Observing Local Oxygen Interstitial Diffusion in Donor-Doped Ceria by

Article pubs.acs.org/JPCC

Observing Local Oxygen Interstitial Diffusion in Donor-Doped Ceria by 17O NMR Relaxometry Ralf Heinzmann,† Ibrahim Issac,† Jens-Peter Eufinger,‡ Gregor Ulbrich,§ Martin Lerch,§ Jürgen Janek,*,‡ and Sylvio Indris*,∥ †

Institute of Nanotechnology, Karlsruhe Institute of Technology, P.O. Box 3640, 76021 Karlsruhe, Germany Institute of Physical Chemistry, Justus Liebig University Gießen, Heinrich-Buff-Ring 58, 35392 Gießen, Germany § Institut für Chemie, Technical University Berlin, Straße des 17. Juni 135, 10623 Berlin, Germany ∥ Institute for Applied Materials−Energy Storage Systems, Karlsruhe Institute of Technology, P.O. Box 3640, 76021 Karlsruhe, Germany ‡

ABSTRACT: The local structure and oxygen ion dynamics in pure CeO2 and CeO2 doped with Ta and Gd is studied by 17O nuclear magnetic resonance (NMR) spectroscopy, Raman spectroscopy, and conductivity measurements. 17O magicangle spinning nuclear magnetic resonance spectra of CeO2 doped with Ta show two oxygen environments. A major peak at 876 ppm represents the oxygen ions on the regular crystal lattice site with four Ce4+ neighbors. A second small peak at 751 ppm (1−2% area fraction) represents point defects that can be ascribed to interstitial oxygen ions. The concentration of these point defects increases with the dopant concentration. Oxygen ion dynamics was studied by variable-temperature spin− lattice relaxation time measurements on pure and Ta-doped CeO2 powders. The obtained results reveal a fast local jump process in Ta-doped CeO2 with a jump rate of about 2.5·108 s−1 at 350 K. This local hopping process is assigned to interstitial oxygen ions trapped by tantalum dopand ions.

1. INTRODUCTION Cerium dioxide (CeO2−x) is a long studied nonstoichiometric phase, which is also one of the best studied mixed ionic/ electronic conductors and oxygen storage materials with excellent chemical and mechanical stability.1−3 Acceptordoped ceria is a superior solid electrolyte4−6 and plays an important role as electrolyte for intermediate temperature solid oxide fuel cells (SOFC). Its fluorite-type structure is depicted in Figure 1.1,7 The cerium ions form a face-centered cubic lattice and the oxygen ions occupy all tetrahedral sites in case of perfect 1:2 stoichiometry, thus the lattice provides only a single cation and a single anion site with high symmetry. Because of the ideal tetrahedral symmetry around the oxygen ions the chemical shielding as well as the quadrupolar interaction would be orientation-independent in the ideal crystal.8 Substitution of cerium ions by dopants like Ta5+, Gd3+, and/or Zr4+ results in distortions of the local symmetry due to differences in the ionic radii, valence state, and electronic configuration (Table 1).9 The crystal structure tolerates chemical reduction (i.e., oxygen deficiency) without phase change in particular at high temperatures, and ceria is therefore often written as CeO2‑x. The investigation of ceria-based materials covers a wide field of applications. As one example for its catalytic applications, CeO2 is a key component in the three-way catalyst (TWC) for the treatment of exhaust gas from automobiles.10,11 The use of ceria in TWC relies on the beneficial effect of precious metal− © 2016 American Chemical Society

Figure 1. Fluorite-type structure of CeO2. Oxygen ions are shown as red spheres, and cerium ions are shown as white spheres.

ceria interaction and on the activity of the redox couple CeIV− CeIII with its ability to change from CeIV under oxidizing conditions to CeIII under reducing conditions and vice versa.1 CeO2 also has the ability to induce secondary processes that can affect the catalyst performance,12 i.e., influence the Received: April 1, 2016 Revised: April 6, 2016 Published: April 7, 2016 8568

DOI: 10.1021/acs.jpcc.6b03341 J. Phys. Chem. C 2016, 120, 8568−8577

Article

The Journal of Physical Chemistry C

CeO233,36 as a consequence of nuclear irradiation. However, the migration of interstitial oxygen has only been subjected in few experimental studies,37,38 but not for fluorite-type oxides. The interstitial diffusion in the latter was only considered theoretically because of a critical drawback: Unfortunately, the distinction between the parallel vacancy and interstitial migration in oxides is neither possible with conductivity nor with tracer experiments. Here we demonstrate that NMR allows identification of interstitial motion. Nuclear magnetic resonance (NMR) is a versatile tool to study local structures as well as dynamics in condensed matter.8,39−43 Local structures can be investigated element- and even isotope-specifically. This includes the local coordination,44,45 variations in the bond lengths/angles,46 the presence of paramagnetic neighbors,47,48 and electric field gradients at the sites of the nuclei.49−51 The techniques to study dynamics in condensed matter include temperature-dependent line shape analysis,45,52 NMR relaxometry,44,53 and field-gradient NMR.54 In this paper we report on 17O NMR experiments investigating the local environments and the dynamics of oxygen ions in undoped and doped CeO2. In general, oxygen NMR investigations are scarce because the only NMR-active isotope of oxygen, 17O, is also the least abundant (0.037%) of all the three stable isotopes (16O: 99.76%, 18O: 0.2%) making 17 O one of the more challenging quadrupolar nuclei (I = 5/2) in NMR spectroscopy. Therefore, isotopic enrichment of 17O was required in this study. In addition, magic-angle spinning (MAS) is often not able to completely remove the secondorder quadrupolar broadening which leads to broad spectral features and complex powder line shapes. Recent studies with (direct or indirect) dynamic nuclear polarization (DNP) showed a strong signal-to-noise enhancement even with natural abundance of 17O NMR in solids. This was achieved via a microwave-induced direct or indirect transfer of polarization from electronic spins in stable radicals to 17O nuclear spins.55 Unfortunately, this technique is only applicable at low temperatures, which is unapplicable in the present case of mobile oxygen at elevated temperature. Previous NMR investigations on CeO2 were restricted to samples doped with Y2O3,56−59 La,60 or Sc.61 In the present paper, we combine results from phenomenological experiments on the long-range charge transport with results from spectroscopic experiments on the local microscopic charge carrier dynamics in pure and Ta-doped ceria (CTO). Because of its significant defect concentration ([Oi″] ≈ 3 × 1019 cm−3) CTO is well suited to study the interstitial oxygen diffusion. 17O NMR visualizes changes in the local (electrical) environments around oxygen ions. Spin−latticerelaxation measurements on Ta-doped CeO2 show oxygen hopping in the vicinity of tantalum ions. As counterpart to donor-doped ceria and for comparison, Gd-doped ceria was examined with the same methods.

