabilization is a result of fluctuations in the sizes and in the net arrival rate of vacancies which enables some clusters to grow further by Oswald ripening processes instead of decaying thermally. The phenomenon of redistribution of atoms between different sinks under a steady supply of equal numbers of Is and Vs during the irradiation can be traced back to a so-called "bias" of these sinks for the capture of interstitials over the capture of vacancies. The capture efficiency (or strength), p\'y, of a given sink can be expressed by the number of atomic volumes per sink which, if visited by a three-dimensional migrating I or V, will lead to its incorporation into the sink. For dislocations p\iV is the number of absorbing atomic volumes in a slab of thickness equal to the Burgers vector perpendicular to the dislocation line. The values pi1 for the capture of single interstitials may be larger than pj1 for the capture of single vacancies, e.g., because of the larger elastic interaction of Ix with S compared to V^ Other reasons for different ps values may be the poisening of a sink by deposited impurities or an applied stress. A stress field causes polarization effects or an anisotropy of the defects moving in different directions with respect to the stress direction. Especially large bias effects or imbalances in the I and V fluxes to dislocations and cavities, respectively, are to be expected if the Is move one-dimensionally via glissile loops \g. This bias has recently been discussed by Trinkaus et al. (1992) and is basically due to the fact that an interstitial loop is captured within much larger distances in the long ranging stress field of a dislocation than at a corresponding cavity. Recently the terms "production bias" (Abromeit and Wollenberger, 1988; Woo and Singh, 1990) and "cascade localization
induced bias" (CLIB) (Yoshiie et al., 1991) have been introduced in the literature. The term "production bias" derives from the fact that a DC introduces different numbers of interstitials and vacancies in clusters. In the steady-state situation discussed above, this effect itself, however, does not lead to an imbalance in the EI/EV fluxes to different sinks in a direct way. The CLIB mechanism assumes that from a DC overlapping with a cavity - on the average more vacancies will end up at that cavity than interstitials. This assumption, however, is based on an oversimplified model of the DC and needs to be tested critically by MDSs before it can be considered as a valid mechanism for radiation induced void growth. Imbalances in the fluxes of I and V to the different types of sinks in a material under irradiation give rise to several important radiation damage phenomena in nuclear materials (see also Table 5-9). Swelling by void growth was soon after its unexpected discovery by Cawthorne and Fulton (1967) associated with the preferential absorption of interstitials at dislocations (Bullough and Perrin, 1969; Harkness and Lie, 1969; and others). Theoretical models based on this concept yield a semiquantitative description of the dependence of the swelling rate on temperature, dose rate and microstructure. No adequate theories exist, however, for the incubation period which precedes the swelling stage and which depends sensitively and in a complex manner on microstructure, gas content, type of irradiation, etc. Irradiation (or in-pile) creep is thought to be caused by the stress-induced preferential absorption (SIPA; Heald and Speight, 1975) of interstitials at dislocations of different orientation with respect to the applied stress. But there are experimental observations which cannot be explained by SIPA alone and thus several
9.5 Defect Production and Reactions at Elevated Temperatures
other mechanisms have been proposed to contribute to irradiation creep (see e.g., Mansur, 1988). The situation is similar for irradiation growth, i.e., an anisotropic dimensional change of noncubic materials under irradiation (for details see e.g., Woo and McElroy, 1988). A more comprehensive discussion of methods, achievements and problems in modeling radiation damage effects is found in a review by Mansur (1988). An analysis of the evolution of the microstructure which takes into account the latest findings in cascade damage is given by Wiedersich (1991) and Trinkaus et al. (1992). Experimental results and modeling efforts on radiation-induced changes in the microstructure of specific alloy systems are found in the pertinent chapters of this Volume and in the following references: Gittus (1978), Nolfi (1983), Garner et al. (1987) and the proceedings of the ASTM International Symposia on Effects of Radiation on Materials and of the International Conferences on Fusion Reactor Materials. 9.5.3 Helium Bubble Formation
The extraordinary properties of helium (and other inert gases) as an impurity in bubble
235
solids are due to its high enthalpy of solution. Values of several eV (e.g., 3 eV in Ni) lead to equilibrium concentrations far below the ppm range even close to the melting point. This means that helium produced in nuclear materials by one of the processes described in Sec. 9.2.3 is present in high supersaturation and has thus a strong tendency to precipitate into bubbles. Experimental results and theoretical considerations (Trinkaus, 1985,1986) have shown that the development of the bubble population during continuous helium build-up under irradiation proceeds in the following way (Fig. 9-39): Immediately after the onset of irradiation and helium generation, the concentration of "solute" (single mobile) helium atoms, cHe increases linearly with time (incubation stage) until clustering of these atoms becomes substantial. Then absorption of diffusing helium by the clusters reduces the rate of increase of cHe. When the gain of "solute" helium by continuing generation is balanced by the loss to clusters, both cHe and the clustering rate CB reach maximum values (nucleation peak at t = t* in Fig. 9-39). Afterwards both quantities decrease and the cluster/ bubble density CB saturates or increases only very slowly. In this stage, the newly
void
incuba- nuclea- gas absorp- vacancy condensation tion tion tion
/ /
m
o
L CHe/
/
X
A
„»
Figure 9-39. Time dependence of solute (mobile) helium concentration cHe, cavity (bubble or void) density CB and average cavity radius rB. The characteristic stages of the cavity evolution under continuous helium production are outlined (schematic, not to scale). The time of maximum nucleation rate t* is usually very short compared to the duration of the following stages.
236
9 Physics of Radiation Damage in Metals
created helium is almost completely absorbed by the existing bubbles (growth stage). Because of the low solubility of helium in solids, bubble nucleation occurs at extremely low helium concentrations, i.e., after very short service times. Although the detailed evolution kinetics of bubble populations is complicated there are two limiting cases, for which explicit expressions for the bubble density reached after the end of the nucleation stage can be derived: (1) At low temperatures (T < 0.4 Tm) and high helium introduction rates (e.g., by aparticle implantation or tritium decay) even dihelium clusters do not decay before capturing an additional helium atom and thus forming stable bubble nuclei. In this case, the evolving bubble number density is given by the ratio of the helium introduction rate and the helium diffusion coefficient. If, in addition, diffusion controlled clustering of thermal or irradiation-induced vacancies is slow, helium bubble formation is only possible if the required space is produced by athermal processes such as self-interstitial emission or dislocation loop punching (see Trinkaus and Wolfer, 1984, for a review). (2) At high temperatures (T>0.4T m ) and low or moderate helium generation rates PHe (i.e., in the parameter range of high temperature helium embrittlement of nuclear materials), only clusters above a certain minimum size are stabilized by the continuing helium supply against decay. Hence, the nucleation process has to overcome a barrier Gc depending strongly on the supersaturation as well as on the type of nucleation site (Fig. 9-40). In this case, the final bubble density C% in the grain interior is proportional to the ratio of the helium generation rate PHe and the helium dissociation rate from the bubble nuclei (dissociation or permeation controlled nu-
easy nucleation
(a)
(b) Figure 9-40. (a) Dependence of the free enthalpy G on the volume V of a cavity (schematic). Vc is the volume of the critical nucleus. For V < Vc a cluster tends to decay whereas for V > Vc it will grow, (b) Examples of bubble sites with decreasing nucleation barrier Gc: 1 in the ideal lattice, 2 at a precipitate in the grain, 3 in a grain boundary and 4 at a grain boundary precipitate (schematic).
cleation). Accordingly, the temperature dependence of Cg is determined by the sum of the activation energies for resolution of helium atoms from bubble nuclei and for the helium mobility. Experimental observations confirm these predictions. Figure 9-41 shows that the data of different investigations in austenitic stainless steel lie within a rather narrow band if normalized to the production rate PHe. The tern-
237
9.5 Defect Production and Reactions at Elevated Temperatures
perature dependence of C^/PHe yields an activation energy of somewhat below 3 eV which agrees with the permeation energy of helium in stainless steel determined by independent helium bubble coarsening experiments (Rothaut et al., 1983). The situation is more complicated for bubble nucleation at extended defects such as dislocations, grain boundaries and precipitate-matrix interfaces, which act as helium absorbers and sites of easy nucleation. The nucleation at such a site is controlled by the helium flux from the grain interior. If this (time dependent) flux were known, the resulting bubble densities and sizes could be estimated on a similar basis as for the matrix bubbles. However, such a flux affects the bubble evolution in the vicinity of the heterogeneity which, in turn, affects the flux. This coupling of bubble formation at the defects and in the nearby matrix
renders a theoretical description very difficult. However, useful analytical expressions have been derived (Trinkaus, 1986; Singh and Foreman, 1989) for different combinations of nucleation mechanisms (diatomic and multiatomic) in the matrix and at grain boundaries, respectively. They are the basis for a theoretical modeling of high temperature embrittlement of nuclear materials by helium-induced intergranular failure. After nucleation has ceased, incorporation of additional helium produced with the passage of time, t, increases the size of the bubbles but does not change their density (Fig. 9-39). At T > 0.4 Tm the bubbles may be assumed to absorb sufficient thermal or irradiation-induced vacancies to keep their internal gas pressure close to its equilibrium value p^ e , given by PL = —
1/mo'V1) Figure 9-41. Bubble concentration Cg normalized to the helium production rate PHe vs. reciprocal temperature 1/T in stainless steel implanted with helium at T (compilation of data from different authors). The solid line corresponds to an activation energy of 2.8 eV, which is close to the permeation energy of He in stainless steel (after Singh and Trinkaus, 1992).
(9-H)
where y is the surface free energy and rB is the bubble radius. The mean radius rB of such equilibrium bubbles increases as (PHe t)112 for ideal gas behavior (Fig. 9-42). A bubble population is generally spending most of its life time in this gas driven growth stage. For a detailed discussion including nonideal gas behavior and grain boundary bubbles see Trinkaus (1986). An applied mechanical stress and/or an irradiation induced vacancy supersaturation changes the gas driven bubble growth into an accelerated growth due to vacancy condensation (Fig. 9-39). For this transition a critical condition must be reached from which a critical radius rB can be defined:
Bff2
32 F v3
< >- = ^ F
with r =
4v
» ?
(9-12)
kB T of helium atoms 3 a in the n is the 27 number critical bubble, Fv = V/rl < 4 JI/3 is a geo-
238
9 Physics of Radiation Damage in Metals
Figure 9-42. Mean radius rB and concentration C£ of helium bubbles in DIN 1.4914 martensitic steel as a function of the helium content cHe. The material was implanted with a rate of lOOappm/h at 873 and 973 K, respectively (from Stamm, 1988).
metrical factor relating the bubble volume V to the radius rB. o can be an actual stress or an effective driving force ("chemical" stress) caused by an imbalance of the fluxes of the irradiation-induced interstitials and vacancies, respectively, to the cavity (biaseffect). Figure 9-43, showing the free energy of a gas-containing cavity as a function of its radius, illustrates this "bubble to void" transition: if a bubble has accumulated a number n of helium atoms during the gas driven growth stage so that no2 = (n&2)c or, equivalently, rB = rB, the minimum in the free energy disappears and unstable growth by vacancy condensation sets in. The concept of a critical radius, rB, where cavity growth changes from bubble to void behavior has been used to analyse irradiation induced swelling (Mansur et al., 1986). In modeling high temperature embrittlement due to helium (Trinkaus, 1985) this concept has been successfully combined with the nucleation theory mentioned above. In these models, the onset of fast unstable growth at rB > rB is equated with the beginning of the linear void swelling regime and with crack initiation at grain boundary cavities, respectively. For a general review on the influence of helium on the properties of nuclear materials see Ullmaier (1984).
9.6 Acknowledgements
-3
Figure 9-43. Free energy F of a bubble containing n atoms of ideal helium gas and subject to an external stress a vs. its radius rB. Fo and rB are the equilibrium values of F and rB, respectively, o — 0. The parameter represents the ratio of n a2 to its critical value (n G2\ for which the bubble becomes unstable, a is compressive for the upper and tensile for the lower curves.
We are grateful to R. S. Averback (Urbana), B. N. Singh (Riso) and H. Wollenberger (Berlin) for a critical reading of the manuscript and for helpful comments. We are particularly indebted to H. Trinkaus (Julich) for many fruitful discussions and valuable suggestions. Finally we thank Mrs. M. Garcia for typing the manuscript and handling the numerous changes with skill and patience.
9.7 References
9.7 References Abromeit, C , Wollenberger, H. (1988), I Mater. Res.
3(4), 640. Alurralde, M., Caro, A., Victoria, M. (1991), / Nucl Mater. 183, 33. ASTM/E 521-89 (1989), Standard Practice for Neutron Radiation Damage Simulation by ChargedParticle Irradiation. Philadelphia, PA: Am. Soc. Test. Mat., p. 167. Averback, R. S. (1986), Nucl Instr. Meth. Phys. Res. B15, 675-687. Averback, R. S., Seidman, N. D. (1987), Mater. Sci. Forum 15-18, 963. Averback, R. S., Benedek, R., Merkle, K. L. (1978), Phys. Rev. B18, 4156. Averback, R. S., Horngming, H., Diaz de la Rubia, T., Benedek, R. (1991), / Nucl. Mat. 179-181, 87. Banerjee, S., Urban, K. (1984), Phys. Stat. Sol. (a) 81, 145. Barnes, R. S. (1965), Nature (London) 206, 1307. Becker, D., Dworschak, E, Lehmann, C , Rie, K. T., Schuster, H., Wollenberger, H., Wurm, J. (1968), Phys. Stat. Sol. 30, 219. Becker, D. E., Dworschak, R, Wollenberger, H. (1972), Phys. Stat. Sol. (b) 54, 455. Beeler, J. R. (1983), Radiation Effects Computer Experiments. New York: North-Holland Phys. Publ. Bender, O., Ehrhart, P. (1983), /. Phys. F13, 911. Biersack, J. P., Haggmark, L. G. (1980), Nucl. Instr. and Meth. 174, 257 (and later versions of TRIM). Brailsford, A. D., Bullough, R. (1981), Phil Trans. Roy. Soc. (London) 302, 87. Bullough, R., Perrin, R. C. (1969), in: Radiation Damage in Reactor Materials. Vienna: IAEA, STI/ PUB/230, p. 233. Bullough, T, English, C. A., Eyre, B. L. (1987), Mat. Science Forum 15-18, 1069. Burger, G., Meissner, H., Schilling, W (1964), Phys. Stat. Sol. 4, 267. Cawthorne, C , Fulton, E. J. (1967), Nature (London) 216, 515. Diaz de la Rubia, T, Guinan, M. W. (1991), UCRLJC-107488. Diaz de la Rubia, T, Averback, R. S., Horngming, H. (1989), /. Mater. Res. 4, 579. Donnelly, S. E., Evans, J. H. (Eds.) (1991), Fundamental Aspects of Inert Gases in Solids. New York: Plenum Press. Dunlop, A., Lesueur, D., Morillo, I , Spohr, R., Vetter, J. (1989), C. R. Acad. Sci. Paris, 309, Serie II, 1277-1282. Ehrhart, P. (1991), in: Landolt-Bornstein, Vol. HI/25, Atomic Defects in Metals: Ullmaier, H. (Ed.). Berlin: Springer-Ver lag. Ehrhart, P., Averback, R. S. (1989), Phil. Mag. A 60, 283. Ehrhart, P., Schlagheck, U. (1975), in: Fundamental Aspects of Radiation Damage in Metals, Proc. Int.
239
Conf: Robinson, M. T, Young, F. W. (Eds.). Oak Ridge, TN: Oak Ridge Report CONF-751006-P2. Ehrhart, P., Robrock, K.-H., Schober, H. R. (1986), in: Physics of Radiation Effects in Crystals: Johnson, R. A., Orlov, A. N. (Eds.). Amsterdam: North Holland Phys. Publ. English, C. A., Jenkins, M. (1987), Mater. Sci. Forum
15-18, 1003. Flynn, C. P., Averback, R. S. (1988), Phys. Rev. B38, 7118. Foreman, A. J. E., English, C. A., Phythian, W. J (1992 a), Phil. Mag. A 66, 655. Foreman, A. J. E., Phythian, W. X, English, C. A. (1992 b), Phil Mag. A 66, 671. Garner, F. A., Perrin, J S. (Eds.) (1985), in: ASTMSTP 870, Vol. II. Philadelphia, PA: Am. Soc. Test. Mat, pp. 863-1163. Garner, F. A., Packan, N. H., Kumar, A. S. (Eds.) (1987), in: ASTM-STP955. Philadelphia, PA: Am Soc. Test. Mat. Garner, F. A., Heinisch, H. L., Simons, R. L., Mann, F M. (1990), Radiation Effects and Defects in Solids 113, 229-255. Gittus, J. (1978), Irradiation Effects in Crystalline Solids. London: Appl. Sci. Publ. Guinan, M. W, Kinney, J. H., Van Konynenburg, R. A., Damask, A. C. (1981), /, Nucl. Mater. 103/ 104, 1217-1220. Greenwood, L. R., Kneff, D. W, Skowronski, R. P., Mann, R M. (1984), /. Nucl. Mater. 122/123,1002. Hamberg, A. (1914), Geol Fonen Stockholm Fohr 36, 31. Hameed, M. Z., Smallman, R. E., Loretto, M. H. (1982), Phil Mag. A46, 707-716. Hardtke, Ch., Schilling, W, Ullmaier, H. (1991), Nucl. Instr. Meth. Phys. Res. B59/60, 377-381. Harkness, S. D., Li, C. Y. (1969), in: Radiation Damage in Reactor Materials. Vienna: IAEA, STI/ PUB/230, p. 189. Harries, D. R. (1966), J. Brit. Nucl Energy Soc. 5, 74. Heald, P. T, Speight, M. V (1975), Acta Metall 23, 1389. Heinisch, H. L. (1990), Radiation Effects and Defects in Solids 113, 53-73. Heinisch, H. L., Singh, B. N. (1992), Phil. Mag., to be published. Hertel, B., Diehl, X, Gotthardt, R., Sultze, H, (1974), in: Applications of Ion Beams to Metals. New York: Plenum Press, pp. 507. Hohenstein, M., Seeger, A., Single, W (1989), / Nucl Mater. 169, 33. Howe, L. M., Rainville, M. H. (1979), Phil Mag. A 39, 195-212. Ishida, L, Sasaki, T, Yoshiie, T, Iwase, A., Iwata, X, Kiritani, M. (1991), J. Nucl Mat. 179-181, 913916. Jager, W, Merkle, K. L. (1988), Phil. Mag, A 57, 479-498. Jaouen, C , Riviere, X P., Delafond, X, (1991), Nucl Instr. Meth. Phys. Res. B 59/60, 406-409.
240
9 Physics of Radiation Damage in Metals
Jenkins, M. L., Wilkens, M. (1976), Phil. Mag. 34, 1155-1167. Johnson, W L. (1986), Progr. in Mat. Sci. 30, 81-134. Jung, P. (1991), in: Landolt-Bornstein, Vol. 111/25, Atomic Defects in Metals: Ullmaier, H. (Ed.). Berlin: Springer-Verlag. Kim, S.-X, Nicolet, M.-A., Averback, R. S., Peak, D. (1988), Phys. Rev. B37, 38. Kinchin, G. H., Pease, R. S. (1955), Rep. Progr. Phys. 18, 1. King, W. E., Benedek, R., Merkle, K. L., Meshii, M. (1982), in: Point Defects and Defect Interaction in Metals: Takamura, I, Doyama, M., Kiritani, M. (Eds.). Tokyo: Univ. Tokyo, p. 789. Kinney, J. H., Guinan, M. W., Munir, Z. A. (1984), /. Nucl. Mater. 122/123, 1028. Kiritani, M., Yoshiie, T, Kojima, S., Satoh, Y. (1990), Rod. Eff. and Defects in Sol. 113, 75. Kirk, M. A., Blewitt, T. H. (1982), J. Nucl. Mater. 108/109, 124-136. Koike, X, Okamoto, P. R., Rehn, L. E., Meshii, M. (1989), /. Mater. Res. 5, 1143-1150. Kulcinski, G. L. (1976), in: Proc. Int. Conf Radiation Effects and Tritium Technology for Fusion Reactors, Gatlinburg, U.S.A.: Watson, J. S., Wiffen, F. W. (Eds.). Springfield, VA: U.S. Dept. of Commerce, p. 1-17. Kulcinski, G. L., Doran, D. G., Abdou, M. A. (1975), Special Technical Publications 570. Philadelphia, PA: Am. Soc. Test. Mat., p. 329. Liu, H. C , Mitchell, T. E. (1983), Acta Metall. 31, 863-872. Luzzi, D. E., Mori, H., Fujita, H., Meshii, M. (1986), Acta Metall. 34, 629-639. Mansur, L. K. (1988), in: Kinetics of Nonhomogeneous Processes: Freeman, G. R. (Ed.). New York: Wiley-Interscience. Mansur, L. K., Lee, E. H., Maziasz, P. X, Rowcliffe, A. P. (1986), /. Nucl. Mater. 141-143, 633. Maziasz, P. X (1988), in: Proc. Internat. Conf on Materials for Nuclear Reactor Core Applications, Vol. 2. Bristol, England. London: BNES. Merkle, K. (1976), in: Radiation Damage in Metals: Peterson, N.L., Harkness, S. D. (Eds). Metals Park, OH: Am. Soc. Met., p. 75. Metzner, H., Sielemann, R., Klaumunzer, S., Hunger, E. (1987), Phys. Rev. B. 36, 9535. Miehle, W., Plewnia, A., Ziemann, P. (1991), Nucl. Instr. Meth. Phys. Res. B59/60, 410-413. Muller, A., Naundorf, V, Macht, M. P. (1988), J. Appl. Phys. 64, 3445, Muroga, T, Ishino, S., Okamoto, P. R., Wiedersich, H. (1984), J. Nucl. Mater. 122/123, 634. Nastasi, M., Mayer, X W. (1991), Materials Science Report 5, 1-51. Nolfi, F. V. (Ed.) (1983), Phase Transformations During Irradiation. London: Appl. Sci. Publ. Norgett, M. X, Robinson, M. T, Torrens, I. M. (1975), Nucl. Eng. Des. 33, 50.
