Nanoscale Heterogeneities and Oxygen Storage Capacity of Ce

J. Phys. Chem. B 2003, 107, 13007-13014

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Nanoscale Heterogeneities and Oxygen Storage Capacity of Ce0.5Zr0.5O2 E. Mamontov,*,†,§ R. Brezny,‡ M. Koranne,‡ and T. Egami†,| Department of Materials Science and Engineering, UniVersity of PennsylVania, Philadelphia, PennsylVania 19104-6272, and W. R. Grace & Co.-Conn., Grace DaVison Research, Columbia, Maryland 21044-4098 ReceiVed: May 28, 2003; In Final Form: August 25, 2003

Nanocrystalline powders of Ce0.5Zr0.5O2 were synthesized and characterized for the application as a catalyst support and oxygen storage medium in automotive catalysis. Detailed pair-distribution function analysis of the neutron diffraction data revealed an unusual structure of the nanocrystallites, consisting of domains of Ce0.4Zr0.6O2 composition of ∼25-30 Å in size in a matrix of Ce0.7Zr0.3O2. The domains cannot be detected by conventional diffraction analysis because the crystallographic structure and orientation of the atomic planes are the same for all ceria- and zirconia-enriched regions within the crystallite, which therefore gives rise to Bragg scattering as a whole. The oxygen storage capacity did not show a correlation with either crystallite size or surface area of the samples. We suggest that the smaller size of zirconia-enriched domains associated with larger interfacial area between ceria- and zirconia-enriched regions is responsible for the increase in oxygen storage capacity.

1. Introduction Alloying CeO2 into ZrO2 or vice versa often leads to improving technologically important characteristics of both oxides. For instance, the toughness of ceria-doped ZrO2 is considerably higher than that of undoped or yttria-doped zirconia.1-3 In automotive emission control catalysis, adding ZrO2 to catalytically active CeO2 improves oxygen storage capacity (OSC) and thermal stability of ceria owing to suppression of thermal sintering4 and stabilization of oxygen defects.5 The compositional dependence of the structure of (Ce1-xZrx)O2 has been studied extensively,6-14 and there are some detailed studies of the structure-property relationships in ceria-zirconia pertinent to its OSC.15 Despite extensive efforts, the phase diagram of the CeO2ZrO2 system is not free from controversy because of the presence of metastable phases depending on the thermal history. At room temperature, CeO2 is cubic, and ZrO2 is monoclinic, with space groups Fm3m (no. 225) and P21/c (no. 14), respectively. Although the only truly thermodynamically stable phases in (Ce,Zr)O2 are cubic and monoclinic, ceria-zirconia phases of intermediate composition are characterized by overall tetragonal symmetry. Three tetragonal phases, all of space group P42/nmc (no. 137), with tetragonality increasing with zirconia contents, are known as t, t′, and t′′.9-11 Precise positions of the boundaries between monoclinic, t, t′, t′′, and cubic phases depend on the thermal history of the system. Although the formation of metastable phases makes obtaining a reliable phase diagram of ceria-zirconia difficult, it provides the opportunity for optimizing the properties of the CeO2-ZrO2 system by stabilizing the desired structure. * To whom correspondence should be addressed. E-mail: mamontov@ seas.upenn.edu. Phone: (215) 898-7944. Fax: (215) 573-2128. † University of Pennsylvania. ‡ W. R. Grace & Co.-Conn. § Present address: NIST Center for Neutron Research, Gaithersburg, MD 20899-8562. E-mail: [email protected]. Phone: (301) 975-6232. Fax: (301) 921-9847. | Present address: Departments of Materials Science/Physics and Astronomy, University of Tennessee, Knoxville, TN 37996-1508.

