An electrical conductivity and X-ray study of a high-temperature transition in U4O9

The high temperature transition in U₄O₉ has been studied by electrical conductivity measurements and X-ray diffraction. From the electrical conductivity measurements, a similar variation of log σ$\text{T}$ with reciprocal temperature to that in the transition range near room temperature is observed in the temperature range from about 300 to 800°C. Like the low temperature transition, a small lattice contraction is also observed in that temperature range by means of X-ray diffractometry, and the transition temperature increases from 530 to 620°C with increasing O/U ratio. After the transition the intensity of $\text{4a}{}_0$ superlattice reflections increases, but that of $\text{8a}{}_0$ superlattice reflections disappears. The mechanism of this high temperature transition is considered to be a second-order transition of the order—disorder type based on the configurational change of U⁴⁺ and U⁵⁺ with the shift of some portions of the lattice oxygen atoms from the lattice sites to the interstitial positions. The phase diagram of U₄O₉ is presented on the basis of the electrical conductivity and X-ray data.     

K-shell X-ray cross-section ratios for 1H, 4He and 6Li projectiles on targets from fluorine to titanium

Experimental K X-ray emission ratios have been measured for nine elements ranging from F to Ti (9 ? Z₂ ? 22) using ¹H⁺, and ⁶Li⁺ projectiles. The experimental ratios are in fair agreement with theoretical predictions for the velocity range studied, $\text{?}{}_{\mathrm{k}}\text{\$}\text{?}\text{1}$.     

Porosity dependence on thermal diffusivity and thermal conductivity of lithium oxide Li2O from 200 to 900°C

The porosity dependence on thermal diffusivity and thermal conductivity of Li₂O was studied in the temperature range of 200 to 900°C. The relationship between thermal diffusivity α_M and porosity P was expressed as α_M = α_T(1?ζP), $\text{}\text{?}$gz being 0.93, and also as $\text{α}{}_{\mathrm{M}}\text{=}\text{α}{}_{\mathrm{T}}\text{(1 + ηP)}$, η being 1.74 (200°C) ~ 1.11 (900°C). The relationship between thermal conductivity k_M and porosity obeyed the generalized Loeb equation k_M = k_T(1?γP), $\text{}\text{?}$gg being 1.70, and obeyed the modified Maxwell-Eucken equation $\text{k}{}_{\mathrm{M}}\text{=}\text{k}{}_{\mathrm{T}}\text{(1?P)}\text{(1 + βP)}$, β being 1.81 (200°C) ~ 1.32 (900°C). Alternative empirical equations were also attempted as α_M = α_T(1?P)^m and k_M = k_T(1?P)ⁿ, m? being 0.91 and $\text{n}\text{?}\text{= 1.06}$. The temperature dependence on thermal diffusivity and thermal conductivity was found to be expressed as α = (A' + B'T)^?1 and k = (A + BT)^?1. The pore correction factors are discussed in terms of porosity and temperature compared to published results.     

Electrical conductivity of UO2+x

Isothermal and isobaric conductivities of UO_2+x nave been measured as a function of oxygen pressure and temperature. Intrinsic disorder predominates at low-oxygen pressures. The pressure dependence of the conductivity can be expressed as in the intermediate oxygen pressure range in which more than one type of charge carrier predominate. At high oxygen pressures, it has been proposed that oxygen vacancy-interstitial trios in association with U⁺⁵ ions faciliate the fast transport of oxygen interstitials in the region.     

