Phase Separation in Fluorite-Related U1–yCeyO2–x: A Re

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Cite This: Inorg. Chem. XXXX, XXX, XXX−XXX

Phase Separation in Fluorite-Related U1−yCeyO2−x: A Re-Examination by X‑ray and Neutron Diffraction David Simeone,*,†,⊥ Philippe Garcia,‡ Audrey Miard,‡ Gianguido Baldinozzi,†,⊥ Florence Porcher,¶ and Jean-Francois Berar§ †

CEA/DEN/DMN/SRMA/LA2M-LRC CARMEN, CEA, Université Paris-Saclay, F-91191 Gif-sur-Yvette, France CNRS/CentraleSupelec/UMR 8085, Grande voie des vignes, F-92290 Chatenay Malabry, France ‡ CEA, DEN, DEC, Centre de Cadarache, F-13108 Saint-Paul-Lez-Durance Cedex, France ¶ CEA/DRF/Iramis/LLB, CEA, Université Paris-Saclay, F-91191 Gif-sur-Yvette, France § Institut NEEL CNRS/UGA UPR2940 25 rue des Martyrs BP 166, F-38042 Grenoble cedex 9, France

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⊥

S Supporting Information *

ABSTRACT: The phase separation observed at low temperature (below circa 600 K) in the U1−yCeyO2−x system and for values of y between roughly 0.34 and 0.5 purportedly involves fluorite structures only. However, for y values above 0.5, an oxygen-deficient C-type bixbyite is also reported. In this work, the phase separation in U0.54Ce0.46O2−x has been reexamined using X-ray and neutron diffraction. Below a critical temperature, the existence of two fluorite related structures in the miscibility gap is confirmed: a stoichiometric U0.54Ce0.46O2 phase and an oxygen-deficient U0.54Ce0.46O2−x phase. Although the former is indeed a fluorite, we show that the other endmember phase has a C-type bixbyite structure. This would suggest that the oxygen-deficient phase can be described as a bixbyite over the entire cerium composition range. that in the remainder of this paper, the term “oxygen-deficient phase” loosely refers to a compound of general formula U1−yCeyO2−x (where x > 0), which, although fluorite related, is not necessarily a fluorite. The literature mentions no ordered oxygen-deficient phases in this composition range. At higher cerium or plutonium compositions,7,10 however, the oxygendeficient end-member phase has been identified as a C-type bixbyite. Interestingly, in analogous binary RO2−x (0 < x < 0.5) systems with only one cation species present in two oxidation states (R3+ and R4+), fluorite-related superstructures are also commonly reported.14,15 This paper is an attempt at understanding how, in these ternary oxides, the oxygen-deficient end-member phase switches from being a fluorite at low aliovalent cation content to a C-type bixbyite at higher cation concentrations. Not only does the questionconstitute an academic challenge to understand and model the structural stability of oxygen-deficient fluorite oxides1 but also it has very practical implications in regard to the stability of mixed oxide fuels, crucial for the safe and sustainable use of nuclear power. To address the issue, two types of experiments have been carried out. The first is an in situ microstructural characterization

1. INTRODUCTION A characteristic property of AO2 compounds such as UO2, CeO2, and PuO2, which crystallize to form fluorite structures, is the readiness with which the cation lattice incorporates a large proportion of aliovalent ions to form “fluorite-type solid solutions”. These A1−yByO2−x systems are often attributed to a fluorite structure1 in which the cation lattice is virtually complete and the anion lattice can be highly defective. The distribution of cations and of anion defects is assumed to be completely random. However, there are a number of examples where these “fluorite-type solid solutions”, while existing at high temperature, give way at lower temperature to ordered fluorite-related superstructures.2,3 For anion-excess systems, the structural mechanisms for the formation of ordered phases seems well established leading to “vernier structures” as in the Zr2O2− Nb2O5, Zr2O2−Ta2O5 and (U, Zr)(O, F)2−x ternary oxides.4 No such consensus exists for anion-deficient systems, but despite the relative dearth of structural information,2 substoichiometric ternary U−Ce5−8 or U−Pu9−13 oxides have been extensively investigated. These two systems have a number of common features. For intermediate plutonium or cerium concentrations (typically 0.15 ≤ y ≤ 0.5 for Pu bearing ternary oxides), two fluorite phases coexist below a critical temperature: a stoichiometric U1−yByO2 (B = Ce or B = Pu) and an oxygendeficient U1−yByO2−x phase also reported to be a fluorite. Note © XXXX American Chemical Society