Table 1. Effective Ionic Radii in 8-Fold Coordinated Environment and Electronic Configuration of Ce4+ and the Dopants Ti4+, Ta5+, Gd3+, and Zr4+ 4+

Ce Ti4+ Ta5+ Gd3+ Zr4+

radius (Å)

electronic configuration

0.97 0.74 0.74 1.053 0.84

[Xe] [Ar] [Xe] 4f14 [Xe] 4f7 [Kr]

dispersion of supported metals, promote the water gas shift and steam reforming reactions, increase the thermal stability of the support, and promote noble metal reduction/oxidation. In addition to its catalytic properties, pure CeO2 is also known as a high temperature electrical conductor with the conductivity proportional to pO21/6. Its point defect properties have been intensively studied and are well understood.13−18 The transference number for oxygen ions of undoped CeO2 is less than 0.002 between 873 and 1573 K, and thus, pure CeO2 is primarily an electronic conductor. The minor ionic conductivity of undoped ceria at 1273 K and pO2 = 10−6 bar is less than 3% of the total conductivity. The n-type electronic conductivity is based on small polaron transport; i.e., the electronic carriers are locally trapped as a consequence of the displacement of adjacent atoms.19 These localized electronic defects move by a thermally activated hopping process as proposed by Tuller et al. The activation energy of 0.4 eV for the electron hopping process increases with increasing oxygen deficiency x in CeO2‑x, i.e., with the increasing concentration of electrons. Investigations on the electrical conductivity of nanocrystalline ceria also show that the grain boundary regions are more easily reduced than the bulk of the crystal.20−22 At 1073 K an electron hole mobility of 10−5 cm2 V s−1 was found in the high oxygen partial pressure region which is lower compared to the electron mobility of 6.1 × 10−3 cm2 V s−1 at low oxygen partial pressures. In general, in undoped ceria, the ionic conductivity is always lower than the electronic conductivity. By doping with divalent or trivalent metals (Ca2+, Gd3+, Sm3+) the ionic transference number can be increased such that applications as oxygen solid electrolyte are possible. The majority charge carriers in these systems are oxygen vacancies,4,6,23 which are formed to compensate the reduced positive charge of the di- and trivalent dopant ions. Less attention has been paid to donor doped ceria (Ta5+, Nb5+) due to the limited applications.24−27 It is a predominant electronic conductor although the majority defects are oxygen interstitials at ambient atmosphere and temperatures below 800 K. Since the mobility of the oxygen interstitials is lower than the electronic mobility, the electronic charge transport is dominating. Ceria doped with 1% TaO2.5 shows an activation energy of 0.2 eV for the electrical conductivity which is significantly lower than that of pure ceria.26 Interstitial oxygen species were found in several oxides, e.g. La2NiO4+δ,28 ZnO,29 Sr6−2xNb2+2xO11+3x,30 Bi2O3,31 and tin-doped In2O3,32 and were proposed in pointdefect models of fluorite-type oxides, too. However, in fluoritetype oxides they have not been detected by means of crystallographic techniques although Bi2O3 and In2O3 crystallize in the bixbyite structure, a defective superstructure of the cubic fluorite structure. Theoretical calculations predict the presence of interstitial oxygen in UO2,33−35 ThO2,33,36 and

2. EXPERIMENTAL SECTION 2.1. Sample Preparation and Phase Analysis. Pure and doped ceria single crystals were grown by the skull-melting technique.62 The respective oxides (CeO2, Chempur 99.9%; Ta2O5, Chempur 99.99%; TiO2, Chempur 99.5%; Gd2O3, Chempur 99.9%) were homogenized in a powder mixer. After melting the premixed powders, single crystals were grown by directed solidification. The homogeneity of dopant ion distribution and the elemental composition (Table 2) were quantified by electron beam microprobe (Gennevilliers Cedex 8569


Article


measured via impedance spectroscopy in a two-electrode setup. Therefore, the as-prepared crystals were cut into rectangular bars with dimensions of approximately 1 × 1 × 3 mm3 on whose ends silver conducting paste was painted to improve the electrical contact. These samples were mounted on a spring loaded sample holder equipped with two parallel silver electrodes inside a high temperature tube furnace. The temperature was varied in a range from 1073 K down to 423 K while at each temperature the sample was allowed to equilibrate for 4 h before the measurement was started. The impedance measurement was performed with a Princeton Applied Research EG&G Potentiostat/Galvanostat (model 283) in combination with a model 1025 Frequency Response Analyzer in the frequency range from 500 kHz to 100 mHz. The excitation amplitude was kept at 20 mV for all measurements. The obtained impedance spectra were evaluated with the program package PowerSine. The conductivity of the samples was then calculated via the eq 1

Table 2. Compositions of the Obtained Crystals sample label

composition

CeO2 10% CGO 0.5% CTO 2.5% CTO

Ce0.991Zr0.009O2‑δ Ce0.889Gd0.105Zr0.006O2‑δ Ce0.9832Ta0.0046Ti0.0007Zr0.0115O2‑δ Ce0.957Ta0.025Zr0.018O2‑δ

