Kinetics of Structural and Microstructural Changes at the Solid

Article pubs.acs.org/JPCC

Kinetics of Structural and Microstructural Changes at the Solid/ Solution Interface during Dissolution of Cerium(IV)−Neodymium(III) Oxides Stéphanie Szenknect,* Adel Mesbah, Denis Horlait, Nicolas Clavier, Sandrine Dourdain, Johann Ravaux, and Nicolas Dacheux Institut de Chimie Séparative de Marcoule, UMR 5257 CNRS/CEA/UM2/ENSCM, Site de Marcoule, Bâtiment 426−BP 17171, 30207 Bagnols sur Ceze Cedex, France S Supporting Information *

ABSTRACT: Improving the understanding of dissolution mechanisms at the solid/solution interface of mixed CeIV1−xNdIIIxO2−x/2 dioxides is a critical step in the frame of generation IV (Gen IV) nuclear fuel elements that integrates minor actinides recycling. In order to give significant insight into the dissolution kinetics of Ce0.4Nd0.6O1.7 sintered samples prepared from the thermal conversion of oxalate precursors, X-ray reflectivity (XRR) and grazing incidence X-ray diffraction (GI-XRD) were used to characterize the microstructural and structural changes of the solid/solution interface and were coupled with solution analysis. The dissolution mechanism occurred in three steps. After a short period of dissolution in which no Nd release was observed, GI-XRD revealed the existence of a 200 nm depth chemical gradient below the surface formed during the early times of the dissolution. A Nd-enriched zone was located at the surface of a Nd-depleted zone. Simultaneously, XRR data showed the development of a surface layer of 40 Å of low electron density. Both results fit well with the presence of a Nd-based thin surface layer phase. Then, the Nd-based surface layer dissolved and developed microporosity, leading to a strong increase of the surface of this layer. Thus, during the initial times, the Nd-based layer mainly contributed to dissolution, which played a crucial role in the kinetics of dissolution and explained why incongruent dissolution was first observed as a transient stage to congruent dissolution. From almost 3% of mass loss, a steady state was reached, corresponding to a congruent dissolution mechanism. The duration of the transient steps depended on the alteration conditions.

1. INTRODUCTION Mixed actinide dioxides such as (U,Pu)O2 are already used as fuels in pressurized-water reactor (PWR) nuclear reactors and stand as potential fuels within the development of several generation IV (Gen IV) concepts such as the sodium-cooled fast reactor (SFR) or gas-cooled fast reactor (GFR).1,2 Moreover, as the reprocessing of actinides coming from nuclear spent fuel into new fuel elements has emerged, the recycling of minor actinides including neptunium, americium, and curium in such mixed-oxides fuels is now often considered. CeO2 is frequently used as a surrogate for plutonium dioxide3,4 since both oxides crystallize in the same fluorite-type structure (face centered cubic, space group Fm3m ̅ ; JCPDS cards 01-081-0792 for CeO2 and 00-051-0798 for PuO2) with close unit cell parameters due to similar ionic radius [VIIIr(Ce4+) = 0.97 Å and VIII r(Pu4+) = 0.96 Å].5 Similarly, neodymium is often used as surrogate for trivalent minor actinides such as americium and curium. In such conditions, ceria-based mixed oxides of general chemical composition Ce1−xLnxO2−x/2 were used as model compounds for advanced nuclear fuels. In addition, Ce4+ and Ln3+ mixed oxides were widely studied as potential solid electrolytes for IT-SOFCs.6 In this aim, © 2012 American Chemical Society

several papers were already dedicated to the structural investigation of ceria-based mixed oxides. The incorporation of trivalent lanthanide in CeO2 fluorite-type structure is then obtained through the direct substitution of Ce4+ by Ln3+ along with the formation of half an oxygen vacancy to counterbalance the lack of positive charge.7,8 The formation of a cubic superstructure (Ia3̅ space group) related to the ordering of oxygen vacancies accumulated in the fluorite structure was reported for x ≥ 0.4, whatever the nature of the doping lanthanide element.9 Furthermore, this structure is maintained up to x = 1 for lanthanide elements ranging from Eu to Lu.10 For lighter lanthanides (such as La, Nd, and Sm), Ln2O3 does not crystallize commonly in the former cubic structure and a secondary phase is formed for higher mole fractions. For instance, Nd2O3 is reported with hexagonal structure (P3̅1m space group; JCPDS file 00-043-1023), and the limit of incorporation of neodymium in CeO2 reaches x ≈ 0.7.9,11,12 These limits of incorporation of Ln3+ in the cubic structure also Received: January 23, 2012 Revised: May 11, 2012 Published: May 12, 2012 12027

