The a-decay of the actinides taking place in the crystal lattice creates an alpha particle and a recoil atom. The recoil atom produced has a range of about 12 nm and causes a dense collision cascade with typi­cally about 2300 displacements (Frenkel pairs) within a short distance, around 7.5 nm in size. The a-particle has a path of about 10 pm, with a cascade of about 265 displacements at the end of its range.1 Although recombination will take place, point defects and eventually extended defects (dislocations, dislocation loops) will survive in the crystal lattice, resulting in changes in the properties of the materials. Computer simulations of the radiation effects in fcc plutonium have shown that the defect recombination stage is much longer than that in other metals and that the vacancies do not seem to form clusters.19 In addition to the radiation damage, helium ingrowth takes place.

As discussed by Hecker and Martz,20 the expan­sion of the lattice of a-Pu is significant due to

2500

2000

2Г

1500 l-T

1000 500

self-irradiation, when held at cryogenic tempera­tures, saturating at about 10vol.%. In contrast, the (Ti-stabilized) p-phase shows a slight contraction and the (Al-stabilized) 8-phase a substantial contrac­tion, the latter saturating at 15 vol.%. Of course this is also reflected in other properties such as electrical resistivity.21,22 The radiation effects recover upon annealing to room temperature, a few percent of the damage remaining. Gorbunov and Seleznev23 observed that a-Pu containing predominantly 239Pu retains its crystal structure after prolonged storage at room temperature. A sample of predominantly shorter lived 238Pu (t1/2 = 87.74 years) contains both the a — and p-forms at immediate examination and additionally the g-, Z-, and e-phases after a similar storage period. Chung et a/.24 showed by X-ray diffraction and dilatometry measurements on 238Pu-doped 8-phase plutonium samples that the lattice expansion by self-irradiation appears to be the primary cause for dimensional changes during
the initial 23 years of aging. Following the initial transient, the density change is primarily caused by a constant helium ingrowth rate as a result of particle decay. The two effects were combined in an equation for the expansion AL/L with an exponential (radia­tion damage) and a linear (helium ingrowth) part:

AL/L ffi A[1 — exp(-Bt)] + Ct [1]

where A, B, and C are constants and t is time.

The self-irradiation is one of the main causes that complicates the study of the heavy actinide metals. For example, berkelium metal (t^2 = 314 days; ^0.2% 249Cf growth per day) shows signs of amorphization (weak and diffuse X-ray spectra) at room temperature, which improved after annealing and thermal cycling, and the samples were found to contain two crystallo­graphic structures at room temperature, double hexag­onal close-packed (dhcp) and fcc, ofwhich the former is the stable form.25 An extreme case is Es; its crystal structure has been resolved only by rapid electron diffraction of thin film material due to the very short half-life of the isotope used.26

The cation diffusion occurs in actinide dioxides via cation vacancies in the vicinity of oxygen vacancy according to Jackson eta/.234 As a consequence, cation diffusion may occur in hyperstoichiometric dioxides MO2 + x (with x > 0). The diffusion coefficient increases drastically with departure from stoichiom­etry as x2 in MO2 + x as follows:

D(x, T)=A, x2exp^—[15]

Recently, Gao et a/.235 obtained the prefactor D0 = 2.341 x 10-2m2s-1 and migration energy E0 = 2.5 eV for UO2 + x.

Knorr et a/.236 mentioned that the effect of the intergranular grain boundary is to change the pre­factor to a linear dependence upon stoichiometry. The diffusion coefficient is expressed as follows:

D = Dox ex^kE^^ [16]

They found for UO2 + x, D0 = 7.5 x 10-^m2s—1 and E0 = 2.47 eV.236

This apparent activation energy is very close to the one obtained by Gao et a/.235 A systematic study of the difference between monocrystals and polycrystals done by Sabioni et a/.220 led the authors to the conclusion that most of the activation ener­gies reported by previous authors probably refer to polycrystals.

2.04.4.2.5.1 Uranium monocarbide UC

Freshly cleaved UC is bright gray with a metallic shine. It promptly darkens in contact with oxygen, due to the formation of a thin oxide layer on the surface.