Okamoto, P. R., Meshii, M., (1989), in: Science of Advanced Materials: Wiedersich, H., Meshii, M. (Eds.). Metals Park, OH: Am. Soc. Met., pp. 3 3 98. Packan, N. H., Stoller, R. E., Kumar, A. S. (Eds.) (1990), ASTM-STP1046, Vol.11. Philadelphia, PA: Am Soc. Test. Mat., p. 5-373. Peisl, X, Franz, H., Schmalzbauer, A., Wallner, G. (1991), Mat. Res. Soc. Symp. Proc, Vol. 209. Chicago, IL: Mat. Res. Soc, p. 271. Phythian, W. X, English, C. A., Yellen, D. H., Bacon, D. X (1991), Phil. Mag. A 64, 821-833. Protasov, V I., Chudinov, V G. (1982), Radiation Effects 66, 1-7. Qaim, S. M. (1981), Handbook of Spectroscopy, Vol. 3. Boca Raton: CRC Press, p. 141. Rauch, R., Peisl, X, Schmalzbauer, A., Wallner, G. (1990), J. Phys. Condens. Matter 2, 9009. Reiley, T. C , Jung, P. (1977), in: Proc. Int. Conf on Radiation Effects in Breeder Reactor Structural Materials, Scottsdale, U.S.A.: Bleiberg, M. L., Bennett, X W. (Eds.). New York: AIME, p. 295. Rehn, L. E., Okamoto, P. (1987), Mater. Sci. Forum
15-18, 985. Rothaut, X, Schroeder, H., Ullmaier, H. (1983), Phil. Mag. A 47, 781. Schilling, W, Burger, G., Isebeck, K., Wenzl, H. (1970), in: Vacancies and Interstitials in Metals: Seeger, A., Schumacher, D., Schilling, W, Diehl, X (Eds.). Amsterdam: North Holland, Phys. Publ. Schulson, E. M. (1979), J. Nucl. Mat. 83, 239-264. Seidman, D. N., Averback, R. S., Benedek, R. (1987), Phys. Stat. Sol. (b) 144, 85. Seitz, F., Koehler, X S. (1956), Solid State Phys. 2, 305. Shimomura, Y, Fukushima, H., Guinan, M. W. (1990), J. Nucl. Mater. 174, 210. Singh, B. N., Foreman, A. X E. (1989), ASTM-STP 1046. Philadelphia, PA: Am. Soc. Test. Mat. p. 555-571. Singh, B. N., Trinkaus, H. (1992), J. Nucl. Mater. 186, 153. Stamm, U. (1988), Ph.D. thesis, RWTH Aachen (Germany). Theiss, U., Wollenberger, H. (1980), J. Nucl. Mater. 88, 121. Thome, L. (1988), in: Nuclear Physics Applications on Materials Science: Recknagel, E., Soares, X C. (Eds.) Dordrecht: Kluwer Academic Publ., pp. 183-208. Trinkaus, H. (1985), /. Nucl. Mater. 133/134, 150. Trinkaus, H. (1986), Radiation Effects 101, 91. Trinkaus, H., Wolfer, W. G. (1984), J. Nucl. Mater. 122/123, 552. Trinkaus, H., Singh, B. N., Foreman, A. X E. (1992), to be published in /. Nucl. Mater. Turnbull, D. (1984), in: Second Israel Materials Eng. Conf: Grill, A., Rokhlin, S. I. (Eds.). Israel: BeerSheva.
9.7 References
Ullmaier, H. (1984), Nuclear Fusion 24, 1039.t Ullmaier, H. (Ed.) (1991), in: Landolt-Bornstein, Vol. 111/25, Atomic Defects in Metals: Berlin: Springer-Verlag. Urban, K., Saile, B., Yoshida, N., Zag, W. (1982), Point Defects and Defect Interaction in Metals: Takamura, I , Doyama, M., Kiritani, M. (Eds.). Tokyo: Univ. Tokyo Press, p. 783. Vetrano, I S., Bench, M. W, Robertson, I. M., Kirk, M. A. (1989), Metall. Trans. 20 A, 2673. von Rossum, M., Cheng, Y-T. (1988), Defect and Diffusion Forum 57/58, 1-32. Wallner, G., Franz, H., Rauch, R., Schmalzbauer, A., Peisl, J. (1989), Mat. Res. Soc. Proc. Vol. 138, 35. Weber, W. I , Mansur, L. K., Clinard, F. W., Parkin, D. M. (1991), J. Nucl. Mater. 184, 1. Wiedersich, H. (1991), Nucl. Instr. Meth. Phys. Res. B 59160, 51. Wiedersich, H. (1992), Materials Science Forum 9799, 59. Wigner, E. P. (1946), J. Appl. Phys. 17, 857. Williams, S., Poate, J. M. (Eds.) (1984), Ion Implantation and Beam Processing. New York: Academic Press. Winterbon, K. G., Sigmund, P., Snaders, J. B. (1970), Mat. Fys. Medd. Dan. Vid. Selsk
37,1.
Wolf, D., Okamoto, P. R., Yip, S., Lutsko, I K , Kluge, M. (1990), /. Mater. Res. 5, 286. Wollenberger, H. (1990), Nucl. Instr. and Methods in Phys. Res. B48, 493-498. Woo, C. H., McElroy, R. J. (1988), J. Nucl. Mater. 159. Woo, C. H., Singh, B. N. (1990), Phys. Stat. Sol. B159, 609 [and (1992), Philos. Mag. A, to be published].
241
Yoshiie, T, Satoh, Y, Kojima, S., Kiritani, M. (1991),
179-181, 954. Zinkle, S. J. (1988), /. Nucl. Mat. 155-157, 1201.
General Reading Atomic Defects in Metals, in: Landolt-Bornstein, Vol. 111/25; Ullmaier, H. (Ed.) (1991). Berlin: SpringerVerlag. Averback, R. S., Seidman, D. N. (1987), Energetic Displacement Cascades and their Roles in Radiation Effects, in: Materials Science Forum, Vol. 15-18. Zurich: Trans Tech Publications. Mansur, L. K. (1988), Mechanics and Kinetics of Radiation Effects in Metals and Alloys, in: Kinetics of Nonhomogeneous Processes: Freeman, G. R. (Ed.). New York: Wiley-Interscience. Physics of Irradiation Effects in Metals, in: Materials Science Forum, Vol. 97-99: Szenes, G. (Ed.) (1992). Zurich: Trans Tech Publications. Physics of Radiation Effects in Crystals: Johnson, R. A., Orlov, A. N. (Eds.) (1986). Amsterdam: North-Holland. Radiation Damage in Metals: Peterson, N. L., Harkness, S. D. (Eds.) (1976). Metals Park, OH: Am. Soc. Met. Ullmaier, H., Schilling, W. (1980), Radiation Damage in Metallic Reactor Materials, in: Physics of Modern Materials, Vol. I. International Atomic Energy Agency-SMR-46/105.
10 Fusion Reactor Materials Dale L. Smith, Richard F. Mattas and Michael C. Billone Argonne National Laboratory, Argonne, IL, U.S.A.
List of Symbols and Abbreviations 10.1 Introduction 10.2 Structural Materials 10.2.1 First Wall/Blanket Structure 10.2.1.1 Austenitic Steels 10.2.1.2 Ferritic/Martensitic Steel 10.2.1.3 Vanadium Base Alloys 10.2.2 Divertor Structural Materials 10.2.2.1 Copper Alloys 10.2.2.2 Molybdenum Alloys 10.2.2.3 Niobium Alloys 10.2.2.4 Summary of Divertor Structure Issues 10.3 Plasma Facing Materials 10.3.1 Introduction . 10.3.2 Graphite and Carbon-Carbon Composites (CCCs) 10.3.2.1 Graphite Base Properties 10.3.2.2 Graphite Sputtering 10.3.2.3 Hydrogen/Tritium Retention 10.3.2.4 Neutron Irradiation Effects 10.3.3 Beryllium 10.3.3.1 Beryllium Base Properties 10.3.3.2 Physical Sputtering 10.3.3.3 Tritium Retention and Release 10.3.3.4 Neutron Irradiation Effects 10.3.4 Tungsten 10.3.4.1 Base Properties 10.3.4.2 Physical Sputtering 10.3.4.3 Neutron Radiation Effects 10.4 Blanket Materials 10.4.1 Introduction 10.4.2 Materials' Properties' Database for Solid Breeders/Beryllium 10.4.3 Materials' Properties' Correlations for Solid Breeders/Beryllium 10.4.4 Database Evaluation for Solid Breeders/Beryllium 10.4.5 Comparison of Solid Breeder/Beryllium Properties 10.4.6 Models for Solid Breeder/Beryllium Performance Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. Allrightsreserved.
245 248 252 253 255 263 267 270 272 272 276 276 278 278 281 281 282 284 287 290 290 291 291 293 294 296 296 296 298 298 300 300 301 305 310
244
10.4.6.1 10.4.6.2 10.4.6.3 10.4.6.4 10.4.7 10.4.7.1 10.4.7.2 10.4.7.3 10.4.7.4 10.4.8 10.4.9 10.5 10.5.1 10.5.2 10.5.3 10.5.4 10.6
10 Fusion Reactor Materials
Thermal Performance Models Mechanical Performance Models Tritium Retention/Release Models Helium-Induced Swelling Models Summary of Solid Breeder/Beryllium Blanket Materials Thermal Performance Tritium Performance Mechanical Performance Chemical Stability/Compatibility Liquid Metal Coolants Summary Insulators Bulk Insulators Windows Optical Fibers Reflectors References
311 314 315 321 323 324 324 325 326 327 332 332 332 333 334 335 336
List of Symbols and Abbreviations
245
List of Symbols and Abbreviations a a1 a2 as C C (a, t) Cp CT D Do De Z) p , dp E EB ED Es ET Ex G G Ga g h hc Hs /, J o J k /ca kh kd fceff /cg kjc kr M m N{ Nu P p p,p 12
effective grain radius accommodation coefficient for steel accommodation coefficient for ceramic breeder specific pore-solid surface area mobile hydrogen concentration time-dependent surface concentration of tritium specific heat at constant pressure trapped hydrogen concentration diffusivity; displacements per atom pre-exponential hydraulic diameter particle diameter; large, small Young's modulus; energy hydrogen-defect binding energy migration energy heat of solution detrapping energy activation energy for an atom to enter the bulk from the surface implant source term due to hydrogen bombardment from the plasma; hot gap; linear distortion tritium generation rate H e content j u m p distance; mass flux radiation conductance term heat transfer coefficient hardness of the softer material (steel) tritium inventory bulk hydrogen flux, bulk tritium flux Boltzmann constant; thermal conductivity adsorption rate constant conductivity of pebble material desorption rate constant effective thermal conductivity H e conductivity fracture toughness molecular recombination rate constant thermal stress factor 2g/Dp cycles to failure Nusselt number Larson-Miller parameter porosity pressure
246 Pf
pg Pn2 *H 2 O
*HTO
Pr Q* R *i
r2 r
12
10 Fusion Reactor Materials
packing fraction He pressure H 2 partial pressure H 2 O moisture pressure HT partial pressure HTO moisture pressure Prandtl number heat of transport ideal gas constant root-mean-square (rms) roughness height (in m) of steel rms roughness height (in m) of breeder or Be
[(rl + rD/lV'2
Re S
Reynolds number solubility tritium solubility st Sievert's law pre-exponential So allowable stress sm allowable time-dependent stress sm T,Th,T0 temperature T/Tm homologous temperature Tx steel surface temperature T2 breeder or Be surface temperature Te plasma edge temperature Tg average gap temperature Tirr irradiation temperature Tm melting temperature Tr reference temperature ^test test temperature V, Vo volume V bulk velocity Z atomic number a am S
molecular sticking coefficient; thermal expansion coefficient mean thermal expansion coefficient tensile ductility
S
kjkh
e s± 82 sc elch 8e es Aet sth A
total strain steel thermal emissivity emissivity of ceramic breeders or Be creep component of total strain thermal creep component elastic c o m p o n e n t of total strain swelling c o m p o n e n t of total strain total strain range thermal expansion component of total strain 4.4919786 x 10" ^ T 2 - 3 3 2 ^
List of Symbols and Abbreviations
ji v Q Q0 Qirr a cr a ae a0 O.5r m -5O°C) 5%
1/3 UTS @ T l%e 1/3 UTS @ T 5%e
100 DPA 150 DPA Neutronic propertiesd (at 1 MW • y/m2) DPA H (appm) He (appm) 2(W/cm 3 )
>750 b >650 b 750 b
125 -200(425 °C)
200(600°C)
205 (500 °C) 195(550°C) 205(500°C) 190 (550 °C) 190(500 °C) 175 (550 °C) 190(400 °C) 150(550 °Q 155 (400 °C) 130(550 °C)
175(500°C) 160(550 °C) 145(500°C) 85(55O°C) 175(500°C) 160(550°C) 198(400°C) 160(550°C) 163(400°C) 150(550°C) 125(400°C) 115 (550 °C)
220(500°C) 235(650°C) 220(500°C) 235(650°C) 220(500°C) 235(650°C)
11.3 594 157 9.8
11.1 450 110 9.8
11.3 240 57 7.1
165(500°C) 165(650°C) 125(500°C) 125 (650 °C)
a
At 400 °C for PCA and HT-9; at 500 °C for VCrTi [references for data given by D. L. Smith et al. (1991 a)]; Estimated values for VCrTi; c Based on calculations by Majumdar reported by D. L. Smith et al. (1991a); d Approximate: values depend on blanket composition [unpublished results from Gohar (1981)].
b
Chemical Compatibility Chemical compatibility with candidate coolants, tritium breeding materials and the hydrogen (DT) plasma are an important consideration in the selection of the
structural material. Compatibility of austenitic steels with the helium coolant does not present a design constraint. Corrosion mass transfer of austenitic steel in water is not a lifetime problem, but radioactive mass transfer may present maintenance
258
10 Fusion Reactor Materials
0
200
400 600 800 1000 Temperature (°C)
1200
Figure 10-8. Thermal conductivity of candidate alloys. Nb 20
900
I
800 — 0
200
I
I
400 600 800 TEMPERATURE, °C
1000
1200
Figure 10-10. Electrical resistivity of candidate alloys.
240
INCO 625 0
200
400 600 800 1000 1200 Temperature, (°C)
-85
Figure 10-9. Specific heat of candidate alloys.
V-15Cr-5Ti Ti - 6242
J 0
200
I
400 600 800 Temperature (°C)
I 1000
1200
Figure 10-11. Elastic moduli of candidate alloys.
259
10.2 Structural Materials
T
I
I
TENSILE DATA
I en
10
Id
P
200
400
10
600
— NO RUPTURE L-M = T (20 + log tr)•, T IN K, t r IN HOURS
J_J
18
L
19 20 21 22 23 24 LARSON-MILLER PARAMETER, PxlO"3
Temperature, (°C)
25
Figure 10-12. Calculated thermal stress factor for selected structural alloys.
Figure 10-13. Larson-Miller plot for selected structural alloys.
difficulties. However, aqueous stress corrosion of the austenitic steels under fusionrelevant conditions is an important consideration. Because of radiolysis, control of free oxygen species in the water at desirable low levels may not be possible. Radiation hardening of the steel and cyclic
operation may also exacerbate the problem. Figure 10-14 illustrates the sensitivity of austenitic steels to stress corrosion cracking as a function of test temperature (Ruther and Kassner, 1993). Increased sensitivity to cracking is observed at temperatures above about 150°C. As shown
300
Temperature (°C) 250 200 150
Temperature (°C) 100
300
10 7
10"7
Types 316NG& 316LSS
o o
m1
101 1.7 1.8 1.9
2
-8 ppm O2 HP. water 2.1 2.3 2.4 2.5 2.6 2.7
1000/T (K-1)
150
100
R =0.7 Tr = 1 0 s
10* -
DC CD O
Heat no. D440104 13198 16850
200
Types 316NG& 316LSS
R -0.7 Tr = 10s DC
250
109 10"1
Heat no. 0 D440104 • 13198 A 16850
1011 1.7 1.8 1.9 2
•
-8 ppm O2 5 ppm Cr
2.1 2.3 2.4 2.5 2.6 2.7 1000/T (K-1)
Figure 10-14. Temperature dependence of crack growth rates (CGRs) of Type 316 SS in oxygenated water.
260
10 Fusion Reactor Materials 1
100-
o o
-
80
1
'
This study
Refs.
• •
0
Type 304 Type 316
CO
CD 6 0
H
40 -
1
1
o
•• • ••
•
I
DO: 32ppm Strain rate: 1.7-2.8x10'V o
•
D
20 -
c0 • , 1Q23
1Q
24
i
1Q26
-JQ25
°%
Neutron fluence (n/m2, E >1 MeV) O.-J dependence . •L$l«— Figure 10-15. Fluence of IGSCC in Type 304 and 316 stainless steels.
Temperature (°C) 650 600 550 500 450 400
3
10
!
!
in Fig. 10-15, Kohyama et al. have also noted that neutron irradiation to modest fluences also affects the stress corrosion cracking of austenitic steels (Kohyama et al. 1992). Cold-worked steel is generally considered to be more susceptible to stress corrosion than steels in the solution annealed condition. The corrosion resistance of austenitic steels to lithium and Pb-Li alloy blanket materials has been investigated under a variety of conditions. The corrosion process is dominated by dissolution of nickel from the alloy. Therefore the corrosion rate of the PCA alloy with the higher nickel content is higher than that for type 316 steel. The Pb-Li alloy is more corrosive than
350
^
f
Type 316 stainless steel TCL and HT-9 TCL and FCL
Weight loss in lithium * f $ ORNL OS70OO ANL WARD • UW A SU H JAERI
eg
E Type 316 stainless— steel "FCL
(0
o CO CO
b 10c
10"
" ICL BJ$ Type 316 stai n I ess 4 PCA steel ^HT-9 FCL • C O Type 3 1 6 stainless steel DType 316 CW stainless steep AAAType 304 stainless steel 0 PCA 52 u V HT-9 €O9Cr-1Mo 1.0
1.1
1.2
1.3 1.4 1000/T (K)
Scatter band Type 316 stainlesssteel FCL
1.5
1.6
Figure 10-16. Corrosion rate data for austenitic and ferritic steels exposed to flowing lithium. 1.7
261
10.2 Structural Materials
lithium. The corrosion data, as correlated by Chopra (Chopra and Smith, 1986) are presented in Figs .10-16 and 10-17. A nominal corrosion limit for a liquid metal system is ~ 20 |im/y. Based on this corrosion rate and a modest velocity, the allowable interface temperatures for austenitic steels in lithium and a Pb-Li alloy are ~ 450 °C and 400 °C, respectively. Stress corrosion problems have not been encountered in alkali metal systems. Limited data by Chopra indicate that the fatigue properties of austenitic steel are unaffected by testing in a relatively pure lithium environment (Chopra and Smith, 1984). Compatibility of the austenitic steels with ceramic breeding materials is not considered to be a serious problem at the allowable operating temperatures. However, significant interaction with beryllium at 600 °C has been reported (Flament et al., 1992).
10
Temperature (°C) 650 600 550 500 450 400
3
350
= r
10 2 O)
ORNL
10 J2 o S2
b 10°
10-
1.0
1.1
ANL A PCA oType 316 Stainless steel D Type 316 CW stainless steel v HT-9 O 9Cr-1 Mo USSR < 16Cr-15Ni-2Mo >13Cr-1 Mo
J
1.2
I
J_
1.3 1.4 1000/7 (K)
J
1.5
I
1.6
1.7
Figure 10-17. Corrosion rate data for austenitic and ferritic steels exposed to flowing Pb-17Li alloy.
SA,PCA,9-15dpa * FFTF.0.5 appm He/dpa - o ORR,18 appm He/dpa ? HF IR,47 appm He/dpa
.£ 2
400 500 600 Irradiation temperature (°C)
700
Figure 10-18. Temperature dependence of swelling of solution annealed (SA) PCA in mixed spectrum and fast reactors.
Irradiation Effects The effects of neutron irradiation on the properties and performance of the austenitic steels present a major constraint to the use of an austenitic steel structure. Of particular significance is the observed difference in the neutron spectrum or the effects of simulated fusion-relevant helium generation rates. Table 10-5 presents a comparison of the He/dpa ratios of several materials exposed to a fission and fusion neutron spectrum. Extensive development of the cold-worked PCA alloy in the fission reactor and early fusion reactor programs indicated substantial improvements in the irradiation damage resistance of PCA compared to the conventional type 316 steel. However, as shown in Fig. 10-18, results from the spectral tailoring experiments (Stoller et al., 1988) indicate that the swelling of austenitic steels is significant even at low damage levels with simulated fusion-relevant He/dpa ratios of ~ 15 He/ dpa. Also, the differences between the PCA and the type 316 steels are small and the cold-worked structure appears to provide only modest improvements.
262
10 Fusion Reactor Materials
Irradiation-induced embrittlement is probably the most critical feasibility issue for all structural alloys. Most alloys exhibit significant increases in yield strength and reductions in ductility after only modest neutron fluences equivalent to a few dpa. This loss of ductility is generally sensitive to the strain rate. Low strain rate (^10~ 4 s~ 1 ) tensile tests provide data relevant to normal loading, while high strain rate tests, e.g., Charpy impact tests, provide data relevant to plasma disruption loading conditions. Data derived from spectrally tailored experiments (to simulate fusion-relevant He/dpa ratios), conducted in the Oak Ridge Research Reactor (Kohyama et al., 1992), are summarized in Fig. 10-19. The yield strengths of both solution-annealed and cold-worked austenitic steels are increased to values of - 8 0 0 MPa at temperatures of 100-400 °C after only 7 dpa. The uniform elongation of solution-annealed materials is drastically reduced at 200-400 °C from - 3 0 % to about 0.3%. The uniform elongation of the cold-worked material is also reduced to < 1 % at r < 4 0 0 ° C ; however, the total
DU
illf|
OSAJ316 • 20%CWJ316 A SAJPCA V15%CWJPCA O 25% CW US-PCA X20%CWUS-316 + 25% CW JPCA
50 - v
•LONGA
o
40
elongation remains at 2 - 3 % at temperatures of 300-400 °C. The effects of irradiation on the fracture toughness of the austenitic steels have been summarized by Odette and Lucas (1992). Most of the data available (Fig. 10-20) are for temperatures of 370-430 °C. Also, only a few tests have been performed at a high strain rate. The effect of irradiation on the high strain rate properties in the critical temperature range of 200-350 °C is not well characterized. The available results show a sharp reduction in fracture toughness after only a few dpa. The results indicate saturation at 10-30 dpa; however, most of the data are for low strain-rate tests at temperatures around 400 °C. The effects of irradiation on weldments are quite limited but generally indicate a lower fracture toughness than for the base metal. The effects of high helium concentrations are not known. Measurement of fatigue properties during irradiation is extremely difficult. Postirradiation fatigue data generally indicate only modest effects of irradiation (Grossbeck and Liu, 1982). The irradiation creep
_x
ORR 7 dpa T'irr ~ — T'test
/ UNIRRADIATED SA '
30
UJ
OC
O u.
z
20
V*-SA • * \
Figure 10-19. Irradiation temperature dependence of uniform elongation of austenitic steels irradiated in ORR.
UNIRRADIATED A 15-25% CW
10 0 100
200
300
400
IRRADIATION TEMPERATURE (°C)
500
600
10.2 Structural Materials 400
10.2.1.2 Ferritic/Martensitic Steel
A Dynamic/Fast/SA(304, 316) A Dynamic/Fast/Weld (304, 308, 316) O Static/Mixed/SA (304, 316, 347, 348) Static/Mixed/Weld (316) • Static/Fast/SA (321, 304, 316) • Static/Fast/Weld (308) Q Static/Fast/CW (316)
300
CO
Q- 200
100
8
10
Figure 10-20. Variation of fracture toughness of austenitic steels with irradiation.
properties of austenitic steels have been investigated by Grossbeck (Grossbeck et al., 1990). Results obtained by the spectral tailoring experiment (Fig. 10-21) indicate higher irradiation creep rates at 60 °C than at higher temperatures.
1.4 1.2
ORR spectral tailoring experiment 7 dpa O
rf
1.0
V
0.8
/ y"
/
0.4 /
0.2
•
.- o
60 °C 200 °C 330 °C 400 °C
100
PCA 25 % Cold worked
/x 200
300
400
500
Effective stress (MPa)
Figure 10-21. Irradiation creep of austenitic steel irradiated in ORR.