In this work, we report the synthesis and structural characterization of nanocrystalline powders of Ce0.5Zr0.5O2 for the application as a catalyst support and oxygen storage medium in three-way automotive catalytic converters. Catalytic converters are designed to achieve simultaneous reduction of NOx and oxidation of CO and hydrocarbons. The important role of ceriazirconia is to stabilize the local oxygen partial pressure at the catalyst surface as the air-to-fuel ratio in the engine exhaust fluctuates around the stoichiometric ratio. Thus, the property of interest in this study is the OSC, which is known to vary among differently prepared ceria-zirconia samples of identical composition and crystallographic structure. The OSC of the samples we studied did not exhibit a correlation with either the crystallite size or surface area. In search of the origin of the OSC variations, we focused on the nanoscale structure of ceriazirconia. The limitations of conventional diffraction analysis preclude accurate assessment of nanoscale compositional variations, which nevertheless can be investigated using the pairdistribution function (PDF) structural analysis. The PDF analysis reveals unusual nanoscale intracrystallite compositional heterogeneities associated with the formation of zirconia-enriched domains in the ceria-enriched matrix. Both the crystallographic structure and orientation are identical for all ceria- and zirconiaenriched regions within the crystallite, which differentiates the domains from those observed previously and makes them impossible to study using conventional diffraction analysis. The spatial extent of the nanoscale compositional heterogeneities varies among samples prepared by different processes, and we try to relate the domain size with the OSC. 2. Experimental Section 2.1. Synthesis. The samples referred below as samples 1, 2, and 3 were synthesized as follows. For sample 1, water and acetic acid were mixed and cerium carbonate was added to form a clear solution of cerium acetate. The mixture was stirred for 48 h to completely dissolve cerium carbonate. Zirconium acetate was added to the cerium acetate and stirred to make a homogeneous solution, which was then spray dried at 110 °C

10.1021/jp030662l CCC: $25.00 © 2003 American Chemical Society Published on Web 11/01/2003

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to form a powder of mixed acetates. The powder was calcined in a muffle furnace at 500 °C for 1 h to form the Ce0.5Zr0.5O2 oxide. Preparation of sample 2 involved dissolving cerium carbonate into zirconyl nitrate and concentrated nitric acid. The resulting solution was stirred to completely dissolve the carbonate. The acid solution was then added to concentrated ammonia to induce precipitation. The precipitate was washed with water and spray dried prior to calcining at 500 °C for 1 h. To prepare sample 3, a cerium nitrate solution was obtained by dissolving cerium carbonate in water and nitric acid. Zirconyl nitrate was added to the cerium nitrate to form a mixed nitrate solution. The nitrate solution was coprecipitated with 5 N ammonia, and the precipitate was filtered, washed, and spray dried. The dried product was calcined at 500 °C for 1 h to form the final Ce0.5Zr0.5O2 product. To simulate the accelerated aging of ceria-zircoinia catalyst supports in an automobile catalytic converter, the samples were calcined at 800 °C for 2 h. After the thermal aging, the BET surface area was 1, 20, and 50 m2/g for samples 1, 2, and 3, respectively. 2.2. Characterization of OSC. The OSC of samples 1-3 aged at 800 °C was determined by thermal gravimetrical analysis (TGA). Before the measurements, the samples were impregnated with 20 ppm of Pd and were held in flowing air at 500 °C for 1 h to remove residual water. The weight loss was measured in the course of flowing a mixture of 10% H2 in nitrogen at 500 °C. The weight loss was then converted into the oxygen contents assuming the initial oxygen concentration in the sample of 100%. The slope of the TGA plot shows the oxygen release rate, and the amount of oxygen desorbed from the sample defines its OSC. 2.3. Structural Characterization and Data Analysis. Neutron diffraction experiments were carried out at the Special Environment Powder Diffractometer at the Intense Pulsed Neutron Source (IPNS) of Argonne National Laboratory using the time-of-flight method. The Rietveld analysis of the diffraction data over the d-spacing range of 0.65-4.5 Å (1.4 Å-1 < Q < 9.7 Å-1) was carried out using the GSAS program16 that performs least-squares fitting of the diffraction data in reciprocal space with a model diffraction pattern. The background was modeled using a polynomial function and was subtracted from the diffraction pattern. The peak profile was modeled using a two-sided exponential function to describe the instrumental contribution, which was convoluted with a linear combination of Gaussian and Lorentzian functions to describe the sample contribution. The overall scale factor, background parameters, peak shape parameters, sample extinction and absorption parameters, unit cell parameters, positions of the oxygen ions, and isotropic thermal factors were refined. The structure factor, S(Q), was obtained from the diffraction pattern by means of subtracting the separately measured background and applying corrections for absorption, inelastic, and multiple scattering.17 Fourier transformation was applied to S(Q) to obtain the PDF in real space:

G(r) ) 4πr[F(r) - F0] )