Bober eta/.166 studied the spectral reflectivity p of liquid UC, using a laser sphere reflectometer with a polarized laser beam at different angles and wave­lengths (458, 514, 647, and 752 nm). The refractive index n and absorption constant к were obtained. n resulted to be around 2 at 458 and 514 nm, and 2.5 at 647 nm, slightly decreasing with temperature in both cases, up to1.7 at 4100 K and 458 nm and 2.1 at 647 nm. The value of n at 752 nm is 1.7, independent of temperature. The same trend was observed for к, which takes the value 2.5 at 458 and 514 nm and 3.1 at 647 nm, decreasing with T whereas к = 2.5 at 752 nm is independent of T

The normal spectral emissivity e2 of UC has been investigated at 650 nm in polycrystalline sam­ples. Results are certainly affected by oxidation of the sample surface. The most complete trend is the one proposed by Bober et a/.,167,168 based on reflectivity measurements on 96.5% th. d. UC (with 0.36 wt% O and 0.02 wt% N) between 300 and 2780 K and in liquid UC between 2780 and 4200 K:

e650 = 0.566 — 2.7209 x 10-6T + 2.7697 x 10-9 T2

— 2.7102 x 10-12T3 + 2.8618 x 10-16T4 [41]

for 300 K < T< 2780 K, and

e650 = 0.45 2 — 4.3 2 47 x 10-6(T — 2780)

+ 3.1967 x 10-9(T — 2780)2

— 1.6784 x 10-12(T — 2780)3

— 4.6641 x 10-16(T — 2780)4 [42]

for 2780 K < T< 4200 K. An experimental error of ±5% should be taken into account, leading to uncertainty bands larger than the proposed emissiv — ity variation as a function of temperature. Even considering those uncertainty bands, a marked dis­continuity of about 0.1 in ei upon melting remains clear. Such a gap is probably dependent on the oxygen-impurity content.169 For practical purposes, it is reasonable to assume for solid UC e650 = 0.55 ± 0.02 and for liquid UC e650 = 0.45 ± 0.02. De Bruycker169 provided the following wavelength depen­dence of solid and liquid UC spectral emissivity for 488 nm < 1 < 900 nm:

Solid UC: £i = 0.75746 — 0.467911 + 0.184912 [43]

Liquid UC: sx = 0.79998 — 0.755451 + 0.3903612 [44]

De Coninck et a/.170 measured the total hemispheri­cal emissivity eT of nearly stoichiometric UC, obtaining eT = 0.45 between 1400 and 2100 K.

2.04.4.2.5.2 Uranium sesquicarbide U2C3

Freshly cleaved uranium sesquicarbide is bright gray with a metallic luster. U2C3 optical functions have never been studied so far.

2.04.4.2.5.3 Uranium dicarbide UC2

a-UC2-x is bright gray with a metallic shine and darkens upon oxidation. The normal spectral emis — sivity of a — and p-UC2-x at 650 and 2300nm was

studied by Grossman171 and by De Croninck eta/.152 The results of this latter work, in agreement with Grossman’s, are shown in Figure 19. The total hemi­spherical emissivity was estimated to be around 0.55 between 1300 and 2350 K. These results suggest that the emissivity of UC2_x does not detectably vary in the visible range upon the a! p transformation, whereas it varies very little in the infrared range.

2.06.3.3.2.1 Preparation

U2F9 (which is called uranium enneafluoride) is prepared by treating UF4 with UF6. The kinetics are faster (about 1 day) when the process is per­formed under higher pressure of UF6 and high temperature. Another way is to start from the penta- fluoride and decompose it into U2F9. It is less conve­nient because it includes the synthesis of pure pentafluoride which is less stable. U2F9 is the more stable intermediate fluoride under argon atmosphere without moisture. It forms black needles.

U4F17 was obtained by maintaining UF4 at 593 K and introducing UF6 at 2.3 x 103 Pa pressure

during 2 days.59 The diffraction pattern of this material resembling that of UF4 but containing uniquely differ­ent features is described as ‘distorted UF4.’ It is a black powder. This synthesis is to be avoided due to the corrosion problem at high temperature under UF6. It is easier to decompose higher intermediate fluorides such as UF5 or U2F9 than synthesize directly.

We have found no physical properties for these compounds in the literature. Mixtures of UF5-U2F9 have a tendency to melt at lower temperature than pure UF5.