The 8-12% Cr martensitic steels have been investigated for use as fuel cladding materials in fast fission reactors and are considered candidates for the first wall/ blanket of a fusion reactor. These alloys are based on HT-9 (a 12% C r - 1 % Mo alloy) and the European alloy MANET (a 12% Cr-0.5% Mo alloy). Modifications currently being evaluated include 8-9% C r - 1 - 2 % Mo alloys and reduced activation versions, viz., 8-9 % C r - 2 % W alloys from which the molybdenum and small amounts of niobium and nickel have been removed. The current trend is toward the 8-9% Cr alloys. The key features of this alloy system relate to swelling resistance, better thermal stress factors than the austenitic steels, better liquid metal corrosion behavior than the austenitic steels, and a substantial database for both baseline and irradiated material properties. The critical design issues for this alloy class relate to more difficult weld characteristics, including postweld heat treatment, irradiation embrittlement, limited operating temperature, and ferromagnetic properties. Fabrication
i
0
263
cold-worked
In order to enhance the properties of these alloys, they are used in the normalized and tempered condition. The properties are quite sensitive to the thermo-mechanical treatment, typically a 950-980 °C anneal for 2 hours, followed by ~1070°C for 0.5 h and finally 750-780 °C for 2 h. Significant variations in properties are obtained by temperature variations of 2040 °C and significant variations in annealing/tempering times. In order to obtain good welds, these alloys typically require some pre-weld heating followed by a controlled post-weld heat-treatment of
264
10 Fusion Reactor Materials
~ 725 °C for 1 h. Modest variations in the post-weld heat treatment produce significant variations in the weld properties. Baseline Properties The baseline physical properties characteristic of this class of alloys are compared with those of other alloy systems in Table 10-5 and Figs. 10-7 to 10-11. As in the case of the austenitic steels, the thermal stress factor and the thermal creep properties are important for comparison of the projected performance capability. Figures 10-12 and 10-13 present a comparison of the calculated thermal stress factor and the thermal creep properties of the martensitic steels with the austenitic steels and vanadium alloys. The relatively low creep strength is the primary operating temperature constraint for this alloy system. Chemical Compatibility The chemical compatibility of the ferritic/martensitic steels with candidate coolants and tritium breeding materials is generally similar to or better than that of the austenitic steels. Compatibility with the helium coolant does not present a design constraint. Corrosion by water is acceptable and stress corrosion is of less concern than for the austenitic steels. As shown in Figs. 10-16 and 10-17, the corrosion resistance of the martensitic steels is superior to that of the austenitic steels in both lithium and the Pb-Li alloy. The allowable interface temperature limits are about 50 °C higher than those for the austenitic steels. The corrosion rates in lithium are affected by the nitrogen concentration in lithium. The effects of the lithium environment on the fatigue properties of HT-9 have also been investigated (Chopra and Smith, 1981). An effect of higher nitrogen content in lithium on the fatigue properties of the
steel has been observed. Compatibility of the martensitic steels with ceramic breeding materials is not considered to be a serious problem at the allowable operating temperatures for the steel. Irradiation Effects The effects of neutron irradiation on the properties and performance of the martensitic steels pose a major concern regarding their use as a first wall/blanket structure. The most critical issue involves irradiation embrittlement, including the helium transmutation effects. Since a high flux, 14 MeV neutron source is not available for testing materials, the effects of fusion-relevant helium production rates must be determined by various simulations. The martensitic steels are highly resistant to irradiation-induced swelling in a fast fission spectrum to damage levels of over 70 dpa at temperatures of 400-650 °C (Gelles and Thomas, 1984). However, significant cavity formation has been observed in alloys irradiated in the High Flux Isotope Reactor (HFIR), particularly in alloys with nickel additions (Maziasz and Klueh, 1989). Neutron reactions with the nickel enhance the helium production rate to levels closer to that of a fusion spectrum. Although the interpretation of these results remains controversial (Klueh, 1992), the enhanced void formation and swelling in alloys containing nickel are generally attributed to the higher helium generation rates. Even with the helium effect, swelling in the martensitic steels is considerably less than that for the austenitic steels. Neutron irradiation has a significant effect on the mechanical properties of the martensitic steels. Low fluence irradiations (~ 7 dpa at 25 °C) produce a significant increase in the yield strength of these steels,
10.2 Structural Materials
as indicated in Fig. 10-22. An additional effect of the helium is also observed. Although the hardening appeared to saturate at low fluences (~ 10 dpa) with low helium concentrations, this saturation was not observed in alloys irradiated in HFIR with higher helium concentrations (Klueh et al., 1986). The effect of irradiation on the ten1400
1
1
1300 1200
(167)
STEEL o o • A A
0
12Cr-1MoVW 12Cr-1MoVW-1Ni (He, appm)
5 10 15 20 25 30 DISPLACEMENT DAMAGE (dpa)
Figure 10-22. Effect of irradiation and helium concentration on the yield stress of martensitic steels.
Steel Heat treatment MANET-I 950-980oC/2h+1075°C/0.5h+750°C/2h Heat No.: 53645
265
sile ductility at low strain rates correlates with the yield strength. The shift in the ductile-to-brittle transition temperature (DBTT) from high strain rate Charpy impact tests after irradiation is the most critical effect for the martensitic steels. The effects of alloy composition, irradiation fluence and temperature, and simulated helium effects have been investigated (Klueh, 1992). Figure 10-23 shows the shift in the DBTT of the MANET alloy as a function of the irradiation temperature. Available data (Klueh and Alexander, 1992) indicate that 12 Cr and 9 Cr alloys show similar behavior, but with lower shifts for the 9 Cr steels. Alloys irradiated in HFIR tend to show a larger shift in the DBTT than the same alloys irradiated in a fast reactor spectrum. Also, alloys with nickel additions to produce higher helium concentrations exhibit significantly larger shifts in the DBTT. Figures 10-24 and 10-25 summarize much of the available data on these effects. The helium effects appear to dominate the shift. Limited data on reduced activation compositions (9 Cr-2 WVTa) exhibit a lower shift in the DBTT at low fluences. It is not clear
Irradiated at Unirr. 300°C 400°C 475°C dpa
S SB
Figure 10-23. Effect of irradiation temperature on the impact properties of MANET steel.
-A—A
-100
100 200 300 Temperature [°C]
400
500
266
10 Fusion Reactor Materials
400
1
• A 350 -- • • A •
300 _-
o
1
HFR (MANET - 5DPA) HFIR SUBSIZE EBR-II SUBSIZE EBR-II FULL SIZE UBR SUBSIZE UBR FULL SIZE
Open Symbols - Low Fluence
I
12Cr- 1MoVW A 40,200 M w ^
(A 40,110 = DPA, appmHe) ( • 26 = 26DPA)
A 40,110
250 --
A A 9,25
X 200 h
1 EQ=
I
1
A10-17 A
150 D
A10-17 • 26 T5M
•
100 -
y. •. .
A 50 -
j|
LxX I
100
fUU
I
350 -
• • •
I
200 300 400 IRRADIATION TEMPERATURE (°C)
i
1
1
500
Figure 10-24. Effect of irradiation temperature and helium on the impact properties of 12Cr martensitic steels.
600
1
9Cr-1MoVNb
HFIR SUBSIZE EBR-II SUBSIZE UBR FULL SIZE
A 40/200
Open Symbols - Low Fluence
300 -
(A 40/35 = DPA, appmHe) ( • 5 = DPA)
-
250 A 40/35 (Oo)
t I C/D H
200 —
150 _
QQ
Q
_ A
100
_t
_
• 50
n
-
Figure 10-25. Effect of irradiation temperature and helium on the impact properties of NCR martensitic steels.
• • 26,13 •5 1
100
1
1
• •
200 300 400 IRRADIATION TEMPERATURE (°C)
• 500
• 600
10.2 Structural Materials
whether this is due to the compositional change or the higher purity alloy with lower nickel and lower helium generation. Effects of simultaneous helium generation on the DBTT shift are critical. The effects of irradiation on fatigue have not been investigated in detail. Results, primarily on the MANET alloy (Klueh, 1992), indicate only modest effects on the post-irradiation fatigue characteristics of this alloy, even with simultaneous helium additions. 10.2.1.3 Vanadium Base Alloys Vanadium alloys are considered an attractive first wall/blanket option because of their higher temperature potential, long radiation lifetime, high heat flux capability and low activation characteristics. Based on previous work in support of the fast breeder fission reactor program, vanadium base alloys with a few percent titanium were shown to exhibit excellent swelling resistance (D. L. Smith et al., 1985). The focus is on solid solution-strengthened alloy compositions with 3 - 7 % Cr and 3 - 5 % Ti. The leading candidate alloy is nominally a V-5Cr-5Ti composition (Loomis and D. L. Smith, 1992). The alloys currently being evaluated are typically in the solution-annealed condition (~1125°C for 1 h); however, further development may find improved treatments. The critical design issues for this alloy system relate to a limited experience base for refractory metals, inert atmosphere requirements for welding, nonmetallic element (O, N, H, C) interactions at elevated temperatures, and irradiation-induced embrittlement similar to the other alloys. The favorable physical and mechanical properties combined with the high swelling resistance suggest that high performance and long lifetime are potentially attainable.
267
Fabrication Resources of this material are adequate and high purity material has been produced (Gupta and Krishnamurthy, 1992). Secondary fabrication processes, including plate, sheet, tubing and wire have been used for several compositions. Since the alloys of most interest exhibit high ductility, these compositions are also the most readily made. The major issue is to avoid environmental contamination during processing since vanadium will react readily with oxygen at an elevated temperature. Although only limited information is available on the welding characteristics of these alloys, they appear to be readily weldable by all fusion processes (D. L. Smith et al., 1985). Since the alloys of interest are simple solid solutions, they do not appear to be sensitive to segregation. All welding must be done in an inert environment. Post-weld heat treatment does not appear to be necessary for thin sections; however, further work is required to determine the weld parameters for thick sections. Baseline Properties The baseline physical and mechanical properties of this alloy system are compared with those of the steels in Table 10-5 and Figs. 10-7 to 10-11. The thermal stress factor and the thermal creep properties of the vanadium alloy system are superior to those of the steels, as shown in Figs. 10-12 and 10-13. The values for a V-5Cr-5Ti alloy are slightly lower than those for the V-15Cr-5Ti alloy shown. If desirable, the strength properties of these alloys may be improved through further development by minor compositional and/or thermo-mechanical treatments. Only limited data have been reported on the fatigue properties. The results in Fig. 10-26 indicate that
268
10 Fusion Reactor Materials 1M
" " l ' I M I I I I I ' ' M ""l r • 550°C (V-15%Cr-5% Ti) a 650°C (V-15 % Cr-5 % Ti) 550°C (20%CW316) 650°C (20%CW316)
c 03
o Average trend curve for 20% C Type 316 stainless steel 0.2 3 i i mini 4 i i i imil 5 i i mini 6 i i mini 7 i i mini 8 10 10 10 10 10 10 Cycles to failure, Nf
Figure 10-26. Fatigue data for V-15Cr-5Ti and Type 316 stainless steel.
the fatigue properties of V-15Cr-5Ti are superior to those of austenitic steel (Liu, 1981). Chemical Compatibility The chemical compatibility of the vanadium alloys is dominated by nonmetallic element (e.g., oxygen) interactions. These alloys are compatible with helium; however, the helium must be highly purified, and effective purification methods must be used to maintain the helium purity. Limited data indicate that certain vanadium alloys are resistant to corrosion by 300 °C pressurized water. The V-15Cr-5Ti alloy appears to be corrosion resistant and resistant to stress corrosion cracking (Diercks and Smith, 1986). Alloys without chromium are much less corrosion resistant. The amount of chromium required for improved corrosion resistance is not well defined. Vanadium is highly corrosion resistant to high purity lithium and probably to the Pb-Li alloy (D. L. Smith et al., 1985). At high temperatures (above 450-500 °C), the liquid metals must be purified to prevent nonmetallic element interactions. The predicted corrosion behavior of vanadium al-
loys in lithium is much lower than that in the steels. At lower temperatures the rate of these reactions will be kinetically controlled such that acceptable corrosion behavior can be obtained with stable surface reaction products, e.g., nitrided surfaces. Further work is required to define acceptable conditions. Irradiation Effects As in the case of the steels, the effects of neutron irradiation on the properties and performance of vanadium-base alloys is a critical issue. The most critical of these effects is the potential for irradiation embrittlement, including the effects of helium. Optimism for acceptable performance is based on the favorable baseline mechanical properties and the possibility of further improvements by optimization of the composition and/or thermo-mechanical treatment. Neutron irradiation to ~ 100 dpa and ion irradiation to > 200 dpa have shown that vanadium alloys containing a few percent of titanium are highly resistant to swelling (Loomis and Smith, 1992). Figure 10-27 shows the swelling as a function of titanium concentration at fluences approaching 100 dpa. Alloys in which helium was preinjected (74-100 appm by tritium decay) showed swelling behavior similar to the same alloys without helium. Alloys with greater than 3 % Ti show improved swelling resistance. Irradiation hardening increases the yield strength and decreases the ductility in most alloys tested (Loomis and Smith, 1992). These effects tend to saturate at about 20 dpa. Figure 10-28 illustrates the uniform elongation of several V-Cr-Ti ternary alloys after irradiation at 420 °C to fluences of up to 84 dpa. The total elongations of all but one of these alloys range
10.2 Structural Materials 8 V-(0-15)Cr-5Ti alloys
g 7 Q.
6
^
5-
600° C
o - 17-21 DPA • - 40-50 DPA
• •
- 73-77 DPA - 84 DPA
• •
—
r-O-i
2 4 6 8 10 12 Chromium concentration (Wt.%)
14
Figure 10-27. Effect of Cr addition on the swelling of V-(0-15Cr)-5Ti alloy after irradiation at 600°C.
from 5 to 15%. The uniform elongations of these alloys are typically about 75 % of the total elongation. The effects of pre-injection of helium on the tensile properties of V-3Ti-lSi, V-20Ti and V-5Ti under a variety of conditions have been reported.
269
(Van Witzenburg and DeVries, 1990). Only modest effects of the helium are observed. As in the case of the martensitic steels, the effects of irrdiation and helium on the high strain rate fracture toughness are the greatest concern. Results obtained at 420 °C after irradiation to 44 dpa indicate significant shifts (-200°C) in the DBTT of V-15Cr-5Ti and V-10Cr-5Ti alloys (Loomis and Smith, 1992). Loomis has also observed that the DBTT in V-Cr-Ti alloys varies considerably with composition, as shown in Fig. 10-29. These data indicate a minimum in the DBTT below - 200 °C for alloys with - 3 - 1 0 % Cr + Ti. For example, the V-5Cr-5Ti exhibits an unirradiated DBTT of less than -200°C. Since the trend with composition of available irradiated data is similar to that of the unirradiated alloys, it is anticipated that alloys with 3-10% Cr + Ti will have DBTTs significantly below room temperature. Experiments are in progress to evalu-
20 VANADIUM ALLOYS FFTF-MOTA Irradiation
18
16 h - V-18Ti 14
- 420 C Irr.
0
100
Figure 10-28. Uniform elongation of vanadium alloys after irradiation to 30-40 dpa.
• , A , I _ 41-46 DPA O, T , D - 28-34 DPA
200 300 400 500 600 700 IRRADIATION/TEST TEMPERATURE (°C)
800
10 Fusion Reactor Materials
o o
24 - 43 dpa eV
Figure 10-43. Maximum allowable graphite surface temperature on the ITER divertor plate as a function of plasma edge temperature.
to be viable for the divertor. The dashed lines indicate how uncertainties in the plasma physics can influence the allowable temperature. Even when physical sputtering is eliminated from the calculation, there is a temperature limit of ~1300°C due to radiation-enhanced sublimation. For the ITER reference case, the surface temperature limit has been set at 1000-1200 °C (Kuroda et al, 1991). This rather modest temperature limits the allowable thickness of the graphite and hence the expected lifetime of the divertor plate. 10.3.2.3 Hydrogen/Tritium Retention During operation of fusion reactors, plasma facing materials will be subjected to bombardment by energetic ions. Tritium ions will be implanted into the surface and can be trapped in the material, leading to potentially high inventories in the plasma chamber. The effect of radiation-induced traps will be to increase the inventory in the material. A high tritium inventory is a safety concern, since this inventory could be released during an accident. Tri-
Hydrogen retention in a material is generally described by a formalism involving hydrogen diffusion in the presence of trapping sites, with a recombination kinetics limited boundary condition. Hydrogen retention results from the trapping of hydrogen at radiation damage or intrinsic defects, and from bulk hydrogen solid solution formation or second phase (i.e., hydride) precipitation. The complete characterization of hydrogen trapping and release must therefore include detailed information on mobility, solubility, and trapping of hydrogen, as well as the hydrogensolid phase diagram. Unfortunately, the database is incomplete. Nevertheless it is important to outline the hydrogen trapping formalism in order to appreciate the available data and the gaps in the database. Hydrogen diffusion in an undamaged lattice at a temperature T is characterized by a diffusivity D =
DoQxp[~ED/(kT)]
(10-1)
where Do is the pre-exponential, Eu is the migration energy, and k is the Boltzmann constant. Likewise, the solubility S of hydrogen in a metal exposed to hydrogen gas at pressure p is given by Sievert's law (10-2) where So is the Sievert's law pre-exponential, and Es is the heat of solution. Hydrogen trapping and release in a material can be described by the diffusion equation 8C _ _ dJ _ 9CT 5J (10-3) dt ~ dx dt
10.3 Plasma Facing Materials
where C is the mobile hydrogen concentration, CT is the trapped hydrogen concentration, / is the bulk hydrogen flux, and G is the implant source term due to hydrogen bombardment from the plasma. Traps are generally characterized by a trap concentration (CT) and a detrapping energy {ET). The detrapping energy is defined as the sum of the migration energy (ED) and the hydrogen-defect binding energy (EB). Further details of the trapping term can be found in other publications (Wilson and Baskes, 1978; Baskes, 1980 a). Second phase precipitation is not included in the present treatment for simplicity, although cases where such a formalism should be included will be discussed when appropriate. The bulk tritium flux J is given by dC dx+
CQ*dT kT1 dx
(10-4)
where hydrogen diffuses from a concentration gradient (the first term) or by a temperature gradient (the second term) driving force. This latter phenomenon is termed the Soret effect, where Q* is the heat of transport. A negative value of Q* means that hydrogen will tend to diffuse to the hot side of a temperature gradient. Values for g* will be given in the discussion of individual materials, when available. Deviations from Fickian diffusion due to the hypothesized interaction of hydrogen with point defects to form mobile complexes (Gorodetsky etal., 1980; Tanabe etal., 1981) are not considered in this study. The surface boundary condition is of the form 0
otherwise
(10-7)
The formalism presented in the previous paragraphs is applicable to a smooth, metallic surface, which can be approached under controlled laboratory conditions. The actual hydrogen trapping and release behavior for a realistic fusion surface with a rough, porous layer of redeposited material is unknown. Present modeling of hydrogen trapping and release varies the molecular recombination rate constant kx to account for surface roughness and contamination. In the Baskes theory for kr, the molecular sticking coefficient a is used as a measure of surface contamination. A clean metallic surface can have a value of a of 0.5 (Boszo et al., 1977), while values of 10~ 3 to 10 ~ 4 are typical for oxide-coated surfaces (Rendulic and Winkler, 1978). Wienhold etal. (1979, 1980) use a surface roughness factor (cr = actual area/geometric area) that multiplies kT in Eq. (10-5) to account for the effects of roughness. Values of a range from 2 for an electropolished surface to >20 for a heavily blistered and sputtered surface (Maeda etal., 1981). In general, surface contamination should reduce the release of hydrogen isotopes at the plasma-side surface, increasing the tritium inventory and permeation problems. Alternatively, surface roughness or interconnected porosity could enhance the release of implanted hydrogen at the plasma-side surface, reduc-
286
10 Fusion Reactor Materials
ing the bulk tritium inventory and permeation rates. A third complication for redeposited material could be a bulk non-interconnected porosity. This porosity would act as an internal sink for hydrogen, especially for materials with low solubility, such as tungsten. The time to steady state permeation would be drastically increased, but the tritium inventory would also increase. Experimental Results of Tritium Retention for Carbon Carbon exhibits a wide variation in trapping behavior, depending on sample temperature. Near room temperature, hydrogen is completely trapped at the end of the range (Scherzer etal., 1976; Langley etal., 1978; Wampler and Magee, 1981; Brice and Doyle, 1981) until saturation is reached. The saturation concentration of ~0.4 D/B ratio is independent of ion energy. Further implantation produces a rapid release of the implanted hydrogen at nearly the implantation rate. The behavior of hydrogen in carbon is often modeled by the 'local mixing model' and its derivatives (Doyle et al., 1980, 1981; Brice and Doyle, 1981; Brice et al., 1982). At higher temperatures, the trapped deuterium concentration in the near surface region decreases and rapid evolution of the implanted hydrogen is observed (Erents, 1975; Erents etal., 1976; Braganza etal., 1978). Thermal desorption of the hydrogen trapped in the near surface region is generally observed in the range of 900 to 1200 K. This phenomenological behavior is observed for virtually all carbon morphologies that have been studied. However, the assessment of the carbon database offers an intriguing dilemma. Direct measurements of diffusivity by means of outgassing after energetic triton recoil injection using nuclear
reactions such as 6 Li(n, a) 3 H, etc. (Walter etal., 1973, Rohrig etal., 1976; Causey et al., 1979) indicate that the migration energy for hydrogen in carbon is ~ 2 . 8 4.6 eV. On the other hand, implantation measurements by Hucks et al. (1980) with atomic hydrogen, and by Sone and MeCracken (1982) using low energy ion beams, indicate that hydrogen may be highly mobile in carbon at room temperature in the absence of radiation damage. It was also observed that there is significantly less retention in tokamak-exposed carbon probes if they are annealed to remove lattice damage prior to exposure. Using data from Hucks etal. (1980), Erents (1975), Erents etal. (1976) and Langley etal. (1978), a migration energy for hydrogen in pyrolytic carbon has been calculated to be ~0.5eV (and a detrapping energy from lattice damage of ~2.5 eV). Whether this order of magnitude difference in migration energy can be attributed to differences in carbon morphology, measurement temperature range, or to some fundamental differences in measurement techniques and interpretation is not known. However, the distinction between an immobile hydrogen atom (ED = 4 eV), and a hydrogen atom that is mobile at room temperature (ED = 0.5 eV) but is readily trapped at radiation damage (i?T = 2.5eV) is more than semantics. In the former case, tritium is confined to the implantation range. The tritium inventory in the carbon will be low, and tritium permeation through the carbon will be zero in a fusion application. In the case of mobile but readily trapped hydrogen, tritium inventory can be much higher since injected tritium can diffuse into the bulk and can become trapped at bulk neutron damage sites (Doyle, 1981); steady-state tritium permeation can also be large, since tritium can readily diffuse through carbon at fusion operating tern-
10.3 Plasma Facing Materials
peratures when the neutron damage traps are saturated. Recent experiments (Causey et al., 1986) performed in the Tritium Plasma Experiment (TPX) at Sandia National Laboratories in Livermore have shown there to be high energy traps for tritium naturally occurring in graphite. The density of these naturally occurring traps is approximately 20 appm. Saturation of these traps with tritium would represent an inventory of approximately 100 grams. While this may seem a large amount of tritium, it may be small compared to the amount that will be retained if neutron damage produces additional traps and these traps are saturated. It is not known what the trap density will be in neutron-irradiated graphite. Earlier studies (Wilson, 1984) with metals have shown the trap density due to radiation damage to be as high as 1 at.%. At this level of trapping, saturation with tritium in the large amount of graphite to be used in a fusion reactor would result in unacceptable tritium inventories. Preliminary experiments (Causey and Doyle, 1989) were performed at Sandia National Laboratories in Livermore and Albuquerque using charged particles to produce the damage. In this study samples of POCO AXF-5Q were bombarded with 1017 C + ions/cm2 at an energy of 6 MeV. The damage was calculated to vary between 0.1 dpa near the surface to a high of around 10 dpa near the end of the particle range. After the irradiation, the samples were exposed to 1 atmosphere of deuterium at a temperature of 1200 °C for 3 h. The samples were then analyzed for deuterium content using nuclear reaction profiling. The tests showed the deuterium concentration to be more or less constant at a level of around 0.06 at.%. Because the DIFFUSE (Baskes, 1983) computer code has been used to determine the temperature, time, and gas pressure
287
necessary to achieve saturation of the traps, this value of 0.06% is also believed to represent the trap density. While this represents a significant increase over the 20 appm trap density found in non-irradiated graphite, it is also much lower than the feared 1 at.%. It cannot be ruled out at this time that neutron damage may result in greater trapping than that seen for the charged particle damaged sample. 10.3.2.4 Neutron Irradiation Effects The properties of graphite are generally degraded by neutron irradiation. First neutron irradiation will result in dimensional changes as a result of densification and swelling. The mechanical strength will also be reduced, and the thermal conductivity will be reduced. Bulk nuclear graphites have been developed for fission reactor applications. The overriding consideration in radiation damage to graphite is the estimated lifetime due to swelling - that is to say, the point at which its mechanical integrity is destroyed due to internal fracture generated by radiation distortion. A typical set of design curves for an isotropic nuclear graphite is given in Fig. 10-44 (D. L. Smith et al., 1982). The classic definition of lifetime is that point in time (fluence) when the graphite distortion returns to its original volume. In actual fact, the mechanical strength persists for some time after this point, as indicated in Fig. 10-45 (D. L. Smith et al., 1982). The three curves presented apply to currently used nuclear graphites, not the high-strain materials such as GraphNOL. Due to the complexities of radiation damage in graphite having primarily to do with the formation of interstitials and their degrees of aggregation, lifetime is a complex function of temperature. There is a
288
10 Fusion Reactor Materials
1
2
3
FLUENCE 0 x 1022 (E > 50 keV)
Figure 10-44. Typical graphite design curves for a nuclear reactor. Graphite is assumed to be isotropic.
maximum in the life at irradiation temperatures in the range 500-550 °C, and minimum at around 1000 °C. Data above 1300°C are scanty and unreliable, but there is reason to believe that the life at 1200-1300 °C is a secondary maximum. Due to vacancy mobilities, a third relative-
120
-
IOO
80
I
ly resistant area should exist at radiation temperatures in excess of 2000 °C. The generally expected changes in thermophysical and mechanical properties that will result from irradiation are summarized in Table 10-12. Specific data are discussed below. The thermal conductivity of irradiated H-451 graphite is shown in Fig. 10-46 (D. L. Smith et al., 1982). The degradation in conductivity occurs over the fluence range 0.1-1.0 x 1025 n/m 2 , with apparent saturation at the higher fluence. Thermal expansivity is again estimated from the behavior of conventional graphites and the data available on AXF and similar materials at 715°C. The latter data are given in Fig. 10-47 (D. L. Smith et al., 1982). More recently, carbon-carbon composites (CCCs) have been considered for fusion applications. Irradiation effects are less well understood in CCCs compared with bulk nuclear graphites, and irradiation experiments are at a relatively early stage. CCCs generally exhibit high thermal conductivity. However, it is expected that neutron irradiation will reduce the thermal conductivity. Recent experiments performed in the U.S. and Japan have shown
I
GRADE ASR-IR O POSITION I • RADIAL POSITION 2 —I
co 60 e> z * 40 20 QC CD
10 15 FLUENCE, EDN (n/m 2 )
20 (xlO25)
Figure 10-45. Strength is retained past end of life for graphite grade ASR.