∫0∞ Q[S(Q) - 1] sin(Qr) dQ

2 π

(1)

where F(r) is the microscopic pair density describing the probability of finding an atom at a distance r beginning from another atom, F0 is the average number density, and Q ) (4π/λ)sin(2θ/2) is the magnitude of the scattering vector. Note that in practice the highest accessible scattering vector Qmax ) (4π/λmin)sin(2θmax/2) and a high-energy probe with low λmin

such as synchrotron X-rays or spallation neutrons from a pulsed neutron source is needed in order to reduce the termination error resulting from truncating the integration in eq 1 at Qmax. The PDF analysis was performed using the PDFFIT program,18,19 which carries out least-squares fitting of the PDF data in a real space with a model G(r). We refined unit cell parameters, positions of the oxygen ions, isotropic thermal factors, and overall scale factor. For the structural models comprising more than one phase, the phases scale factors were refined independently, thereby yielding the partial phase concentrations. The neutron scattering lengths used in the Rietveld and PDF analyses were 0.484 × 10-12 cm for cerium, 0.716 × 10-12 cm for zirconium, and 0.580 × 10-12 cm for oxygen. The agreement factor, χ2, used as a measure of goodness of fit of the experiment with the model is defined in a similar manner in both Rietveld and PDF analyses N

χRietveld2 )

w(Qi)[Iexp(Qi) - Imod (Qi)]2 ∑ i)1 Nobs - Npar

(2)

N

χPDF2 )

w(ri)[Gexp(ri) - Gmod (ri)]2 ∑ i)1 Nobs - Npar

(3)

where Nobs and Npar are the numbers of independent observations in the data and independent parameters of the model, respectively. The residuals of the diffraction pattern intensities, I(Qi) and PDF intensities, G(ri) are taken with the appropriate weights, w(Qi) and w(ri). In PDF analysis of the data collected up to a maximum scattering vector value of Qmax performed for Rmin < r < Rmax, the number of independent observations can be evaluated as (Rmax - Rmin)Qmax/π. 20,21 In Rietveld analysis, the number of independent observations is estimated as the range of Q chosen for the analysis divided by the instrument resolution. 3. Results The results of TGA measurements using 10% H2 in N2 at 500 °C to determine the OSC of the Ce0.5Zr0.5O2 samples are presented in Figure 1. In the absence of a precious metal catalyst, the dissociative chemisorption of hydrogen onto the surface of the oxide particles is the rate-limiting step.22 For oxide samples of low surface area, the rate of oxygen release is so low that it precludes measuring their true OSC. Adding a precious metal such as rhodium,22 or in the present study palladium, greatly facilitates the dissociative chemisorption of hydrogen and decreases the reduction time. This allows the comparison of the OSC among samples of greatly different surface area, even though the rate of the oxygen release may still be dependent on the sample surface. Indeed, the initial rate of the oxygen release (the slope of the curves in Figure 1) clearly exhibits correlation with the surface area, being highest for sample 3 and lowest for sample 1. However, unlike the reduction rate, the OSC determined as the total amount of oxygen desorbed from the sample shows no correlation with the surface area, being highest for sample 2 and lowest for sample 3. The total amount of the desorbed oxygen was 2.13, 2.23, and 1.82 g (O2) per mole of Ce0.5Zr0.5O2 for samples 1, 2, and 3, respectively. Given the fact that the composition of the samples is identical, the difference in the OSC is likely associated with their structure. On the basis of the phase diagram,10,11 the crystal structure of

Nanocrystalline Powders of Ce0.5Zr0.5O2

J. Phys. Chem. B, Vol. 107, No. 47, 2003 13009

Figure 1. Oxygen contents as a function of time in the TGA measurement of OSC using 10% H2 in N2 at 500 °C. The inset displays the expanded data measured during the first 150 s. Figure 3. Rietveld refinement of the crystallographic structure. Measured (symbols) and fitted (solid line) data are shown along with the difference plots (below) and agreement factors (above).

Figure 2. Atomic structure of tetragonal (Ce,Zr)O2 phases, space group P42/nmc. Arrows show displacements of oxygen ions from their positions in cubic (Ce,Zr)O2.

Ce0.5Zr0.5O2 is expected to be tetragonal, with space group P42/ nmc, as shown in Figure 2. The structure of tetragonal ceriazirconia phases, known as t, t′, and t′′, is closely related to the fluorite structure of cubic (Ce,Zr)O2. Compared to the fluorite structure, the oxygen ions surrounding 8-coordinated cations are displaced from their original positions along the [001] direction. This is because the ionic radius of Zr4+ (0.84 Å) is too small compared to that of Ce4+ (0.97 Å) to accommodate eight oxygen ions.23 Thus, Zr4+ normally prefers to be 7-coordinated.24,25 The correlated oxygen displacements are associated with elongation along the [001] direction in the t and t′ phases