Received: May 21, 2019

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DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry

furnace was purged using highly purified helium prior to the start of the experiment. The temperature of the sample was determined using a methodology exposed in reference.20 Diffraction patterns of a LaB6 NIST reference powder were collected at room temperature to characterize the instrumental aberrations of the setup in a Bragg−Brentano configuration. Reliability factors (Rwp = 8.76%, RB = 6.88%) resulting from Rietveld refinements of the LaB6 diffraction patterns allows an estimate Δa to be made of the cell parameter accuracy ( a < 10−4 )21 and of the instrumental broadening of the XRD setup. Neutron patterns were collected on the 3T2 beamline (λ = 0.123 nm), a high resolution two-axis diffractometer dedicated to powder diffraction at the LLB facility. A NaCaAlF standard was used to determine the instrumental background and broadening. All diffraction patterns were refined using the XND (X-ray and neutron diffraction) software.22 2.3. Annealing Strategy Relative to the in Situ X-ray Diffraction Study: Temperature and Oxidation Cycles. The Xray diffraction experiments are run in an open system, which means that the samples are exposed to a flowing gas mixture. The stoichiometry of the solid is controlled by allowing the oxygen activity in the solid to equilibrate, via a buffering reaction, with that of the gas phase. A mixture of hydrogen and water vapor was used to control the oxygen pressure of the gas phase hence the oxygen content of the solid. Our strategy was to reduce the sample at high temperature in the single fluorite phase region and maintain the reducing conditions. Equilibration was carried out at a temperature of 1173 K and under an oxygen potential of −428 kJ/mol until the microstructure had stabilized. This oxygen potential (ΔG(O2) = RgTLn(pO2), where Rg is the ideal gas constant and T and pO2 are respectively the temperature and the oxygen partial pressure of the sample) can be imposed since the setup is equipped upstream and downstream from the sample holder with an electrochemical oxygen pump and pressure gauge system.23,24 This enables the control and measurement of the oxygen partial pressure and the fact that the input and output oxygen pressure values are identical, guarantees that the system is buffered by the hydrogen and water vapor mixture which is used. According to Markin,5 these thermodynamic conditions correspond to an oxygen to metal (O/M) ratio of roughly 1.95. Once at equilibrium, the sample was then quenched down to 573 K, with a ramp rate of 5 K s−1, and further diffraction patterns were collected at constant temperature every 50 K down to room temperature (303 K). A 15 min period was allowed for the system to thermally equilibrate prior to beginning the data collection.

of the phase separation in U0.56Ce0.44O1.95 using X-ray powder diffraction. The second is a structural study of the oxygendeficient end-member using neutron diffraction. The (U0.56Ce0.44O1.94) sample preparation conditions were indeed chosen to reproduce, based on Markin’s5,6 data, the composition of the end-member. The paper is organized as follows: in section 2, we describe the materials studied along with the characteristics of the X-ray and neutron setups and the experimental strategy adopted; in section 3, we study the phase separation as a function of temperature from the analysis of in situ X-ray diffraction patterns. Neutron diffraction patterns collected at room temperature are used to reveal the presence of superstructures in the oxygen-deficient end-member phase. In the Discussion, we essentially analyze the consistency between X-ray and neutron data, and we suggest a possible mechanism for explaining the phase separation in the ternary U−Ce−O substoichiometric system.