dx.doi.org/10.1021/jp3007602 | J. Phys. Chem. C 2012, 116, 12027−12037

The Journal of Physical Chemistry C

Article

laboratory bulk rate measurements defines the other end of the scale: that of the sample, from hundreds of micrometers to centimeters in the present case. This comparison also requires further averaging of the processes over a representative surface area, the so-called “normalization step”. It is commonly accepted that BET surface area is representative of the reactive surface area and that BET-normalized dissolution rates are appropriate far from equilibrium at low pH.33 However, BET surface area measurement is not possible in the case of sintered dioxides pellets and other methods must be developed to quantify the accurate surface area of such samples. Fischer et al.30 obtained a linear correlation between BET surface area and LSM data. They concluded that the largest pores of natural iron oxides contributed mainly to the BET surface area. That means that the quantification of surface area requires a method with high spatial resolution (typically submicrometric) and a large field of view. From this point of view, X-ray reflectivity (XRR) seems to be promising, as it gives access to physical properties of the solid surface, as electron density and roughness of the entire illuminated sample. Focused on the solid/solution interface, this work thus presents a new approach of the dissolution kinetics for sintered Ce1−xNdxO2−x/2 samples. XRR and grazing incidence X-ray diffraction (GI-XRD) were used to characterize microstructural and structural changes of the solid/solution interface and were coupled to the analysis of elementary releases in solution in order to gain thorough understanding of the dissolution mechanism.

differ depending on the chemical route of synthesis and/or the heat treatment conditions. Among them, oxalic precipitation was frequently considered for the preparation of precursors of such solid solutions since it led to quantitative precipitation of lanthanides, homogeneous distribution of cations,13 and improved the properties of the final mixed oxides regarding their sintering capability14−16 or chemical durability.17 The evaluation of the chemical durability of Ce1‑xNdxO2‑x/2 as a surrogate of nuclear fuel arises not only in the frame of the reprocessing operations planned in Gen IV type reactors, but also in view of nuclear spent fuel repository. However, the dissolution kinetics of CeO2 in nitric acid media was rarely examined.18 CeO2 usually appears to be strongly resistant to corrosion in oxidizing media. Moreover, the corrosion resistance of powdered Ce1−xNdxO2−x/2 mixed oxide prepared from oxalate precursors was only recently reported.19−21 The influence of physicochemical properties like crystallite size, specific surface area, and porosity on the normalized dissolution rates under various conditions (temperature, nitric acid concentration) was also highlighted.19 The results suggested that dissolution was mainly driven by surface reactions occurring at the solid/solution interface. The key role of the neodymium content (and thus of the presence of oxygen vacancies) was also underlined. Indeed, the chemical durability of Ce1−xNdxO2−x/2 solid solutions was reduced by almost one decade for each 9 mol % of neodymium incorporated. It was attributed to the decrease of the energy of cohesion of the solid associated with the substitution of Ce4+ by Nd3+ in the structure.22 However, the sintering step is expected to modify some of the properties that had a significant impact on the dissolution kinetics of powdered samples, i.e., obviously surface area and density, but also pore size distribution, grain size, and/or occurrence of the grain boundaries. Indeed, the influence of the crystallite size was evidenced recently for thorium−uranium,17 thorium−cerium,23 and cerium−neodymium dioxides.19 As an example, the average crystallite size (controlled through heating temperature) contributes to modifications of the normalized dissolution rates of thorium−cerium dioxides in the same order of magnitude than conventional parameters such as the leachate acidity. Grain boundaries could also present a different reactivity than the bulk material, as was evidenced for sintered thoria pellets.24 Guo and Waser25 also showed that the grainboundary cores of Nd-doped ceria were positively charged, probably owing to the oxygen vacancy enrichment therein. Grain boundaries of sintered pellets of cerium−neodymium oxides may thus constitute preferential dissolution zones. As noted by many other workers, both chemical and physical properties of the solid surface can affect the mechanism of precipitation and dissolution. Bulk measurements of the rates were thus improved by studies of the mineral surface after reaction by means of techniques such as atomic force microscopy (AFM), 26,27 scanning electron microscopy (SEM), and light optical microscopy methods, mainly, vertical scanning interferometry (VSI)28,29 and confoncal laser scanning microscopy (LSM).30 Through the use of a combination of techniques, it is thus possible to track the microstructural evolution of the solid/solution interface during dissolution and, combined with the elementary releases in solution, to give insight on the dissolution kinetics observed at the macroscopic scale.31 The changes in mineral surfaces exposed to solutions occur on a variety of spatial scales. Ultimately, the smallest limiting scale is that of chemical bonds.32 The comparison with