2.06.3.3.2.2 Crystal structure

U2F9 as a body-centered cubic structure, space group I43m — Td with a = 8.4716(5) /A (Figure 12) and a theoretical density pth = 7.06gcm—3 62,63 This sym­metry was confirmed by high-resolution neutron powder diffraction data.54

U4F17 has a deformed UF4 lattice. The symmetry is monoclinic, space group C6h, C2/c, with a = 12.09 (0.08) A, b = 10.81(2) A, c = 8.29(4) A, and b = 128.0(8).12

C. Gueneau, A. Chartier, and L. Van Brutzel

Commissariat a I’Energie Atomique, Gif-sur-Yvette, France

© 2012 Elsevier Ltd. All rights reserved.

2.02.1 Introduction

Owing to the wide range of oxidation states +2, +3, +4, +5, and +6 that can exist for the actinides, the chemistry of the actinide oxides is complex. The main known solid phases with different stoichiome­tries are shown in Table 1.

Actinide oxides mainly form sesquioxides and dioxides. The +3 oxides of actinides have the general formula M2O3, in which ‘M’ (for metal) is any of the actinide elements except thorium, protactinium, ura­nium, and neptunium; they form hexagonal, cubic, and/or monoclinic crystals.

Crystalline compounds with the +4 oxidation state exist for thorium, protactinium, uranium, nep­tunium, plutonium, americium, curium, berkelium, and californium. The dioxides MO2 are all isostruc­tural with the fluorite face-centered cubic (fcc)
structure. Most of these actinide compounds can be prepared in a dry state by igniting the metal itself, or one of its other compounds, in an atmosphere of oxygen. The stability of the dioxides decreases with the atomic number Z. All dioxides are hypostoichio — metric (MO2 _ x). Only uranium dioxide can become hyperstoichiometric (MO2 + x). The thermodynamic properties of the dioxides vary with both temperature and departure from the stoichiometry O/M = 2.

Only uranium, neptunium, and protactinium form oxide phases with oxygen/metal ratio >2. An oxida­tion state greater than +4 can exist in these phases. The +6 state exists for uranium and neptunium in UO3 and NpO3. Intermediate states are found in U4O9 and U3O8 arising from a mix of several oxida­tion states (+4, +5, +6).

Detailed information on the preparation of the binary oxides of the actinide elements can be found in the review by Haire and Eyring.

The absence of features at the Fermi level in the observed XPS spectra indicates that all the dioxides are semiconductors or insulators.2

Systematic investigations of the actinide oxides using first-principles calculations were very useful to explain the existing oxidation states ofthe different oxides in relation with their electronic structure. For example, Petit and coworkers3,4 clearly showed that the degree of oxidation of the actinide oxides is linked to the degree of f-electron localization. In the series from U to Cf, the nature of the f-electrons changes from delocalized in the early actinides to localized in the later actinides. Therefore, in the early actinides, the f-electrons are less bound to the actinide ions which can exist with valencies as high as +5 and +6 for uranium oxides, for example. In the series, the f-electrons become increasingly bound to the actinide ion, and for Cf only the +3 valency occurs. With the same method, Andersson et al.5 studied the oxidation thermodynamics of UO2, NpO2, and PuO2 within fluorite structures. The results show that UO2 exhibits strong negative energy of oxidation, while NpO2 is harder to oxidize and

PuO2 has a positive or slightly negative oxidation energy. As in Petit and coworkers,3’4 the authors showed that the degree of oxidation is related to the position of the 5f electrons relative to the 2p band. For PuO2, the overlap of 5f and 2p states suppresses oxidation. The presence of H2O can turn oxidation of PuO2 into an exothermic process. This explains clearly why hyperstoichiometric PuO2 + x phase is observed only in the presence of H2O or hydrolysis products.6

Solid actinide monoxides ‘MO’ were reported to exist for Th, Pu, and U. According to the experimen­tal characterization of plutonium oxide phases by Larson and Haschke,7 these phases are generally considered as metastable phases or as ternary phases easily stabilized by carbon or/and nitrogen. From first-principles calculations, Petit et a/.3 confirmed that the divalent configuration M2+ is never favored for the actinides except maybe for EsO. On the contrary, the monoxides of actinide MO(g) are stable as vapor species that are found together with other gas species M(g), MO2(g), MO3(g) which fraction depends on oxygen composition and temperature when heating actinide oxides.