2. EXPERIMENTAL SECTION 2.1. Material Manufacturing and Characteristics. One of the sources of scatter relative to studies involving mixed U−Ce or U−Pu oxides5,7,16 stems from the cation homogeneity of the oxide powder, which is very rarely reported. Powder metallurgy routes are generally not considered as being the most appropriate method for producing homogeneous cation distributions because the low cation interdiffusion coefficients require repeating the grinding and sintering stages of the manufacturing process and often prolonged high temperature annealing. In contrast, liquid routes offer a greater guarantee of obtaining mixed oxide powders which are homogeneous down to the atomic scale.17 These processes generally involve dissolution of the initial oxide products in ad hoc proportions followed by either direct denitration or, alternatively, formation of an ammonium nitrate salt18 or oxalate precipitate19 containing both actinide or lanthanide cations. In this work, we have chosen to manufacture products from oxalate precursors. Chemical analyses using inductively coupled plasma mass spectrometry (ICP-MS) of the initial oxalates revealed a cerium proportion of 0.439. Following calcining, the oxide powders were pressed into pellets and sintered for 4 h at 1700 °C in a flowing argonbased gas mixture containing 5% hydrogen and 0.12% water vapor. The final hydrostatic density of the sintered pellets was approximately 98.5% of the theoretical density. The proportion of cerium with respect to the total cation content was derived from an energy dispersive X-ray (EDX) analysis performed at different locations at the surface of a pellet. This proportion was estimated at 43.9 ± 1.4%, identical to the value obtained from the characterization of the original oxalate. In regard to the X-ray diffraction study, approximately 50 mg were sampled from one of the pellets and then ground using an agate mortar and pestle so as to produce a powder with a typical particle size of a few microns. For the room temperature neutron experiment, four pellets (total mass 12.8 g) were reduced at high temperature (950 °C) and rapidly cooled so that their composition corresponded roughly to that of the end-member phase in Markin’s phase diagram,5 i.e., U0.56Ce0.44O1.93. The average oxygen to metal ratio was determined by weighing pellets prior to and following the reduction annealing. This quantity, averaged over five pellets, was of 1.937 with a standard deviation of 10−3. Following reduction, samples were sealed in a gastight vanadium canister in an atmosphere of highly purified argon. 2.2. Diffraction Setup. XRD patterns were collected at the LEFCA facility with a D8 advance X-ray diffractometer, which uses a Cu Kα radiation from a tube source at 40 kV and 40 mA. The system is equipped with a LynXEye fast-counting position sensitive detector (PSD) with an aperture of 3° and a goniometric radius of 275 mm. It also comprises a heating stage made up of a tantalum heater and a Mo ribbon for direct heating of the sample.12 The data collection was carried out over 2θ angles ranging between 18° and 140°. The U0.56Ce0.44O2−x powder was deposited on the Mo ribbon, and the

3. ANALYSES OF DIFFRACTION PATTERNS To characterize the temperature dependent structural and microstructural changes that occur during the phase separation process in U0.56Ce0.44O1.95, Rietveld refinements are carried out of X-ray diffraction patterns collected during the in situ experiment (section 3.1). However, in contrast with X-rays that are essentially only sensitive to uranium atoms, neutrons are practically equally sensitive to U, Ce, or O atoms since the coherent neutron scattering lengths are of a similar order of magnitude (8.417, 4.840, and 5.803 fm for U, Ce, and O atoms respectively). Consequently, the positions and occupancies of these three types of atoms can be accurately determined from a Rietveld analysis of neutron diffraction patterns. This analysis we provide for a U0.56Ce0.44O1.94 sample and diffraction patterns collected at room temperature (section 3.2). The consistency between the analyses using X-rays and neutrons is discussed in section 3.3, and some underlying phase separation mechanisms involved are suggested. 3.1. X-ray Diffraction. The examination of diffraction patterns reveals that only a single fluorite phase exists at and B

DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry above Tc = 573 K. Below this temperature, two different phases with distinct unit cell parameters appear in the diffraction pattern in agreement with previous studies.5,7 The phase separation temperature derived here agrees fairly well with the one measured by previous authors.5 Assuming both phases are fluorites at this cerium content,5,7 it is possible to determine the variation of their respective unit cell parameters as a function of temperature as is illustrated in Figure 1.

Figure 2. Variation of the size of coherent diffracting domains (CDD) of the stoichiometric (black squares) and oxygen-deficient (red circles) low temperature phases as a function of temperature. Sizes of CDDs are roughly equal to 100 nm and do not change with temperature.