(c/a ) 1.018 and 1.010, respectively). On the other hand, the quasicubic t′′ phase with c/a ) 1 is frequently viewed as “cubic” in X-ray diffraction measurements. The t′′ phase is readily distinguished from the true cubic phase by neutron diffraction which is more sensitive to the positions of oxygen ions. Correlated displacements of oxygen ions along the [001] direction result in the presence of extra peaks in the neutron diffraction pattern. In particular, a weak peak at about 3 Å-1 in Figure 3 is indicative of correlated oxygen displacements as found in t′′ quasicubic ceria-zirconia. This peak is absent in the neutron diffraction pattern of true cubic ceria-zirconia.5 Figure 3 shows the results of the Rietveld analysis of the crystallographic structure for the three samples. Good agreement between the experiment and a t′′ structural model (the agreement factor χ2 ∼ 3-4) confirms that all of the Ce0.5Zr0.5O2 samples belong to the quasicubic t′′ phase, with space group P42/nmc and c/a ) 1. No peak splitting that would indicate the presence of two phases could be detected, and therefore, the diffraction pattern demonstrates the formation of a single, solid-solutionlike ceria-zirconia phase with lattice constants of 5.29, 5.28, and 5.27 Å for samples 1, 2, and 3, respectively. The crystallite sizes, as determined from the width of the diffraction peaks, are 160, 115, and 105 Å for samples 1, 2, and 3, respectively. There is no correlation between the OSC, as shown in Figure 1, and the crystallite size. The lack of correlation between the crystallite size and oxygen storage properties (steady-state reactivity for CO oxidation) in ceria and ceria-zirconia was observed previously by Gorte and co-workers.26 In summary, our Rietveld analysis suggests that the crystallographic structures of the samples are very similar, and provides no explanation of the varying OSC. To examine the structural information in further detail, the PDFs were obtained from the diffraction data. The PDF is obtained by the Fourier transformation of the total scattering factor S(Q) that includes not only Bragg peaks but also diffuse scattering intensities. Thus, the information about local distor-


Mamontov et al.

Figure 4. PDF analysis of the nanoscale structure using a constrained c/a ratio and single-phase structural model. Measured (symbols) and fitted (solid line) PDF data are shown along with the difference plots below. The agreement factors with one standard deviation are presented.

Figure 5. PDF analysis of the nanoscale structure using a constrained c/a ratio and two-phase structural model. Measured (symbols) and fitted (solid line) PDF data are shown along with the difference plots below. The agreement factors with one standard deviation are presented.

tions that give rise to diffuse scattering is retained, thereby making PDF much more sensitive to local structural deviations.27 To investigate the structural correlation over different ranges of interatomic distances, the PDF analysis was performed independently over several 10 Å wide slices of distances covering the range of 1.5-41.5 Å. At first, we used a singlephase structural model similar to the one determined by the Rietveld analysis. We allowed the c/a ratio to assume the values of 1.018, 1.010, or 1, which are characteristic values of t, t′, and t′′ ceria-zirconia phases, respectively,12-14 and chose the best fit. We found that the model with the c/a ratio of 1.018 never yielded a good fit to the data. This is consistent with the fact that the t phase (c/a ) 1.018) is observed for low concentrations of ceria, quite far away from the composition of our samples, Ce0.5Zr0.5O2. The results of the PDF fitting using the single-phase solid-solution model are presented in Figure 4 along with the difference curves and agreement factors with one standard deviation. The agreement at low r is rather poor as suggested by the high values of χ2. This prompted us to test a different structural model comprising two tetragonal phases belonging to P42/nmc space group with independent sets of lattice parameters, a1 and c1 for the first phase and a2 and c2 for the second phase. This model represents the scenario in which two separate phases coexist and form a physical mixture, in contrast to the single-phase model describing the formation of a solid solution. Because lattice parameters depend on the (Ce/Zr) ratio due to different ionic size of Ce4+ and Zr4+, the difference in lattice parameters characterizes compositional variation between two phases. As in the case of the single-phase modeling, the c1/a1 and c2/a2 ratios were allowed to assume the values of 1.018, 1.010, or 1, and the best fit was chosen. Similar to the results of the single-