Solutions of lithium salts in liquid lithium: The interaction of hydride and of nitride with lead and with tin

Electrical resistivity studies of Li-Pb-H(N) ternary solutions ($\text{x}{}_{\mathrm{Li}}\text{? 0.96}$; $\text{T = 675 K}$) have shown that lead does not interact chemically with either LiH or Li₃N in liquid lithium solutions; similar studies of Li-Sn-H(N) ternary solutions ($\text{x}{}_{\mathrm{Li}}\text{? 0.96}$; $\text{T = 775 K}$) have shown that tin is also inert to both LiH and Li₃N in these solutions. Solubility data for LiH in liquid lithium-lead solutions ($\text{x}{}_{\mathrm{Pb}}\text{= 0.01}$; $\text{675 ? T(K) ? 735}$) have also been determined and compared with those for LiH in pure lithium. Enhanced solubilities are observed for the lithium-lead solutions; they are attributed to a decreased LiH activity in the ternary solutions vis-àvis the binary solutions. The significance of the results to the use of lead as a neutron multiplier in reactor lithium is discussed.     

The effect of void density on the dislocation bias

Recent experimental work on the void-swelling characteristics of FV548 steel during irradiation is analysed, and it is shown that the dislocation bias for interstitial condensation $\text{p}$ is a function of the void density $\text{N}{}_{\mathrm{v}}$, decreasing from 4&;#x0303;0% when $\text{N}{}_{\mathrm{v}}$ is 10¹⁹/m³ to 0&;#x0303;.5% when $\text{N}{}_{\mathrm{v}}$ is 10²³/m³. This dependence on $\text{N}{}_{\mathrm{v}}$ is responsible for most of the apparent temperature dependence of the bias. In addition to affecting $\text{N}{}_{\mathrm{v}}$, the temperature imposes a maximum void growth rate, which decreases as the temperature is reduced, and because of this, high swelling rates cannot be obtained at low temperatures even when $\text{N}{}_{\mathrm{v}}$ is low. The full effect of the void-density dependence of $\text{p}$ is therefore only visible at high temperatures.     

Diffusion of carbon in zirconium and some of its alloys

Diffusion of carbon in zirconium, zircaloy-2 and Zr- 2.5% Nb has been studied in the temperature range 873–1523K for zirconium and zircaloy-2 and 753–1523K for Zr-2.5% Nb alloy, using the residual activity technique. The diffusivities (in m²/s) in the α and β phases could be represented by $\text{D}{}_{\mathrm{C/\alpha -Zr}}\text{(873–1123K) = (2.00 ± 0.37) × 10}{}^{\mathrm{?7}}\text{exp [}\text{?(151.59 ± 2.51)}\text{RT}\text{]}$$\text{D}{}_{\mathrm{C/\alpha -Zircaloy-2}}\text{(873–1043K) = (1.41 ± 0.32) × 10}{}^{\mathrm{?7}}\text{exp [}\text{?(158.99 ± 3.14)}\text{RT}\text{]}$$\text{D}{}_{\mathrm{C/\alpha -Zr-Nb-alloy}}\text{(753–873K) = (4.68 ± 0.88) × 10}{}^{\mathrm{?7}}\text{exp [}\text{?(159.98 ± 2.91)}\text{RT}\text{]}$$\text{D}{}_{\mathrm{C/\beta \; Zr}}\text{((1143–1523K) = (8.90 ± 1.60) × 10}{}^{\mathrm{?6}}\text{exp [}\text{?(133.05 ± 1.46)}\text{RT}\text{]}$$\text{D}{}_{\mathrm{C/\beta \; Zircaloy-2}}\text{(1263–1523K) = (2.45 ± 0.61) × 10}{}^{\mathrm{?5}}\text{exp [}\text{?(150.29 ± 1.72)}\text{RT}\text{]}$$\text{D}{}_{\mathrm{C/\beta \; Zr-Nb\; alloy}}\text{(1143–1523) = (1.70 ± 0.42) × 10}{}^{\mathrm{?5}}\text{exp [}\text{?(158.20 ± 2.09)}\text{RT}\text{]}$The activation energies are given in kJ/mole. In the phase transition region, the diffusivities could be represented by the empirical relation: D = D_α^Cα · D_β^Cβ, where C_α, C_β are the concentrations of the two phases in the alloy and D_α, D_β are the extrapolated values of diffusion co-efficients in the α and β phases respectively.The results have been explained in terms of the interstitial mechanism of diffusion.     