which are roughly equal to 100 nm). This suggests that the phase separation does not result from the growth of coherent oxygendeficient domains in a stoichiometric fluorite matrix or vice versa. At each temperature, slopes of Hall−Williamson plots are shown to be small, indicating that no local fluctuation of inter-reticular distances, i.e. no microstrains, exists in the low temperature phases. In an attempt to shed light on the mechanisms responsible for demixing in the U−Ce−O system, Rietveld refinements were first carried out assuming all phases are fluorite as reported in the literature. 7−9 At and above T c , where only a single U0.56Ce0.44O1.95 phase exists, the Bragg reliability factor RB, which is a measure of the quality of the structural model, is satisfactory (4.5%). Below Tc, this factor deteriorates and remains systematically greater than 7%. A comparison of experimental and simulated diffraction patterns shows that intensities of (111), (311), (331), and (333) Bragg reflections of the oxygen-deficient end-member phase are accurately calculated whereas (200), (220), (222), and (400) reflections are systematically badly calculated. A graphical example of such an analysis is plotted in Figure 4. Peaks at 28° and 56° are associated with the (111) and (311) reflections and peaks at 32.5° and 46.8° are associated with the (200) and (220) reflections. It appears that both fluorite (green line) and bixbyite (red line) diffraction patterns fit the (111) and (311) peaks whereas only the c-type bixbyite correctly reproduces the collected X-ray intensity of the (200) and (220) peaks as explained in the caption of Figure 4. This constitutes a strong motivation for seeking an alternative structural model. By analogy with the binary Ce2O3 sesquioxide,25,26 a second attempt was made at describing the X-ray diffraction patterns below Tc assuming that the oxygen-deficient end-member phase has a C-type bixbyite structure (space group Ia3̅) with a unit cell parameter double that of fluorite (aB = 2aF). In our model, uranium and cerium cations can occupy distinct 8b and 24d Wyckoff positions. Oxygen atoms are also allowed to occupy two Wyckoff sites: the 16c and the 48e. Note that in an ideal C-type bixbyite, 48e sites are fully occupied and 16c sites are empty. Figure 3 describes the crystallographic relationships between the Ia3̅ and Fm3̅ m space groups whence the following constraints on the reduced coordinates of oxygen atoms are

Figure 1. Variation of the unit cell parameter of the stoichiometric (black squares) and oxygen-deficient low temperature phase (red circles) as a function of temperature, derived from the analyses of X-ray diffraction patterns. At and above Tc = 573 K, the two phases merge in an unique fluorite structure (blue triangles).

Note that a linear dependence upon temperature of lattice parameters of both phases reflects the fact that we are not dealing with a Martensitic transformation (no change upon temperature in macroscopic strains). This linear feature also indicates that the temperature ramp rate used in our study is sufficient to avoid any exchange of oxygen with the surrounding gas phase during cooling. The slope of the unit cell vs temperature curve relative to the high temperature phase (above Tc), is estimated at 9.7 × 10−6 nm K−1, which is reasonably close to the values published by Markin et al. (this slope varies, depending on stoichiometry, between 8.1 × 10−6 and 9.4 × 10−6 nm K−1 as displayed in Figure 5 from ref 5). Below Tc, the change upon temperature of the unit cell parameter of the stoichiometric phase (full line in Figure 1) is similar to that determined by Lorenzelli and Touzelin7 for stoichiometric mixed oxide containing the same quantity of cerium. Further, the unit cell parameter of the oxygen-deficient end-member at room temperature is equal to 0.5486(2) nm, which is very close to the value of 0.5487 nm extrapolated at the appropriate cerium content from.5,7 These various observations illustrate the consistency of our results with those of previous authors and highlight the fact that under our in situ experimental conditions, reaction kinetics are too low to induce a substantial change in the oxygen content of the material. To characterize the microstructure of the low temperature phases, a Hall−Williamson analysis is performed at each temperature. This provides an indication of the size of coherent diffracting domains (CDD) and of the potential presence of microstrains. Figure 2 shows there is no discernible change with temperature in the size of CDDs for either the stoichiometric (black squares) or oxygen-deficient (red circles) phase (both of C

DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry

• (400) reflections ((200) reflection in the fluorite description) are mainly sensitive to the cation occupation numbers. • (440) reflections ((220) reflection in the fluorite description), are sensitive to the relative occupancy of oxygen atoms in both 16c and 48e positions. • (444) reflections ((222) reflection in the fluorite description) are sensitive to the shift of oxygen atoms with respect to their ideal fluorite position (ξ). Figure 4 shows a comparison between our experimental data and different Rietveld refinements performed at room temperature. Two models are tested. The first assumes the oxygendeficient phase to have a fluorite structure (green curve) and the second uses our alternative model (red curve). The set of refined structural parameters and reliability factors are provided in Table 1. The following comments may be formulated: • The oxygen occupancy relative to the 16c site is estimated at 0.82 (1), which puts the oxygen to metal ratio of the oxygen-deficient phase at 1.910(5), in agreement with the value extrapolated from Markin’s phase diagram5 at the appropriate cerium content (1.92). • The refined value of ξ is equal to 0.010(5), indicating that the bixbyite phase for this cerium content can be viewed as a relatively small distortion of the fluorite phase. It corresponds to a displacement of oxygen ions from their fluorite position of roughly 0.19 Å.