2. EXPERIMENTAL SECTION Mixed oxide samples of Ce0.4Nd0.6O1.7 were obtained by calcination at 1000 °C during 16 h of oxalate precursor prepared according to the synthesis protocol described recently by Horlait et al.9 The powdered samples were then ground during 30 min at 60 Hz in a jar mill and then pressed under 250 MPa in a 3 cm diameter tungsten carbide die. The obtained green pellets were thus sintered at 1400 °C for 10 h in air atmosphere. The surface of the pellets was polished to optical grade for the needs of XRR experiments. The surface of each sintered pellet was then observed by environmental scanning electron microscopy (ESEM) with a FEI Quanta 200. The images were processed and binarized with ImageJ 1.45 m software.34 Two data were extracted from such analysis: a 2D porosity defined by the area of the pores (black pixels) divided by the total area of the investigated domain and a 2D pore size distribution using the “analyse particles” plugin implemented in ImageJ. The image of the surface of each pellet and the corresponding 2D pore size distribution are gathered in Figure S1 (Supporting Information). The specific surface area of each pellet was then deduced from the pore size distribution, assuming the pores were cylindrical.35 The precise mole ratio of each cation in the mixed oxide was determined from energy dispersive spectrometry (X-EDS) analyses using a Bruker AXS X-flash 5010 detector coupled to the ESEM. About 20 measures randomly dispersed at the surface were considered in order to estimate the average molar ratios. The main characteristics of the sintered pellets including X-EDS results are gathered in Table 1. The dissolution experiments were performed in PTFE containers using static conditions. They were carried out in 25 mL of 4 M HNO3 and 5 mL of 0.1 M HNO3 for P1 and P2, respectively. The dissolution vessels were placed in a oven regulated at 60 °C. Aliquots of 20−100 μL of solution were 12028



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reflection was performed according to Salah et al.37 Corrected intensities (I) were then normalized to the incident intensity (I0). Reflectivity curves showed the evolution of the logarithm of the normalized intensity, log(R = I/I0) versus the wave vector transfer, q = (4π sin θ/λ) (Å−1). The experimental XRR curves were calculated using the matrix formalism38 depending on the following parameters of the layers: thickness, τ (Å); electron density, ρe (e− Å−3); roughness, σ (Å). Electron density of the layer is linked to the critical wave vector transfer, qc, at which the plateau of total external reflection exhibits a strong dip:

Table 1. Main Characteristics of Prepared Sintered Pellets (referred to as P1 and P2 throughout) diameter, d (cm) thickness, e (cm) mass, m0 (g) geometric density, ρg (g cm−3) Nd molar fraction, xNd (−) Ce molar fraction, xCe (−) calcd density, ρs (g cm−3) relative density (%) geometric porosity, ε (%) surface porosity, ε2D (%) specific surface area (m2 g−1)

P1

P2

2.25 (1) 0.30 (2) 6.804 (1) 5.76 (5) 0.63 (2) 0.37 (2) 6.76 (1) 85 (2) 15 (1) 16 (2) 0.015 (3)

2.16 (3) 0.11 (1) 2.410 (1) 5.98 (5) 0.59 (2) 0.41 (2) 6.77 (1) 88 (2) 12 (1) 10 (2) 0.019 (3)

ρe =

qc 2 16πre

(2) −15

where re = 2.85 × 10 m is the classical radius of the electron. It can be used without any bias to determine the global porosity ε of the layer since ρ ε=1− e ρe (3)

sampled at regular intervals, centrifuged at 13000 rpm for 1 min, and diluted in 0.2 M HNO3 (analytical grade) prior to evaluating the elementary releases by inductively coupled plasma atomic emission spectrophotometry (ICP-AES, Spectro Arcos). Simultaneously, the pellet was removed from the solution, washed three times in 50 mL of deionized water (18.2 MΩ cm) to stop dissolution, and gently dried on an absorbant towel. The pellet was then weighed before characterization of the solid solution interface by XRR and GI-XRD. In order to compare with the chemical durability of powdered samples, the release of an element i from the material was described by its normalized mass loss, NL(i) (g m−2), and by the associated normalized dissolution rate, RL(i) (g m−2 h−1), calculated as follows:36 NL(i) =