10.3 Plasma Facing Materials
Table 10-12. Extrapolated property values for irradiated GraphNOL N3M.
p CO
E 0.25, -^0.20
Fluence ( " ^ ^ ^ ( E >0.18
fo.15 " g 0.10-
H-451 Axiai
^T ^^~—i
o _
—-
-—
—-=
10.0
j 0.05
i
i
i
I
I
i
i
i
I
I
10.20 to
i
~~
"o!P^
^
- ^—^IliS—-
-=0 05 500
10.0 i
i
' i
I
I
i
600 700 800 900 1000 1100 1200 Irradiation temperature (°C)
Figure 10-46. Thermal conductivity of H-451 nearisotropic graphite at irradiation temperature as a function of irradiation temperature. Conversion: 418 x [cal/(cm s °C)] = [W/(m K)].
large reductions in thermal conductivity (Tanabe, 1991; Burchell, 1991). CCCs having a thermal conductivity above 300 W/ mK at room temperature were irradiated to a level of 0.1 dpa in Japanese reactors. The resulting thermal conductivity was measured to be only in the range 10-50 W/ mK. In the U.S. a 3D CCC (FMI222) was
< o
Saturates at about 30 W/(m K)
Thermal expansion
Rapid rise and fall within first third of life
Maximum at perhaps 50% over initial value. Saturates at about 75% of initial value
Tensile strength
Linear fall off with fluence
About 50% of initial value at end-of-life
Strain-tofailure
Gradual fall off saturating at about half-life
Saturates at about 40% of initial value
Moduli Gradual in(Young's and crease, satushear) rating at about half-life
Saturates at 2.5-3 times initial value
i
= = = = = _ — — — — _
|I0.10-
I
Thermal Monatomic exconductivity ponential decrease, saturating at about half-life
Radial"
.^0.15 1
8
289
Radiation creep
No information; Assume behaves as other nuclear appears to be relatively insen- graphites sitive to graphite grade
irradiated to a level of 3 dpa at 600 °C in the HFIR test reactor. The thermal conductivity was measured at a level of 4060 W/mK compared with an unirradiated
10 20 30 40 (x1021) 2 Fluence (neutrons cm" , E > 50 keV)
Figure 10-47. The average coefficients of thermal expansion from 20 to 600 °C versus fluence accumulated at 715 °C for graphite grades AXF, AXF-5QBG-3, H-395, and P-03.
290
10 Fusion Reactor Materials
value of ~180W/mK at room temperature (Burchell, 1991). Hence the high thermal conductivities for CCCs will very likely be degraded rapidly in a fusion reactor. 10.3.3 Beryllium The primary issues associated with the use of beryllium in fusion devices are: (1) physical sputtering and erosion by plasma particles, (2) swelling due to the generation of helium during neutron irradiation, and (3) neutron degradation of mechanical properties. 10.3.3.1 Beryllium Base Properties Beryllium has a hexagonal-close-packed (h.c.p.) crystal structure throughout its temperature range of useful properties. The room-temperature density is 1.85 g/ cm3 and the melting temperature 1284°C. Poisson's ratio at room temperature depends on the anisotropy of the product material. Hot-pressed block has a value of 0.085, and values for sheet range from 0.070 to 0.093 (Brown and King, 1974). Several of the key physical properties of beryllium are given in Table 10-13 (Baker et al., 1980). All of these properties should be considered as nominal for beryllium, as most vary considerably from product to product, being sensitive to purity, density, and preferred orientation. The product-to-product variations in the properties of beryllium show up most dramatically in the mechanical properties reported for this material. The strength and ductility properties are a strong function of the metal purity, porosity, shape and distribution of inclusions, grain size, and preferred orientation or texture. In some cases these effects can be used to tailor the properties to the application; how-
Table 10-13. Properties of beryllium. Property Atomic number Atomic weight Density (g/cm3) Crystal structure Melting temperature (°C) Boiling temperature (°C) Vapor pressure (Pa (°C))
Value 4 9.01 1.85 h.c.p. 1284 2970 10~4(850) 10" 2 (993) 10° (1192) 1Wjn^ 11(\ jw*^ i\\ j j l 111
Heat of fusion (J/g) Heat of vaporization (J/g) Heat capacity (J/g °C) 500 °C 1000°C 1500°C Coefficient of thermal expansion (10" 6 /°Q 2 5 - 100°C 2 5 - 500 °C 25-1000°C Thermal conductivity (W/(m K)) 50 °C 300 °C 600 °C Electrical resistivity (|iQ cm) 400 °C
1083 24,790 2.25 2.92 3.59(1)
11.6 15.9 18.4 150 125 96 15
ever, in evaluating the literature to judge a new application, they introduce a significant level of confusion. The modulus of elasticity (Young's modulus) is 290 GPa at room temperature. Ultrasonic measurement of the modulus yields a value of 314 GPa. The range of properties measured in room temperature tensile tests is shown in Table 10-14 for a number of different beryllium product forms. In most cases these are manufacturers guaranteed minimum properties. The temperature dependence of the tensile properties of a typical high-strength grade of hot-pressed beryllium, HP21, is shown in Figs. 10-48 and 10-49 (Borch,
291
10.3 Plasma Facing Materials
Table 10-14. Typical room temperature tensile properties of various grades of berylliuma. Material purity and grade
Mechanical properties Strength (MPa) Total i
Yield Ultimate tensile
elongation (%)
Normal purity (1) Vacuum hot-pressed block Standard Structural Structural, high ductility Structural, instrument Instrument Thermal Optical High purity
207 241 207 242 310 166 172 173
276 310 290 311 345-414 255 241 242
1.0 1.5 3.0 2.0 1.0 1.5 2.0 1.0
(2) Wrought products Plate b Sheet Extrusions
207/311 345 207
414/449 483 483
3.0 10.0 5.0
—
20-140
0.1
276
345
1.5
290
414
5.5
241 414
345 550
3.0 4.0
10.3.3.2 Physical Sputtering The physical sputtering characteristics of beryllium are similar to those of graphite and are illustrated in Fig. 10-50 (Roth etal., 1979; Fetz and Oechsner, 1963; Laegreid and Wehner, 1961; Borders et al., 1978). The self-sputtering of Be->Be remains below unity, and therefore it can be used under all plasma energy conditions without reaching a point of runaway selfsputtering. Be does not exhibit the chemically enhanced sputtering observed with graphite. 10.3.3.3 Tritium Retention and Release Tritium particles will bombard the surface of beryllium during operation and will be injected into the material. In addition,
(3) Plasma sprayed0 As-deposited (87-89) d Deposited and sintered (90-92) Deposited and sintered (99)
Ultimate tensile strength 350 -
High purity (isopressed) Standard High strength0
300
a
From Fullerton-Batten and Hawk (1977); Dunmur et al. (1977); and Dunmur (1979) (refs. 8-10); b lower values: 11.4-15.2 mm, higher values: 6.35-11.4 mm thickness; c Typical, not guaranteed, mechanical property values; d density, in percent.
\ \
250
\ CO
Lower yield strength 200
1979). These data suggest an upper temperature limit for application, based on loss of strength, no higher than 575— 600 °C. If ductility were not affected by irradiation, service near this limit would take advantage of the best ductility properties.
E = 0.0385 / min — Transverse — — Longitudinal 150 300
i
400
i
i
500 600 700 Temperature K
800
900
Figure 10-48. Temperature dependence of the strength of a typical high-strength, hot-pressed beryllium, HP21.
292
10 Fusion Reactor Materials 1
10
i
1
28 —
Be e = 0.0385 / min
24
20 —
.i
—
-
/
Transverse-
-
1 v
16 /
T8
/
c
•§ 12
I
/
8 —
^-Longitudinal
-
/ /
300
/
1 i 500 700 Temperature K
Laegreid DSPUT (D) DSPUT (BeO) 4 T, He) IPP
i
900
Figure 10-49. Temperature dependence of the total elongation of HP21 beryllium.
tritium is produced from a nuclear reaction with beryllium. Thus a large inventory of tritium can potentially build up in beryllium in a similar manner to graphite. The modeling of tritium trapping is outlined in the previous section on graphite. Recent tests indicate that relatively high temperatures must by achieved before tritium is released from beryllium (Billone, 1991). Tritium release has been investigated for 100% dense and 80% dense beryllium following irradiation. Release was measured as a function of temperature in a post-irradiation test. In the case of 100% dense beryllium, a temperature of 600 °C was needed to release 90% of the tritium. Release in the 80% dense sample occurred at somewhat lower temperatures, with 85% of the tritium being released at 500 °C. The release data as a function of temperature are shown in Fig. 10-51.
i mnl 10 2 103 Ion energy (eV)
10 4
Figure 10-50. Normal incidence sputtering yields for Be targets. These data are similar to results obtained for a BeO target and may be representative of the oxide surface.
0.001
. 200
400
.
.
600 800 Time (h)
.
. 1000
. 1200
Figure 10-51. Tritium release from 80%-dense coldisostatically-pressed/sintered Be with 55 appm tritium, 780 appm He, and 0.9 wt.% BeO.
10.3 Plasma Facing Materials
10.3.3.4 Neutron Irradiation Effects Older reviews of irradiation effects in beryllium (Bush, 1965; Kangilaski, 1971) based on relatively low fluence experimental results conclude that the metal is intrinsically resistant to purely displacement damage events. They conclude that observed effects of irradiation at temperatures above the cryogenic range are due primarily to transmutation helium, rather than to point defects or defect clusters. Several experiments confirm the resistance of beryllium to defect cluster formation under irradiation. Measurements of resistivity (Blewitt, 1958; Williams et al, 1972) and thermal conductivity in material irradiated at 20 or 77 K showed that recovery was complete, with no residual damage on annealing to 270 K. While changes at low temperature were large for both properties, the complete recovery suggests that neither of these point-defect-sensitive properties would be much affected by higher temperature irradiation. More direct evidence was developed by Carpenter and Fleck (1977) using high-energy electron bombardment and direct TEM observation of the damage microstructure. They found that, although visible damage resulted in bombardments near room temperature, no visible damage could be developed for bombardments producing 15dpaat 300 °C. Mishima etal. (1977) showed that for irradiation at 550 °C in a fission reactor, where both displacement damage and helium are produced, visible damage does result. They observed that cavities, presumably equilibrium helium bubbles, were produced by the irradiation. Considerable helium mobility at this temperature can be inferred from the preferred location of bubbles on grain boundaries and at dislocations.
293
The thermal conductivity of beryllium at elevated temperatures should be little affected by irradiation. As mentioned earlier, irradiation at ~ 77 K resulted in a reduction of the thermal conductivity from 340 to 70W/mK (Williams etal., 1972); however, full recovery was achieved by annealing to 270 K. The results further indicated that the thermal conductivity was mainly electronic (Williams etal., 1972). As a result, the coarse distribution of bubbles for irradiation at projected fusion service conditions will have little or no effect on the conductivity at those temperatures. Helium accumulation in neutron-irradiated beryllium results in swelling and degradation of the mechanical properties, especially loss of ductility. Observation of the preferred location of the bubbles suggests that, when helium is created, it is probably located first in interstitial positions and moves with high mobility to trapping sites such as dislocations and grain boundaries. Bubbles formed when sufficient helium has collected at these sites are relatively immobile, and helium release does not occur on thermal annealing until much higher temperatures are reached. At temperatures where beryllium can be used, the average migration distance for helium precipitation will be of the order of a few hundred angstroms. It is therefore unlikely that the as-fabricated porosity could collect any significant fraction of the helium, and is thus not a viable mechanism for suppressing swelling or preventing grain boundary embrittlement. Swelling correlations have been formulated to predict helium driver swelling in fusion reactors (D. L. Smith et al., 1991 a). The data were taken from previous fission reactor irradiations for temperatures iMev m.5xio21nvt>iMev -
A
50 —
o A ^
o
40 —
A u
/
>
30 /
_
AX^MINIMUM ELONGATION rA \ VALUES OF UNIRRADIATED \BERYLLIUM —
/
20 —
3.8xio21nvt>iMev
\
/
o
O2.4xl020-
\
A
' ^
/ A
\A
•
0
ro—• 200
60
I
1 stress
40
*
—a—Ln.
400 600 TEMPERATURE, °C
I
I
I
Yield stress,irradiated s
50 a. |
Figure 10-53. The influence of irradiation and tensile testing temperature on the elongation of beryllium.
/
10
< X
o
* S
30 JO
295
Yield
800
I (a)
stress,unirradiated
\M N V ^v
Elongation, unirradiated
c 30 %
A
g> 20 Ay
20 CD
10
Elongation,irradiated a^^*^&k I I I I f J 100
70
200 300 400 500 Temperature (°C)
1
60
1
1
1
i
Figure 10-54. Effect of irradiation of 350 °C on the tensile properties of beryllium for fast fluences of (a) 2 x 1020 n/cm2, (b) 6 x 1020 n/cm2, and (c) 8 x 1020 n/cm2.
10 0
600
T
I (b)
Yielc stress, irradiated
v
i
l
l
Yield stress,irradiated
I (c)
Yield stress,unirradiated 50
30
boo b
40
\
-
^ A
20
\
Yield stress, unirradiated
3^
Elongation, unirradiated
30
c 20 iS
A
N.
00: Be(80%TD, t)
\
^ ^ Li 2 0 (b) Li2ZrO3 (b)
\
Li4SiO 4 (br *- -~ 10:
0
LiA102 ( b T ^ ^ ^ ^—-^ - _
"" '" ""
100 200 300 400 500 600 700 800 900 T/C
Figure 10-60. Bending (b) strength of 80%-dense, 10 |im grain size breeder materials.
1000
& 10 E ?o 1 £
0.1
/
oBe n BE(80 %) APCA oHT9 + Li2O x LiA10 2 • Li4SiO4
/
/
/
/
'
'
/ /
(0
.E as
0.01 /
£_ 0.001 / O
s
0.0001
(a)
10~ 5
1000
-5T 100
10
t (0
o
1
2
0.1
Be BE(80 %) Li2O LiA10 2 Li2ZrO3 Li 4 Si0 4
\ \
/
/
1 MeV). However, the primary driving force for the swelling is helium production. Most of the He is produced from the (n, 2n) reaction which has a threshold energy of 2.7 MeV and a cross-section which peaks at ~0.66 b at 3.25 MeV. Thus considering the vastly different fast flux spec-
tra for thermal and fast fission reactors and fusion blankets, it is insufficient to develop correlations versus fast fluence. The effective cross-section is quite different for test and design applications. It is more useful to have a swelling correlation as a function of He content. The database for swelling of hot-pressed and extruded Be is reviewed by Nardi (1991). Figure 10-71 shows the swelling data versus the helium content for hotpressed Be, for which the relationship between helium retention and fast fluence has been measured. These data sets include Be irradiated in ATR at T&75°C (Billone et al. 1991), ATR-irradiated Be subjected to one hour post-irradiation anneals (ATR-A) at 200 < T< 500 °C (Billone et al. 1991), BR2-irradiated Be at 7 ^ 4 5 °C (Sannen and De Raedt, 1991), and Be irradiated in EBR-II at 4 2 7 < r < 4 8 7 ° C (Billone et al., 1991). All samples were 100% dense with similar micro structures (~25 mm grain diameter) and oxygen impurity contents (1.5-2.0 wt.% BeO).
322
10 Fusion Reactor Materials
oBR2(45°C) • ATR(75°C) AATR-A(200 X ) oATR-A(300 ' O + ATR-A(400 °C)
V
K B R 2 - A ( 4 0 0 ° C) H V
_
x
T
¥ 4
• EBR-IK427" C) • EBR-11(457° C) A EBR-11(467° C) • EBR-11(477° C) o EBR-11(482° C) a EBR-11(487° C) HATR-A(500
•o
vBR2-A(600 C) • BR2-A(800 C) X
o
a
3 X
o
o ° o
• •
o
OO
° ° u Z
I* ° 0
5
Figure 10-71. Swelling data for 100%-dense hot-pressed beryllium. 10
15 20 Ga(1000 appm)
The correlation developed to describe He-induced swelling in Be is AF/F 0 = 0.115(Ga/103) • [ l + ( 3 . 0 x l 0 - 3 ) ( G a / 1 0 3 ) 0 5 T15 •exp(-3940/T)] (10-14) where AV/V0 is volumetric swelling in %, Ga is He content in appm (
>o OR) c r 3 "o^Q#
0 CD
o°
-.5
°
o o
' oo
°o
a
o
-1
o T < 100° C
Figure 10-72. Difference between Be swelling correlation and data as a function of He content.
>200° c 1.5
D
10
15 20 , 1000 appm
work needs to be done to extend the validity of the correlation to higher temperatures, as well as to higher (p > 0) as-fabricated porosity volume fractions. As successful as the correlation is in matching most of the Be swelling database, it does not include fabrication variables (e.g., grain size, porosity and overall size), a coupled He release model, and grainboundary versus grain matrix bubble distributions, such as are found to be important in modeling swelling in fission reactor fuels. Models which have been applied to describe noble-gas-induced swelling in ce-
-1.5 100 200 300 400 500 600 700 800 900 T, °C
Figure 10-73. Difference between Be swelling correlation and data as a function of temperature.
25
30
ramie and metallic fission reactor fuels should be adapted to Be and the fusion blanket ceramics. 10.4.7 Summary of Solid Breeder/Beryllium Blanket Materials
The database, materials' design correlations, and models for the performance of breeder ceramics and beryllium have been discussed. For some of the properties, the database is adequate for developing correlations to describe the performance of the material throughout the anticipated range of fabricational and operational parameters. In a number of cases, data are available for only one temperature or only one porosity value, and extrapolation is necessary. In these cases, trends in the behavior of similar materials were used to develop 'zeroth' order approximations, and the generation of additional data is suggested. In other cases, particularly with regard to tritium and helium behavior, the phenomena are too complicated to be described by a single correlation. For these phenomena, the database is reviewed along with the status of the modeling effort. Based on this critical review, the following areas are highlighted for additional study.