CAMECA SAS). After grinding the large crystals (obtained in a range between 1 × 1 × 3 mm3 and 5 × 5 × 20 mm3) to powders they were characterized by X-ray powder diffraction (Siemens D5000) supporting exclusively the presence of the expected fluorite-type phases. As the skull-melting method itself needs a small piece of metal for ignition, every single crystal contains small amounts of zirconium impurities (1−2%, see Table 2). The sample denoted as 0.5% CTO was attempted to prepare with an additional Ti concentration. However, the solubility of Ti turned out to be very small, and thus, only a small content of Ti was found in the respective crystal. In the following, this crystal will be considered as “pure” (Ce,Ta)O2 as the defect properties resulting from the very small Ti content are far dominated by the defects formed due to Ta as majority donor dopant. Furthermore, since Zr4+ as well as Ti4+ are isovalent to Ce4+, their influence on the defect properties is practically negligible. 17 O enrichment was performed by heating a quartz tube containing a small quantity of the sample to 873 K in a 17O2 environment (Sigma-Aldrich, 90 at%) for 96 h. XRD patterns were measured on an STOE STADI P diffractometer (Mo Kα radiation, curved germanium (111) monochromator, Debye− Scherrer geometry) in sealed glass capillaries. 2.2. Raman Spectroscopy. In order to further investigate the local vibrational modes of the single crystals, Raman spectroscopy was applied. The as-prepared single crystals were partially reduced due to the high temperature during the growth process. The partial reduction did become apparent from a color darkening. To obtain a uniform oxidation state all single crystal samples were oxidized in oxygen atmosphere at 873 K for 12 h. Raman measurements were carried out at room temperature and ambient atmosphere with a Bruker Senterra Raman microscope. The excitation source was a green Ar laser with a wavelength of 532 nm operated at a power of 2 mW. All obtained spectra were normalized to the very intense F2g mode of CeO2 at 463−465 cm−1. 2.3. 17O MAS NMR Spectroscopy. 17O NMR measurements were performed on pure and doped CeO2 powders after 17 O enrichment. These measurements were carried out on a Bruker Avance 600 MHz spectrometer (B0 = 14.1 T) using 2.5 mm zirconia rotors in dry nitrogen atmosphere. 17O-enriched water was used in all NMR measurements as the reference for the chemical shift of 17O (0 ppm). The typical values for the recycle delay and the π/2 pulse length were 10 s and 5 μs, respectively. 17O MAS NMR experiments were performed at 298 K and a spinning speed of 25 kHz with a single-pulse sequence. 17O variable-temperature NMR relaxation time measurements were accomplished on a Bruker Avance 300 MHz spectrometer (B0 = 7.1 T) between 220 and 673 K under static conditions. A saturation-recovery as well as an inversion− recovery pulse sequence were used to measure spin−lattice relaxation times (T1). The recycling delay for all relaxation time experiments was set to 5T1. 2.4. Electrical Conductivity. The total electrical conductivity of the pure and doped Ceria single crystals was

σ=

l A×R

(1)

where σ denotes the total conductivity, l the electrode distance, A the electrode area, and R the measured resistance. The obtained total electrical conductivity is the sum of the partial ionic and electronic conductivities, and further information is required to distinguish both contributions. Therefore, we additionally determined the oxygen ion transference number of selected sample compositions by using a galvanic concentration cell method.27,63,64 2.5. Ionic Transference Number. Measuring the open cell voltage of a mixed ionic/electronic conductor exposed to a chemical potential gradient is a common method for determining the ionic transference number. According to Wagner’s transport theory65 the open circuit voltage E for a given chemical potential difference of oxygen is given by E=−

1 4F

∫μ

μOII

2

I O2

t O2− dμO

2

(2)

where μO2 denotes the chemical potential of oxygen, tO2− the transference number of oxygen ions, and F the faraday constant. If the chemical potential difference of oxygen is small (to be ensured by small ratios of oxygen partial pressures during the experimental procedure) the transference number tO2− can be considered as constant and eq 2 is approximated into the following well-known equation: E = −t O2− ×

pOII RT ln I 2 4F pO 2

(3)

In case of a sample with a transference number smaller than unity, ionic and electronic currents will flow through the sample, and redox reactions will take place at the electrodes in order to level out the chemical potential difference. Therefore, highly reversible electrodes are required as they have to be well equilibrated with the surrounding gas atmosphere. In general, porous platinum meets this demand. Since the obtained single crystals were too small to form a membrane between two gas compartments, the as prepared single crystals were embedded into an alumina matrix. This procedure, the sample preparation for the EMF measurement and the experimental setup has been reported in a previous publication.27 8570


Article

The Journal of Physical Chemistry C The oxygen partial pressure was fixed at a value of pIO2 = 0.69 × 105 Pa in the left compartment and pIIO2 = 0.072 × 105 Pa in the right compartment by using appropriate Ar/O2 mixtures leading to a (logarithmic) mean oxygen partial pressure of pO̅ = 0.22 × 105Pa . The temperature was stepwise varied in a 2

range of 1100 ≥ T/K ≥ 600 while monitoring the open cell voltage and the oxygen sensor signals. When a steady state for one temperature step was reached, the transference number of oxygen ions was calculated from the respective voltage values.

3. RESULTS AND DISCUSSION 3.1. Phase and Structure Analysis. XRD measurements were performed on the samples before and after enrichment with 17O (Figure 2). The XRD patterns show no changes confirming that the crystal structure of the samples is maintained.

Figure 3. Magnified Raman spectra of ceria single crystals with different dopants measured at ambient atmosphere and room temperature. Inset: Detail section of the very intense F2g mode around 463−465 cm−1 present for all compositions. All spectra were normalized to this F2g mode.