Δmi fi A

and

RL(i) =

dNL(i) dt

s

in which ρes (e− Å−3) is the calculated electron density of the solid. [N.B., electron density and mass density are linked through ρes = ρsNA(∑jxjZj/∑xjMj), where xj (dimensionless) is the mole ratio of element j in the solid, Zj (e−) the atomic number, Mj (g mol−1) the molar mass, and NA (mol−1) Avogadro’s number.] GI-XRD measurements were performed using the same Bruker D8 diffractometer and primary optics adopting the Bragg−Brentano geometry. Secondary optics were composed of long Sollers slits of 0.12° and of a point detector. GI-XRD patterns were collected between 15° and 60° (2θ mode) with a 0.02° step. The angular domain was limited at higher angles by the beam knife-edge that equipped the reflectivity stage. The incident angle, θi, was fixed to a value ranging from 0.2° to 1°. By increasing the incident angle, it was possible to scan a layer of increasing thickness at the surface of the material. The penetration depth, z (nm), of X-rays (E = 8.048 keV) as a function of θi is linked to the material absorption coefficient, μ (cm2 g−1), and the material density, ρ (g cm−3). For a compound material, the mass absorption coefficient is obtained from the sum of the absorption cross sections of the constituting atoms by

(1)

where Δmi (g) is the mass of element i released in solution, A (m2) is the solid surface area, and f i (dimensionless) is the mass fraction of i in the solid. Thereafter, the reaction of dissolution was considered to be congruent when all the normalized dissolution rates were identical (i.e., when all the elements were released with the same ratios as the stoichiometry of the initial material). For a congruent dissolution, the mass of solid dissolved, Δm (g), can thus be calculated from the concentration of any constitutive element measured in solution. The mass of solid dissolved was thus estimated from the total concentrations of Ce and Nd measured in solution and compared to the mass loss evaluated by weighing the pellet. XRR measurements were performed using a Bruker D8 Advance diffractometer equipped with a motorized reflectivity stage that allows vertical translation of the sample. The complete primary optics setup for XRR analysis was composed of a Cu Kα1,2 (λ = 1.541 84 Å) source, a Göbel mirror, a motorized divergence slit, a fixed 0.2 mm slit, an automatic absorber, a fixed 0.1 mm slit after the absorber, and 2.5° Sollers slits. The secondary optics included a motorized antiscattering slit, a graphite monochromator, 2.5° Sollers slits, a 0.05 mm receiving slit, and a point detector. Standard θ−2θ scans were used for data collection. Step sizes and counting times were adapted to the intensity collected by the detector. The resolution of this setup was estimated to 0.03°. The background was determined for each XRR measurement using the off specular scan method37 starting at 2θ = 0.7° with δ(θ) = 0.2° and subtracting from the collected intensity. Geometric correction in the region of the plateau of total external

μ = NA

∑j xjσj ∑j xjMj

(4) −1

where σj (cm atom ) is the total atomic absorption cross section derived from CXRO X-ray Data Booklet.39 As the density of the solid decreased during the dissolution experiments, the penetration depth of X-rays as a function of θi was also updated. The qualitative analysis of recorded patterns confirmed the presence of a single Ia3̅ phase attributed to the Ce1−xNdxO2−x/2 solid solution with x = 0.59 and 0.63. Further refinements were performed using the Thomson−Cox−Hastings pseudo-Voigt function convoluted with an axial divergence asymmetry function40 after the extraction of the instrumental function using pure silicon. In the Le-Bail refinement (profile matching) using the Fullprof Suite program,41 all the profile parameters were considered, including microstructural broadening, and lead to good agreement factors. As an example, the refinement 2

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concentrations and the associated normalized mass loss obtained for each pellet is presented in Figure 2. Ce and Nd were clearly released congruently during leaching tests of P1 in 4 M HNO3 at 60 °C (Figure 2a). The three different estimations of NL (from the weighings of the pellet and from the elementary analysis of Ce and Nd) gave almost the same value of the normalized dissolution rate of the solid (RL = 0.33 ± 0.01 g m−2 h−1) that remained almost constant during all the dissolution experiment (i.e., 107 h). Moreover, this value appeared in good agreement with that determined by Horlait et al.21 when leaching powdered sample of Ce0.41Nd0.59O1.70 under the same conditions (i.e., 0.15 ± 0.01 g m−2 h−1) despite the uncertainties associated with the specific surface area determination for sintered pellets. A total of 34.5% of P1 was dissolved after 107 h in 4 M HNO3 at 60 °C. This observation suggested that neither precipitation of neoformed phases nor formation of protective layers acting as diffusion barrier for the elementary releases occurred in such experimental conditions. The results obtained in 0.1 M HNO3 at 60 °C, are presented in Figure 2b. In this medium, the normalized dissolution rate of the solid calculated from the evolution of the normalized mass loss of the pellet reached RL = 4.8 ± 0.3 × 10−3 g m−2 h−1. Under these conditions, only 1.1% of the mass of P2 was dissolved after 126 h of leaching test. This value was significantly lower than the RL value obtained in 4 M HNO3,

of the XRD pattern obtained for pellet P1 before the dissolution experiment is plotted in Figure 1.

Figure 1. Observed and calculated patterns of the pellet P1 at t = 0 h.