phase modeling, a good fit was never obtained with c/a ) 1.018 for either of the two phases. The results of the PDF analysis based on the two-phase model are presented in Figure 5 along with the difference curves and agreement factors with one standard deviation. The fitting yields much lower agreement factors compared to the single-phase model for shorter interatomic distances up to 21.5 Å, especially in the range of 1.511.5 Å. At higher r, the difference between the agreement factors for the single- and two-phase models is statistically indistinguishable. Note that in the computing of the agreement factors the number of independent parameters, which is different for single- and two-phase models, has been accounted for as shown in eq 3. The relatively high standard deviations of the agreement factors preclude precise determination of the range over which the two-phase “physical mixture” model better describes the structure. A more precise determination can be achieved by analyzing the r dependence of the lattice parameters as shown in Table 1. Also given in the table are the phase fractions as determined from the two-phase analysis. The lattice constants of the two phases obtained in the modeling are clearly different at low r, but this difference decreases at higher interatomic distances. At the highest values of r (31.5-41.5 Å), one of the two phases seems to deviate from this trend. However, this is simply an artifact because the data in this range is adequately described by a single-phase model. Note the similarity of the lattice parameters between the single-phase model and phase 2 in the two-phase model over the range of 31.5-41.5 Å. Furthermore, the fraction of phase 1 drops almost to zero. These demonstrate that at high interatomic distances the two-phase model invariably converges into a single-phase model for all of the samples.

Nanocrystalline Powders of Ce0.5Zr0.5O2

J. Phys. Chem. B, Vol. 107, No. 47, 2003 13011

TABLE 1: Results of the PDF Analysis Using a Constrained c/a Ratioa fitting range

1.5-11.5 Å

11.5-21.5 Å

21.5-31.5 Å

31.5-41.5 Å

5.3024 ( 0.0004 5.3024 ( 0.0004

5.2855 ( 0.0003 5.3383 ( 0.0003

a, Å c, Å

5.3128 ( 0.0013 5.3128 ( 0.0013

Sample 1, Single-Phase Model 5.3069 ( 0.0004 5.3069 ( 0.0004

fraction of phase 1 a1, Å c1, Å fraction of phase 2 a2, Å c2, Å

59% 5.2470 ( 0.0060 5.2990 ( 0.0060 41% 5.3390 ( 0.0050 5.3930 ( 0.0050

Sample 1, Two-Phase Model 59% 5.2620 ( 0.0060 5.3140 ( 0.0060 41% 5.3210 ( 0.0050 5.3750 ( 0.0050

52% 5.2683 ( 0.0009 5.3210 ( 0.0009 48% 5.3045 ( 0.0009 5.3576 ( 0.0009

3% 5.1580 ( 0.0040 5.1580 ( 0.0040 97% 5.2863 ( 0.0005 5.3391 ( 0.0005

a, Å c, Å

5.2812 ( 0.0015 5.3340 ( 0.0015


5.2715 ( 0.0005 5.3242 ( 0.0005

5.2722 ( 0.0005 5.3249 ( 0.0005


71% 5.2420 ( 0.0060 5.2940 ( 0.0060 29% 5.3580 ( 0.0060 5.3580 ( 0.0060


29% 5.2580 ( 0.0050 5.2580 ( 0.0050 71% 5.2830 ( 0.0070 5.3350 ( 0.0070

3% 5.2190 ( 0.0010 5.2190 ( 0.0010 97% 5.2731 ( 0.0040 5.3258 ( 0.0040

a, Å c, Å

5.3049 ( 0.0016 5.3049 ( 0.0016


5.2982 ( 0.0005 5.2982 ( 0.0005

5.2817 ( 0.0008 5.3345 ( 0.0008


60% 5.2340 ( 0.0080 5.2860 ( 0.0080 40% 5.3550 ( 0.0060 5.3550 ( 0.0060


45% 5.2670 ( 0.0070 5.2670 ( 0.0070 55% 5.3000 ( 0.0050 5.3530 ( 0.0050

4% 5.2124 ( 0.0014 5.2645 ( 0.0014 96% 5.2823 ( 0.0008 5.3352 ( 0.0008

a Lattice parameters are shown with one standard deviation, and the overlapping values of lattice constants at high r are typed in boldface. For sample 2, the values of lattice parameters of the single- and two-phase models in the range of 21.5-31.5 Å overlap not with each other but with the corresponding values in the range of 31.5-41.5 Å.

Although for all of the samples this convergence is practically complete in the range of 31.5-41.5 Å, sample 2 demonstrates the onset of the convergence at shorter interatomic distances (21.5-31.5 Å) exhibiting the most significant reduction of the fraction of phase 1. Besides, sample 2 is the only one for which the lattice constants in this range are already similar to those in the range of 31.5-41.5 Å. These suggest that in sample 2 the convergence between single- and two-phase models occurs at somewhat lower interatomic distances compared to those of samples 1 and 3. Because the PDF peaks are broadened by thermal and quantum oscillations, efforts have to be made in the PDF modeling to find a physically sound structural model rather than the one that merely yields the lowest agreement factor. When we repeated the same set of the single- and two-phase modeling procedures as described above but using nonconstrained c/a ratios, we persistently obtained a ratio of c/a < 1 at low values of r. Despite this problem with the nonconstrained modeling at low interatomic distances, at higher r, the results agreed with the trend observed in the constrained modeling. The detailed results of the nonconstrained modeling are available in the Supporting Information. In particular, in the nonconstrained modeling for sample 2, the crossover between the single- and two-phase models occurs in the range of 31.5-41.5 Å. At these interatomic distances, the lattice parameters of the dominant phase in the two-phase model become similar to those of the single-phase model, and the fraction of the second phase drops to just 1%. In samples 1 and 3, the lattice parameters for singleand two-phase models do not yet converge in the range of 31.541.5 Å, the highest r range of this study. Furthermore, the