investigated compositions and are assigned to other Ce−O vibrational modes as described by Weber et al. in more detail.67 Since Raman spectroscopy is very sensitive to local disorder and deviations from the ideal crystal lattice, small concentrations of dopant or impurity ions will have significant effect on some features. Group theoretical calculations predict one additional mode at 272 cm−1, experimentally found at 264 cm−1.67 Neither in the pure CeO2 nor in the doped crystals this mode is present. It may be superimposed by other additional modes at 277 cm−1, only appearing in crystals with Ta5+ doping (CTO), and 240 cm−1 which is present in all samples. To our knowledge these modes rather arise from a change in the phonon-dispersion due to the small dopant concentration of Ta5+ and Zr4+ than from a local mode of a Ta−O or Zr−O vibrations. But regardless of their origin, the modes at 277 and 240 cm−1 can be considered as “fingerprint” for Ta and Zr in CeO2, respectively. In the spectrum of the crystal doped with Gd (10% CGO) two observations can be made. First, the main F2g mode is shifted down (red-shift) to 463 cm−1 with respect to the undoped CeO2. As doping with Gd3+ introduces oxygen vacancies into the lattice the average coordination number of Ce is reduced also affecting the force constant. Therefore, a slight red-shift of the F2g mode can be observed.66 Second, in the range from 100 to 300 cm−1 three broad bands at 123, 190, and 250 cm−1 on a high background can be noticed. According to Nakajima et al. these are defect-induced Raman bands involving oxygen vacancies.69 Under consideration of metalvacancy and oxygen-vacancy defect clusters of the type M4Ov (oxygen vacancy surrounded by four nearest-neighbor metal ions) and O6Ov (oxygen vacancy surrounded by six nextnearest-neighbor oxygen ions, ignoring the surrounding four metal ions) the observed broad bands can be well explained. Finally, it is worth to emphasize that tantalum has significant impact on the local structure and environment of the oxygen ions. The observation of separate signals in Raman spectra only for Ta-doped CeO2 may also imply interstitials as a different oxygen species. 3.2. 17O MAS NMR Measurements. 17O MAS NMR experiments performed on nominally pure CeO2 and Ta-doped CeO2 reveal a very sharp symmetric NMR signal at 877 ppm (Figure 4a) and thus a highly symmetric environment around

Figure 2. XRD patterns of CeO2 doped with Ta before and after enrichment with 17O in comparison to the database pattern.

Ancillary to the routine crystal structure and phase analysis via X-ray diffraction (XRD) Raman spectroscopy offers high sensitivity to local as well as to long-range ordering in both the cation and the anion sublattices. In particular, changes in the local oxygen vacancy distribution cannot be resolved by XRD, but are apparent in Raman spectroscopy.66 Thus, Raman spectroscopy is very useful to probe small impurities and small dopant contents in mixed electrolyte systems. CeO2 crystallizes in the cubic fluorite-type structure, which only reveals one very strong triply degenerated first-order Raman active mode at 463−465 cm−1 with F2g symmetry (Figure 3). This vibrational mode can be regarded as a symmetric three-dimensional stretching of oxygen around each metal cation center.66−68 Moreover, several second-order features with low intensities and primarily A1g symmetry are found in the spectra. Phonon modes at about 365, 557, 590, 674, and 1173 cm−1 appear in samples with almost all 8571


Article


Figure 4. 17O MAS NMR spectra of doped and pure CeO2 powders. (a) full range including the spinning sidebands, (b) extended view of the isotropic peaks. The spinning sidebands are marked with an asterisk. The arrow in part a marks the smaller peaks that are shown in detail in part b.

Figure 5. Typical impedance spectra of CGO (A) and CTO (B) in the Nyquist representation.

the 17O nuclei, as expected from the fluorite-type structure (Figure 1). The width of this peak is about 4 ppm for pure CeO2 and about 6 ppm for both Ta-doped samples. The chemical shift belongs to oxygen ions on regular crystal sites in the CeO2 structure and it is in good agreement with previously reported data.70 In addition to this major peak, there are additional peaks with smaller positive shift (Figure 4b). For nominally pure CeO2, there are two more peaks at 750 ppm (0.45% area fraction) and 699 ppm. These can be assigned to defects in the crystal structure, e.g., impurity atoms or intrinsic defects such as oxygen vacancies associated with Ce3+ ions. The samples doped with Ta also reveal the small peak at 750 ppm, but with significantly increased intensity in comparison to pure CeO2. Its relative intensity amounts to 1.4% for 0.5% CTO and to 1.8% for 2.5% CTO. This peak is therefore assigned to interstitial oxygen ions in the vicinity of Ta dopant ions. For the sample doped with Gd the isotropic peak at 876 ppm and also its spinning sidebands are much broader than for the other samples. This is due to the unpaired electron spin density of the paramagnetic Gd3+ ions (Table 1) which is transferred to the neighboring 17O probe nuclei. The width of the isotropic peak is about 40 ppm in this case. Because of the overlapping of

the isotropic peak and the spinning sidebands the smaller signal at 750 ppm could not be detected. 3.3. Total Conductivity and Transference Number. The temperature dependence of the total conductivity was measured with electrochemical impedance spectroscopy (EIS). Exemplary impedance spectra in the Nyquist representation (-Im (Z) vs Re (Z)) are shown in Figure 5. The spectra for CGO (A) and CTO (B) look quite different: The CGO sample exhibits only one depressed semicircle, which is shifted on the real axis. In terms of an equivalent circuit this spectrum arises from a resistor (R1) in series with a parallel RC-circuit (R2/ CPE). The resistor in the high frequency regime (R1) belongs to the sample bulk resistance whereas the parallel RC-circuit can be assigned to the symmetric platinum electrodes. As such electrodes often are not ideal electrical elements the capacitance can be better described by a constant phase element instead of a capacitor. The CTO sample shows only one single point on the real axis. This point originates from a single resistor which represents the sample bulk in this case. The absence of any capacitance in the spectrum is an indication for predominant electronic conductivity, which will be discussed later. 8572


Article


exhibits by far the lowest conductivity. This finding can be attributed to the low charge carrier concentration with neither significant acceptor nor donor doping. While the donor-doped CTO samples achieve the highest values the acceptor doped CGO and YSZ samples reside in the medium conductivity range. In comparison to YSZ, CGO has a higher conductivity and is in good agreement with reported values for polycrystalline samples.5,71,72 According to literature reports the majority charge carriers of acceptor doped electrolytes (CGO and YSZ) can be clearly assigned to oxygen vacancies.4,6,23,71,72 Considering the transference number measurement (Figure 7), it turns out that 0.5% CTO shows almost pure electronic

In order to interpret the conductivity results, we have to consider defect formation by aliovalent substitution. According to Kröger−Vink notation the introduction of Gd2O3 as acceptor leads to the following equation: x Gd 2O3 + 2CeCe + OOx → 2Gd′Ce + V •• O + CeO2