The vaporization and thermal stability of lithium molybdates

Two lithium molybdates, δ-Li₄MoO₅ and Li₂MoO₄, were evaporated and measured by high temperature mass spectometry. Various lithium and molybdenum oxide ions were observed, and their partial pressures were obtained. From the thermochemical calculation of evaporation, the heats of formation of the molybdates were obtained for the following reactions, $\text{Li}{}_2\text{O(c) +}\text{1}\text{2}\text{MoO}{}_3\text{(c) =}\text{1}\text{2}\text{δ-Li}{}_4\text{Mo0}{}_5\text{(c), ΔH}{}_{\mathrm{r.298}}{}^{\mathrm{o}}\text{= ? 120.4 kj.mol}{}^{\mathrm{?1}}$, and $\text{Li}{}_2\text{O(c) + MoO}{}_3\text{(c) = Li}{}_2\text{MoO}{}_4\text{(c), ΔH}{}_{\mathrm{r.298}}{}^{\mathrm{o}}\text{= ? 154.7 kj.mol}{}^{\mathrm{?1}}$. Thermochemically, $\text{Li}{}_2\text{MoO}{}_4$ is less stable than $\text{δ-Li}{}_4\text{MoO}{}_5$.     

Solubility of nitrogen in liquid lithium and thermal decomposition of solid Li3N

The solubility of nitrogen in liquid lithium was determined from 195 to 441°C by direct sampling and equilibrium nitrogen pressure over solid Li₃N were measured at eight temperatures between 660 and 778° C. The solubility data may be represented by log₁₀S = 3.323 ? 2107 T^?1, where S is in mol% Li₃N and T is in K. From a thermodynamic analysis of the combined solubility and decomposition data, the standard free energy of formation of solid Li₃N was estimated to be ΔG°_f(kcal/mol) = 33.2 × 10^?3T ? 39.1. For dilute solutions of Li₃N in lithium, the Sieverts' law constant, $\text{K}{}_{\mathrm{S}}\text{=}\text{N}{}_{\mathrm{Li3N}}\text{p}{}^{\mathrm{<mtext>1</mtext><mtext>2</mtext>}}$, is given by In $\text{K}{}_{\mathrm{S}}\text{(}\text{atm}{}^{\mathrm{<mtext>?1</mtext><mtext>2</mtext>}}\text{) = ? 13.80 + 14 590 T}{}^{\mathrm{?1}}$, where N_Li3N is the mole fraction of Li₃N and p is the nitrogen pressure. The system Li-Li₃N appears to conform to a simple eutectic diagram in which the eutectic point occurs at ≈ 0.05 mol % Li₃N and 180.3° C. The melting point of Li₃N was found to be (813 ± 1)° C. Implications of the results of this study regarding the compatibility of liquid lithium with the strucutral materials of interest to fusion reactors are discussed.     

51Cr diffusion in α-Zr single crystals

The volume diffusion coefficients of the fast-diffusing solute ⁵¹Cr have been obtained in oriented α-Zr single crystals, in the directions parallel and perpendicular to the c axis. The dependence of these diffusion coefficients on temperature was also measured between 750°C and 848°C.Single crystals were grown by thermal cycling through the α α β transformation temperature (862°C). Diffusion coefficients were measured using the “ thin film” method. In some experiments non-Gaussian penetration profiles were obtained and this behaviour is also analyzed.The diffusion of ⁵¹Cr is faster in the c axis direction, with Q_∥(153 kj/mol) < Q_⊥ (163 kj/mol). The anisotropy factor $\text{f}{}_{\mathrm{a}}\text{=}\text{D}{}_{\mathrm{\parallel }}\text{D⊥}\text{≈ 3}$. This factor, the activation energies and frequency factors are discussed in comparison with those of other solutes.     