324
10 Fusion Reactor Materials
10.4.7.1 Thermal Performance The thermal properties of unirradiated ceramics and Be are well characterized in terms of temperature and porosity dependence. There is no strong evidence that the thermal properties are degraded with irradiation, beyond the increased porosity effects of helium-induced swelling. Thermal conductivity data and correlations/models are available for both sintered products and pebble beds. In the case of a high conductivity material such as Be, pebble-bed data are desirable because the models presented and their validation pertain to low conductivity ceramics. The solid/solid conductance term in these models may not adequately describe the behavior of a Be bed. The same is true for the model which describes interfacial heat transfer. It has been developed and validated primarily for ceramic/metal interfaces, rather than metal/metal interfaces. For both the sintered product and the pebble-bed, uncertainties in thermal transport exist because of the feedback of deformation and possible cracking on the thermal performance. 10.4.7.2 Tritium Performance Sophisticated models have been developed to describe tritium retention in, and release from, ceramic breeder materials. Also, there is a large database from laboratory tests and in-reactor purge-flow tests with on-line tritium monitoring. The database for Li2O is the most extensive. However, most of the tests are relatively shortterm in duration, resulting in low burnup and low fast neutron exposure. There are unresolved issues, particularly with regard to the ternary ceramics, as to the burnup effects on tritium performance. These mainly concern changes in stoichiometry and chemical form with lithium burnup. The BEATRIX-II test results will extend
the database for Li2O up to about 5 at.% burn up. It is important to demonstrate good tritium performance for burnups as high as 20-30%. While the database for tritium performance of lithium-based ceramics is quite large, there have been problems in interpreting these data and in separating out systems' effects and instrumentation response from ceramic breeder performance for the on-line transient release data. The most direct way of validating the models and codes is to compare predictions to end-of-test tritium inventory data. For Li 2 O, there are only four such data points available. The BEATRIX-II test results will more than double that number. Fewer direct inventory data are available for the ternary ceramics. Thus there is still some question as to whether enough data are available and enough validation of models has been done to answer two important design questions with confidence. The first question concerns the minimum local temperature that the breeder can operate at without building up excessive inventory. No inventory data are available below a minimum local temperature of 425 °C. The second question concerns the amount of purge protium needed to ensure a low tritium inventory. Typically 100:1 protium/ tritium ratios are recommended for design from the inventory perspective. However, the higher that this ratio is, the more difficult it is to separate out the tritium from the purge for reprocessing purposes. Although the protium: tritium ratio has been varied as an experimental parameter over the range of ~ 0 to 1000, the results are ambiguous. The Be multiplier material will also generate tritium. Although the generation rate in the Be is about 1/100 of that in the breeder, kilograms of tritium will be generated in the Be over the life of the fusion
10.4 Blanket Materials
reactor. Thus tritium retention during normal operation and tritium release during overheating transients are still issues for the Be. There have been no on-line experiments with Be to monitor tritium release with changing operating conditions. The limited post-irradiation database available suggests that tritium is essentially "trapped" in Be below a temperature To (300-600 °C for dense Be depending on the He content). Between To and an upper temperature Th (500-900 °C for dense Be, again depending on the He content), some of the tritium is released. At and above Th, the tritium is released in a burst mode after several hours. However, the same tests performed on porous Be (about 80% dense) indicate low-temperature, in-reactor release, a gradual release with increasing temperature, and no significant burst release. These results suggest that as-fabricated porosity and porosity created by Heinduced swelling have a significant effect on enhancing release and reducing the inventory. Other parameters which may be important are BeO content and distribution, grain size, and overall sample size. While separate-effects tests have been unsuccessful in generating rate constants (e.g., diffusivity) and solubility, enough data are available for a preliminary validation of models (e.g., those used to determine gas release and swelling for fission reactor fuels) which contain mechanisms for porosity and swelling effects on gas release. 10.4.7.3 Mechanical Performance While the materials presented in this section are not structural materials, their stress, strain, deformation, and fracture behavior may influence the thermal performance of the blanket and blanket lifetime. Below is a rather specific list of me-
325
chanical properties which should be determined for breeder ceramics and Be in order to allow an assessment of the impact of these materials on thermal performance and lifetime. The fracture toughness of Li 4 SiO 4 has been determined for unirradiated material over a wide range of grain sizes and porosities at room temperature. High temperature data are needed for this material. The fracture toughnesses of the other three materials are needed as a function of porosity, grain size and temperature. Bending strength is another useful parameter. This needs to be determined for Li2O as a function of porosity, grain size and temperature. More bending-strength data are available for the ternaries. However, the database needs to be broadened. Both fracture toughness and bending strength are responses of the material to applied mechanical loads. The data can be used to predict the response of the material to thermal loads. However, there is a great deal of uncertainty in this calculation. Usually it gives a lower limit to the temperature gradient which will cause fracture. It would be very useful to measure the thermal shock resistance of the ceramics directly in tests with temperature gradients. The database for out-of-reactor thermal creep is reasonably well developed at high temperatures (>700°C) for all of the ceramics, except Li 2 ZrO 3 . However, most designs call for a significant fraction of the breeder to operate well below this temperature. Extrapolation of the creep correlations to in-reactor temperatures below 700 °C is complicated by the unknown effects of radiation on the deformation rates of these materials and by the changing of thermal creep mechanisms (e.g., matrix creep versus grain boundary sliding). Lower temperature data are needed under inreactor conditions in order to predict the
326
10 Fusion Reactor Materials
mechanical response of the breeder/cladding system after contact has been established. The more the porous ceramic creeps and hot-presses in response to load, the less the stress will be on the cladding. The swelling of the breeder is important from two perspectives. Swelling is a driving force for breeder/cladding contact and stresses on the cladding. Also, differential swelling within the breeder can cause internal stresses which may crack the breeder. Results from the FUBR-1A experiment have provided the primary estimate of breeder swelling. For the two materials which exhibited the highest swelling rates (i.e., Li2O and Li 4 Si0 4 ) more data are needed in the temperature range of 400600 °C under controlled moisture conditions. In FUBR-1A, the Li2O actually experienced grain growth and sintering at 500 °C after long exposure. Such behavior is often associated with the formation of LiOH(T) due to possibly high moisture levels in the closed FUBR-1A capsules. Many experimental results are reported in the literature on the mechanical properties of unirradiated and irradiated Be. The mechanical properties, particularly the fracture and ductility properties, are dependent on the method of fabrication, which affects the resulting oxygen impurity content and the grain size. Initial porosity and generated He also have a strong influence on the mechanical properties. It is important to first summarize the Be properties for anticipated design fabrication methods and then to recommend additional tests to determine the ductility, fracture strength, and yield strength as functions of temperature, fabricated porosity, and He content. This work is in progress within the U.S. ITER R & D program. Helium-induced swelling in beryllium is important to its mechanical performance,
as well as its tritium release characteristics. Recent progress has been made in determining the He content of irradiated Be and in quantifying swelling as a function of annealing temperature in post-irradiation tests. However, it is also important to conduct in-reactor tests with simultaneous heat and helium generation. Also, more data and better modeling are needed for quantifying the effects of fabrication variables (e.g., grain size and porosity), impurity concentrations (e.g., BeO), and overall dimensions on swelling. 10.4.7.4 Chemical Stability/Compatibility The chemical stability and compatibility of the breeder ceramics are important for establishing upper temperature limits for the bulk of the breeder, and for breeder interfaces with other materials. A related performance parameter is Li mass transport. Based on calculations and data, it appears that mass transport may be an issue for Li2O and for Li4SiO4. The baseline data for Li2O comes from two data sources with consistent results. Basically, the data show that if the only source of oxygen in the system comes from Li burnup, then mass transport can be minimized by proper control of the purge flow system. However, these findings need to be substantiated on a component-scale level with prototypical temperature gradients. For Li4SiO4 material, baseline data on mass transport are still needed. With regard to compatibility, the key issues appear to be Li2O/Be compatibility both in- and out-of-reactor for the anticipated temperature range of 400 to 700 °C, and Be/steam reactions for overheating events. The thermodynamic driving force for these chemical reactions is high, but there is some question about the kinetics of the reactions. The reaction between Li2O
10.4 Blanket Materials
and stainless steel, while apparently faster than for the other ceramics, is still reasonably slow at anticipated design interface temperatures. However, there are some recent results which suggest that for relatively pure Li2O the reaction is even slower than previously thought. Additional data are desirable to resolve this point. While the interaction between the other breeders and stainless steel and Be appears to be reasonably slow at the anticipated operating temperatures, in-reactor compatibility data are desirable. 10.4.8 Liquid Metal Coolants In a recent review, Malang et al. (1991) discussed critical issues for liquid metal breeding blankets. In order to perform design analyses for fusion blankets with either breeding or non-breeding liquid metal coolants, it is necessary to have a property database for physical, thermal, hydraulic, electromagnetic, chemical stability/compatibility, and tritium generation/transport. Included in the reports of the design studies summarized in Table 10-17 are assessments of the metallic coolant databases with emphasis on Li and Li 17 Pb 83 . In addition, Maroni etal. (1973) have critically reviewed the properties of liquid Li. Metallic coolant properties are also summarized in a variety of handbooks and texts. However, as with the solid breeder and Be properties, there are a number of mistakes which propagate from the original reference to the handbook to the design report. For example, electrical resistivity is properly given in units of Qm. Sometimes it is incorrectly listed as Q/cm, which, when converted to SI units, is then in error by a factor of 104. Thus there is a need to review critically the properties' database for liquid metal coolants in as much detail as was done for solid-breeder/
327
Be materials. This was done partially for Li 17 Pb 83 at the ITER Specialists' Meeting on Blanket Materials' Database. In developing the material which follows, a number of sources, in addition to the ones mentioned above, were used. These include: the Materials Handbook for Fusion Energy Systems (Davis, 1980— 1986), the Reactor Handbook (Vol. 1, Materials) edited by Tipton (1960), Argonne National Laboratory technical reports, the Purdue University Thermophysical Properties of Matter by Touloukian et al. (1970), Vapor Pressure of the Elements by Nesmeyanov (1963), and open literature publications. Attempts were made to resolve discrepancies within each source and among different sources. However, it is important that this exercise be repeated with wider peer review. Table 10-27 lists the basic physical properties for liquid metal coolants. Figures 1074 to 10-80 show the variations of density, specific heat, thermal conductivity, viscosity, vapor pressure, and electrical resistivity and temperature. These graphs should facilitate a critical review of properties' correlations, as well as guiding the user in the proper application of correlations. They are based on correlations to the data and are shown only within the database. In comparing liquid metal coolants for design applications, there are several groups of parameters which are of interest. For static liquid metals, the thermal conductivity (k) is important. Figure 10-76 shows a comparison of the thermal conductivity of the coolants of interest. For flowing coolants, the product of the density (Q) and the specific heat at constant pressure (Cp) is important in determining temperature increase along the coolant flow path (see Fig. 10-80). The temperature rise between the bulk coolant and the coolant/solid interface is a function of the Nusselt (Nu)
328
10 Fusion Reactor Materials
Table 10-27. Basic physical properties for liquid metal coolants.
Coolant
Melting temp. (°C)
Boiling b temp. (°C)
Latent heats (kJ/kg)
Density at Tm (kg/m3)
Fusion Ga K Li Li 17 Pb 83 a Na
1983 760 1317
30 64 181 235 98 -11
Based on atom fraction;
883 784 b
6200 829 518 9600 928 871
80.2 61.1 66.2 33.9 113
Volume exp. at melting (%)
Vaporization 4246 2077 19595 4208
-3.1 2.41 1.5 3.5 2.5 2.5
at one atmosphere
10000 9000
Figure 10-74. Densities of liquidmetal coolants. 800
Figure 10-75. Heat capacities of liquid-metal coolants. 800
10,4 Blanket Materials
329
Figure 10-76. Thermal conductivities of liquid-metal coolants. 800
Figure 10-77. Viscosities of liquidmetal coolants. 100
200
300
400
500
600
700
800
T(°C)
Figure 10-78. Vapor pressures of liquid-metal coolants. 400
600
800
1000
7(°C)
1200
1400
1600
1800
330
10 Fusion Reactor Materials
Figure 10-79. Electrical resistivities of liquid-metal coolants. 100
200
300
400
500
600
700
Figure 10-80. Heat capacity times density for liquid-metal coolants.
400
800
number. For turbulent flow through a pipe with constant heat flux through the walls, the heat transfer coefficient (hc in Wm~ 2 K" 1 ) is given by hc = (k/De) Nu
(10-15)
where Nu = l + 0.025 (Re Pr)0-8 Re = VDeQ/fi= VDJv Pr =Cpfi/k
800
(10-15a) (10-15b) (10-15 c)
V = bulk velocity D e = hydraulic diameter ( = diameter for a tube) Q = density ju = viscosity v = fi/g = kinematic viscosity C p = specific heat at constant pressure and k
= thermal conductivity.
10.4 Blanket Materials
For liquid metal flow through a magnetic field, the heat transfer coefficient is more complicated. While solutions to this problem are geometry, fluid, and magnetic-field dependent, a conservative (i.e., pessimistic) approach is to assume laminar flow for design scoping studies. The Nusselt number for developed laminar flow through a pipe is (10-15 d)
Nu = 4.36
In general, the laminar-flow Nusselt number varies between two and five for different geometries of interest.
331
The Reynold's number (Re) is also important in determining the pressure drop in a non-magnetic field. Thus from a thermal-hydraulic view, coolants with low kinematic viscosity (Fig. 10-81) and high Prandtl number (Pr in Fig. 10-82) are desirable. For Li and Li 17 Pb 83 , other important considerations are the tritium diffusivity and solubility, as these materials are tritium breeders as well as coolants. Figure 10-83 shows the diffusivity of tritium as a function of temperature for these materials, as well as for solid blanket materials.
Figure 10-81. Kinematic viscosities for liquid-metal coolants. 100
200
300
400
500
600
700
800
T (°C)
Figure 10-82, Prandtl numbers for liquidmetal coolants. 800
332
10 Fusion Reactor Materials
0.0001 ^0
* ^ \ (46.9) \
V
| 1.00E-6
o Li2O o UA1O 2 ABeO o A1 2 O 3 + Li17Pb83 « Li • Y-Fe A Be
(28.4) -____^
X
(19.5)
f 1.00E-8 CO
^ ^ X \ (126.SC) \> X\
1 1.0E-10 \ CD r>
\
•§ 1.0E-12 CO
X 1.0E-14 0.5
\
(95.1,SC)
\\(220,SC) \ \ 1
(102)
programs in place to fill some of the gaps in the database. In the case of liquid metals, an extensive database is available within both the fission and fusion research communities. However, there is a need to peer-review these data, develop design correlations where needed, and document this work at the same level of detail as has been done for solid breeder/beryllium materials.
(239,SC) 1.5 2 1000 K/(T)
2.5
Figure 10-83. H-isotope diffusion in blanket materials.
These properties are needed to assess permeation losses within the blanket and ease of tritium recovery outside the blanket region. The high solubility of tritium in lithium leads to low permeation losses within the blanket, and some difficulty in tritium extraction outside the blanket. The situation is reversed for Li 17 Pb 83 . Chemical compatibility with the structural material and the gas and liquid environments, as well as MHD effects, are also important considerations when selecting a liquid-metal coolant. While these are beyond the scope of this review, the reader is referred to Malang etal. (1991), which contains brief descriptions of these issues and an excellent set of references to further pursue these issues. 10.4.9 Summary
The database for solid breeder materials and beryllium is well developed. It has been critically reviewed by members of the fusion design and materials' communities. Models and properties' correlations have been developed to facilitate the application of this database to design analyses. Areas for further development have been identified and prioritized, with on-going R & D
10.5 Insulators 10.5.1 Bulk Insulators
Leading candidates for bulk insulators include A12O3, MgO, MgAl 2 O 4 (spinel), and BeO. Degradation of these insulators can be permanent, as from neutron damage; or transient, as a result of absorption of ionizing energy or generation of a high concentration of defects during irradiation. Transient damage, which takes the form of a temporary decrease in insulating properties during irradiation, is of special concern because such degradation can occur as soon as the fusion device begins to generate radiation fields. The effect can be large: Farnum et al. (1992) have observed a dramatic increase in room-temperature ac conductivity in a single-crystal of A12O3 during irradiation with 3 MeV protons, as ionizing energy produces an excess number of charge carriers. Much of the increase dissipates with time as the damage dose builds up; this effect is attributed to trapping and recombination of electrons and holes at defect sites. Hodgson (1991) has identified another, possibly more serious problem: degradation of electrical and perhaps even structural properties when insulators are irradiated while under an imposed electric field. This phenomenon, which is not well un-
10.5 Insulators
derstood, is currently under study at several laboratories around the world. The lifetime neutron fluence in nearterm fusion systems is not sufficient to cause significant structural degradation in most ceramics. However, Macor machinable glass-ceramic, one of a family of materials that lends itself to inexpensive production of complex insulator configurations, has been shown to swell significantly at 4 x l 0 2 2 14MeV n/m 2 (Coghlan and Clinard, 1991). Since the net swelling observed is attributed to growth of the crystalline mica phase and densification of the glassy phase, a large reduction of strength due to internal stresses is likely. At the much larger, lifetime neutron fluence characteristic of long-term fusion systems, many ceramics can undergo significant swelling and strength loss. Fission reactor irradiation tests of BeO conducted at ~100°C by Hickman (1966) have shown that 80 % of the strength is lost at a fluence about fifty times less than that characteristic of the ITER lifetime fluence. Strength loss results from anisotropic swelling of individual grains in this non-cubic ceramic, lead to severe internal stresses and microcracking. A similar problem can affect polycrystalline A12O3, although a usable strength may be retained to the end-oflife fluence (Tucker etal., 1986). Cubic MgAl 2 O 4 spinel can actually strengthen under high-dose fission neutron irradiation (Hurley et al., 1981). However, Zinkle and Kojima (1991) have shown that cavitation is observed in grain boundaries of spinel at fusion-relevant He/dpa (appm generated He/displacements per atom) levels, so that the usefulness of this ceramic in a fusion environment is uncertain. It should be noted that a significant fraction of absorbed ionizing energy results from stopping within the material, ions knocked from their lattice positions by
333
neutron bombardment. A useful approximation by Dell and Goland (1981) for A12O3 at the first wall is that 6.1 x 1015 n/ m 2 deposits 1 Gy. This comprises about a tenth of the ionizing energy absorbed from accompanying gamma rays in a fusion reactor, although damage effects are not necessarily comparable, since energy is deposited more locally in the case of neutrons. 10.5.2 Windows Candidate window materials for fusion systems will be subjected to radiation that can degrade the transmission characteristics, induce luminescence, cause dimensional changes, and alter the mechanical properties. Proposed candidate materials for windows are, according to Taylor (1988): - crystalline quartz for the far infrared region - ZnSe for the infrared - fused silica for the visible and near ultraviolet - sapphire and magnesium fluoride for the ultraviolet. Approaches to the development of windows for near-term designs include the selection of materials that exhibit the greatest resistance to radiation damage, the operation of windows at elevated temperatures ( J T > 3 0 0 O C ) where irradiation-induced defects can continuously anneal, periodic annealing of defects following lowtemperature operation, and the judicious design of windows to accommodate material changes without failure. The following discussion emphasizes the effect of dimensional changes in SiO2, with optical damage in this material being discussed in the section on optical fibers.
334
10 Fusion Reactor Materials
Primak and co-workers (1955, 1958, 1962, 1981) have shown that both fused silica and crystalline quartz change dimensions when irradiated with low fluences of fast neutrons or gamma rays at ambient temperatures. Three mechanisms have been identified that contribute to dimensional changes in fused (vitreous) silica during reactor irradiations: densification due to collision damage, contraction but later expansion due to gamma rays, and radiation-induced stress relaxation. Figure 10-84 shows trend lines for dimensional changes in these materials. Inhomogeneities in the original material can result in localized stress centers as well as macroscopic dimensional changes. Small dimensional changes in window materials can result in high stress levels in window assemblies. It is estimated that the upper limit for the radiation-induced interference strain for window assemblies of current design is ~ 5 x 10 ~4; design optimization may improve this limit by a fac-
tor of two (Taylor, 1988). This leads to the conclusion that windows of current design should not be subjected to a fluence in excess of ~ 5 x 1021 fusion n/m 2 . Larger dimensional changes could be tolerated if sliding seal assemblies can be developed.
10.5.3 Optical Fibers Silica-based optical fibers are the leading candidates for fusion system diagnostic applications operating in the range ~ 0 . 4 2.0 jim. These materials are mechanically viable, thermally stable, and less susceptible to radiation-induced attenuation than are other types of glasses. To preserve their mechanical and thermal characteristics these fibers must be metal-coated rather than polymer-jacketed, but the technology to do this exists. Radiation-induced color centers can never be eliminated, but as a result of recent research these defects are becoming understood, and it is known that
Fusion Neutrons (n m ) 22
21
10
10 Contraction
vitreous silica ^ n irradiation /
O) v-3
c 10
_ Expansion
crystalline ^
: :
/ / ^~^_// /
N ^ "**O
quartz / / \ ^ / n irradiation// 7^
CD
OA
yS
CO
I
I ioo CO
10"
_
/ /
/
/
/
Figure 10-84. Fractional contraction of vitreous silica and expansion of crystalline quartz as a function offluenceof fission neutrons and ionization dose, at ambient temperature. The upper scale shows equivalent fluence of fusion neutrons. The term 'OA' refers to optical axis.
/
y
Contraction vitreous silica e" irradiation 1
10
100 6
Fluence scale, 10 Gy or 10 21 damaging neutrons nrr2
1000
335
10.5 Insulators
they can be controlled to a degree. (For a review, see Griscom, 1991.) Figure 10-85 shows the spectral dependence of induced absorption in a commercial, high-purity, fused silica of low water content, after irradiation with 2 MeV electrons (Friebele et al., 1987). Also shown are Gaussian resolutions of several of the distinct color center bands that have been isolated in the course of correlated optical and electron spin resonance (ESR) experiments. The latter technique gives information about the atomic-scale structure of the color centers (Griscom, 1990). A striking dependence on dose rate has been reported for the 5.85 eV band (Palma and Gagoz, 1972), indicating the importance of in-situ experiments and tests employing rates characteristic of fusion reactors. Onset of saturation for this band is near 1010 Gy (Pfeffer, 1985), or 1000 times the ionizing dose represented in Fig. 10-85, thus transmissivity of the SiO2 fibers may be expected to be severely degraded by the expected fusion fluences. Other absorption bands in the infrared region have been tentatively ascribed to alkali and hydrogen impurities (Friebele etal., 1985). If this interpretation is correct, pick up of hydrogen isotopes from a fusion reactor environment could prove to be a problem. Construction of an optical fiber involves core and cladding glasses of different compositions, so that ~ 10% dopants must be introduced into one or the other in order to achieve waveguiding. With the exception of fluorine, the usual dopants employed for this purpose lead to greater radiationinduced attenuation. Thus a pure-silicacore, F-doped-silica-clad fiber would seem to be the best for fusion diagnostic applications, although telecommunications fibers with a Ge-doped silica core and clad with pure silica should also be tested.
Wavelength (nm) 800 600 500
400
300
3
4
250
Energy (eV) Figure 10-85. Optical absorption in Suprasil W2 fused silica after irradiation with 2 MeV electrons to 10 MGy.
Finally, in addition to the radiation-induced attenuation bands, there are also radiation-induced luminescence bands that may be excited either by the radiation itself or by light propagating in the fiber. The spurious light could mask the true optical data from the plasma. In particular, an emission band near 2.7 eV is intrinsic to pure silica and may arise from any or all of the following: self-trapped excitons (Itoh et al., 1988), oxygen vacancies (Tohmon etal., 1989), or two-coordinated silicon ions (Skuja et al., 1984). In diagnostic applications, it is likely that the luminescence emissions will have to be removed from the data stream by computer processing, perhaps by making use of the light output of reference fibers exposed to the same radiation fluxes but prevented from propagating the signal light. Similar approaches may be employed for the optical absorption bands. 10.5.4 Reflectors Reflectors (mirrors) are specified for use at wavelengths from 0.3 to 11 fim and 0.1 to 5 mm, and at temperatures up to 350 °C.
336
10 Fusion Reactor Materials
These diagnostic elements may be made of a variety of materials or materials' combinations, depending on the wavelength to be reflected. Possibilities include: all metal, all ceramic, or all glass; ceramic with metal coating; ceramic or glass with dielectric coating; metal with dielectric coating; metal with metal coating. Transient damage may be in the form of temporary luminescence or darkening. Permanent degradation could result from darkening deposition of contamination layers, sputtering, microcracking, debonding, spalling, swelling, or structural damage. Some of these topics have previously been discussed with respect to SiO2 in the sections on windows and optical fibers. Many of the structural problems may also be encountered with metals, although materials with significant ductility should show adequate resistance to brittleness-related degradation mechanisms such as microcracking and spalling. Metals are of course also immune to luminescence and darkening effects.
10.6 References Abdou, M., Baker, C , Brooks, X, DeFreece, D., Ehst, D., Mattas, R., Morgan, G. D., Smith, D., Trachsel, C. (1982), A Demonstration Tokamak Power Plant Study (DEMO), Argonne National Laboratory Report ANL/FPP-82-1. Abdou, M., Gierszewski, P., Tillack, M., Sze, D. K., Bartlit, J., Berwald, D., Grover, X, McGrath, R., Puigh, R., Reimann, X (1985), FINESSE Phase I Report, University of California (Los Angeles), Report PPG-909, UCLA-ENG-85-39. Almen, O., Bruce, G. (1961), NucL Instrum. Methods 11, 257, 279. ASM (1979), Metals Handbook, Vol. 2: Properties and Selection: Nonferrous Alloys and Pure Metals, 9th ed. Metals Park, OH: American Society for Metals. Attaya, H., et al. (1992), J. NucL Mater. 191-194, 1469. Badger, B., et al. (1975), UWMAK-II, Univ. of Wisconsin, UWFDM-122.