3. RESULTS AND DISCUSSION 3.1. Dissolution Tests: Ce and Nd Releases Observed at the Macroscopic Scale. The evolution of Ce and Nd

Figure 2. Evolution of the concentrations of Ce and Nd and the associated normalized mass losses, NL(Ce) and NL(Nd), determined during the dissolution of Ce0.37Nd0.63O1.68 pellet in 4 M HNO3 at 60 °C (a) and of Ce0.41Nd0.59O1.70 pellet in 0.1 M HNO3 at 60 °C (b). 12030



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showing the dependency of the normalized dissolution rate on the proton activity. On the basis of the data reported by Horlait et al.,20,21 the normalized dissolution rate should reach 1.15 ± 0.1 × 10−2 g m−2 h−1 when leaching powdered Ce0.4Nd0.6O1.7 in 0.1 M HNO3. Considering the uncertainties associated with the measurement of RL value, the result obtained in this work appears in good agreement with those obtained on powdered samples, even though the proton activity effect seems to be more pronounced for sintered samples. In addition, the results obtained in 0.1 M HNO3 (Figure 2b) showed different evolutions of total concentrations and normalized dissolution rates with the dissolution time. During 70 h, a steady state was observed for the normalized mass loss determined from Nd, and then the normalized dissolution rate for Nd increased with RL(Nd)/RL(Ce) > 3, as a consequence of an incongruent dissolution mechanism. This specific behavior of Nd could have suggested that the release of trivalent lanthanide incorporated in the solid solution was enhanced. However, if the dissolution mechanism was unchanged, incongruent dissolution should have been also observed in 4 M HNO3, i.e., for higher reaction progress. The existence of a very short step (t < 1.5 h) of incongruent dissolution could be also suspected in 4 M HNO3, but this is hard to evidence. In this sense, a higher normalized dissolution rate RL(Nd) compared to RL(Ce) at the beginning of the dissolution could indicate that the surface of the solid was initially neodymium-enriched compared to the bulk material. Bellière et al.42 observed such chemical separation within a crystalline particle of cerium−lanthanum oxides prepared by the solidstate route by combining electron energy-loss spectroscopy (EELS) with scanning transmission electron microscopy (STEM). In this case, the solid-state route led to separation between enriched and depleted phases regarding the loading element. The formation of an Nd-enriched surface layer could also have occurred during the first 60 h of the dissolution. Similar observations were made during the dissolution of wollastonite under acidic conditions.31 Using electron microprobe measurements, the authors evidenced the formation of a Ca-depleted layer at the wollastonite/water interface, although they do not favor a particular mechanism of formation of this layer: migration through the depleted crystal structure or surface reprecipitation. An accurate analysis of the structure of the solid/solution interface will provide clues to explain this apparent specific behavior of Nd at the beginning of the dissolution. 3.2. GI-XRD Measurements: Structural Changes at the Surface of Dissolving Ce0.4Nd0.6O1.7 Mixed Oxides. GIXRD is the most appropriate method to evidence dynamic structural changes at the surface of crystallized materials.43 GIXRD patterns with θi = 1° were then recorded regularly during the dissolution of each pellet. The patterns recorded for P1 (4 M HNO3) and P2 (0.1 M HNO3) are presented in parts a and b of Figure 3, respectively. The unit cell volume corresponding to each pattern was extracted from the profile matching procedure (see Supporting Information, Table S1). For θi = 1°, a regular decrease of the unit cell volume with the dissolution time was recorded whatever the alteration conditions. Qualitatively, it could be associated with the neodymium depletion in the solid solution. Indeed, the change in the unit cell parameter of Ce1−xNdxO2−x/2 solid solutions, when plotted with respect to the Nd loading, follows a second-order polynomial law in both fluorite (F-type CeO2) and C-type structure, i.e., the oxygen-

Figure 3. GI-XRD patterns recorded with θi = 1° for P1 (a) and P2 (b) samples and for various dissolution times.

vacancy superstructure of the former one (with a double unit cell parameter). 9,12 Ikuma et al. 12 derived from their experimental results the following empirical relation between the unit cell parameter a (Å) and the Nd mole ratio x: a = −0.0847x2 + 0.2047x + 5.4111. This relation was selected to calculate the Nd mole ratio from the unit cell volume, because the temperature used to prepare the oxides (1400 °C) was similar to the sintering temperature for P1 and P2 samples. The evolution of Nd and Ce mole fractions in the pellet subsurface is presented in Figure 4. The reported thickness, z, is equal to the penetration depth of X-rays for θi = 1°. The initial Nd and Ce mole fractions measured by GI-XRD and X-EDS were not significantly different. Consequently, the initial average composition of 200 nm of the pellets subsurface was close to the composition of the bulk material (X-EDS analysis depth reached 1 μm). The density of the leached layer decreased significantly during dissolution in 4 M HNO3 at 60 °C (XRR results, Figure 7a). Thus, the penetration depth of Xrays (z) increased strongly for P1 with the dissolution time (Figure 4a), as z reached 0.9 μm after 110 h of dissolution. The interpretation of the results reported in Figure 4a was thus 12031