presence of the minor phase in the two-phase modeling is still appreciable (8% for sample 1 and 10% for sample 3). To summarize, the PDF modeling with constrained c/a ratio shows the crossover where a two-phase “physical mixture” structural model converts into single-phase “solid-solution” model for all of the samples. For sample 2, this crossover is attained at lower values of r. In the nonconstrained PDF modeling, the crossover apparently shifts to higher r for all of the samples. The crossover is clearly observed only in sample 2, whereas samples 1 and 3 exhibit signs of approaching convergence between the two- and single-phase models as demonstrated by the decrease in the fraction of the second phase. The overall conclusion is that regardless of the details of the modeling procedure all of the samples demonstrate the presence of two phases at low r but appear single-phase at higher r. In sample 2, the crossover between two-phase and single-phase structural models is attained at lower interatomic distances. 4. Discussion The existence of two (Ce,Zr)O2 phases with different lattice parameters detectable only at short interatomic distances suggests a nanodomain structural model of Ce0.5Zr0.5O2 crystallites as shown in Figure 6. Within the 100-150 Å crystallites, there exist CeO2- and ZrO2-enriched regions. The lattice constants of these regions are slightly different owing to the difference in the ionic radii of Ce4+ and Zr4+. However, the crystallographic orientation of the atomic planes is unchanged within the crystallite as depicted by the grid of the lattice planes in the figure. The crystallite thus gives rise to the Bragg scattering as a whole, resulting in diffraction peaks characteristic of a single phase.


Figure 6. Proposed structure of nanocrystallites in the Ce0.5Zr0.5O2. The grid shows a set of atomic planes. Because of the difference in ionic radii of Ce4+ and Zr4+ the lattice constants of Ce1-X1ZrX1O2 and Ce1-X2ZrX2O2 are slightly different. However, both the crystallographic structure and orientation of the atomic planes are the same for all of the ceria- and zirconia-enriched regions within the crystallite.

Unlike the diffraction pattern, the PDF is sensitive to the local variations of the lattice constants. If a particular pair of atoms lies within a region, either ceria- or zirconia-enriched, the distance between such atoms is proportional to the lattice constant of the material in this region. On the other hand, if two atoms lying across the domain boundary are probed, the distance between them is some weighted average of the lattice constants of ceria- and zirconia-enriched regions, depending on how far from the domain boundary the atoms lie. In our study, at small values of r, the PDF predominantly samples the pairs of atoms contained within the ceria-enriched or zirconia-enriched regions, thus making the material appear a physical mixture of the two independent phases with different sets of lattice constants. As r is increased, the PDF is progressively more likely to probe atom pairs spanning across domain boundaries that leads to progressive convergence of the lattice constants of the two phases in PDF modeling. At even larger values of r, the PDF exclusively probes the distances between atoms that belong to different regions, and the PDF pattern starts looking like that of a single phase, with an average lattice constants of the “solid solution”. The position of the crossover between the two- and single-phase models is a measure of spatial extension of the compositional heterogeneities, which was found to be smaller in sample 2 compared to those of samples 1 and 3. To estimate the degree of ceria-zirconia compositional variation, we calculated the unit cell volume using the lattice constants in Table 1 over the range of 1.5-11.5 Å and compared it with the compositional dependence of the unit cell volume in (Ce1-xZrx)O2 available from the literature.15 For all of the samples, the composition of the zirconia-enriched phase was found to be approximately Ce0.4Zr0.6O2, whereas the composition of the ceria-enriched phase was found to be approximately Ce0.7Zr0.3O2. Using the molar fraction of the phases from Table 1, one calculates the overall composition of Ce0.52Zr0.48O2 for sample 1, Ce0.49Zr0.51O2 for sample 2, and Ce0.52Zr0.48O2 for sample 3, which is in agreement with the nominal composition of Ce0.5Zr0.5O2. Although the PDF analysis does not directly yield information on the topology of ceria- and zirconia-enriched regions, the fact that the fraction of the zirconia-enriched phase in the two-phase modeling decreases continuously at larger distances is consistent with a model of Ce0.4Zr0.6O2 domains embedded in the Ce0.7Zr0.3O2 matrix. The domain size can be derived from the position of the crossover between two- and single-phase models and is estimated as ∼25 Å in sample 2 and ∼30 Å in samples 1 and 3. Whereas for all of the samples the models fully