Phase relation and defect structures of nonstoichiometric U4O9±y and UO2+x at high temperatures

Phase relations in the composition range from $\text{UO}{}_{\mathrm{2+x}}$ to $\text{U}{}_3\text{O}{}_{\mathrm{8?z}}$ were studied by electrical-conductivity measurements and X-ray diffraction in the ranges $\text{1025°C ? T ? 1140°C}$ and . The plot of log σ versus log showed straight lines with distinct slopes, which corresponded to four regions ($\text{UO}{}_{\mathrm{2+x}}$, $\text{U}{}_4\text{O}{}_{\mathrm{9?y}}$, $\text{U}{}_4\text{O}{}_{\mathrm{9+y}}$ and $\text{U}{}_3\text{O}{}_{\mathrm{8?z}}$). The existence of the hyperstoichiometric $\text{U}{}_4\text{O}{}_{\mathrm{9+y}}$ phase was suggested in the temperature range from 1025 to 1126°C. The peritectoid temperature ($\text{U}{}_4\text{O}{}_{\mathrm{9\pm y}}\text{= UO}{}_{\mathrm{2+x}}\text{+ U}{}_3\text{O}{}_{\mathrm{8?z}}$) was estimated to be present between 1126 and 1131°C. The partial free enthalpies and entropies for the two-phase equilibrium ($\text{U}{}_4\text{O}{}_{\mathrm{9+y}}\text{+ U}{}_3\text{O}{}_{\mathrm{8?z}}$, and $\text{U}{}_4\text{O}{}_{\mathrm{9?y}}\text{+ UO}{}_{\mathrm{2+x}}$) were calculated and compared with previous results. From the dependence of the electrical conductivity on the oxygen partial pressure the nonstoichiometric defect structures of $\text{UO}{}_{\mathrm{2+x}}$ and $\text{U}{}_4\text{O}{}_{\mathrm{9\pm y}}$ were interpreted as consisting of doubly charged oxygen interstitials (O_i'') and doubly charged oxygen vacancies (V_O''). At room temperature, the homogeneity range of the U₄O₉ phase was investigated with a Debye-Scherrer camera.     

Existence of hypostoichiometric chromium sesquioxide at low oxygen partial pressures

The nonstoichiometric composition of $\text{Cr}{}_2\text{O}{}_{\mathrm{3\pm x}}$ was measured by means of thermogravimetry in the range of $\text{1173 ≦ T/K ≦ 1318}$ and . The compositional deviation from stoichiometry, $\text{x}$, in the hyperstoichiometric Cr₂O_3+x phase was observed to be smaller than 2 × 10^?4, irrespective of temperatures, provided that the hyperstoichiometric Cr₂O_3+x exists. The existence of the hypostoichiometric Cr₂O_3?x phase was first established in this study in the region of low oxygen partial pressure below 10^?5 Pa. From the oxygen partial pressure dependence of $\text{x}$ in Cr₂O_3?x, the defect structure was discussed with the neutral chromium interstitials in the composition near stoichiometry and with the triply charged ones far from stoichiometry. The partial molar enthalpy and entropy of oxygen of Cr₂O_3?x showed the complex compositional dependences, suggesting the change of the type of the predominant defect.     

Chemical thermodynamics of fusion reactor breeding materials and their interaction with tritium

Liquid lithium, lithium alloys (solid and liquid) and ceramic lithium compounds are candidate breeding materials for (D,T) fusion reactors.Besides their tritium breeding capability, which results from neutron capture, their thermochemical properties and their interaction with tritium are of particular interest. A good knowledge of the physical and chemical properties of liquid lithium exists; and the systems Li-LiH, Li-LiD and Li-LiT have been studied in great detail. For dilute solutions of D₂ in liquid lithium, Sieverts' law was found to be valid down to an atom fraction of $\text{x}{}_{\mathrm{D}}\text{= 10}{}^{-6}$; in the vapor, lithium polymers up to Li₄ and lithium deuterides are found.In the system liquid Li-Pb, the solubility of D₂ was measured as a function of temperature and alloy composition, and correlated with the activities of the constituent metals. The solubility of D₂ was found to obey Sieverts' law at low concentrations, and is many orders of magnitude smaller than that in liquid lithium. This holds also for solid “Li₇ Pb₂”.Vaporization studies yielded data on the thermal stability of the oxides: Li₂0, γ-LiAlO₂, β-Li_sAlO₄, LiAl₅O₈, Li₂ZrO₃, Li₄ZrO₄, Li₈ZrO₆, Li₂SiO₃ and Li₄SiO₄. Tritium diffusivity was studied in Li₂O, γ-LiAlO₂, β-Li₅AlO₄ and Li₄SiO₄. A large number of gaseous lithides were detected during these studies.     