In Sections 2.02.2 and 2.02.3, the phase diagrams of the actinide-oxygen systems, the crystal structure data, and the thermal expansion of the different oxide phases will be described. The related thermodynamic data on the compounds and the vaporization behavior of the actinide oxides will be presented in Sections 2.02.4 and 2.02.5. Finally, the transport properties (diffusion and thermal conductivity) and the thermal creep of the actinide oxides will be reviewed in Sections 2.02.6 and 2.02.7.

Some data on the Gibbs free energy of formation for actinide nitrides exist. In this section, the Gibbs free energy of formation of uranium nitride, UN, pluto­nium nitride, PuN, uranium and plutonium mixed nitride, (U, Pu)N, neptunium nitride, NpN, and americium nitride, AmN are summarized.

The values of AfG for UN were calculated from the values of AH2g8 and thermal functions for uranium and nitrogen.

Source: Hayes, S. L.; Thomas, J. K.; Peddicord, K. L. J. Nucl. Mater. 1990, 171,300-318; Matsui, T.; Ohse, R. W. High Temp. High Press. 1987, 19, 1-17; Cordfunke, E. H. P.; Konings, R. J. M.; Potter, P. E.; Prins, G.; Rand, M. H. Thermochemical Data for Reactor Materials and Fission Products; Elsevier: Amsterdam, 1990; p 667; Chase, M. W.; Curnutt, J. L.; Prophet, H. JANAF Thermochemical Tables; Dow Chemical Co.: Midland, MA, 1965.

Curium has a few long-lived isotopes, some of which are strong a-emitters (particularly 242Cm and 244Cm). Their formation in U-Pu nuclear fuel, there­fore, increases the radiotoxicity of the waste. How­ever, even after a few years of irradiation at burnups >100 GWdton—1, the total Cm concentration in MOX was observed to be very low, ^10—3 at.%.228 In this scenario, the rare literature studies on Cm carbides have an essentially academic profile.

The substantially covalent (a) Cm-C bonding in gaseous curium-carbon complexes obtained by laser metal-polymer coablation was studied in the late 1990s.229

More recently, Radchenko et a/.230 prepared the first samples of curium carbides by high-vacuum high-temperature condensation of metallic 244Cm onto an iridium support coated with amorphous car­bon. Cm2C3 and Cm3C, isostructural to Am2C3 and Sm3C, respectively, were identified by XRD. Cm2C3 has a bcc crystal lattice of the 143d space group with a = 839.04 ± 0.05 pm. Cm3C has fcc Fe4N-like lattice (already observed for some lanthanides) with a = 517.2 ± 0.2 pm.

Since no other carbides were detected at any carbon concentration, these authors concluded that Cm2C3 and Cm3C are the only existing Cm carbides.

2.07.4.1 Alloying Elements and Phase Diagrams

Like any metal, pure Zr exhibits rather poor engi­neering properties. To improve the properties of a given metal, the metallurgical engineering proce­dures are always the same: It consists in finding addi­tions, any species of the periodic table could be considered, with significant solubility, or heat treat­ments producing new phases that could improve the properties. The relative solubility ofthe various alloy­ing elements in the a — and р-phases is therefore one basis for the choice of additions, as well as for devel­oping the heat treatments, for microstructure control.

For the nuclear applications, neutron physics requirements restrict the possibilities, by rejection
of the isotopes having high interaction cross-sections, or isotopes that would transmute to isotopes of high capture cross-section or having high irradiation impact (Co). Elements such as Hf, Cd, W, and Co have therefore not been considered for alloy devel­opments. With low nuclear impact, O, Sn, and Nb have been selected (Al and Si having also low nuclear impacts were not retained because of degradation in corrosion resistance), while other TMs (Fe, Cr, Ni, etc.) can be accepted up to limited concentrations (below 0.5% total).

The additions have to improve the engineering properties. The main properties to be improved are the corrosion behavior in hot water and the mechan­ical strength (yield stress, ductility, and creep). As described below, Sn and Nb are added for corrosion resistance, and elements forming secondary phases (Nb and Fe, Cr, and Ni) or solid solutions are also used for increasing the mechanical properties.

Last, the microstructure obtained after the ther­momechanical processing should not change without control under irradiation. Therefore, hardening obtained by precipitation or strain hardening can be considered only if the irradiation-induced evolution of the initial microstructure will be compensated by the development of irradiation-induced microstruc­tural defects. In this respect, the evolution of preci­pitates in Zircaloys is of high importance for corrosion behavior and geometrical integrity. These points are discussed in Chapter 5.03, Corrosion of Zirconium Alloys and Chapter 4.01, Radiation Effects in Zirconium Alloys.