Baker, C. C , Abdou, M. A., etal. (1980), STARFIRE-A Commercial Tokamak Fusion Power Plant Study, Argonne National Laboratory Report ANL/FPP-80-1. Baker, C. C , Brooks, X N., Ehst, D. A., Smith, D. L., Sze, D. K. (1985), Tokamak Power Systems Studies - FY1985, Argonne National Laboratory Report ANL/FPP-85-2. Baldwin, D. L. (1992), in: Proc. Int. Workshop on Ceramic Breeder Blanket Interactions, Clearwater Beach, FL, November 22-23, 1991, p. 43. Baldwin, D. L., Billone, M. C. (1993), Presentation at 6th Int. Conf. on Fusion Reactor Materials (ICFRM-6). Barabash, V. R., etal. (1992), /. NucL Mater. 191194, 411. Baskes, M. I. (1980 a), SAND80-8201. Baskes, M. I. (1980b), J. NucL Mater. 92, 318. Baskes, M. I. (1983), DIFFUSE 83, SAND83-8231. Bates, X R, Johnston, W. G. (1977), Radiation Effects in Breeder Reactor Structural Materials. New York: AIME, p. 625. Billone, M. C. (1991), in: Beryllium Technology Workshop, Clearwater Beach, FL: Longhurst, G. (Ed.), EGG-FSP-10017. Billone, M. C , Baldwin, D. L. (1992), in: Proc. Int. Workshop on Ceramic Breeder Blanket Interactions, Tokyo, Japan. Billone, M. C , Grayhack, W. T. (1988), Summary of Mechanical Properties Data and Correlations for Li2O, Li4Si04, LiAlO2, and Be, Argonne National Laboratory Report ANL/FPP/TM-218. Billone, M. C , Grayhack, W. T. (1989), Adv. Ceram. 25, 15. Billone, M. C , Lin, C. C , Baldwin, D. L. (1991), Fusion Technol. 19, 1707. Billone, M. C , Attaya, H., Kopasz, X P. (1992), Modeling of Tritium Behavior in Li2O, Argonne National Laboratory, ANL/FPP/TM-260. Billone, M. C , Dienst, W, Flament, T., Lorenzetto, P., Noda, K., Roux, N. (1993), ITER Solid Breeder Blanket Materials Database, Argonne National Laboratory Report ANL/FPP/TM-263, to be published. Blewer, S. R., Whitley, X B. (1983), in: Proc. 10th Symp. on Fusion Engineering: Philadelphia, PA: IEEE, p. 1027. Blewitt, T. H. (1958), ORNL-2614, 64. Borch, N. R. (1979), in: Beryllium Science and Technology, Vol. 1: Webster, D., London, G. X (Eds.). New York: Plenum Press, Chap. 8. Borders, X A., Langley, R. A., Wilson, K. L. (1978), J. NucL Mater. 76, 168. Boszo, R, Ertl, G., Grunze, M., Weiss, M. (1977), Appl. Surf. Sci. 1, 103. Braganza, C. M., Erents, S. K., McCracken, G. M. (1978), J. Nucl Mater. 75, 2201. Brice, D. K., Doyle, B. L. (1981), /. NucL Mater. 103-104, 503.
10.6 References
Brice, D. K., Doyle, B. L., Wampler, W. R. (1982), in: Proc. 5th Int. Conf. on Plasma Surface Interaction in Controlled Fusion Devices, May 3-5, 1982, Gatlinburg, TN. Brooks, J. N. (1983), Nucl. Technol./Fusion 4, 33. Brown, W. R, Jr., King, B. (1974), Aerospace Structural Metals Handbook. Belfour Stulen. Burchell, T. D. (1991), in: Proc. of the US-Japan Workshop Q-142 on High heat Flux Components and Plasma-Surface Interactions for Next Devices, SAND92-0222. Bush, S. H. (1965), Irradiation Effects in Cladding and Structural Material. Metals Park, OH: American Society for Metals. Carpenter, G. J. C , Fleck, R. G. (1977), in: Beryllium 1977. Paper 26. Causey, R. A., Doyle, B. L. (1989), unpublished. Causey, R. A., Elleman, T. S., Verghese, K. (1979), Carbon 17, 323. Causey, R. A., Baskes, M. I., Wilson, K. L. (1986), /. Vac. Sci. Technol. A4, 1189. Chopra, O. K., Smith, D. L. (1981), J. Nucl. Mater. 103-104, 651. Chopra, O. K., Smith, D. L. (1984), /. Nucl. Mater. 122-123, 1213. Chopra, O. K., Smith, D. L. (1986), /. Nucl Mater. 141-143, 566. Claudson, T. X, Pessl, H. J. (1965), Irradiation Effects on High Temperature Reactor Structural Materials, BNWL-23. Battelle Pacific Northwest Laboratory. Coghlan, W. A., Clinard, R W, Jr. (1991), J. Nucl. Mater. 179-181, 391. Conn, R., et al. (1994), ARIES-II UCLA Report, in press. Dalle Donne, M., Sordon, G. (1990), Fusion Technol. 17, 597. Dalle Donne, M., et al. (1991), Status Report, KfK Contribution to the Development of DEMO-relevant Test Blankets for NET/ITER, Part 2: BOT Helium Cooled Solid Breeder Blanket, Kernforschungszentrum Karlsruhe Reports KfK 4928, 4929. Davis, J. W. (1980-1986), Materials Handbook for Fusion Energy Systems, U. S. Department of Energy Report, DOE/TIC-10122. Davis, I W, Smith, D. L. (1979), J. Nucl. Mater. 85-86, 71. Dienes, G. X, Damask, A. C. (1961), /. Nucl. Mater. 3(1), 16. Diercks, D. R., Smith, D. L. (1986), /. Nucl. Mater. 141-143, 617. Dell, G. R, Goland, A. N. (1981), Radiation Damage Parameters in Multicomponent Nonmetals, Brookhaven National Laboratory Report, BNL 29615. Dobson, R. L., Whitley, J. B. (1987), Erosion Corrosion of Copper in a High Velocity Water Environment, Sandia National Laboratories, SAND870312. Doyle, B. L. (1981), SAND81-0622; and in: US INTOR, INTOR/81-1, Vol. 1, Chap. VII.
337
Doyle, B. L., Wampler, W. R., Brice, D. K., Picraux, S. T. (1980), J. Nucl. Mater. 93-94, 551. Doyle, B. L., Wampler, W. R., Brice, D. K. (1981), J. Nucl. Mater. 103-104, 513. Dunmur, I. W. (1979), Beryllium Science and Technology, Vol. 2, Chap. 9. Dunmur, I. W, et al. (1977), in: Beryllium 1977: 4th Int. Conf. on Beryllium. The Metals Society, paper 35. Edwards, D. X, et al. (1992), /. Nucl. Mater. 191-194, 416. Ehst, D. A., etal. (1986), Tokamak Power Systems Study, ANL/FPP/86-1. Erents, S. K. (1975), in: Proc. Int. Conf. on Applications of Ion Beams to Materials, University of Warwick, U.K. Erents, S. K., Braganza, C. M., McCracken, G. M. (1976), /. Nucl. Mater. 63, 399. Evans, R. R. V., Weinberg, A. G., Van Thyne, R. X (1963), Acta Metall. 11, 143. Fabritsiev, S. A., Gosudarenkova, V. A., Potapova, V A., Rybin, V. V, Kosachev, L. S., Chakin, V. P., Pokrovsky, A. S., Barabash, V. R. (1992), /. Nucl. Mater. 191-194,426. Farnum, E. H., Kennedy, X C , Clinard, R W, Frost, H. M. (1992), /. Nucl. Mater. 191-194, 548. Federici, G., Wu, C. H., Raffray, A. R., Billone, M. C. (1992), J. Nucl. Mater. 187, 1. Fetz, H., Oechsner, H. (1963), in: Proc. 6th Int. Conf. Phenomenes d'lonisations dans les Gaz, Paris, p. 39. Flament, T, et al. (1992), /. Nucl. Mater. 191-194, 163. Friebele, E. X, Long, K. X, Askins, C. G., Gingerich, M. E., Marrone, M. X, Griscom, D. L. (1985), in: Critical Reviews of Technology: Optical Materials in Radiation Environments Proc. SPIE, Vol. 541, Levy, P., Friebele, E. X (Eds.) Bellingham, WA: SPIE, 70. Friebele, E. X, Higby, P. L., Tsai, T. E. (1987), Diffus. Defect Data 53, 203. Fukumoto, K., Kinoshita, C, Abe, H., Shinohara, K., Kutsunada, M. (1991), /. Nucl. Mater. 179181, 935. Fullerton-Batten, R. C , Hawk, X A. (1977), in: Beryllium 1977: 4th Int. Conf. on Beryllium. The Metals Society, paper 49. Gelles, D. S., Thomas, L. E. (1984), Ferritic Steels for Use in Nuclear Energy Technologies. Warrendale, PA: TMS/AIME p. 559. Gohar, Y. (1981), ANL, unpublished results. Gorodetsky, A. E., Zakharov, A. P., Sharapov, V W, Alimov, V. Kh. (1980), / Nucl Mater. 93-94, 588. Gorynin, I. V, Fabritsiev, S. A., Rybin, V. V, Kasakov, V. A., Pokrovsky, A. S., Barabash, V. R., Prokofiyev, Y. G. (1992 a), J. Nucl Mater. 191194, 401. Gorynin, I. V, etal. (1992b), J. Nucl Mater. 191194, 421.
338
10 Fusion Reactor Materials
Grieger, G., Mori, S., Stacey, W M., Jr., Kadomtser, B. B., et al. (1988), International Tokamak Reactor, Phase Two A, Part II, Report of the International Tokamak Reactor Workshop held in five sessions in Vienna from 1985 to 1987. Vienna: International Atomic Energy Agency. Griscom, D. L. (1990), in: Glass: Science and Technology, Vol. 4B: Uhlmann, D. R., Kreidl, N. X (Eds.). Boston: Academic Press, p. 151. Griscom, D. L. (1991), /. Ceram. Soc. Jpn. 99, 923. Grossbeck, M. L. (1991), /. Nucl Mater. 179-181, 568. Grossbeck, M. L., Liu, K. C. (1982), J. Nucl. Technol. 58, 538. Grossbeck, M. L., et al. (1990), American Society for Testing and Materials, Special Technical Publication ASTM-STP 1046, Vol. II. p. 537. Gupta, C. K., Krishnamurthy, N. (1992), Extractive Metallurgy of Vanadium. Amsterdam: Elsevier. Hall, R. O. A., Martin, D. G. (1981), J. Nucl Mater. 101, ill. Hechtl, E., Bohdansky, X, Roth, J. (1981), /. Nucl. Mater. 103-104, 333. Hesketh, R. V. (1967), Collapse of Vacancy Cascades to Dislocation Loop, Battelle National Laboratory, BNL 50083, p. 389. Hickman, B. S. (1966), in: Studies in Radiation Effects, Series A, Physical and Chemical, Vol. 1: Dienes, G. X (Ed.). New York: Gordon and Breach, p. 72. Hodgson, E. R. (1991), /. Nucl. Mater. 179-181, 383. Holdren, I P., et al. (1989), Environmental, Safety and Economic Aspects of Magnetic Fusion Energy, UCRL-53766. Hollenberg, G. W. (1987), in: Fusion Reactor Materials Semiannual Progress Report for Period Ending September 30, 1986, DOE/ER-0313/1, 373-380. Hucks, P., Flaskamp, K., Vietzke, E. (1980), / Nucl. Mater. 93-94, 558. Hull, A. B., Purdy, L, Loomis, B. A. (1992), in: Proc. Fifth Int. Conf. on Fusion Reactor Materials (ICFRM-5), Clearwater FL. Hurley, G. F, Kennedy, J. C , Clinard, F. W, Jr., Youngman, R. A., McDonnell, W. R. (1981), /. Nucl. Mater. 103-104, 761. IAEA (1991), ITER Tokamak Device, ITER Documentation Series, No. 25. Vienna: International Atomic Energy Agency. Itoh, C , Tanimura, K., Itoh, N. (1988), /. Phys. C: Solid State Phys. 21, 4693. Jones, P. M. S., Gibson, R. (1967), /. Nucl. Mater. 21, 353. Jones, R. H., et al. (1992), /. Nucl. Mater. 191-194, 75. Kalinin, G. M. (1991), /. Nucl. Mater. 1193-1198. Kangilaski, M. (1971), Radiation Effects Design Handbook, Section 7, Structural Alloys, NASA CR-1873. Kirk, M. A., Robertson, I. M., Vertano, I. S., Jenkins, M. L., Funk, L. L. (1986), in: Radiation Induced
Changes in Microstructure: Garner, F. A., Packan, N. M., Kumar, A. S. (Eds.). Philadelphia: ASTM, pp. 48-69. Klueh, R. L. (1992), J. Nucl. Mater. 191-194, 116. Klueh, R. L., Alexander, D. J. (1992), /. Nucl. Mater. 191-194, 896. Klueh, R. L., Maziasz, P. J. Vitek, J. M. (1986), J. Nucl. Mater. 141-143, 960. Kohyama, A., et al. (1992), /. Nucl. Mater. 191-194, 37. Kopasz, J. P., Seils, C. A., Johnson, C. E. (1993), in: Int. Workshop on Ceramic Breeder Blanket Interactions, Tokyo, Japan, October 26-29, 1992. Kuroda, T, et al. (1991), ITER Plasma Facing Components, ITER Documentation Series, No. 30. Vienna: International Atomic Energy Agency. Laegreid, N., Wehner, G. K. (1961), /. Appl. Phys. 32, 365. Langley, R. A., Blewer, R. S., Roth, J. (1978), J. Nucl. Mater. 76-77, 313. Lauritzen, T, et al. (1981), Irradiation Induced Embrittlement of Some High Nickel Alloys, General Electric Rpt. GEFR-00576. Liu, K. (1981), / Nucl. Mater. 103 and 104, 913. Loomis, B. A., Smith, D. L. (1992), /. Nucl. Mater. 191-194, 84. Maeda, S., Mohri, M., Hashiba, M., Yamashina, T, Kaminsky, M. (1981), /. Nucl. Mater. 103-104, 445. Makin, M. J., Minter, F. X (1959), Acta. Metall. 7, 361. Malang, S., Leroy, P., Casini, G. P., Mattas, R. F , Strebkov, Yu. (1991), Fusion Eng. Des. 16, 95. Maroni, V. A., Cairns, E. J., Cafasso, F. A. (1973), A Review of the Chemical, Physical, and Thermal Properties of Lithium that are Related to Its Use in Fusion Reactors, Argonne National Laboratory Report ANL-8001. Materials Advisory Board (1963), Evaluation Test Methods for Refractory Metal Sheet Specimens, Report MAB-192-M. Washington, DC: Materials Advisory Board, National Academy of Sciences. Mayerhofer, U. (1986), Diploma Thesis, Technical University of Munich. Maziasz, P. X (1985), /. Nucl. Mater. 133'/134, 134. Maziasz, P. X, Klueh, R. L. (1989), American Society for Testing and Materials, Special Technical Publication, ASTM-STP-1046, p. 35. Mishima, Y, Ishino, S., Shiozawa, S. (1977), in: Beryllium 1977, Paper 25. Najmabadi, F., Conn, R. W (1991), The ARIES I Tokamak Reactor Study, UCLA-PPG-1323. Nardi, C. (1991), Status of Knowledge About the Beryllium Swelling by Neutron Irradiation, ENEA Report, ISSN/1120-5598. Nesmeyanov, An. N. (1963), Vapor Pressure of the Elements, translated by X I. Carasso, New York: Academic Press. Noda, K., Billone, M. C , Dienst, W, Flament, T, Lorenzetto, P., Roux, N. (1990), Summary Report
10.6 References
for the ITER Specialists' Meeting on Blanket Materials Data Base, ITER Technical Memorandum. Odette, G. R., Lucas, G. E. (1992), J. Nucl Mater. 191-194, 53. Palma, G. E., Gagoz, R. M. (1972), /. Phys. Chem. Solids 33, 111. Peterson, D. T. (1982), U.S./DOE Report ER-0045/8, 304. Peterson, D. T, Hull, A. B., Loomis, B. A. (1992), in: Proc. 5th Int. Conf on Fusion Reactor Materials (ICFRM-5), Clearwater, FL. Petukhov, V. A., Chekhovskoi, V. Ya. (1972), High Temp.-High Pressures 4, 611. Pfeffer, R. L. (1985), J. Appl Phys. 57, 5176. Piet, S. X, Cheng, E. T, Porter, L. X (1990), Fusion Technol, 636. Pionke, L. X, Davis, X W. (1979), Technical Assessment of Niobium Alloys Data Base for Fusion Reactor Applications, USDOE RPT COO-4247-2. Primak, W. (1958), Phys. Rev. 110, 1240. Primak, W. (1981), Radiation Damage in Diagnostic Window Materials for TFTR, ANL/FPP/TM-146. Primak, W, Edwards, E. (1962), Phys. Rev. 128, 2580. Primak, W, Fuchs, L. H., Day, P. (1955), J. Am. Ceram. Soc. 38, 135. Rendulic, K. D., Winkler, A. (1978), Surf. Sci. 74, 318. Rohrig, H. D., Fischer, P. G., Hecker, R. (1976), J. Am. Ceram. Soc. 59, 316. Rosenberg, D., Wehner, G. K. (1962), / Appl. Phys. 33, 1842. Roth, X Bohdansky, X, Blewer, R. S., Ottenberger, W, Borders, X (1979), /. Nucl. Mater. 85-86,1077. Roth, X, Bohdansky, X, Martinelli, P. A. (1981), in: Proc. Int. Conf on Ion Beam Modification of Materials, Budapest. Roth, X, Bohdansky, X, Wilson, K. L. (1982), in: Proc. 5th. Int. Conf on Plasma Surface Interactions, Gatlinburg, TN. Roth, X, Bohdansky, X, Ottenberger, W. (1989), /. Nucl. Mater. 165, 193. Ruther, W. R, Kassner, T. F. (1993), USDOE Rpt. ER-0313/14, 395. Rybin, V. V., Smith, D. L. (1992), /. Nucl. Mater. 191-194, 30. Sannen, L., De Raedt, Ch. (1992), in: 17th SOFT Conference, Rome, Italy. Scherzer, B. M. U., Behrisch, R., Eckstein, W, Littmark, U., Roth, X, Sinha, M. K. (1976), J. Nucl Mater. 63, 100. Skuja, L. N., Streletsky, A. N., Pakovich, A. B. (1984), Solid State Commun. 50, 1069. Smith, D. L., et al. (1982), US FED-INTOR Activity, Chapter VII-Impurity Control and First Wall Engineering, Materials Data Base, USAFED-INTOR/ 82-1. Smith, D. L., Morgan, G. D., et al. (1984), Blanket Comparison and Selection Study, Argonne National Laboratory Report, ANL/FPP/84-1; Fusion Technol. 8, 10 (1985).
339
Smith, D. L., et al. (1985), / Nucl. Mater. 135, 125. Smith, D. L., et al. (1990), U. S. Contribution on Plasma Facing Components, ITER-TN-PC-1-O-U2, Garehing, Germany. Smith, D. L., Altovsky, I. V., Barabash, V. R., Beston, X, Billone, M., Boutard, X L., Burchell, T, Davis, X, Fabritsiev, S. A., Grossbeck, M., Hassanein, A., Kalinin, G. M., Lorenzetto, P., Mattas, R., Noda, K., Nygren, R., Odintsov, N. V., Rybyn, V. V., Takatsu, H., Vinokurov, V. P., Watson, R., Wu, C. (1991 a), ITER Blanket, Shield and Material Data Base, Part B (Material Data Base), ITER Report No. 29, IAEA/ITER/DS/29. Vienna: International Atomic Energy Agency. Smith, D. L., Antipenkov, A., Baker, C , Billone, M., Daenner, W, Gohar, Y, Kuroda, T, Lorenzetto, P., Maki, Y, Mori, S., RafFray, A., Shatalov, G., Sidorov, A., Simbolotti, G., Sviatoslavsky, I., Takatsu, H., Yoshida, H. (1991 b), in: Proc. of the 13th Int. Conf on Plasma Physics and Controlled Nuclear Fusion Research, International Atomic Energy Agency, Washington D. C , 1-6 October 1990. Smith, X N., Meyer, C. H., Layton, X K. (1977), /. Nucl. Mater. 67, 234. Smolik, G. R., Merrill, B. X, Wallace, R. S. (1992), in: Proc. 5th Int. Conf on Fusion Reactor Materials (ICFRM-5), Clearwater, FL, U.S.A., November 17-22, 1991, Part A, p. 153. Sone, K., McCracken, G. M. (1982), in: Proc. 5th Int. Conf on Plasma Surface Interactions in Controlled Fusion Devices, May 3-5, 1982, Gatlinburg, TN. Southern Research Institute (1966), Report on the Mechanical and Thermal Properties of Tungsten and TZM Sheet Produced in the Refractory Metal Sheet Rolling Program - Part 1, Southern Research Institute, Rpt. No. AD-6386-31. Steichern, X M. (1976), J. Nucl. Mater. 60, 13. Stoller, R. E., et al. (1988), /. Nucl. Mater. 155-157, 1328. Suiter, D. X (1983), Lithium Based Ceramics for Tritium Breeding Applications, McDonnell Douglas Astronautics Company Report MDC E2677 (UC20). Sullivan, X D., Brayman, C. L., Verrall, R. A., Miller, X M., Gierszewski, P. X, Londry, F. (1991), Fusion Eng. Des. 17, 79. Tanabe, T. (1991), in: Proc. of US-Japan Workshop Q-142 on High Heat Flux Components and PlasmaSurface Interactions for Next Devices, SAND920222. Tanabe, T, Saito, N., Etoh, Y, Imoto, S. (1981), J. Nucl. Mater. 103-104, 483. Taylor, A. (1988), Report on the Survivability of Diagnostic Windows for the CIT Reactor, Argonne National Laboratory, Fusion Power Program Internal Document. Tietz, T. E., Wilson, J. W. (1965), Behavior and Properties of Refractory Metals. Standard, CT: Stanford University Press.
340
10 Fusion Reactor Materials
Tipton, C. R., Jr. (Ed.) (1960), Reactor Handbook, Vol. I, New York: Interscience. Tohmon, R., Shimogaichi, Y, Mizuno, H., Ohki, Y, Nagasawa, K., Hama, Y. (1989), Phys. Rev. Lett. 62, 1388. Touloukian, Y S., Buyco, E. H. (1970), in: Thermophysical Properties of Matter, Vol. 4: New York: IFI/Plenum. Touloukian, Y S., Powell, R. W, Ho, C. Y, Klemens, P. G. (1970), in: Thermophysical Properties of Matter, Vol. 1: New York: IFI/Plenum. Touloukian, Y S., Powell, R. W, Ho, C. Y, Nicolaou, M. C. (1973), in: Thermophysical Properties of Matter, Vol. 10, New York: IFI/Plenum. Tucker, D. S., Zocco, T, Kise, C. D., Kennedy, I C. (1986), /. Nucl. Mater. 141-143, 401. Van Witzenburg, W, Bruyne, H. J. (1993), Effects of Irradiation on Materials, ASTM-STP-1175. Van Witzenburg, W, DeVries, E. (1990), Effects of Irradiation on Materials, ASTM-STP-1125. Walter, K. H., Kienberger, K. H., Lange, G. (1973), /. Nucl. Mater. 48, 287. Wampler, W R. (1984), /. Nucl. Mater. 122/123,1598. Wampler, W R., Magee, C. W. (1981), J. Nucl. Mater. 103-104, 509. Watson, R. D. (Ed.) (1989), ITER Divertor Engineering Design, Summary of June-October Joint Working Session, ITER-TN-PC-8-9-1.