For reason for charge neutrality the incorporation of two Gd3+ ions into the crystal lattice creates one oxygen vacancy. These mobile oxygen vacancies are responsible for a high ionic conductivity in acceptor doped electrolytes. In contrast, the incorporation of Ta2O5 as donor has to be explained by two equations depending on the respective pO2 and temperature, as it was described earlier for the analogous system Nb2O5−CeO2.24,25 At low and medium oxygen partial pressures and high temperatures eq 4a holds and Ta5+ is compensated by electrons, localized on Ce sites (polarons). x Ta 2O5 + 4CeCe → 2Ta•Ce + 2Ce′Ce + 2CeO2 + 1/2 O2

(4a)

At high oxygen partial pressures and temperatures below 800 K, Ta5+ is rather compensated by formation of oxygen interstitials eq 4b. x Ta 2O5 + 2CeCe + V ix → 2Ta•Ce + O″i + 2CeO2

(4b)

Under variable oxygen partial pressures and temperatures a transition between electron and interstitial compensation takes place. With increasing oxygen partial pressure and decreasing temperature, oxygen is incorporated into the bulk at the expense of electrons, and interstitials become the dominating negative defect: x 2Ce′Ce + 1/2 O2 + V ix → 2CeCe + O″i

(4c)

At an oxygen partial pressure corresponding to air and a temperature below 800 K, the equilibrium is shifted to the right side of eq 4c; i.e., a considerable concentration of oxygen interstitials is formed, but a small concentration of electrons is still present. Since the electron mobility is much higher than the ionic one, the total conductivity is still dominated by electron transport in this regime. According to the well-accepted polaron model for ceria-based materials, electronic charge transport takes place via hopping of one electron from a Ce3+ to a neighboring Ce4+ site.19 Figure 6 presents the conductivities of all investigated samples as a function of reciprocal temperature, where CeO2

Figure 7. Conductivity and ionic transference number as a function of reciprocal temperature for CeO2 and 0.5% CTO.

conductivity (t(O2−) < 0.01) in the high temperature range. Below 600 K a maximum with predominant ionic charge transport was observed, which indicates oxygen interstitial transport in the light of the discussion above. Unfortunately, t(O2−) becomes more and more inaccurate at low temperatures on account of higher cell resistance. But the observed maximum was still considered as significant. In contrast, CeO2 shows a transition from being predominantly an ionic conductor at low temperature to a mixed conductor at high temperatures. As mentioned in the introduction, pure CeO2 is mainly an electronic conductor. As a consequence we conclude that the single crystal sample used in this work contains small impurities of M3+/M2+ elements acting as acceptor. Therefore, the ionic conductivity is increased and along with this also the transference number of oxygen ions. The activation energy EA for charge transport was determined by an Arrhenius plot according to the equation: EA (4) k·T All investigated samples have different activation energies and moreover several different regimes with respect to temperature (Figure 8). The activation energy depends on the type of the majority charge carriers, on possible formation ln(σT ) = ln(A) −

Figure 6. Conductivity as a function of reciprocal temperature for CeO2, CGO, CTO, and reference materials. 8573


Article


Figure 8. Arrhenius plot of the total conductivity and activation energy for the charge transport.

of defect associates, on contributions of defect formation energies, and on ion size effects. The 10% CGO sample shows only one activation energy of EA = 0.71 eV which is in good agreement with values reported in literature.72 The mixed conducting CeO2 sample presents a first region at low temperature (EA = 0.83 eV) which is dominated by ionic charge transport (refer to Figure 7). The second high temperature region (EA = 1.32 eV) can be interpreted as follows: With increasing temperature, defect formation will enter into the activation energy together with the migration energy resulting in the highest activation energy of all samples. As a consequence, more oxygen vacancies and excess electrons are formed. Since the mobility of electrons is higher than that of oxygen vacancies, the electronic contribution to the total conductivity increases with increasing temperature. The 0.5% CTO sample with mainly electronic conductivity reveals three different regions with different values for EA: (1) The low temperature regime is determined by mixed charge transport as proven by the EMF transference measurement. (2) In the medium temperature regime a transition between the majority defects eq 4a and eq 4b takes place. With increasing temperature, which resembles more reducing conditions, the equilibrium in eq 4c is shifted to the left side whereby more electrons are formed. (3) In the high temperature region the slope becomes smaller again and the lowest value of EA = 0.57 eV with respect to all samples is achieved. In this regime the electron concentration is fixed by the constant tantalum concentration (Ta•Ce = CeCe ′ ). The measured low activation energy is therefore only determined by polaron charge transport and decoupled from defect formation enthalpies. However, conductivity measurements only give access to the long-range charge transport rather than to short-range charge carrier dynamics on the microscopic length scale of few atoms. Here, the NMR measurements add valuable information. 3.4. Oxygen Dynamics in Pure and Donor-Doped CeO2. In order to investigate the oxygen ion dynamics in doped and pure CeO2, we performed temperature-dependent spin−lattice relaxation time (T1) experiments. These experiments were carried out in the temperature range from 220 to 673 K. For all samples and all temperatures, the magnetization transients could be well described with a single-exponential function. For pure CeO2, the relaxation rate (T1−1) shows a weak temperature dependence at temperatures below 400 K (Figure 9). For higher temperatures, T1−1 is strongly increasing

Figure 9. 17O NMR spin−lattice relaxation rate vs inverse temperature for pure and doped CeO2.

with increasing temperature, which indicates an Arrhenius-like behavior. This contribution is diffusion-induced42 with a maximum in T1−1 occurring at temperatures above the temperature range accessible in our experiments (T > 673 K). From the high-temperature slope, an activation energy of 0.4 eV can be roughly estimated. In line with the total conductivity and the oxygen ion transference number the diffusing species is assumed to be oxygen vacancies. The activation energy in the low temperature regime was found to be 0.83 eV which is considerably higher than the one estimated in NMR. At first glance this mismatch seems surprising but it is well-known that activation energies determined by NMR are underestimated by a factor of approximately 2−4 compared to conductivity measurements.42,60 In contrast, for the samples doped with Ta, the spin−lattice relaxation rate T1−1 shows a pronounced maximum at temperatures between 300 and 400 K. In general, long-range motion of charge carriers in electrolytes is “frozen-in” at such low temperatures (see Figure 4). Therefore, the occurrence of a maximum is a clear evidence for a local hopping process of oxygen ions which causes this additional contribution to the spin−lattice relaxation rate. Interestingly, in pure CeO2 this local jump process does not occur. In donor-doped ceria oxygen vacancies are minority defects.27,73 At high oxygen partial pressure and low temperatures (θ < 800 K), as applied in our experiment, oxygen interstitials are the predominant oxygen defects. This interstitial oxygen is probably trapped at tantalum sites under formation of charged defect associates: • Ta•Ce + O″i → (TaCe O″i )′