Figure 3. Crystallographic relationship between the Ia3̅ and the Em3̅m space groups, their metrics and Wyckoff positions. In a C-type bixbyite (space group Ia3̅ origin at 3̅) 8b, 24d cation and 48e anion positions have a site occupancy of 1. The 16c Wyckoff sites are empty. By imposing (0, 0, 1/4) to be the coordinates of 24d positions, the value of ξ at 0 and allowing the full occupation of 16c Wyckoff sites, the fluorite structure is recovered with a unit cell parameter half that of the bixbyite.

assumed: X16c = Y16c = Z16c = 1/8 − ξ, X48e = Z48e = 1/2 − X16c = 3/8 + ξ, and Y48e = 1/4 − X16c = 1/8 + ξ. Note that in our model, the first reduced coordinate of cations in 24d positions was set to zero. For ξ = 0, the fluorite structure is indeed recovered. This model thus describes the transformation from a bixbyite to a fluorite in a continuous fashion. We have calculated explicitly the structure factor F(hkl) as a function of cation and anion occupations and ξ. In so doing, we have found that

Figure 4. Zooms on calculated X-ray diffraction patterns resulting from different Rietveld refinements. The experimental pattern (black line) is refined assuming that the oxygen-deficient end-member phase is fluorite (green lines in graphs a and b: Rwp = 10% and RB = 9%) or C-type bixbyite (red line in graphs c and d: Rwp = 7% and RB = 4%). Whereas both models fit (111) and (311) reflections at 2θ = 28° and 2θ = 56° adequately, the fluorite model badly fits the intensities of the (200) and (220) reflections at 2θ = 32.5° and 2θ = 46.8°, suggesting that oxygen vacancies in the oxygen-deficient endmember phase may not be randomly spread in the anion sublattice as expected for a fluorite structure. D

DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry

and 35.3° can be observed, which cannot be indexed if the oxygen-deficient phase is assumed to be a fluorite. Indeed, the (420) peak which defines a cation bearing plane, could possibly be associated with a fluorite reflection but with twice the lattice parameter expected for the fluorite (the (210) reflection would not show up). As for the two latter reflections, which correspond to oxygen bearing planes, their existence precludes an Fcentered lattice as is the case for a fluorite. They can however be indexed as (420), (314), and (215) reflections assuming a Ctype bixbyite structure. On the strength of these observations, a Rietveld refinement was carried out assuming two phases are present: a stoichiometric fluorite and a C-type bixbyite. The resulting structural parameters are summarized in Table 2. One notes that the reliability factor Rwp = 4.4%, which reflects the quality of the data, is about half the one published in the only equivalent study (i.e., a neutron diffraction study of a substoichiometric U−Ce mixed oxide).8 The principle conclusions to be drawn from this table are

Table 1. Structural Parameters Determined from the Rietveld Refinement of X-ray Diffraction Patterns Collected at Room Temperature (See Text)a atom

W.P.

P.S.

coordinates

occ.

U Ce O

4a 4a 8c

U Ce U Ce O

8b 8b 24d 24d 16c

O

48e

m3̅m (0, 0, 0) 0.55(6) m3̅m (0, 0, 0) 0.43(9) 4̅3m (1/4, 1/4, 1/4) 0.99(1) Phase I: Fm3̅m, a = 5.4464(3), RB = 7.00% 0.3̅. (1/4, 1/4, 1/4) 0.56(1) 0.3̅. (1/4, 1/4, 1/4) 0.43(4) 2.. (0, 0, 1/4) 0.55(7) 2.. (0, 0, 1/4) 0.44(1) 0.3. (0.115(5), 0.115(5), 0.82(1) 0.115(5)) 1 (0.385, 0.135, 0.385) 1.00 Phase II: Ia3̅ , a = 10.9718(4), RB = 3.33%

Uiso (A2) 0.040(3) 0.040(3) − 0.004(1) 0.004(1) 0.020(5) 0.020(5) − −

a

Column 1: nature of atoms. Column 2: Wyckoff positions, column 3: point symmetry of Wyckoff positions. Column 4: reduced atomic positions. Column 5: occupation numbers. Column 6: isotropic Debye−Waller terms). Standard deviations are given in parentheses. The reliability factor Rwp is equal to 6.34%.