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Figure 4. Evolution of Ce and Nd mole ratios determined at the surface of the pellet dissolved in 4 M HNO3 (a) and 0.1 M HNO3 (b). GI-XRD measurements were performed for θi = 1°. The corresponding X-rays penetration depths are shown for (a) P1 and (b) P2, respectively. For X-EDS, the analysis depth reached 1 μm.

layer was not sufficient to influence significantly the composition averaged over 1 μm determined by X-EDS. In order to examine the surface of the P2 pellet, the Nd profile was built by performing GI-XRD measurements with lower grazing incident angles (0.2° ≤ θi ≤ 1°) and with the same dissolution time. The evolution of the surface composition during dissolution is reported in Figure 5. The results showed the existence of a composition gradient over 200 nm from the surface of the pellet. The Nd-enriched zone was located at the surface of the Nd-depleted zone evidenced by GIXRD measurements with θi = 1°. However, this separation was not evidenced in the initial Nd mole fraction profile for P2 sample. Thus, the chemical gradient was formed during the first times of the dissolution and then evolved toward a steady state. GI-XRD measurements revealed Nd enrichment at the surface of the sample. However, each data reported in Figure 5 corresponded to the average composition of the whole layer probed by GI-XRD. The existence of such a chemical gradient could be explained by the presence of a thin Nd-rich surface layer above the Nd-depleted zone. XRR is the appropriate

complicated due to the GI-XRD profile matching procedure that averaged the composition of a layer increasing in thickness. However, the general trend corresponded to progressive Nd depletion down to the composition of the leached layer (x ≈ 0.53). At the most advanced stage of dissolution, the thicknesses probed by both techniques were almost the same (≈1 μm) and both GI-XRD profile matching and X-EDS analyses gave almost the same composition for the leached layer. Using less aggressive conditions (0.1 M HNO3, Figure 4b), the variation of the density of leached layer was less important. For θi = 1°, z slightly increased from 190 to 220 nm when increasing the dissolution time and Nd depletion was also observed in this zone. The two last data points showed a slowdown in the composition evolution, meaning that the steady composition was approached after 120 h of dissolution (x ≈ 0.51). This slowdown was also detectable from the evolution of the normalized mass loss of P2 (Figure 2b). After 100 h of dissolution, the thickness of the Nd-depleted leached 12032



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technique to characterize thin films and thus gave additional information on the properties of this surface layer. 3.3. XRR Measurements: Surface Microstructural Changes of the Sintered Pellets during Dissolution. Combined with others techniques, XRR was reported to be a useful technique to determine some porous properties in mesoporous silica thin films,44 density and thickness of La2O3− HfO2 thin films deposited on Si wafers during their dissolution,45 and protective properties and dissolution ability of the gel formed during nuclear glass alteration.46 To our knowledge, this technique was never used up to the present to characterize the evolution of surface microstructure of sintered ceramics during their dissolution. Indeed, in this case either structural and microstructural changes occurred at the solid/ solution interface, due to the loss of mass of solid during dissolution. The evolution of density, pore size distribution, and surface roughness is suggested by ESEM images in Figure 6a. From this point of view, XRR seemed to be the most appropriate technique, as it gives access to a 2D-characterization of the solid/solution interface properties. The XRR curves monitored during dissolution of Ce1−xNdxO2−x/2 in 4 M

Figure 5. Evolution of Nd mole fraction profiles during the dissolution of P2 pellet in 0.1 M HNO3.

Figure 6. (a) Experimental XRR curves obtained during the dissolution of Ce1−xNdxO2−x/2 sintered samples in 4 M HNO3 (P1) and ESEM micrographs of the surface exposed to solution for t = 0 and t = 92 h. (b) Experimental (symbols) and simulated (lines) XRR curves of the P2 pellet dissolved in 0.1 M HNO3. 12033