The diffusion process observed at ∼350 K in NMR is therefore most likely a local jump of trapped oxygen interstitials around Ta5+ ions. This explanation is supported by a “sizeeffect” as known for other ceria systems. Since Ta5+ has a significantly smaller ionic radius than Ce4+ (see Table 2) there is more space around tantalum for localized diffusion of interstitial oxygen. To the best of our knowledge, interstitial oxygen diffusion has never been proven experimentally in doped ceria. Only few theoretical studies deal with the migration of interstitial oxygen 8574


Article

The Journal of Physical Chemistry C in ceria and reported activation energies of 0.56−0.8 eV.33,34 The activations energies obtained by NMR are 0.25 ± 0.05 eV for both samples. These values give a good correlation to the calculated migration data taking the above-mentioned factor (2−4 times lower) into account. For temperatures above 550 K, the relaxation rate T1−1 starts to increase again for the samples doped with Ta (Figure 9), in a way similar to that of pure CeO2. This increase at higher temperatures might be attributed to the long-range transport of oxygen ions associated with higher energy barriers. We believe that the trapped interstitial oxygen ions are released and now contribute to the long-range charge transport. This assumption correlates well with the course of the oxygen ion transference number which has a maximum in a short-range of temperature. Here, the electronic contribution is still small while the thermal energy is sufficient for the release of interstitial oxygen from the tantalum site. At higher temperatures, electronic charge transport dominates again. It should be noted that the static 17O NMR spectra of these samples show no temperature dependence, i.e. a so-called motional narrowing does not occur. This is due to the fact that the static spectra are dominated by the contribution at 877 ppm associated with the majority of oxygen ions in a regular environment with no dopants/vacancies in the direct environment. This confirms that the local hopping observed by the local maxima in the relaxation time measurements is associated with oxygen interstitials in the direct environment of the dopant atoms which is also corroborated by the small activation barriers. The increase in the T1−1 data above 500 K for all samples indicates the onset oxygen ion transport on longer length scales. The observed course of the data in Figure 9 can be mathematically described as follows: Local hopping over activation barriers EA results in motional correlation times given by ⎛E ⎞ τ = τ0 × exp⎜ A ⎟ ⎝ kT ⎠

Ta-doped CeO2 samples. eq 6 is valid for dipolar as well as for quadrupolar relaxation mechanisms which both result in the same temperature dependence for the spin−lattice relaxation rate and differ only in their prefactors.41 eq 7 together with the given Larmor frequency ωL can be used to determine the motional correlation time at the temperature where the maximum occurs. Fitting eqs 5 and 6 to the experimental data (dashed lines in Figure 9) yields a hopping rate of τ−1 = 2.5 × 108 s−1 at temperatures of 365 and 340 K for 2.5% CTO and 0.5% CTO, respectively. In GCO, where long-range O2− ion transport is the dominating charge transport process in this temperature regime, local hopping processes cannot be observed due to the strong paramagnetic contributions to the 17O NMR relaxation rate.

4. CONCLUSIONS The electrical conductivity of doped and undoped CeO2 single crystals was measured by impedance spectroscopy. It turns out, that the electrical conductivity of donor-doped CeO2 is higher than that of pure CeO2 and shows regimes with different activation energies. 17O MAS NMR experiments reveal differences in defect concentrations for different doped CeO2 samples. A fast local oxygen hopping process associated with small activation barriers is found for the Ta-doped samples, in addition to a slower one that is present in all samples. It is proposed that below a certain temperature local hopping of interstitial oxygen around a tantalum site is observed by NMR. With increasing temperature these “trapped” interstitial oxygen ions are released and contribute to the charge transport. To our knowledge, interstitial oxygen transport in ceria has never been experimentally observed. Future studies may be extended to other oxides which contain interstitial oxygen.

■

*(S.I.) E-mail: [email protected]. Telephone: +49-721-60828508. *(J.J.) E-mail: [email protected]. Telephone.: +49 + 641-99-34500.

(5)

where τ0 is a pre-exponential factor, k is the Boltzmann constant, and T is the temperature. In the following, we identify the motional correlation time with the mean residence time of the oxygen ions. The mean residence time differs from the correlation time only by a factor of the order of unity.42,74 Increasing the temperature T results in a monotonic decrease of the mean residence time τ. The hopping of the oxygen ions results in fluctuations in the local fields around the oxygen probe nuclei. These fluctuating fields can be described by a spectral density J(ω). The components of J(ω) are high at frequencies around the average hopping rate τ−1. The field fluctuations induce transitions between the Zeeman levels and thus a very effective spin−lattice relaxation T1−1 ∝

τ 4τ + 1 + ωL 2τ 2 1 + 4ωL 2τ 2

Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We are grateful to the German Federal Ministry of Education and Research (BMBF) for financial support. We thank Prof. Peter J. Klar and co-workers (Department of Physics, University Giessen) for fruitful discussions on Raman spectroscopy. This study was supported by the DFG within Projects Ja648/21-1 and Le781/14-1.

■

REFERENCES

(1) Trovarelli, A. Catalytic Properties of Ceria and CeO2-Containing Materials. Catal. Rev.: Sci. Eng. 1996, 38, 439−520. (2) Mogensen, M.; Sammes, N. M.; Tompsett, G. A. Physical, Chemical and Electrochemical Properties of Pure and Doped Ceria. Solid State Ionics 2000, 129, 63−94. (3) Etsell, T. H.; Flengas, S. N. Electrical Properties of Solid Oxide Electrolytes. Chem. Rev. 1970, 70, 339−376. (4) Steele, B. C. H. Appraisal of Ce1‑yGdyO2‑y/2 Electrolytes for ITSOFC Operation at 500 °C. Solid State Ionics 2000, 129, 95−110. (5) Chaubey, N.; Wani, B. N.; Bharadwaj, S. R.; Chattopadhyaya, M. C. Physicochemical Properties of Rare Earth Doped Ceria

(6)

if the contribution of the spectral density at the Larmor frequency J(ωL) is large or, in other words, if the hopping rate τ−1 is of the order of the Larmor frequency ωL, i.e.