Most of the binary phase diagrams with Zr are already known and many ternary or higher-level
diagrams of industrial interest are now known.17 The need for a better control of the processing of the current alloys and the aim of finding new alloys and structures without too much experimental work have been a driving force for the modern trend in numeri­cal simulation for material science. It is now also possible to extrapolate the binary data to multicom­ponent systems. In that respect, a thermodynamic database for Zr alloys, called ZIRCOBASE, has been developed under the Calphad methodology.18 This database contains 15 elements and is frequently updated. The most complex ternary or quaternary phase diagrams available are optimized or computed using this database, and, in the case of missing basic thermodynamic data, with the contribution of ‘ab — initio’ computations.19 The phase diagrams presented in this review were obtained according to this procedure.

Oxygen is highly soluble in the a-phase, and stabi­lizes at high temperature (Figure 4). Oxygen has to be considered as an alloying element. This use of oxygen for strengthening is rare in metallurgy, com­pared to the use of nitrogen. However, the use of nitrogen for strengthening would severely deteriorate the corrosion resistance, and nitrogen is removed as much as possible. The purpose of oxygen additions is to increase the yield strength by solution strength­ening, without degradation of the corrosion resis­tance. The O content is not specified in the ASTM standards, but usually it is added to concentrations in the range of 600-1200 ppm, and this has to be agreed between producer and consumers. High O concen­trations (O > 2000 ppm) reduce the ductility of the alloys; therefore, O additions above 1500 ppm are not recommended. In addition, O atoms interact with the dislocations at moderate temperatures,
leading to age-strengthening phenomena in temper­ature ranges depending on strain rate.20 The oxygen in solid solution in a-zirconium is an interstitial in the octahedral sites. In the Zr-O system, the only available stable oxide is ZrO2. A monoclinic phase is stable at temperatures up to about 1200 °C, above which it transforms to a tetragonal structure. The impact on corrosion of the different phases of ZrO2, according to temperature and pressure is dis­cussed in Chapter 5.03, Corrosion of Zirconium Alloys.

Tin tends to extend the a-domain, and has a maximal solubility in the hcp Zr of 9 wt% at 940 °C (Figure 5). It was originally added at concentrations of 1.2-1.7% to increase the corrosion resistance, especially by mitigating the deleterious effects of nitrogen. The amount of Sn needed to compensate the effect of 300 ppm of N is about 1% of Sn. How­ever, in N-free Zr, Sn has been observed to deterio­rate the corrosion resistance. Therefore, the modern trend is to reduce it, but only slightly, in order to maintain good creep properties.21

Iron, chromium, and nickel, at their usual concentra­tions, are fully soluble in the р-phase (Figure 6). However, in the a-phase, their solubilities are very low: in the region of 120 ppm for Fe and 200 ppm for Cr at the maximum solubility temperature.22 In the pure binary systems, various phases are obtained: ZrFe2 and ZrCr2 are Laves phases with cubic or hexagonal structure, while Zr2Ni is a Zintl phase with a body-centered tetragonal C16 structure. These precipitates are called the Second Phase Par­ticles (SPPs).

In the Zircaloys, the Fe substitutes for the corresponding TM and the intermetallic compounds found in Zircaloy are Zr2(Ni, Fe) and Zr(Cr, Fe)2.

The formation of these precipitates, and more complex ones in industrial alloys, is analyzed in detail for the control of the corrosion behavior of the Zircaloys. Indeed, a strong correlation has been observed between precipitate size distributions and corrosion kinetics, the behavior being opposite for BWRs and PWRs. A better uniform corrosion resis­tance is obtained for Zircaloys used in PWRs if they contain large precipitates, while better resistance to the localized forms of corrosion is seen in BWRs in materials that have finely distributed small precipi­tates 23,24 With an increase in the particle diameters from 0.05 to 0.1 pm or higher, the in-pile corrosion of Zircaloy cladding diminishes appreciably. However, nodular corrosion may occur in BWR cladding with a further increase in the particle diameters above about 0.15 pm25 (Figure 7).