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the leached layer. The porosity of both surface layer and leached layer was determined from their electron density (Figure 8b). The results showed that the surface layer was also dissolving. After a transient step, the rates of formation and dissolution of the surface layer reached a steady state, leading to the stabilization of the layer thickness and porosity. However, the porosity was still slowly increasing within the leached layer, and the spatial resolution of the ESEM apparatus did not allow observing such evolution. 3.4. Discussion: Coupled Structural and Microstructural Evolution of the Dissolving Surface. Previous studies devoted to the kinetics of dissolution of mixed dioxides were performed on powdered samples and underlined the effect of the nature of the trivalent cation and the incorporation rate20,21 in the solid solutions, pH, temperature,47,48 or complexing reagents.49 Few explored the effect of the structure and microstructure of sintered samples. Hingant et al.17 and Claparede et al.,19 however, pointed out the key role played by the size of the crystallite, the density of grain boundaries, the porosity, and the grain size on the resistance of the materials to aqueous alteration. In the present work, the techniques used allowed studying simultaneously the dynamics of structural and microstructural changes of sintered polycrystalline materials under various alteration conditions. It showed the joint influence of these changes on the macroscopic dissolution rate of the solid. The understanding of the evolution of the solid/solution interface properties was thus improved. A dissolution mechanism of the studied pellets is proposed in Figure 9. On the basis of the results obtained through the different techniques involved in this work, it includes the description of the surface evolution with three steps and presents the corresponding schematic reconstructions of the interface during dissolution. In order to slow down the progress of the dissolution reaction, experiments were performed in 0.1 M HNO3 at 60 °C. Under these conditions, it was possible to focus the study on the two first steps (1 and 2) of the dissolution mechanism proposed. Such results were thus representative of the very beginning of the dissolution process. Under these conditions, the comparison of the normalized mass losses NL(Ce) and NL(Nd) suggested that the dissolution was initially incongruent, with a preferential neodymium release (step 2), after a period where no release of Nd was observed (step 1). Surface structural characterization by GI-XRD and microstructural characterization by XRR gave some key information to improve the understanding of such unexpected results. The first step (mass loss lower than 1%) was associated with the development of a surface layer of 40 Å of low electron density at the solid/ solution interface (XRR data). Given its very thin thickness and the angular resolution of the diffractometer, it was not possible from GI-XRD results to identify the nature of this Nd-based surface layer. Below this layer, the Nd mole ratio sharply decreased. From 150 nm to less than 1 μm, the solid was Nddepleted compared to the raw material. The mechanism responsible for the formation of the Nd-based surface layer is not resolved. Migration through the depleted crystal structure and formation of a diffusion barrier at the solid/solution interface was envisaged. However, surface precipitation of neoformed phases such as hydrated Nd oxides or (oxo)hydroxydes could not be excluded due to their very low solubility [for example, log Ks°(Nd(OH)3) = −23.34 ± 0.5550] and considering that local high activity values could be

HNO3 (P1) and in 0.1 M HNO3 (P2) are presented in parts a and b of Figure 6, respectively. In 4 M HNO3, neither the second critical angle nor oscillations that would be associated with the presence of a surface layer were observed in these curves. The “nanoscale” roughness of the solid/solution interface was modified, leading to the rapid decrease of the intensity of the reflected signal. After 57 h of dissolution time, the plateau of total reflection was not observed and XRR measurement became impossible. Besides, the rapid alteration of the pellet surface gave rise to a decrease of the critical angle, related to the decrease of the electron density at the surface due to the grain boundaries and pore openings. The data were fitted with a model accounting for the evolution of the electron density and roughness of the leached layer (Figure 7). The porosity, ε, was obtained from eq

Figure 7. Evolution of the porosity and roughness determined by XRR of the altered surface of Ce1‑xNdxO2‑x/2 (P1) during dissolution in 4 M HNO3 and comparison with the porosity obtained from ESEM images analysis.

3, in which ρe is the fitted electron density of the leached layer and ρes is the value calculated from the crystallographic data. ε increased strongly during the first 20 h of dissolution, then the evolution slowed down. The two last values were rather uncertain. The porosity of the altered surface could reach a constant value after 60 h of dissolution. The resolution of ESEM was not sufficient to evidence such evolution of the leached layer porosity. The solid/solution interface evolved more slowly when using less aggressive conditions. It was thus possible to observe microstructural changes in more detail. However, as the time scale of the experiment increased, these observations concerned only the beginning of the dissolution process. After 61 h in 0.1 M HNO3, XRR curves revealed the apparition of an oscillation between 0.06 and 0.12 Å−1, signing the development of a thin layer at the surface of the pellet (Figure 6b). At the same dissolution time, GI-XRD measurements show the formation of a Nd-enriched zone at the surface of a Nd-depleted zone. The presence of a thin surface layer of a Nd-bearing phase could explain both results. The evolution of the parameters fitted on XRR curves during dissolution in 0.1 M HNO3 is presented in Figure 8a. The thickness, τ (Å), of the low electron density surface layer increased within the first 60 h of dissolution and then reached a constant value of 40 Å. Simultaneously, the electron density of this surface layer decreased, as well as the electron density of 12034