ωL × τ ≈ 1

AUTHOR INFORMATION

Corresponding Authors

(7)

Therefore, such a local jump process will result in a spin− lattice relaxation rate that reveals a maximum as a function of (inverse) temperature (cf. eq 6) as observed in Figure 9 for the 8575


Article


(28) Jorgensen, J. D.; Dabrowski, B.; Pei, S.; Richards, D. R.; Hinks, D. G. Structure of the Interstitial Oxygen Defect in La2NiO4+δ. Phys. Rev. B: Condens. Matter Mater. Phys. 1989, 40, 2187−2199. (29) Hausmann, A.; Schallenberger, B. Interstitial Oxygen in ZincOxide Single-Crystals. Z. Phys. B: Condens. Matter Quanta 1978, 31, 269−273. (30) Li, M.-R.; Hong, S.-T. Interstitial Oxygen in Perovskite-Related Sr6−2xNb2+2XO11+3X. Chem. Mater. 2008, 20, 2736−2741. (31) Boyapati, S.; Wachsman, E. D.; Jiang, N. Effect of Oxygen Sublattice Ordering on Interstitial Transport Mechanism and Conductivity Activation Energies in Phase-Stabilized Cubic Bismuth Oxides. Solid State Ionics 2001, 140, 149−160. (32) Inerbaev, T. M.; Sahara, R.; Mizuseki, H.; Kawazoe, Y.; Nakamura, T. Interstitial Oxygen and Dopant Atoms Arrangement in Tin-Doped Indium Oxide. Mater. Trans. 2007, 48, 666−669. (33) Middleburgh, S. C.; Lumpkin, G. R.; Grimes, R. W. Accommodation of Excess Oxygen in Fluorite Dioxides. Solid State Ionics 2013, 253, 119−122. (34) Bai, X.-M.; El-Azab, A.; Yu, J.; Allen, T. R. Migration Mechanisms of Oxygen Interstitial Clusters in UO2. J. Phys.: Condens. Matter 2013, 25, 015003. (35) Janek, J.; Timm, H. Thermal Diffusion and Soret Effect in (U,Me)O2+δ: The Heat of Transport of Oxygen. J. Nucl. Mater. 1998, 255, 116−127. (36) Xiao, H. Y.; Zhang, Y.; Weber, W. J. Stability and Migration of Charged Oxygen Interstitials in ThO2 and CeO2. Acta Mater. 2013, 61, 7639−7645. (37) Leon-Reina, L.; Losilla, E. R.; Martinez-Lara, M.; Bruque, S.; Aranda, M. A. G. Interstitial Oxygen Conduction in Lanthanum OxyApatite Electrolytes. J. Mater. Chem. 2004, 14, 1142−1149. (38) Jeon, S. Y.; Choi, M. B.; Im, H. N.; Hwang, J. H.; Song, S. J. Oxygen Ionic Conductivity of La2NiO4+δ via Interstitial Oxygen Defect. J. Phys. Chem. Solids 2012, 73, 656−660. (39) Abragam, A. The Principles of Nuclear Magnetism; Oxford University Press: Oxford, U.K., 1999. (40) Slichter, C. P. Principles of Magnetic Resonance; Springer: Berlin, 1996. (41) Spiess, H. W. In Dynamic NMR Spectroscopy; Diehl, P., Fluck, E., Kosfeld, R., Eds.; Springer: Berlin, 1978; pp55−214. (42) Heitjans, P.; Indris, S. Diffusion and Ionic Conduction in Nanocrystalline Ceramics. J. Phys.: Condens. Matter 2003, 15, R1257− R1289. (43) Böhmer, R.; Jeffrey, K. R.; Vogel, M. Solid-State Li NMR with Applications to the Translational Dynamics in Ion Conductors. Prog. Nucl. Magn. Reson. Spectrosc. 2007, 50, 87−174. (44) Hain, H.; Heinzmann, R.; Scheuermann, M.; Wünsche, L.; Hahn, H.; Indris, S. Study of Local Structure and Li Dynamics in Li4+xTi5O12 (0 ≤ x ≤ 5) using 6Li and 7Li NMR Spectroscopy. Solid State Nucl. Magn. Reson. 2012, 42, 9−16. (45) Indris, S.; Heitjans, P.; Grey, C. P. Multinuclear NMR Spectroscopic Studies of Structure and Dynamics in Hydrous NaAlSi3O8 and Ca0.5AlSi3O8 Glasses. J. Non-Cryst. Solids 2012, 358, 2862−2867. (46) Indris, S.; Scheuermann, M.; Becker, S. M.; Šepelák, V.; Kruk, R.; Suffner, J.; Gyger, F.; Feldmann, C.; Ulrich, A. S.; Hahn, H. Local Structural Disorder and Relaxation in SnO2 Nanostructures Studied by 119 Sn MAS NMR and 119Sn Mossbauer Spectroscopy. J. Phys. Chem. C 2011, 115, 6433−6437. (47) Grey, C. P.; Dupré, N. NMR Studies of Cathode Materials for Lithium-Ion Rechargeable Batteries. Chem. Rev. 2004, 104, 4493− 4512. (48) Indris, S.; Cabana, J.; Rutt, O. J.; Clarke, S. J.; Grey, C. P. Layered Oxysulfides Sr2MnO2Cu2m‑0.5Sm+1 (m = 1, 2, and 3) as Insertion Hosts for Li Ion Batteries. J. Am. Chem. Soc. 2006, 128, 13354−13355. (49) Indris, S.; Heitjans, P.; Hattendorf, J.; Zeitz, W.-D.; Bredow, T. Local Electronic Structure in MgB2 from 12B β-NMR. Phys. Rev. B: Condens. Matter Mater. Phys. 2007, 75, 024502.