Heat capacity and thermal decomposition of lithium peroxide

The heat capacity of Li₂O₂ was measured by adiabatic scanning calorimetry from 301 to 566 K resulting in $\text{C}{}_{\mathrm{p}}\text{= 59.665 + 52.123 × 10}{}^{\mathrm{?3}}\text{T + 5.0848 × 10}{}^5\text{T}{}^{\mathrm{?2}}\text{(J/K · mol)}$. The thermal decomposition of Li₂O₂ was studied by continuous calorimetric measurements. An endothermic decomposition was observed above 570 K. The enthalpy of the thermal decomposition was determined to be 25.8 kJ/mol. The effects of atmosphere and heating rate of the decomposition were also studied for the powder and the compact specimens.     

Electrical conductivity of UO2-ThO2 solid solutions

The electrical conductivities of UO_2+x. ThO₂ and their solid solutions, in thermodynamic equilibrium with the gas phase, were measured as a function of temperature, and of oxygen partial pressure in the temperatnre range 800 to 1200°C. The slope of the plot log α versus 1/$\text{T}$ for UO_2+x and UO₂-rich solid solutions exhibits a single region, whereas in the ThO₂-rich solid solutions it exhibits two regions. The pressure dependence of the conductivity (σ) in the UO₂-rich solid solutions can be represented by σ ∝ [O_i] ∝ in the range of $\text{0.01 < x < 0.1}$. Here, O_i is an interstitial oxygen and the partial pressure of oxygen, and it varies with the ThO₂ content. At greater deviation from stoichiometry ($\text{x ? 0.1}$) the presence of U₄O₉ or (Th U)₄O₉ phases influences the conductivity data. In ThO$\text{2}$ or ThO₂-rich solid solutions. P-type conduction at high oxygen pressures is interpreted as arising from the incorporation of excess oxygen into oxygen vacancies.     

Electrical conductivity measurement and thermogravimetric study of pure and niobium-doped uranium dioxide

The electrical conductivity and nonstoichiometric composition of UO_2+x and (U_1?yNb_y)O_2+x (y = 0.01, 0.05 and 0.10) were measured in the range $\text{1282 ≦ T ≦ 1373}$ K and Pa by tie four inserted wires method and thermogravimetry, respectively. The electrical conductivity of (U_1?yNb_y)O_2+x plotted against the oxygen partial pressure indicated a minimum corresponding to the transition between $\text{n}$- and $\text{p-}\text{type}$ cone uction. The band-gap energy of (U_1?yNb_y)O_2+x was calculated to be $\text{(248 ± 12)}$ kJmol.^?1, independent of niobium content, which is nearly the same as that of UO_2+x. From the oxygen partial pressure dependences of both the electrical conductivity and the deviation $\text{x}$ of UO_2+x and (U_1?yNb_y)O_2+x, the defect structures in these oxides were discussed with the complex defect model consisting of oxygen vacancies and two kinds of interstitial oxygens.     