Due to the low solubility of the transition metals (Fe, Cr and Ni) in the Zr matrix, coarsening of the precipitates, after the last p-quench, occurs at very
low rates, during the intermediate annealing heat treatments, following each step of the rolling process. Therefore, the precipitate growth integrates the ther­mal activation times of each recovery, and their tem­peratures and durations can be used to control the size of SPPs. This integrated coarsening activation time is referred as the ‘A ’ or ‘SA ’ parameter.

The A-parameter calculates the integral of the activation processes for the different anneal durations and temperatures. The annealing parameter is defined as A = S, (t, exp(—Q/RT)), where t, is the time (in hours) of the ith annealing step, at tempera­ture T (in K); Q/T is the activation temperature of the process involved. The activation energy for the process should have been taken as the one controlling the coarsening, that is, the diffusion. However, as the early studies were undertaken with the aim of improving the corrosion resistance, an unfortunate practice has been induced to take 40 000 K as the value of Q/T. A more correct value would be

32 000 K, which fits very well with the recrystalliza­tion kinetics. The influence of the A-parameter on the corrosion of Zircaloy is discussed in more detail in Chapter 5.03, Corrosion of Zirconium Alloys. High resistance to uniform corrosion in PWR is obtained for the A-parameter close to (1.5—6.0) x 10~19h. In BWR, the A-parameter value for the Zircaloy-2 cladding in BWR has to be in the range (0.5-1.5) x 10~18h (Figure 7).25 This corresponds to precipitates larger than 0.18 pm. The SA approach has been developed for the Zircaloys and is clearly not applicable for other alloys, such as the Zr-Nb alloys.

Niobium (columbium) is a p-stabilizer that can extend the bcc domain to a complete solid solution between pure Zr and pure Nb at high temperatures (Figure 8). A monotectoid transformation occurs at about 620 °C and around 18.5 at.% Nb. The solubil­ity of Nb in the a-phase is maximal at the monotectic temperature, and reaches 0.65%.

Water p-quenching of small pieces leads to the precipitation of a’ martensite supersaturated in Nb. Tempering at intermediate temperature results in p-Nb precipitation within the a’ needles and subsequent transformation of a’ into a. When quenching is performed from an a + p region, a uniform distribution of a — and p-grains is obtained, and the Nb-rich p-phase does not trans­form. By aging at temperatures in the range of 500 °C, the metastable Nb-rich p-phase can be decomposed into an hcp ю-phase. This gives a sharp increase in mechanical strength because of the fine micro­structure obtained by the p-ю transformation.26 In the usual form of the Zr-2.5% Nb, the cold work condition after a + p extrusion and air-cooling, the microstructure consists of Zr grains with layers of p-Nb rich phase (close to eutectoid composition). Owing to the affinity of Fe for the p-phase, most of this element is found in the minor p-grains. These p-grains are metastable and decompose, upon aging, to a mixture of a-Zr and pure p-Nb. The Nb dis­solved in the a-hcp Zr phase is itself metastable and the irradiation-induced precipitation of the supersat­urated Nb solid solution is believed to be the origin of the improvement in corrosion resistance under irradiation of these alloys.27

In the case of Zr-1% Nb used for VVER and RBMK, or M5® in PWRs, the concentration of Nb in the Zr matrix after processing corresponds to the maximum solubility near the monotectoid tempera­ture, which is higher than the solubility at the service temperature. Owing to the slow diffusion of Nb, the equilibrium microstructure cannot be obtained ther­mally. However, the irradiation-enhanced diffusion allows precipitation of fine p-Nb needles in the grains after a few years in reactors.28

Sulfur has recently been observed to be extremely efficient in improving the creep resistance, even at concentration as low as 30-50 ppm. This chemical species, formerly not considered as important, is now deliberately added during processing to reduce the scatter in behavior and to improve the high tempera­ture mechanical properties.29 The efficiency of such low concentrations on the creep properties has been explained by the segregation of the S atoms in the core of the dislocations, changing their core configurations. It does not affect the corrosion properties.30

In the case of complex alloys, other thermody­namical interactions are expected and intermetallic compounds including three or four chemical ele­ments are observed. The chemistry and the crystal­lography of these phases may be rather complex.

200 nm

Two examples will be given of the complex structure

and behavior of these intermetallics.