Wehner, G. K. (1957), Phys. Rev. 108, 35. Wehner, G. K. (1962), General Mills Report No. 2309. Wienhold, P., Ali-Khan, I., Dietz, K. J., Profant, M., Waelbroeck, F. (1979), /. Nucl. Mater. 85-86,1001. Wienhold, P., Profant, I, Waelbroeck, F , Winter, J. (1980), /. Nucl. Mater. 94, 866. Wiffen, F. W. (1973), Nuclear Metallurgy, Vol. 18, The Metallurgical Society - AIME, p. 176-197. Wilkinson, W. D. (1969), Properties of Refractory Metals. New York: Gordon and Breach. Williams, J. M., Hinkle, N. E., Eatherly, W. P. (1972), ORNL/TM-3917. Wilson, K. L. (1984), Nucl. Fusion Special Issue 1984, Langley, R. A. et al. (Eds.), p. 28. Wilson, K. L., Baskes, M. I. (1978), /. Nucl. Mater. 76-77, 291. Yang, W J. S., Hamilton, M. L. (1984), /. Nucl. Mater. 122/123, 748. Yih, S. W. H., Wang, C. T. (1979), Tungsten: Sources, Metallurgy, Properties and Applications. New York: Plenum Press. Yoshida, H. (1990), Experimental Studies on Blanket in JAERI, ITER Rpt. ITER-IL-BL-5-0-2. Zinkle, S. I, Kojima, S. (1991), /. Nucl. Mater. 179181, 395.
11 Mixed Oxide Fuel Pin Performance Alvin Boltax Reactor Materials Technology (Formerly with Westinghouse Advanced Energy Systems) Pittsburgh, PA, U.S.A.
List of 11.1 11.2 11.3 11.3.1 11.3.2
Symbols and Abbreviations Introduction Fuel Pin and Assembly Design Steady-State Mixed Oxide Fuel Performance Initial Fuel Performance Results (1965 to 1972) Irradiation Behavior of First Generation Liquid Metal Reactor Mixed Oxide Fuel 11.3.2.1 Development Testing 11.3.2.2 Fuel Cladding Mechanical Interaction 11.3.2.3 Breached Fuel Pins 11.3.3 Prototype Fuel Behavior 11.3.4 Cladding Breach in Prototype Reactor Fuel 11.3.5 Development of Second Generation Liquid Metal Reactor Fuel 11.3.5.1 Advanced Austenitic Steels 11.3.5.2 High Nickel Alloys 11.3.5.3 Ferritic and Martensitic Stainless Steels 11.3.6 Axially Heterogeneous Fuel 11.4 Transient Behavior of Mixed Oxide Fuel 11.5 Fuel Pin Performance Codes 11.5.1 LIFE Code Development History 11.5.2 LIFE Code Description 11.5.3 LIFE Code Predictions 11.6 Summary 11.7 References
Materials Science and Technology Copyright © WILEY-VCH Verlag GmbH & Co KGaA. All rights reserved.
342 344 346 348 349 352 353 354 356 362 367 368 369 371 372 374 374 379 380 381 383 386 387
342
11 Mixed Oxide Fuel Pin Performance
List of Symbols and Abbreviations
D, Do F L, Lo P Q T V X
cladding diameter, cladding diameter under zero load condition fuel volume fuel pin length, fuel pin length under zero load condition plenum volume linear power temperature volume axial position on fuel pin
AES ANL ANS BNES BOF CDE CDF CRBR DFA DFR DOE dpa EBR-II EFR EOL FCCI FCMI FFTF FMS Fs GCR HCDA IFR LMR LWR ODS O/M PFR PIE PNC PPS PRISM
Atomic Energy Society of Japan Argonne National Laboratory American Nuclear Society British Nuclear Energy Society bottom of fuel core demonstration experiment cumulative damage fraction Clinch River Breeder Reactor driver fuel assembly Dounreay Fast Reactor Department of Energy (U.S.) displacement per atom Experimental Breeder Reactor-II European Fast Reactor end of life fuel-cladding chemical interaction fuel-cladding mechanical interaction Fast Flux Test Facility ferritic/martensitic steel fissium gas cooled reactor hypothetical core disruptive accidents Integral Fast Reactor liquid metal reactor light water reactor oxide dispersion strengthened oxygen to metal ratio prototype fast reactor post-irradiation examination Power Reactor and Nuclear Fuel Development Corporation (Japan) plant protective system power reactor innovative small module
List of Symbols and Abbreviations
RBCB SOL TD TOF TREAT
run beyond cladding breach start of life theoretical density top of fuel transient reactor test
343
344
11 Mixed Oxide Fuel Pin Performance
11.1 Introduction This chapter focuses on the performance of nuclear reactor fuels with emphasis on mixed oxide, (U,Pu)O 2 , fuel behavior in sodium-cooled liquid metal reactors (LMRs). Performance refers to the mechanical and chemical behavior of fuel pins as a function of design and operating conditions. The behavior of nuclear reactor fuels, generally, and fuel pins consisting of cylindrical pellets of fissionable material in metallic cladding, in particular, exhibit similar characteristics. In all nuclear fuels, the fission event typically produces two fission product atoms with mass numbers centered near 93 and 140. Figure 11-1 shows the mass distribution of fission products for the thermal fission of U-235. (The fission yields for U-235 and Pu-239 in fast reactors show similar trends.) The generation of high energy fission products in nuclear fuels is responsible for the following important phenomena: - Fuel growth and swelling - Enhanced fuel plasticity - Release of gaseous fission products (Xe and Kr) - Accumulation of non-volatile fission products
84
96
108 132 Mass number
144
Figure 11-1. The mass distribution of fission products from thermal fission of 235 U.
- Release of volatile fission products (e.g., Cs, I, Te, etc.) - Change or modification of the fuel microstructure due to temperature and temperature gradients Dependent on reactor type and function, nuclear fuel is almost always contained in a non-fissionable structural material. Examples of reactor types, functions, fuel types and structural containment materials (claddings) are given in Table 11-1. The LWRs are the most widely used reactors, accounting for 10 to 75 % of the electrical generating capacity in the major industrialized countries (e.g., 20%
Table 11-1. Nuclear reactor types, functions, fuel and cladding. Reactor Light water reactor (LWR) Liquid metal reactor (LMR)
Gas cooled reactors (GCR)
Coolant
Function
Fuel
Cladding
Pressurized water
Electric power
UO 2
Zircaloy
Sodium
Electric power
Stainless steel
Lithium
Space electric power
(U, Pu)O 2 (U, Pu)C (U, Pu)N U-Pu-Zr UN
Helium
Electric power
Hydrogen
Propulsion
uo2 uoc uc2
Refractory metals Pyrolytic carbon Pyrolytic carbon
345
11.1 Introduction
in the US and 75% in France). The status of the development of Zircaloy clad UO 2 fuel for water-cooled nuclear power reactors was recently published (Simnad, 1989; see also Chap. 3 of this Volume). Currently there are 12 operating LMRs in the world. Table 11-2 provides a list of these reactors with a brief description of the designs of the fuel pin. Mixed oxide fuel is the most widely used in the liquid metal reactors. The selection of mixed oxide fuel follows from the successful deployment of LWR technology including extensive experience with oxide fuel fabrication, irradiation testing and fuel reprocessing. In the US, LMR development effort is currently focussed on metal fuel, specifically the ternary U-Pu-Zr fuel system and the integral fast reactor (IFR) (Till, 1990). The IFR satisfies US criteria regarding passive safety characteristics, fuel cycle
and diversion resistance, and radioactive waste considerations. Chapter 1 provides a review of the status of development work on metal fuel. The presentation of information on LMR mixed oxide fuel will start with a discussion of fuel pin and assembly design (Section 11.2). Steady-state fuel performance (Section 11.3) is the largest section and it includes a historical perspective starting from the early LMR fuel irradiation tests with preproduction fuel pellets and commercially available cladding materials to the present situation with high burnup prototypic fuel with specially-designed cladding materials. The remaining sections of the chapter include discussions of the transient behavior of mixed oxide fuel pins (Section 11.4), the development of fuel pin performance codes (Section 11.5) and finally, conclusions and views on
Table 11-2. Planned and operating liquid metal reactors. Reactor
Fuel
Cladding
Pin diameter in mm
EBR-II FFTF (PRISM)b
U-Fs (U, Pu)O 2 U-Pu-Zr
316 316 HT-9
5.8 5.8 6.7
Phenix Superphenix (EFR) b
UO 2 ,(U,Pu)O 2 (U, Pu)O 2 (U, Pu)O 2
316 316 + Ti 15.15 Ti or PE16
6.6 8.5 8.5
PFR
(U, Pu)O 2
316
5.8
Germany
KNK-II SNR300
(U, Pu)O 2 (U, Pu)O 2
15.15 Ti 15.15Ti
6.0 6.0
CIS
BR-10 BOR-60 BN-350 BN-600
PuO 2 UO 2 UO 2 ,(U,Pu)O 2 UO 2 ,(U,Pu)O 2
316 + Nb 316 + Nb 316 + Nb 316 + Nb
5.0 6.0 6.1 6.9
Japan
JOYO MONJU a (DFBR)b
(U, Pu)O 2 (U, Pu)O 2 (U, Pu)O 2
316 PNC316 ferritic or HiNi austenitic
6.3 6.5
India
FBTR
(U, Pu)C
316
5.1
Country U.S.
France
U.K.
a
Under construction; b in design phase
346
11 Mixed Oxide Fuel Pin Performance
Plan view
Top end cap
Tag gas capsule
Cladding Primary control Shutdown control
Radial shield
Upper axial blanket pellets
0.4 cm Lower axial blanket pellets
f Wire wrap
Bottom end cap
cm
Keyway
_ J L— 0.89 cm (B)
(C)
Figure 11-2. LMR core (A), fuel assembly (B) and fuel pin (C).
the future development of LMR fuel (Section 11.6).
11.2 Fuel Pin and Assembly Design The core of a LMR consists of a closepacked array of hexagonal cross-section fuel and blanket assemblies, as shown in
Fig. 11-2. The triangular arrangement of fuel pins in hexagonal ducts maximizes the fuel volume fraction. The fuel pins in a prototype LMR assembly consist of a one meter column of mixed oxide fuel pellets (typically 75% UO 2 plus 25% PuO 2 ) enclosed in stainless steel cladding. Above and below the fuel column are UO 2 pellets which form the upper and lower axial blankets. Free space is provided within the
11.2 Fuel Pin and Assembly Design
fuel pin to store released fission gas (Xe and Kr). The volume of the free space is of the order of the fuel pellet volume. Table 11-2 provides examples of fuel pin designs. Fuel pins are located or spaced within a stainless steel duct by either a grid structure or a wire-wrap on each fuel pin. The spacing of LMR fuel pins is critical because of the potential for local overheating and possible breach of the cladding in regions where fuel pins are in contact or in close proximity. The fuel assembly ducts provide the following functions: - The ducts provide mechanical support for the fuel pins and a means to load the pins into the core as a unit. - The ducts force the sodium to flow past the fuel pins and permit individual assembly orificing to control coolant outlet temperatures. - The ducts provide safety barriers to the potential propagation of pin ruptures to adjacent fuel assemblies. The irradiation behavior of fuel pins and assemblies vary considerably depending on specific design and operating parameters. The key fuel pin design parameters include: the composition and microstructure of the cladding material, the cladding diameter and wall thickness, the fuel composition (Pu content and O/M ratio), density and dimensions (solid vs. annular pellets) of the fuel, and the plenum volume (free space). The key fuel pin operating parameters are: linear power, peak cladding temperature, fast neutron flux and operating lifetime (fuel burnup and fluence). A typical design of an LMR fuel pin involves 20% cold-worked Type 316 (CW316) stainless steel cladding with a 6 mm OD and a 0.4 mm wall thickness. Solid mixed oxide fuel pellets are 90 to 95% dense with a diametral gap of 0.15 mm between pellets and cladding. The
347
mixed oxide fuel is a solid solution of 75 % UO 2 and 25 % PuO 2 with an overall oxygen to metal (U + Pu) ratio of 1.96 to 1.99. As noted previously, the plenum-to-fuel volume ratio is near or above unity giving rise to end-of-life (EOL) fission gas pressures of less than 10 MPa. Fuel pins typically operate at peak linear power levels of 30 to 45 kW/m with cladding temperatures ranging from 370 °C (inlet) to 650 °C (outlet). The peak fast neutron flux levels range between 2 and 4 x l O 1 5 (n/cm2)/s (E>0.1 MeV). For a three year operating lifetime, peak fuel burnups are near 10% (fission of 10% of the U plus Pu atoms) with fast neutron fluence levels of 2.0xl0 2 3 n/cm 2 . A neutron dose level of 2 x 10 23 n/cm2 is equivalent to the displacement of each atom from its lattice position 100 times during the course of a three year irradiation. Thus, a fast neutron dose of 2 x 10 23 n/cm2 (E>0.1 MeV) corresponds to 100 displacements per atom (dpa). Such high dose levels can give rise to a wide range of radiation-induced phenomena in cladding materials including void swelling, in-reactor creep and changes in cladding mechanical properties (yield and tensile strength, ductility and stress-rupture properties). Chapter 6 and later sections of this chapter provide additional information on these phenomena. The basic function of the fuel pin is to provide fission energy to the coolant while remaining intact and relatively dimensionally stable during normal and off-normal reactor operations. The fuel assembly is designed to perform satisfactorily with fuel pin diametral growth of the order of 10% and axial growth of a few percent. Failures or breaches of the fuel pin cladding will release stored fission gas and eventually permit entry of coolant (sodium) into the pin. Sodium and mixed oxide fuel react at
348
11 Mixed Oxide Fuel Pin Performance
low temperatures (10%) have been reported for both wire-wrapped (Washburn and Weber, 1986) and grid (Levine et al., 1986) assemblies. Several examples of cladding breach data obtained in EBR-II will be presented. This data provides the basis for a quantitative method for predicting cladding breach. Cladding breach data obtained in prototype reactors will be discussed in Section 11.3.4. A significant number of the pre-1980 cladding breaches in EBR-II have been evaluated and categorized (Weber et al., 1979). Table 11-4 provides a summary of the key fuel pin data. Approximately half of the cladding breaches have been attributed to local cladding over-temperature where the wire-wrap spacer did not limit pin-to-pin or near pin-to-pin contact. The combined effects of large temperature gradient induced bow of peripheral pins, the use of insulated ducts and numerous reconstitutions of the assemblies to permit periodic examination of the fuel pins contributed to displacement of pins from nor-
357
11.3 Steady-State Mixed Oxide Fuel Performance
Table 11-4. Cladding variables and irradiation conditions of breached fuel pins. Pin
Cladding
Type/ condition
PNL 5A-1 PNL5B-17 PNL 10-14 PNL 11-39 N/E/N122 E/F-N054 P-12AA-63K P-12AB-11B P-23B-1A P-23A-37 P-23B-73E P-14-29 P-41-A3R P-41-B10 P-42-B71 P-40-D91
304/S/A 304/S/A 316/20% CW 316/20% CW 316/20% CW 316/20% CW 316/30% CW 316/20% CW 316/20% CW 316/20% CW 316/20% CW 316/20% CWC 316/20% CWC 316/20% CWC 316/20% CWC 316/20% CWC
Fuel burnup
Cladding fluence
O.d./wall thickness (mm)/(mm)
EOL a %
EOL n/cm2 x 1023 (£>0.1MeV)
6.35/0.41 6.35/0.41 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 5.84/0.38 6.86/0.28 6.86/0.28 6.86/0.28 5.84/0.25
111 13.7 6.6 10.7 4.4 3.9 3.4 7.1 5.3 10.5 9.6 10.4 1.7 1.7 1.5 12.1
0.85 0.96 0.50 1.03 0.31 0.29 0.23 0.52 0.5 1.0 0.95 0.78 0.19 0.19 0.12 0.9
Internal Calculated cladding pressure midwall temperature at top of fuel BOL b /EOL EOL MPa
°C
8.27 9.82 3.93 7.34 2.62 2.48 2.96 6.00 2.96 7.01 5.87 7.20 2.77 1.02 0.88 4.01
475/457 480/471 564/540 536/534 598/585 602/578 590/642 622/628 627/612 595/575 706/660 602/578 568/565 568/564 541/525 614/581
EOL: end of life; b BOL: beginning of life; c pins are clad with FFTF cladding vs. N-lot in prior listings.
mal positions. Estimates of local cladding temperatures at the breach sites, by comparing the cladding microstracture with isothermal annealing of unirradiated cladding, indicated values in the range 580 to 780 °C. These local temperatures are 20 to 160°C above the nominal values. All of the EBR-II breached pins demonstrated a generally benign nature. The large volumes of fission gas released from the high burnup pins had no apparent effect on neighboring pins. Furthermore, the breached fuel pins operated satisfactorily for months after breaching. The continuing operation of breached pins, referred to as run-beyond-cladding-breach (RBCB) is described in detail in Chapter 3. During the 1980's, major effort continued on the study of cladding breaches. Two EBR-II fuel irradiation programs will be highlighted in the following discussion.
The first involved a program to develop advanced oxide fuel pins possessing increased linear power and burnup capability concomitant with a high fuel volume fraction in the core, as compared to the FFTF core design (Lawrence et al., 1986). The fuel pin parameters of interest included: pin diameter, fuel smeared density and cladding thickness. The second was a collaborative DOE/PNC program involving the irradiation of a wide variety of fuel pin designs under duty cycle and transient overpower conditions (Boltax et al., 1990 b). One of the objectives of this program was to extend the development of cladding breach criteria for normal and off-normal operating conditions. Figure 11-15 summarizes the breached fuel experience for the Advanced Oxide test series (Lawrence et al., 1986). Low burnup breaches were obtained with fuel
358
11 Mixed Oxide Fuel Pin Performance
9392-
e 91|
90-
S 8 9 - Cladding thickness/diameter • 0.041-0.043 "§ 88- ^ 0.055 • 0.066 cu
I 87£
S
O Breached pins due to FCMI • Other breached pins
H
86854
6 8 10 12 14 Peak burnup (at. % )
16
18
Figure 11-15. Breached fuel pin experience in the advanced oxide test series as a function of fuel smeared density and relative cladding thickness (after Lawrence et al., 1986).
pins fabricated using FFTF type cladding (20% cold-worked Type 316), high smeared density fuel ( > 90 % TD) and thin walled cladding (cladding thickness/diameter ratio of 0.043). A contributing factor to some of the breaches was the low O/M of the fuel (1.92 to 1.97). The increased fuel thermal expansion behavior of low O/M fuel may have contributed to the FCMI. Additional information on the effects of FCMI were obtained from profilometry data. Step changes in cladding strain were found at the bottom ends of the fuel columns increasing with fuel smeared density. The strain increases were consistent with a solid fission product volumetric swelling rate of 0.3 % per percent burnup. The results of the Advanced Oxide test series clearly demonstrated the effects of FCMI on fuel pin behavior. The work indicated that fuel pins with low swelling cladding would be expected to experience FCMI at lower exposures than corresponding fuel pin designs with higher swelling cladding materials.
During the early 1980's, in support of the design and licensing of CRBR (Clinch River breeder reactor), major effort was concentrated on the development of cladding breach criteria. Ex-reactor and in-reactor stress-rupture data were evaluated and were compared to the behavior of breached mixed oxide fuel pins (Biancheria and Brizes, 1983). The conclusions from this effort are summarized below: - Stress-rupture testing (using pressurized tubes) in the laboratory and in-reactor yielded similar results. Thus, the net effect of neutron irradiation (including swelling, irradiation creep and addition of helium via n,oc reactions) on stressrupture behavior is small. - In contrast, stress-rupture tests carried out in a hot cell after irradiation or in thermal reactors often showed significant degradation of stress rupture capability compared to unirradiated material. These results are not considered to be applicable to fuel pins operating in fast reactors. - Analysis of the available cladding breach data in the mid-80's indicated that a methodology based on cumulative damage fraction (CDF) using stressrupture data provided a useful approach for correlating cladding breaches whereas evaluations based on cladding mechanical strain were less satisfactory. - Using the CDF methodology, an apparent difference was found between the breach probabilities for in- and ex-reactor stress-rupture tests and the breach of mixed oxide fuel pins irradiated in pin bundles. - Figure 11-16 shows the breach probability vs. CDF for stress-rupture tests (by definition a CDF of 1 corresponds to a 50% probability of breach) and cladding breaches in wire-wrapped assemblies irradiated in EBR-II. Cladding
11.3 Steady-State Mixed Oxide Fuel Performance
breach in wire-wrapped assemblies correspond to a 50% probability at a CDF of 0.2 to 0.25. Note that the calculation of CDF includes normal peaking factors and FCCI (fuel-cladding chemical interaction) effects which account for cladding thinning. - The differences between the unirradiated stress-rupture results and the behavior of wire-wrapped fuel pins were believed to be due to irradiation effects, particularly the effects of helium bubbles on stress-rupture behavior. - For design applications with wirewrapped assemblies (and for grid assemblies as noted in Section 11.3.4), the correlation B in Fig. 11-16 was recommended as an upper design limit based on the behavior of the small population of fuel pins operating of 2 a conditions. During the development of the cladding breach criteria referred to above, requirements for additional fuel pin performance data were identified (Boltax and Sackett, 1981). These requirements provided the 99-
50-
B Breaches in wire-wrapped bundles
Unirradiated sress-rupture data
1-
0.01
0.1 1.0 CDF (steady-state)
10.0
Figure 11-16. Breach probability vs. CDF for stainless steel cladding.
359
basis for a collaborative DOE/PNC program on operational transient testing of mixed oxide fuel in EBR-II. The objectives of this program were numerous including providing input for reactor licensing, calibration of fuel pin design codes (LIFE and CEDAR), development of improved fuel pin design criteria as well as other objectives related to fuel and cladding development. Relative to cladding breach criteria, the following aspects of the DOE/PNC program are pertinent (Boltax et al, 1985): - Two types of irradiation test vehicles were used in the program. Wirewrapped fuel pins were irradiated in pin bundles and in assemblies made up of individual shroud tubes. The shroud tubes were thermally insulated from one another by a gas gap and each shroud tube was individually orificed to provide the desired cladding temperature. - The objective of the two types of fuel pin tests (pin bundles vs. individual pins in shroud tubes) was to determine the lifetime capability of individual wirewrapped fuel pins, thermally and mechanically insulated from neighboring pins vs. that of wire-wrapped pin bundles. Any significant difference in lifetime, presumably a shorter lifetime in a pin bundle, could then be attributed to pin bundle effects. - The duty cycle operating conditions included: steady-state operation, steadystate plus periodic (15 %) overpower operation and alternate EBR-II cycles at 100 and 60% power and flow. - The fuel pin designs included three cladding types, two pin diameters and a range of design and operating conditions covering aggressive, moderate and conservative conditions. - The aggressively designed pins (high fuel smeared density, high linear power and high cladding temperatures) were ex-
360
11 Mixed Oxide Fuel Pin Performance
pected to breach at low burnup (< 5%) based on the stress rupture criteria given by the B correlation in Fig. 11-16. The Phase I portion of the DOE/PNC operational transient program involved irradiation of over 100 mixed oxide fuel pins to peak burhups of 10% and fast neutron dose levels of 8 x 10 22 n/cm2 (Boltax et al., 1990 b). Figure 11-17 shows examples of the power histories for aggressively designed fuel pins, none of which breached. Figure 11-18 shows the results obtained on an aggressive fuel pin with 20% coldworked D9 cladding (see Table 11-3 for cladding composition). The 7.0 mm diameter fuel pin contained fuel with a 90 % TD smeared density and operated at a time-averaged linear power of 47 kW/m with peak ID cladding temperatures of 655 °C to a peak burnup of 7.2 % and a fast neutron dose of 6.3 x 10 22 n/cm 2 . LIFE code calculations for several of the DOE/PNC fuel pins are given in Table 11-5. Measured and calculated cladding strains (there was essentially no swelling of the cladding) were in generally good agreement with the exception of the moderate D 9 clad fuel pins operating with alternate cycles at 100 and 60% power and flow. In this latter case, enhanced Cs migration occurred in the high temperature region of the pins due to the combination of high power and moderate fuel smeared density (86.5% TD for the moderate pins vs. 89.1 % TD for the aggressive pins at similar power and cladding temperature). The key feature of Table 11-5 is the high CDF values calculated for unbreached fuel pins. Note that the two cases with CDF values above 1.0 were re-evaluated and sensitivity studies indicate that a CDF value near 0.5 is more appropriate. Section 11-5 provides additional details of these sensitivity studies. The important result obtained from these tests is that individu-
8 % burnup 6.5x1022 n/cm2
7 % burnup 6x1022 n/cm 2
TOP-7
2000
4000
6000 8000 Time,in h
8 % burnup 8x1022n/cm2 10000
12000
14000
Figure 11-17. Duty cycle power histories for mixed oxide fuel pins (after Boltax et al., 1990 b).