Mamontov et al. converge at 31.5-41.5 Å, the models begin to converge at shorter distances (21.5-31.5 Å) for sample 2, but not for samples 1 and 3. The smaller domain size in sample 2 means that over the range of 21.5-31.5 Å the PDF probes more atomic pairs including one atom in a Ce0.4Zr0.6O2 region and the other in a Ce0.7Zr0.3O2 region. Thus, over the range of 21.5-31.5 Å, the PDF pattern of sample 2 is somewhat more solid-solutionlike compared to those of sample 1 and 3. We would like to emphasize that the intracrystallite domain structure we found in this study is rather unusual and different from the matrix-inclusions structure previously observed in multicomponent zirconia-containing systems.28-31 For example, inclusions of the monoclinic phase of CaZr4O9 were observed in cubic (Ca,Zr)O2.28,29 Following the annealing at T > 1000 °C, the size of inclusions grew from ∼50 to ∼300 Å. Further annealing of (Ca,Zr)O2 was associated with the formation of the monoclinic or tetragonal zirconia inclusions about 1000 Å in size.30,31 In all of the previous studies, the crystallographic structure of the inclusions was different from that of the matrix, and the inclusions manifested themselves in extra peaks in the X-ray diffraction pattern or extra spots in the electron diffraction pattern. In the present study, the domains could not be detected by the diffraction measurements because they were only different in composition but not in structure or lattice orientation. The domains found in this study may be better described as regions of nanoscale compositional heterogeneity rather than inclusions. The fact that the crystallographic structure and lattice orientation are identical for all of the CeO2- and ZrO2-enriched regions within the crystallite follows from the presence of a single set of diffraction peaks corresponding to the average structure rather than two sets of diffraction peaks. Furthermore, if the domains were not structurally aligned with the matrix the PDF would become grossly amorphous at the distances exceeding the domain size but less than the crystallite size. Instead, the coherence of the lattice planes was obviously preserved across domain boundaries, even at high r, where the loss of lattice constant correlation led to the appearance of a single solid-solution-like set of lattice constants. Therefore, not only the crystallographic structure but also the atomic plane orientation was the same within the crystallites. It is worthwhile to note that this study clearly illustrates the strength of PDF analysis as a technique for studying the structural features at the nanoscale level, where conventional diffraction techniques fail due to intrinsic limitations. In earlier studies,we have proposed the existence of pure CeO2 and ZrO2 domains in the nanocrystallites of Ce0.7Zr0.3O2.32 The alleged domain structure in the Ce0.7Zr0.3O2 sample was associated with a higher OSC compared to that of solid-solution-like Ce0.7Zr0.3O2 and a physical mixture of 70% of CeO2 and 30% of ZrO2. The results of the present study are consistent with our earlier findings and are more quantitative and detailed. It is thus reasonable to attribute the highest OSC exhibited by sample 2 to the differences in the domain structure, in particular, the smaller size of zirconia-enriched domains. A decrease in the domain size results in an increase in the interfacial area between the ceria- and zirconia-enriched regions. At a given volume fraction of inclusion phase, the interfacial area is inversely proportional to the domain size. Thus, a decrease in domain size from 30 to 25 Å corresponds to a 20% increase in the interfacial area leading to an increase of the fraction of oxygen near the interface. A larger interfacial area may enhance both the short-time OSC related to the oxygen kinetic and total OSC that depends on the thermodynamic of reduction. The increase in the short-time,