• For the Zr-Cr and Zr-Ni binary alloys, the stable forms of the second phase are Zr2Ni or ZrCr2. These phases are effectively the ones observed in the Zircaloys, with Fe substituting for the corresponding TM. Therefore, the general formu­lae ofthe intermetallic compounds in Zircaloys are Zr2(Ni, Fe) and Zr(Cr, Fe)2. The crystal structure of the Zr(Cr, Fe)2 precipitates is either fcc (C15) or hcp (C14), depending on composition and heat treatment. Both structures are Laves phases, with characteristic stacking faults as seen in Figure 9. The equilibrium crystallographic structure is dependent upon the Fe/Cr ratio, cubic below 0.1 and above 0.9, and hexagonal in the middle. Under irradiation, these precipitates transform to amor­phous state and release their Fe in the matrix, with strong impact on corrosion behavior under irradiation.31

• In the Zr-Nb-Fe ternary, other intermetallic com­pounds can be observed (Figure 10): the hexago­nal Zr(Nb, Fe)2 phase and the cubic (Zr, Nb)4Fe2.32 Although of apparent similar composition, the two phases are indeed different: Nb can substitute Fe in the hexagonal phase, while it will substitute Zr in the cubic phase. In these alloys, due to the slow diffusion of Nb, metastable phases are often present and the equilibrium microstructure after industrial heat treatments may be far from the stable one. Therefore, the final microstructure is strongly dependent on the exact thermomechanical history.

In addition, the low solubility of these elements at operating temperatures drastically reduces the dif­fusion kinetics and requires more than a year to reach equilibrium at 450 ° C, in the absence of irradiation.3

Other minor constituents are often found in the form of precipitates. Among them are the carbide fcc — ZrC and silicides or phosphides of various stoichio­metries (Zr3Si, ZrSi2, ZrP, and Zr3P) that act as nucleation sites for the p! a-phase transformation during quenching and, therefore contribute to con­trol the a-platelets thickness and density.

The actinide sesquioxides can crystallize with three different forms: a hexagonal close-packed (a), a mono­clinic (b), or a cubic (c) structure. The hexagonal form is in most of the cases the stable phase at room
temperature. The cubic phase may be considered as a fluorite structure from which 1/4 of the oxygen ions have been removed. The crystal data on the actinide sesquioxides are listed in Table 6.

Few experimental data are available concerning the thermal expansion coefficients of actinide sesqui — oxides. The thermal expansion can be fitted with the following equation (in percentage):

DL

= a0 + a1 x T + a2 x T2 [9]

L0

Konings43 has extracted from experiments the ther­mal expansion of monoclinic B-Cm2O3. We have done the same for Pu2O3 (from Taylor76). In the case of Am2O3, we have used the data obtained by Uchida et a/.65 by MD calculations. A summary is available in Table 7.

2.02.3.2

Other Actinide Oxides

As mentioned in Section 2.02.3.1, the fluorite struc­ture of UO2 has empty octahedral sites that can be occupied by O2_ ions to form UO2 + x The phase diagram data show that the maximum oxygen content corresponds to x = 0.25 (or U4O9) (Figure 1). From his interpretation of neutron diffraction data on UO213, Willis100 found that the interstitials tend to
aggregate to form clusters made of oxygen intersti­tials interacting with normal oxygen anions.101,102

The so-called cluster 2:2:2 is composed of two oxygen vacancies and four interstitials. Below 1400 K, these clustered excess oxygens tend to form an ordered phase with the composition U2O9 _y.

U4O9 is a narrowly hypostoichiometric phase (U4O9 _y) and exists with three different forms:

a-U4O9 _^ (at T< 353 K), p-U4O9 _^ (at 353 K < T< 823 K), and g-U4O9 _(at 823 K < T< 1400 K). The structure of the p-U4O9 phase was studied by Bevan et at}03 who showed that this phase is a superlattice structure based on the fluorite structure of UO2 with a unit cell 64 times the volume of the UO2 cell. The additional O atoms are arranged in cuboctahedral clusters. According to the later analy­sis by Cooper and Willis,104 the centers of the clusters are unoccupied, whereas they are occupied by single O-ions according to Bevan eta/.103 U4O9 decomposes at 1400 K into UO2 + x (disordered with x ~ 0.25) and U3O8 (see Figure 1).

U3O8 is a mixed valence compound with U(V) and U(VI) cations. U3O8 exists in several forms as a function of temperature. At room temperature, a-U3O8 is orthorhombic and transforms to a pseudo­hexagonal structure p-U3O8 at 483 K. Heat capacity measurements by Inaba et a/.105 showed other phase transitions at 568 and 850 K.