• The proportion of cerium ions present on 8b sites is twice the average value for this material, demonstrating most emphatically that the oxygen-deficient phase is not a fluorite.

• Cation sites are insensitive to the cerium content; i.e., cation proportions on 8b and 24d sites are identical to the macroscopic cerium or uranium proportions determined from chemical analyses.

• Cations in 24d positions are slightly displaced with respect to the ideal fluorite position (0, 0, 1/4). • Similarly, oxygen ions in both 16c and 48e positions are displaced, albeit slightly, away from their ideal fluorite positions. In addition, the crystallographic relationships between the coordinates of ions lying in 48e and 16c positions, which were assumed in section 3.1, are now confirmed by the neutron pattern refinement. This confirms that a single ξ parameter can model the displacement all oxygen ions.

3.2. Neutron Diffraction. To test the validity of our structural model, neutron powder diffraction patterns of a sample of general formula U0.56Ce0.44O1.937 (see section 2.1) were collected at room temperature on the 3T2 high resolution diffractometer at LLB. These patterns are plotted in Figure 5 (black line). Three very weak reflections located at 29.5°, 33.1°,

• The oxygen to metal ratio of the oxygen-deficient phase (i.e., 1.935 ± 0.005) is consistent with the value (roughly 1.92) extrapolated from Markin’s room temperature data5 and our X-ray refinement, i.e., 1.91 ± 0.001. From the scale factors, the volume fractions of each end-member phase are computed and the average oxygen content of our material is estimated at roughly 1.94 ± 0.01. This value is entirely consistent with that obtained from weighing our samples, i.e., 1.937 ± 10−3 (see section 2.1). All of these points provide strong evidence of the consistency between the XRD and neutron diffraction data and of the quality of our structural model. 3.3. Discussion. The most salient conclusion to be drawn from the preceding sections is that the oxygen-deficient endmember phase in our U−Ce−O ternary oxide does not have a fluorite structure. In fact, our refinement shows unambiguously that anion vacancies and cations are not randomly distributed. This ordering phenomenon induced by the decrease in temperature has long been observed in many anion-deficient fluorites (e.g., ref 1, chapter 8). Ordering has not been reported in previous studies of the anion-deficient U−Ce−O system in our composition range. Weak bixbyite superlattice reflections are indeed difficult to pick up using X-rays only, because their intensities are extremely low: X-ray intensities relative to the (420), (314), and (215) reflections specific to bixbyite represent respectively 0.19%, 0.27%, and 0.03% of the most intense reflection (i.e., the (222) reflection) for a U0.56Ce0.44O1.93 oxide. For neutron diffraction

Figure 5. Comparison between the refined (red line) and experimental neutron (black line) patterns for the U0.56Ce0.44O1.937 sample collected at room temperature assuming a stoichiometric fluorite (Bragg angles: red sticks) and an oxygen-deficient C-type bixbyite (Bragg angles: green sticks) (Rwp = 4.4%, RBragg = 3.9%). The relative error dY/wY resulting from the Rietveld refinement is plotted in the top graph. Three peaks at around 29.5, 33.1, and 35.3° are visible in the inset. These peaks can only be indexed assuming the oxygen-deficient end-member phase is a C-type bixbyite. E

DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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Inorganic Chemistry

Table 2. Structural Parameters Determined from the Rietveld Refinement of Neutron Diffraction Patterns Collected at Room Temperaturea Atom

W.P.

P.S.

U Ce O

4a 4a 8c

m3̅m m3̅m 4̅3m

U Ce U Ce O O

8b 8b 24d 24d 16c 48e

0.3̅. 0.3̅. 2.. 2.. 0.3. 1

coordinates

U11

U22

U33

0.0073(2) 0.0073(2) 0.0130(7)

0.0073(2) 0.0073(2) 0.0130(7)

0.0073(2) 0.0073(2) 0.0130(7)

0.0105(1) 0.0105(1) 0.0024(1) 0.0024(1) 0.0217(3) 0.0250(2)

0.0105(1) 0.0105(1) 0.0158(5) 0.0158(5) 0.0217(3) 0.0105(3)

0.0105(1) 0.0105(1) 0.0118(2) 0.0118(2) 0.0217(3) 0.0105(3)

occ.