Nanocrystalline Powders of Ce0.5Zr0.5O2 kinetic OSC may result from a higher mobility of oxygen ions near the boundary between CeO2- and ZrO2-enriched regions due to a higher concentration of lattice strain at this interface. As for the total, thermodynamic OSC, it was found both theoretically33,34 and experimentally35-38 that the reduction enthalpy in ceria is much lower at the surface/grain boundary compared to the bulk. Thus, it is likely that a higher interfacial area due to a smaller domain size is associated with larger amount of structural oxygen which is not merely more mobile but less stable thermodynamically. Indeed, as one can see in Figure 1, sample 2 loses the largest amount of oxygen, thus exhibiting the highest thermodynamic OSC. During the initial stage of reduction shown in the inset of Figure 1, the kinetic OSC seems to correlate with the surface area of the samples rather than domain size. After a few tens of seconds, the oxygen loss by sample 2 becomes somewhat faster compared to that by sample 3 despite the larger surface area of the latter, whereas the oxygen loss by low surface area sample 1 remains slow. It is known that a high surface area CeO2 shows two peaks in temperature-programmed reduction measurements, and the low-temperature peak originates from the surface reduction.39 The presence of a low-temperature peak in the temperatureprogrammed reduction profile of a high surface area Ce0.5Zr0.5O2 has been also reported and attributed to the surface reduction.40 Thus, the initial oxygen loss in our TGA experiment is likely due to surface reduction. The bulk oxygen is not involved, and its mobility is irrelevant to the reduction kinetic. As the reduction progresses into the bulk, the mobility of the bulk oxygen becomes important, and sample 2 possessing larger interfacial area starts showing a faster reduction kinetic. For low surface area sample 1, the surface processes such as dissociative adsorption of hydrogen or water desorption apparently remain the rate-limiting step, and the bulk oxygen mobility cannot be deduced from the reduction rate. In view of the previous finding41 that in pure CeO2 bulk oxygen mobility does not limit reduction at T > 650 K, a faster reduction rate exhibited by sample 2 after the initial surface reduction is quite surprising. Although it would be logical to anticipate a correlation between thermodynamic OSC and the interfacial area, our results are insufficient to establish such a correlation. In particular, our structural analysis cannot distinguish between sample 1 and 3, which show substantial difference in the OSC. In principle, analyzing the data over narrower slices should provide more precise information on the domain size. This would make it possible to find if the domain size in sample 1, which shows an intermediate thermodynamic OSC, falls between those in sample 2 and 3. In practice, however, the decreasing number of independent observation in the PDF pattern makes it more difficult to distinguish between different structural models. Our choice of 10 Å wide slices was a compromise between the accuracy of determination of the crossover position and the sufficient information contents of the PDF. The only conclusion that we can draw from the data analysis is that the sample with the smallest domain size, that is, the largest interfacial area, shows the best OSC. This finding suggests that minimization of the size of intracrystallite domains should benefit oxygen storage properties of nanocrystalline ceria-zirconia. Because a catalyst support is subjected to severe thermal aging, it is especially important to impede the growth of the domains at high temperature to sustain OSC. Thus, the design of an ideal ceria-zirconia material for the application in automotive catalysis should additionally include a mechanism to suppress domain growth.

J. Phys. Chem. B, Vol. 107, No. 47, 2003 13013 5. Conclusion Nanocrystalline powders of Ce0.5Zr0.5O2 were synthesized by different techniques, and their OSC was measured using TGA analysis. The structure of the samples was characterized using pulsed neutron scattering. Although diffraction analysis suggested the formation of a complete ceria-zirconia solid solution, detailed PDF analysis revealed a unique intracrystallite, nanodomain structure consisting of domains of Ce0.4Zr0.6O2 composition of ∼25-30 Å in size embedded in a matrix of Ce0.7Zr0.3O2. The domains are better described as regions of nanoscale (Ce/Zr) compositional heterogeneity rather than second phase inclusions because the crystallographic structure and the orientation of the atomic planes remain unchanged within the crystallites. As a result the crystallite gives rise to Bragg scattering as a whole, which makes it impossible to detect the domains by conventional diffraction analysis. On the other hand, the PDF is sensitive to the domain structure due to the local variations of the lattice parameters caused by the difference in Ce4+ and Zr4+ ionic radii. The OSC did not correlate with the surface area or the crystallite size. Instead, our data suggest that the highest OSC is associated with the smaller size of zirconia-enriched domains and may result from the larger interfacial area between Ce0.4Zr0.6O2 and Ce0.7Zr0.3O2 regions. Acknowledgment. This work has benefited from the use of the Intense Pulsed Neutron Source at Argonne National Laboratory which is funded by the U.S. Department of Energy, BESMaterials Science, under Contract W-31-109-ENG-38. The authors are thankful to anonymous referees for raising some interesting questions concerning reductive behavior of ceriazirconia. Supporting Information Available: The PDF analysis using nonconstrained c/a ratio: single-phase structural model (Figure 1S), two-phase structural model (Figure 2S), and summary of the results (Table 1S). This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) (2) (3) (4)

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Nanoscale Heterogeneities and Oxygen Storage Capacity of Ce

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