Steady-state creep of zircaloy-4 fuel cladding from 940 to 1873 K

Steady-state creep rates of as-received zircaloy-4 fuel cladding have been determined from 940 to 1073 K in the α-Zr range, from 1140 to 1190 K in the mixed (α + β) phase region and from 1273 to 1873 K in the β-Zr phase region. Strain rates of between 10^?6 and 10^?2/s were determined under constant uniaxial load conditions. Assuming that creep rates can be described by a power law-Arrhenius equation, the creep rate for α-phase zircaloy-4 is given by: $\text{ge}{}_{\mathrm{ss}}\text{?}\text{= 2000 σ}{}^{5.32}\text{exp}\text{(?284 600/kT) s}{}^{\mathrm{?1}}$; for the β-phase zircaloy-4 by: $\text{ge}{}_{\mathrm{ss}}\text{?}\text{= 8.1 σ}{}^{3.79}\text{exp}\text{(?142 300/kT) s}{}^{\mathrm{?1}}$; and for the mixed (α + β) phase of zircaloy-4 (for creep rates ?3 × 10^?3 s^?1) by: $\text{ge}{}_{\mathrm{ss}}\text{?}\text{= 6.8 × 10}{}^{\mathrm{?3}}\text{σ}{}^{1.8}\text{exp}\text{(?56 600/kT) s}{}^{\mathrm{?1}}$. For the both the α and β phases, the activation energies for creep are in agreement with those of self-diffusion. For the mixed (α + β) phase region, the low creep rate range is controlled by grain boundary sliding at the α/(α + β) phase boundary.     

Die kinetik der wasserstoffabgabe von yhx im vakuum zwischen 550 und 950°C

The time-dependence of the dehydration of $\text{YH}{}_{\mathrm{x}}$ was determined by means of vacuum-heat extraction, for the composition range $\text{0 ≤ x ≤ 2}$. It was shown that the rate of dehydration is determined by the molecular recombination of hydrogen adsorbed at the sample surface, or else by the desorption of H₂ molecules from the surface. Hydrogen release, which was measured for spherical, plate-shaped or cylindrical samples, follows zero-order kinetics in the $\text{YH}{}_{\mathrm{2?y}}\text{and (α-H + YH}{}_{\mathrm{y?2}}\text{)}$ phase fields. In the two-phase field, as compared with the $\text{YH}{}_{\mathrm{2?y}}$ field, the dehydration rate is greater by a factor 4.7; the activation energy is 65 kcal/mole. The hydrogen release from α-Y solid solutions can be described by second-order kinetics if the condition of a low surface occupancy by hydrogen applies. The temperature dependence leads to an ‘activation energy’ of 67 kcal/mole.     

Compressive deformation of polycrystalline uranium at low temperatures

Polycrystalline α-uranium was tested in compression from 20 to 300 K and, with the exception of the 20 K tests, no cracking was observed during extensive plastic deformation ($\text{? ? ? 3.0}$). Alpha uranium was found to be a strongly work hardening material; the dislocation and twinning structures developed, however, are not stable arid resulted in unexpectedly weak material when samples were tested at temperatures above the initial deformation temperature. On the other hand, the results obtained suggest that large strain deformation at warm temperatures should lead to high yield strength uranium at room temperature. The stress-strain curves for annealed polycrystalline α-U from 78 to 300 K can be predicted accurately from the phenomenological relation: $\text{σ = σ}{}_{\mathrm{s}}\text{? exp[-(N?)}{}^{\mathrm{p}}\text{] · (σ}{}_{\mathrm{s}}\text{? σ}{}_{\mathrm{y}}\text{)}$, where $\text{N}$ and $\text{p}$ are material constants equal to 0.613 and 0.58, respectively, and σ_y is the yield strength. The saturation flow stress, σ_s, is predicted to be 6200 MPa (900 ksi) at 78 K and 2900 MPa (420 ksi) at room temperature.     

The enthalpy of formation of uf4 by fluorine bomb calorimetry

The energies of combustion in fluorine of two samples of uranium tetrafluoride, UF₄, were measured in a bomb calorimeter. Based on these measurements, the standard enthalpy of formation of crystalline UF₄, $\text{ΔH}{}^0{}_{\mathrm{f}}\text{(UF}{}_4\text{, c, 298.15 K) = ?1910.6 ± 2.0 kJ mol}{}^{-1}$ was derived.