(0, 0, 0) 0.54 (0, 0, 0) 0.46 (1/4, 1/4, 1/4) 1 Phase I: Fm3̅m, a = 5.4458(5), RB = 8.00% (1/4, 1/4, 1/4) 0.18(6) (1/4, 1/4, 1/4) 0.82(6) (0.006(1), 0, 1/4) 0.66(2) (0.006(1), 0, 1/4) 0.34(2) (0.122(2), 0.122(2), 0.122(2)) 0.87(1) (0.375(2), 0.126(1), 0.371(1)) 1.00 Phase II: Ia3̅ , a = 10.9838(4), RB = 3.23%

a

Column 1: nature of atoms. Column 2: Wyckoff positions. Column 3: point symmetry of Wyckoff positions. Column 4: reduced atomic positions. Column 5: occupation numbers. Columns 6−8: symmetry adapted harmonic Debye−Waller terms). Standard deviations are given in parentheses. The reliability factor Rwp is equal to 4.4%.

The low temperature ordered C-type bixbyite structure can be seen as resulting from a composition and displacive modulation along the ⟨111⟩ direction of the high temperature fluorite phase as suggested by previous authors for similar oxygen-deficient ternary oxides.1,2 This description prompts us to propose a mechanism for the phase separation. As the temperature falls below some critical value Tc, a mutual pinning of oxygen vacancies and electronic defects in the form of Ce3+ ions takes place which leads to a non random occupation of sites by cations and anion vacancies. This mechanism is entirely consistent with the temperatures (below 573 K) at which the phase separation occurs. Although, on the grounds of our experimental data, the above scenario seems plausible, transmission electron microscopy and neutron diffraction studies are required to demonstrate unambiguously the existence of such a modulation. Attention will now focus upon expanding our studies in these areas.

(λ = 0.123 nm), these intensities are respectively 10 times, thrice, and twice more intense. The neutron diffraction patterns presented in the previous section reveal there are basically two classes of reflections. The first comprises reflections which are sharp (high intensity, narrow peaks) and corresponds to the underlying fluorite structure. Reflections belonging to the second class such as the (420), (314), and (215), are broad and of low intensity. The CDDs associated with this second class of reflections (estimated at around 4 nm from the Hall−Williamson plots of neutron diffraction patterns), may constitute the signature of a commensurate phase modulation as observed in similar oxygen-deficient compounds.2 Such a modulation along ⟨111⟩ directions would justify the crystallographic constraints proposed and could result from the mutual pinning of oxygen vacancies in 16c sites and charge compensating Ce3+ ions in 8b sites. This is not surprising since clustering of charged defects, accompanied in this particular case by a phase transition, is known to significantly reduce the Coulomb contribution to the internal energy of a system. The scenario could therefore be that, as the temperature is lowered, the polaron hopping process, which occurs with a low migration barrier in the high temperature phase, ceases to ensure the high level of disorder required for the fluorite phase to exist. At this point, negative cerium ions and positive oxygen vacancies pin each other on their respective sites. Understandably, this occurs at the relatively low temperature of 573 K (see section 3.1), consistent with low charge hopping or vacancy migration energy barriers.27,28

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.inorgchem.9b01469. Details and Rietveld refinement performed on the neutron diffraction pattern (PDF)

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AUTHOR INFORMATION

Corresponding Author

*(D.S.) E-mail: [email protected]. ORCID

David Simeone: 0000-0002-3689-6414 Gianguido Baldinozzi: 0000-0002-6909-0716

4. CONCLUSIONS A study combining Rietveld analyses of X-ray and neutron patterns of samples of general formula U0.56Ce0.44O2−x (x ≈ 0.06) reveals that the oxygen-deficient end-member phase observed in the miscibility gap for a cerium content of around 50%, referred to in the literature as being a fluorite,5,7,8 can in all probability be described as having a C-type bixbyite structure. Rietveld refinements of this phase assuming a body-centered cubic structure (Ia3̅ space group) show that cerium atoms are not randomly distributed over all cation sites which precludes its description as a fluorite. This observation reconciles the thermodynamic and the crystallographic view points in the sense that it can ensure a continuous description of the oxygendeficient phase over the entire cerium composition range.

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Romain Vauchy and Alexis Joly (CEA, DEN, DRMC) for helping us carry out the XRD experiments we report here and Florent Garel and Anne Charlotte Robisson (CEA, DEN, DEC) for help in manufacturing the materials.

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DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.inorgchem.9b01469 Inorg. Chem. XXXX, XXX, XXX−XXX