350 400
Figure 11-18. Cladding strain and cesium profile for a D9 clad fuel pin in the TOP-4A assembly (after Boltax etal., 1990 b).
361
11.3 Steady-State Mixed Oxide Fuel Performance
Table 11-5. Comparison of PIE measurements and LIFE code calculations (Phase I DOE/PNC operational transient program). Operating conditions
Steady-state
Steady state plus periodic overpower Alternating cycles
Pin type a
Burnup
Dose
%
10 22 n/cm 2
Measured
Calculated
D9,A PNC, A
7.0 7.4
6.1 6.4
1.8 0.5
1.5 0.45 e
1.25 (0.5)d 0.09e
D9f A
7.1
6.0
0.6
1.0
0.6
PNC, A
7.3
6.2
0.3
0.35
0.04
D9,A
8.4
8.0
1.0
1.15
0.5
C
Peak cladding strain, %
Peak CDF value
at 100 and 60%
D9,Mb
8.5
7.9
1.9
1.8
1.25 (0.5)d
power and flow
PNC, A
8.4
8.0
0.35
0.3
0.04
a
Cladding type and design, A = aggressive and M = moderate; b moderate pins with 86.5% TD full smeared density; c enhanced cladding strain model required to account for large Cs migration at peak strain region; d sensitivity studies including cladding temperature ( — 20 °C) and increased in-reactor creep at high cladding temperature (650 °C) indicate CDF values are 0.5 rather than 1.25;e the PNC cladding exhibits lower in-reactor creep and has higher stress-rupture capability than the D9 cladding.
ally tested fuel pins can obtain CDF values of 0.5 without cladding breach. In the Phase II portion of the DOE/ PNC operational transient program, fuel pin burnups were extended to 15% with a peak fast neutron dose of 1.5 x 10 23 n/cm2 (Boltax et al, 1991). Cladding breaches were noted for aggressive D 9 clad pins at about 11 % burnup with sibling pins exhibiting peak cladding strains of 4.5%. The calculated CDF values at the time of breach were close to 1.0. In contrast, aggressive fuel pins with PNC cladding performed satisfactorily to 15 % burnup without cladding breach (Boltax et al., 1992). Figure 11-19 shows the currently available data for the Phase I and II irradiation tests. Significant differences in behavior of the two claddings are evident at high burnup. The average cladding strain-rate for the PNC clad fuel pins is 75% of that for the D 9 clad pins.
i
r
i
r
D9 Avg. slope = 0.8%/% burnup
I PNC 316 Avg. slope = 0.6%/% burnup
Phase I : H 4
6 8
t
10
Phase I I 12 14 Burnup, %
16
18
20
Figure 11-19. Cladding strains vs. burnup for single pin tests (Boltax et al., 1991).
362
11 Mixed Oxide Fuel Pin Performance
The similarity in stress-rupture behavior of unirradiated cladding tubes and individually shrouded fuel pins (Curve A in Fig. 11-16) indicates that the factor of 4 to 5 reduction in stress-rupture lifetime, shown by Curve B in Fig. 11-16, is not due to irradiation effects (e.g., helium bubbles on grain boundaries) but instead is due to unaccounted for differences in the behavior of individual pins and bundles of pins. The unaccounted for effects in pin bundles are believed to be associated with local regions with abnormal cladding hot spots or mechanical interaction. Accordingly, Curve B in Fig. 11-16 should be used when calculating the lifetime of pin bundles, whereas Curve A describes the behavior of individual fuel pins. The absence of a significant effect of irradiation on the stress-rupture behavior of D9 cladding is generally consistent with the EBR-II and FFTF irradiation data reported by Lovell et al. (1981) and Puigh and Hamilton (1987), respectively. This result and the major differences in the burnup capability (Fig. 11-19) of mixed oxide fuel with D 9 and PNC cladding (Boltax et
0.245 •
38.1 I
I
al., 1991) do not agree with the conclusions of Wassilew et al. (1987). Wassilew et al. (1987) believe that the in-reactor rupture life of austenitic steels is controlled by helium behavior and is independent of alloy type. 11.3.3 Prototype Fuel Behavior
Prototype fuel performance data from FFTF (Makenas, 1987), Phenix (Ratier et al., 1986) and PFR (Swanson et al., 1986) are generally consistent with each other and with data obtained from test reactors. Figure 11-20 shows the evolution of cladding dimensional changes for FFTF driver fuel pins during the first three operating cycles. At a peak burnup of 8.8% and a fast neutron dose of 1.3 x 10 23 n/cm2, the peak diameter increase was 6%. Immersion density data confirmed that the diametral increases were primarily due to cladding swelling. Figure 11-21 shows examples of FFTF data (Makenas, 1987; Makenas et al., 1991) for CW316 and CWD9 clad fuel pins. For the high swelling CW316 clad-
Axial location in millimeters from bottom of fuel 190.5 342.9 495.3 647.7 I I I j I I 1
6.22
0.241 -
-6.12
0.237 ~
-6.02
i
-5.92
. 1 0.233 •
3 0.229
-1.5
I
I
4.5
7.5
r 10.5
I
I
I
13.5
16.5
19.5
I 22.5
25.5
I
I
28.5
31.5
I 34.5
5.82
37.5
Axial location fn inches from bottom of fuel
Figure 11-20. Diameter change vs. axial position for 20% CW 316 SS clad FFTF driver pins (after Makenas, 1987).
11.3 Steady-State Mixed Oxide Fuel Performance
-38.1
0.26-
0.25 -
114.3
266.7 i
419.1
Distance from bottom of fuel (mm) 571.5 723.9 876.3 -38.1 114.3 266.7
316 SS fuel pins
419.1
571.5
723.9
363
876.3
-6.60
-635
16.1x1022 n/cm2 (£>0.1MeV)
e
-6.09
S 0.23
5.84 22.5 28.5 34.5 -1.5 Distance from bottom of fuel (in.)
Figure 11-21. Diameter measurements from CW 316 SS and CW D9 clad fuel pins irradiated in FFTF (after Makenas et al, 1991).
ding, irradiation in FFTF is limited to 10 to 1 1 % burnup due to pin bundle and duct distortion and, occasionally, cladding breach (Washburn and Weber, 1986). The moderate swelling CWD9 cladding permits higher burnup (~18%), but it too suffers burnup limitations due to pin bundle distortion. A key test of mixed oxide fuel pins with CWD9 cladding involved the C-l test in
FFTF with relatively high late-in-life cladding temperature (650 °C). This test reached a burnup over 11 % without breach (Baker et al., 1991). Figure 11-22 shows the results of diametral and gamma scan measurements. Similar to that reported previously in Fig. 11-18, cladding strain and Cs concentrations are noted at the high temperature end of the fuel pin. The cladding strain at the top ends of the C-l
o-
Figure 11-22. Diameter change and cesium gamma scan from high temperature C-l fuel pin (after Baker et al., 1991).
364
11 Mixed Oxide Fuel Pin Performance
fuel pins has not been observed in other related tests in FFTF using CWD9 at lower cladding temperatures. This behavior is believed to be related to the combined effects of high cladding temperature and fuel-cesium reactions. Figures 11-23 and 11-24 show examples of European data for mixed oxide fuel irradiated in Phenix (Dupuoy, 1982) and PFR (Swanson et al., 1986). The US and European fuel pin experience followed similar evolutionary patterns with initial driver fuel showing excessive swelling and replacement fuel clad with 20 % cold-worked Type 316 steel with titanium additions, showing moderate swelling. The titanium modified Type 316 steels permit peak burnups and doses of 12% and 2 x 10 23 n/cm2 (100 dpa), respectively. Further improvements in fuel performance, using low swelling second generation cladding materials, are described in Section 11.3.5. Additional information on the behavior of prototype mixed oxide fuel as a function of burnup is shown in Figs. 11-25 and 11-26 (Hales and Baker, 1986). The effect of burnup on fuel microstructure is manifested through a maturing of the restructured fuel zones and the continued growth of the central hole. The fuel grows radially, as the cladding swells, to maintain thermal contact between fuel and cladding. The radial fuel and central hole growth are due to the combined effects of fuel swelling, due to solid and gaseous fission products, and the migration of fission gas bubbles and porosity to the central void. The expansion of radial fuel cracks near the fuel-cladding interface with increasing burnup (Fig. 11-25) provides evidence for fuel swelling in the hotter interior fuel regions. Fission gas release data (Fig. 11-27) are in general agreement with predictions based on EBR-II experience. Some differences were noted for KM (Kerr-McGee) vs.
1976 (316 S.A)
7 6-
54-
3-
2-
1-
I
500
1000
I 1200
I 1400
i
I m.
1600 mm axis
Figure 11-23. Evolution of Phenix mixed oxide fuel performance at ~10% burnup and 1.5 x 1023 n/cm2 (after Dupouy, 1982).
Peak burnup/dose 10.7 % / 7 9 dpa NRT Peak burnup/dose 7 . 9 % / 5 7 dpa NRT Peak burnup/dose 6 . 3 % / 4 5 dpa NRT
Upper breeder
Fuel column
Lower breeder
Plenum
Figure 11-24. Diameter swelling profiles for PFR fuel pins clad in CW.M316 (after Swanson et al., 1986).
B & W (Babcock and Wilcox) fuels related to minor differences in as-fabricated fuel structures. At high burnup ( > 8 % ) , essentially 100% of the fission gas produced was released. The Japanese reported on an interesting high burnup test of mixed oxide fuel irradiated in Phenix (Kashihara et al., 1990). The test involved modified (P and B addi-
365
11.3 Steady-State Mixed Oxide Fuel Performance
KerrMcGee
20
10
30 Peak burnup ,MWd/kg M
Figure 11-25. FFTF fuel restructuring as a function of burnup (after Hales and Baker, 1986). KM: Kerr-McGee. B & W: Babcock and Wilcox. Predicted Measu Equiaxed o Columnar grain • Central void O Distance from bottom of fuel, mm 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1002.5 Kerr-McGee Q o Equiaxed fuel pm 80>2.0
Columnar grain \
"\
^ 40H
300Measured gas release
s:260H -1.5 -1.0
20-/
100-
0
•
I
I
B+W fuel pin
I
1
Q
J
L__J
I
L-2.5
D Equiaxed p
^
-500 %
0.5 4 8 12 16 20 24 28 32 36 Distance from bottom of f u e l , inches Distance from bottom of fuel,mm 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9
-600
K-M
E2204
1800
-650 .51
O B +W
,1*0-
Total fission gas / generated / 100 % release /
- 4 0 0 175 o
100 -
-300 I
60-
-200 I
20-
-1001
80-
2() 0
4 8 12 16 20 24 28 32 Distance from bottom of fuel,inches
36
Figure 11-26. Comparison of predicted and measured FFTF fuel restructuring at lOOMWd/kgM burnup (after Hales and Baker, 1986).
40 60 M 100 110 Peak burnup (MWd/kgM)
Figure 11-27. Volume of fission gas released from FFTF driver fuel as a function of burnup (after Hales and Baker, 1986) (STP: standard temperature and pressure; 1 psi(g) = 6.895 • 103 Pa).
366
11 Mixed Oxide Fuel Pin Performance
tions) 20% CW 316 cladding, low (1.95) and high (1.98) O/M fuel and low (85% TD) and high (93% TD) pellet densities (corresponding to fuel smeared densities of 80.2 and 87.8% TD respectively). The peak burnup and fast neutron fluence were
Diametral change 105000 MWd/t 1.9x10 2 2 n/cm 2
Diametral change
11.5% and 1.9 x 10 23 n/cm2, respectively. The maximum linear power was 45 kW/m with a peak mid-wall cladding temperature of 650 °C. Figure 11-28 shows the diametral and swelling profiles for one of the fuel pins (Mizoo et al., 1989). Figures 11-29 to 11-31 show the effects of O/M on fission gas release, fuel swelling and depth of cladding inner corrosion. The lower O/M fuel offers advantages with respect to fuel swelling and cladding corrosion. Figure 11-32 shows that the creep strain was not sensitive to the difference in
Z 2.01.0-
100 Bottom
Axial location
TOP
O/M CI adding temperature (°C) - 6 0 0 --550 " 1 9 8 - 1.99 — •-mr1 9 5 - 1.96 — s
75
Figure 11-28. Measurement of a PNC mixed oxide fuel pin irradiated in Phenix (after Mizoo et al, 1989),
m
m
/ 50
•
/ 5100"
t
25
80"
D O
o
D r\
25
S 6040.2 20-
6
Symbol • O • D
0/M Density(%TD) 1.98 93 1.95 93 1.98 85 1.95 85
50 75 100 125 Local burnup (x1Q3 MWd/t)
150
Figure 11-31. Maximum depth of cladding inner corrosion as a function of local burnup (after Kashihara etal., 1990).
10 20 30 40 50 60 70 80 90 100 110 Pin-average burnupHO3 MWd/t)
1.2
Figure 1.1-29. Fission gas release rate measured by pin puncturing tests (after Kashihara et al., 1990).
1.0
Symbol Density • 93%TD O 85%TD
E 0.8 c ro
Symbol 0/M Density (%TD) • 1.98 93 O 1.95 93
Jt 0.6 to
EX
\
*
0.2 h
20 40 60 80 100 Local burnup(x 103 M W d / t )
Figure 11-30. Relation between fuel swelling rate and fuel O/M ratio (after Kashihara et al., 1990).
•
I 0/M=1.98 |
0
10 20 30 40 50 60 70 Distance from core bottom (cm)
Figure 11-32. Creep strain of cladding with different pellet densities (after Kashihara et al., 1990).
367
11.3 Steady-State Mixed Oxide Fuel Performance
fuel smeared density. This result is not unexpected considering the relatively large total cladding strain shown in Fig. 11-28. Approximately 75% of the cladding strain is due to void swelling. 11.3.4 Cladding Breach in Prototype Reactor Fuel
During the first ten years of FFTF operation, only one driver fuel assembly (DFA) sustained a cladding breach, however, several breaches occurred in experimental fuel assemblies (Baker et al., 1991). To obtain an estimate of the overall reliability of the FFTF DFAs, data from experimental assemblies taken to higher burnup were used. This resulted in a data base of over 4300 fuel pins (20 assemblies) at burnups above 11%. Using this data base, the fuel pin reliability was established as a function of peak burnup as shown in Fig. 11-33. 0.1 1.0
-
10.0
-
50.0
-
No. of pins = 4342 No. of breaches = 5
AAD-6
90.0 95.0
The results indicate an FFTF driver pin reliability of 99.99% at a peak burnup of 11%. A detailed analysis was conducted of one of the five breaches noted in Fig. 11-33. The RTCB-4 breach occurred in a driver fuel pin irradiated in a grid assembly at 11.5% burnup and 1.6 x 1023 n/cm2 (Levine et al., 1986). Measurements taken during the final irradiation cycle showed that the coolant outlet temperature of the RTCB-4 assembly was 25 to 30 °C higher than anticipated, indicative of coolant flow reductions due to cladding diametral growth. Table 11-6 shows the changes in flow rate and coolant outlet temperature during the final irradiation cycle. Figure 11-34 shows the CDF analysis of the RTCB-4 fuel pins operating at peak and overheated conditions. For the initial 7000 hours of operation, the fuel-cladding gap was probably closed with a very low level of FCMI. Between 7000 to 8000 hours, the gap would be expected to open slightly due to the void swelling of the cladding. For the peak pins, the CDF values would be expected to level off at 0.05 to 0.06, well below the level required to cause breach. However, the peak overheated pins would be expected to attain CDF values of 0.13 to 0.15. For a breach criteria based on 50 % probability of breach at a CDF of 0.25 (Fig. 11-16) for wire-wrapped
'•g 99.0 99.5
Table 11-6. Measured RTCB-4 assembly coolant outlet flowrates and temperatures.
RTCB-4 99.90-
Driver
99.95-
Irradiation time in days
I MW-4 98.0 MWd/kgM
99.99-
V+DE-9 100
120 140 Burnup (MWd/kgM)
160
Figure 11-33. FFTF driver fuel reliability (after Baker et al., 1991).
335 365 448 468
Measured outlet flowrate in 1/min
Measured outlet coolant temperature in°C
1620 1616 1563 1480
560 559 558 564
368
11 Mixed Oxide Fuel Pin Performance
Table 11-7. Low swelling cladding and duct alloys. type 316 x/L =0.9 0.10-
30°C Overheating
Breeder program
Gap —-ClosedlOpen-
US Peak
European
I Open — Gap
Japan 0.018 10 Time, 1000 h
12
14
a
Cladding
Duct
HT-9
HT-9
PE-16 15.15Tia
FV448 EM-10
PNC 1520 a FMS ODS
FMS
20% cold worked.
Figure 11-34. Effect of overheating of RTCB-4 fuel pins on CDF values (after Levine et al., 1986).
or grid-spaced fuel assemblies, CDF values of 0.13 to 0.15 are consistent with one breached pin if 20 fuel pins were operating at peak conditions. The CDF analysis of the RTCB-4 breach is, in general, consistent with the breach experience obtained from EBR-II fuel assemblies. However, the critical experimental data and associated analysis have not yet been completed. It is anticipated that the cladding breach experience on second generation cladding materials (Section 11.3.4) will be required to provide a breach data base free of complications related to excessive void swelling. The cladding breach experience of driver fuel in Phenix and PFR is generally similar to that of FFTF (Leclere et al., 1986). As of 1986, only two breaches of DFAs occurred in Phenix with over 130000 pins irradiated with some at 15% burnup. By 1990, Phenix sustained 12 DFA breaches and PFR sustained a similar number (Languille et al., 1990).
reactors (Bleiberg and Bennett, 1977; Brager and Perrin, 1982). By the start of the 90s, the list of candidates had been reduced considerably with major attention focussed on the alloys noted in Table 11-7. Table 11-8 lists the compositions of the current group of candidate alloys. The alloys can be grouped in three categories: Advanced austenitic - 15/15 Ti (1.4970) - PNC 1520 High nickel alloy1 - PE16 Ferritic/martensitic - HT-9 - FMS - FV 448 (1.4914) - EM-10 - ODS (oxide dispersion strengthened) In the following discussion, examples of mixed oxide fuel pin performance data with second generation cladding will be described. Additional discussion of the irradiation behavior of second generation alloys can be found in Chapter 5.
11.3.5 Development of Second Generation Liquid Metal Reactor Fuel
In the early 1980s, a large number of candidate low swelling cladding and duct alloys were tested in small and prototype
1 PNC is also developing high nickel alloys with 25 and 40% Ni. The PNC alloys contain Nb and V nitride for precipitation strengthening.
369
11.3 Steady-State Mixed Oxide Fuel Performance
Table 11-8. Second generation cladding and duct alloys. Element
PNC 1520 15.15 Ti
Cr Ni Mo Mn Si C Ti Nb + Ta P B S V
w
Y2O3 Al Country a Applications a
b
PE-16
HT-9
FMS
EM-10
15 20 2.5 1.7 0.4 0.06 0.25 0.1 0.025 0.004 0.001 0.01 -
15 15 1.2 1.5 0.4 0.1 0.5 < 0.015 0.005 -
17 44 3.3 0.1 0.2 0.07 1.3 0.001 0.1) 650 °C 480 W/cm CW 361 Ti
20000 _
10000
-
Figure 11-42. Behavior of an axially heterogeneous mixed oxide fuel pin in Phenix (after Boidron et al., 1986). Distance from pin base in cm
tive accidents (HCDA). Furthermore, reactor safety testing involves exploration of the post-breach fuel behavior during severe accidents. The following discussion will be limited to the topic of operational transient testing. A recent conference on fast reactor safety (ANS, 1990) provides an excellent status report on safety related transient irradiation tests. The requirements for operational transient testing derive directly from project needs (such as FFTF, CRBR, PRISM, etc.). However, in the following discussion, generic requirements will be presented, recognizing that significant overlap exists
among the specific reactor designs. Accordingly, the general objectives of an operational transient program are to (Boltax and Sackett, 1981): - Demonstrate the ability of a core component to meet the most challenging events and event sequences permitted by the plant protection system (PPS). - Establish the margin between component failure and the capability of the PPS for a range of events and event sequences. - Confirm the validity of design procedures (criteria, analytical methods, etc.) to describe component damage over the
11 Mixed Oxide Fuel Pin Performance
376
design operating range (normal operation and operational transients). The US operational transient testing program has been conducted in the TREAT and EBR-II facilities. TREAT testing of mixed oxide fuels has been in progress for about 30 years with most of the testing devoted to safety related matters (Wright et al., 1990). Operational transient testing of mixed oxide fuel in EBR-II was conducted intermittently in the 1960s and 1970s and more significantly in the 1980s with the initiation of the DOE/PNC Operational Transient Testing Program (Tsai et al., 1985; Boltax et al., 1985). The operational transient testing in TREAT was performed on fuel pins previously irradiated in EBR-II and FFTF. The EBR-II/TREAT tests on mixed oxide fuel pins with 20% cold-worked Type 316 stainless cladding involved 21 single pin capsule tests and a series of loop tests (Weber et al., 1981). Figure 11-43 provides a summary of the results in terms of the margin to failure as a function of ramp rate. For the fast ramp rate transients (50 0/s and 3 $/s), a substantial overpower margin to failure was demonstrated. At lower ramp rates ( < 10 0/s which corresponds to
iou~ -500
160-
Path to events
3$/s
140-
\100-
j
O
I
•
PPS termination Failure threshold
-400 -300^
* 80-200
604020-
f
s
/ 50(/s 3
-100
$/ s
o
.
A
0
2
4 6 Time .seconds
8
10
Figure 11-43. Mixed oxide fuel pins: Margin between PPS terminated transients and failure thresholds (after Weber et al., 1981).
10000
=
1000
I Life code calculations - 1 TREAT dara
/
/ /m/
7
-
100 -
1
I. - 9 0 % TD), the LIFE-4 code tends to slightly underpredict mechanical strain. 11.5.3 LIFE Code Predictions Selected examples of LIFE-4 (Rev. 1) code predictions will be presented. The examples will include evaluations of fuel pin
)
56 54
/
52
0
A
50
0 0 i°
48 00
46
°A
000
44 0
- 42 40
V
38 36
383
/0 36
/°
/
0
0 38
40
42 44 46 48 50 52 Measured Q (melt) in kW/m
54
56
58
Figure 11-50. Linear-residual-corrected computed powers to melt as a function of measured power to melt (after Sundquist, 1990).
384
11 Mixed Oxide Fuel Pin Performance
0.50 = midplane (fuel pin) 1 = blanket pin 2 = bottom sect, (fuel pin) 3 = top axial sect.(fuel pin)
0.4-
0.3-
Figure 11-51. Error in predicted mechanical diametral strain as a function of burnup (after Sundquist, 1990). Burnup in %
A
0 1 2 3
/
0.H-
3
03
= = = =
2
midpla ne (fuel pin) blanks t pin bottor n sect, fuel pin) top a)(ial sec Ufuel pin)
0 0
A7
3 0 3
—,
o
~
~
—
•
—
_
2
-0 1
3
.oo3