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Figure 8. (a) Parameters obtained from the fit of XRR data. (b) Evolution of the porosity of the surface layer and leached layer deduced from the values of the electron densities fitted on XRR curves.

encountered at the solid/solution interface in the case of low solution renewal. During the second step (mass loss lower than 3%), the Ndenriched surface layer was also dissolved and developed porosity at very small scale, leading to a strong increase of the surface contributing to dissolution in the Nd-rich surface layer. The Nd-rich surface layer thus mainly contributed to dissolution, which explains that incongruent dissolution was observed as a transient stage to congruent dissolution. The third step (mass loss higher than 3%) was mainly evidenced during dissolution experiments in 4 M HNO3 at 60 °C (P1). It corresponded to the establishment of the steady state. The determination of the normalized mass loss showed that the dissolution was congruent regarding the release of Nd and Ce in solution. The normalized dissolution rate under such aggressive conditions was high and comparable to that determined on powdered samples of the same composition. The microstructure of the solid/solution interface evolved

quickly with the formation of corrosion pits and large macropores of a few micrometers (ESEM) that rapidly made surface characterization by XRR impossible. However, it seemed that the surface porosity stabilized to a high value and that the dissolution progressed in depth by gravity and capillarity within the porous network. Structural changes in the first layer of grains were also evidenced thanks to GI-XRD. The Nd and Ce mole ratios evolved substantially at the time scale of the experiment. After 100 h of dissolution in 4 M HNO3, the composition of the first layer of grains reached a steady value: xNd ≈ 0.53 and xCe ≈ 0.47. As this composition was Nddepleted compared to the raw material, the dissolution rate was lower in this zone and explained why the dissolution progressed in the underneath raw material. As the dissolution occurred mainly in a zone of the solid where the composition was not modified, the rate of dissolution was comparable with the results obtained on powdered samples for xNd ≈ 0.6 and xCe ≈ 0.4. As the leached layer above the alteration front dissolved 12035



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Figure 9. Proposed schematic structural and microstructural evolution of the solid/solution interface during dissolution of Ce0.4Nd0.6O1.7 sintered samples by the proton-promoted surface mechanism.

Combination of GI-XRD and XRR experiments revealed that the dissolution mechanism occurred in three steps. After a short dissolution period, the composition of the surface of the Ce0.4Nd0.6O1.7 sintered pellet was clearly Nd-enriched compared to the bulk material. This result had a crucial influence on the dissolution kinetics under weakly acidic conditions and explained the apparent incongruent dissolution observed as an initial transient step. However, this effect disappeared as soon as the porosity of the material increased. It seems now important to gain understanding of the mechanism responsible for the Nd-enrichment of the material surface, in order to control the surface composition at the beginning of the dissolution.

slower, no sensitive reduction in thickness of the pellet was detected on the time scale of the dissolution experiment.

4. CONCLUSION From a methodological point of view, the kinetic monitoring of the dissolution of sintered ceramic pellet by GI-XRD and XRR is relevant when the surface of the pellet is not too damaged. For a porous sample, this condition is fulfilled at the beginning of the dissolution experiment, i.e., for a very short period under aggressive conditions. Indeed, beyond 3% of dissolved material, the geometry of the solid/solution interface of porous material becomes too complex to be probed by these techniques. However, in the case of a surface proton promoted dissolution mechanism, the dissolution rate can be strongly reduced in less acidic dissolution media. Thus, structural and microstructural changes at the solid/solution interface can be monitored. These techniques gave access to surface composition, porosity, and roughness during dissolution. Surface topography and surface porosity are linked with the surface area exposed to solution and, thus, indicate the dynamics of surface area changes during dissolution. Mole fractions in the leached layer and surface area are necessary to accurately estimate normalized dissolution rates and to compare the chemical durability of materials of different composition and microstructure.

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ASSOCIATED CONTENT

S Supporting Information *

Figure S1 and Table S1. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest. 12036



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ACKNOWLEDGMENTS The authors would like to thanks Diane Rebiscoul (CEA/ DEN/DTCD) for helpful discussions concerning XRR data acquisition and treatment, Bruno Corso, Valérie Magnin, Olivier Diat (ICSM/LDD), and Nicolas Massoni (CEA/ DEN/DTCD) for assistance in the use and tuning of the diffractometers, and Véronique Dubois (ICSM/LTSM) for ICP-AES analysis. This work benefited from financial support of the French National Research Agency (ANR, project # ANR-08-BLAN-0216) and from the MATINEX French Research Group (Innovative Materials in Extreme conditions, CEA/CNRS/AREVA/EDF/French Universities) included in the PACEN program.

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Kinetics of Structural and Microstructural Changes at the Solid

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