UO3 can crystallize in six forms. The stable form at room temperature a-UO3 is orthorhombic.

Partial information on crystal data of plutonium and curium intermediate oxides with O/metal ratio below 2 is given in Table 8.

232Th, the only natural Th isotope, can absorb ther­mal neutrons to produce fissile 2 3U and is therefore used as fertile material in breeder reactors. Nowa­days, the thorium fuel cycle is mostly envisaged in India, which has about one-fourth of the total world thorium resources, but this option is kept open in other countries such as Norway and Australia, which also have abundant Th ores.33 Thorium dicar­bide is a candidate fertile material for the Generation IV high-temperature reactor (HTR) and VHTR systems, and it is also exploitable for accelerator — driven system (ADS) burners. Solid solutions of UC2-ThC2 were candidate fuels for the Dragon High Temperature Reactor-coated particle fuels.36 However, thorium-based fuel is difficult to recycle because of the radioprotection issues generated by the hard g-emission of 208Tl (2.6 MeV), formed in the 232Th-233U spent fuel.

2.04.2.1 Phase Relationships

Atmospheric pressure phase equilibria in the Th-C system are reported in Figure 4.

Thorium metal has an fcc (a) structure below 1633 K and a bcc structure (p) at higher temperatures. The first can accommodate carbon atoms as intersti­tials, resulting in the formation of thorium monocar­bide without any lattice change.5 The ThC1±x fcc solid solution range, extending from pure Th to ThC196 at high temperatures, is stable between ThC067 and ThC097 below 1300 K. The exact high carbon limit is still under debate.37 A miscibility gap seems to exist in the ThCi_x phase field, between ThC0.06 around 1000 K,38 ThC0.30 at 1413 (±40) K,39 and ThC0 67 at 1150 K,2 probably extending to room temperature with approximately the same composi­tion boundaries. At higher temperatures, single car­bon interstitials can be replaced by C2 groups up to ThC1.96. Thus, only two compounds have been observed in the Th-C system at atmospheric pres­sure: the fcc monocarbide with its broad nonstoichio­metry range and the dicarbide, more often observed

Figure 4 The Th-C phase diagram.

as hypostoichiometric (ThC2—x). Thorium sesqui — carbide Th2C3 has been observed only at pressures above 30 kbar.27 At low temperatures (below 1500 K), ThC2—x is a monoclinic line compound (a) with composition ThCi.94,40 observed in equilibrium with ThC098 at 1528 (±40) K in the presence of oxygen.41 Around 1528 (±40) K, ThC2—x converts eutectoidally to a tetragonal phase (p) with a homogeneity range between C/Th = 1.66 at 1528 K and 1.96 at 1713 K, the temperature at which the a! p ThC2 phase transition occurs at its C-rich phase boundary.40 Pialoux and Zaug42 reported a different phase diagram, with higher C/Th ratios for the Th-rich p-ThC2 phase boundary, extending from 1.96 at 1570 K to 1.85 at 1743 K. This phase diagram does not include the eutectoid decomposition of p-ThC2, but rather a a! p-phase transition in the line com­pound at 1570 K. All authors agree on the formation of a cubic fcc ThC2—x modification (g) as the tem­perature is raised above 1763 (±45). A solid miscibil­ity gap has been observed by Bowman et a/.40 in the ThC-ThC2—x domain, with a maximum at 2123 (±40) K and C/Th = 1.22. The same maximum was observed by Pialoux and Zaug42 at 2173 (±40) K and C/Th = 1.95. There exists a ThC2-C eutectic of proba­ble composition ThC238 and temperature 2718 K. Obvi­ously, some questions on the ThC2—x phase boundaries are still open, often in relation to the large uncertain­ties in the reported transition temperatures.

The commonly accepted melting point of pure Th is 2020 ± 10 K.6 In the low-carbon domain, a eutectic

(or peritectic) isotherm around 1980 K in the compo­sition range of 0.06 < C/Th < 0.13 has been observed.

Two congruent melting points were observed in the solid solution region with 0.13 < C/Th < 1.96, the first at T = 2773 ± 35 K and C/Th = 0.97 ± 0.05, the second at T = 2883 ± 35 K and C/Th = 1.90 ± 0.06.

The boiling point of ThC2 was extrapolated to be 5400 K at 1 atm.43