Ce0.9Ln0.1O1.95 (Ln = Nd, Sm, Gd) as an Electrolyte Material for ITSOFC/SOEC. Solid State Sci. 2013, 20, 135−141. (6) Park, J. H.; Blumenthal, R. N. Electronic Transport in 8 Mole Percent Y2O3-ZrO2. J. Electrochem. Soc. 1989, 136, 2867−2876. (7) Inaba, H.; Tagawa, H. Ceria-Based Solid Electrolytes − Review. Solid State Ionics 1996, 83, 1−16. (8) Duer, M. J. Introduction to Solid-State NMR Spectroscopy; Blackwell Publishing: Oxford, U.K., 2004. (9) Shannon, R. D.; Prewitt, C. T. Effective Ionic Radii in Oxides and Fluorides. Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem. 1969, B25, 925−946. (10) Di Monte, R.; Kašpar, J. Heterogeneous Environmental Catalysis − A Gentle Art: CeO2-ZrO2 Mixed Oxides as a Case History. Catal. Today 2005, 100, 27−35. (11) Trovarelli, A. Structural and Oxygen Storage/Release Properties of CeO2-Based Solid Solutions. Comments Inorg. Chem. 1999, 20, 263−284. (12) Kašpar, J.; Fornasiero, P.; Graziani, M. Use of CeO2-Based Oxides in the Three-Way Catalysis. Catal. Today 1999, 50, 285−298. (13) Tuller, H. L.; Nowick, A. S. Defect Structure and ElectricalProperties of Nonstochiometric CeO2 Single-Crystals. J. Electrochem. Soc. 1979, 126, 209−217. (14) Panlener, R. J.; Blumenthal, R. N.; Garnier, J. E. Thermodynamic Study of Nonstochiometric Cerium Dioxide. J. Phys. Chem. Solids 1975, 36, 1213−1222. (15) Xiong, Y.-P.; Kishimoto, H.; Yamaji, K.; Yoshinaga, M.; Horita, T.; Brito, M. E.; Yokokawa, H. Electronic Conductivity of CeO2: Its Dependence on Oxygen Partial Pressure and Temperature. Electrochem. Solid-State Lett. 2010, 13, B21−B24. (16) Xiong, Y.-P.; Kishimoto, H.; Yamaji, K.; Yoshinaga, M.; Horita, T.; Brito, M. E.; Yokokawa, H. Electronic Conductivity of Pure Ceria. Solid State Ionics 2011, 192, 476−479. (17) Keating, P. R. L.; Scanlon, D. O.; Morgan, B. J.; Galea, N. M.; Watson, G. W. Analysis of Intrinsic Defects in CeO2 Using a Koopmans-Like GGA Plus U Approach. J. Phys. Chem. C 2012, 116, 2443−2452. (18) Zacherle, T.; Schriever, A.; De Souza, R.; Martin, M. Ab Initio Analysis of the Defect Structure of Ceria. Phys. Rev. B: Condens. Matter Mater. Phys. 2013, 87, 134104. (19) Tuller, H. L.; Nowick, A. S. Small Polaron Electron-Transport in Reduced CeO2 Single-Crystals. J. Phys. Chem. Solids 1977, 38, 859− 867. (20) Chiang, Y.-M.; Lavik, E. B.; Kosacki, I.; Tuller, H. L.; Ying, J. M. Defect and Transport Properties of Nanocrystalline CeO2‑x. Appl. Phys. Lett. 1996, 69, 185−187. (21) Chiang, Y.-M.; Lavik, E. B.; Blom, D. A. Defect Thermodynamics and Electrical Properties of Nanocrystalline Oxides: Pure and Doped CeO2. Nanostruct. Mater. 1997, 9, 633−642. (22) Tuller, H. L. Ionic Conduction in Nanocrystalline Materials. Solid State Ionics 2000, 131, 143−157. (23) Rührup, V.; Wiemhöfer, H.-D. Ionic Conductivity of Gd- and YDoped Ceria-Zirconia Solid Solutions. Z. Naturforsch., B: J. Chem. Sci. 2006, 61, 916−922. (24) Naik, I. K.; Tien, T. Y. Electrical Conduction in Nb2O5-Doped Cerium Dioxide. J. Electrochem. Soc. 1979, 126, 562−566. (25) Yashiro, K.; Suzuki, T.; Kaimai, A.; Matsumoto, H.; Nigara, Y.; Kawada, T.; Mizusaki, J.; Sfeir, J.; Van Herle, J. Electrical Properties and Defect Structure of Niobia-Doped Ceria. Solid State Ionics 2004, 175, 341−344. (26) Ananthapadmanabhan, P. V.; Menon, S. B.; Patil, D. S.; Venkatramani, N.; Rohatgi, V. K. Electrical Conductivity of Cerium Dioxide Doped with Tantalum Pentoxide. J. Mater. Sci. Lett. 1992, 11, 501−503. (27) Eufinger, J.-P.; Daniels, M.; Schmale, K.; Berendts, S.; Ulbrich, G.; Lerch, M.; Wiemhöfer, H.-D.; Janek, J. The Model Case of an Oxygen Storage Catalyst − Non-Stoichiometry, Point Defects and Electrical Conductivity of Single Crystalline CeO2-ZrO2-Y2O3 Solid Solutions. Phys. Chem. Chem. Phys. 2014, 16, 25583−25600. 8576


Article

The Journal of Physical Chemistry C (50) Indris, S.; Heitjans, P.; Uecker, R.; Bredow, T. Local Electronic Structure in a LiAlO2 Single Crystal Studied with 7Li NMR Spectroscopy and Comparison with Quantum Chemical Calculations. Phys. Rev. B: Condens. Matter Mater. Phys. 2006, 74, 245120. (51) Bräuniger, T.; Jansen, M. Solid-State NMR Spectroscopy of Quadrupolar Nuclei in Inorganic Chemistry. Z. Anorg. Allg. Chem. 2013, 639, 857−879. (52) Heitjans, P.; Indris, S. Fast Diffusion in Nanocrystalline Ceramics Prepared by Ball Milling. J. Mater. Sci. 2004, 39, 5091−5096. (53) Indris, S.; Heitjans, P.; Uecker, R.; Roling, B. Li Ion Dynamics in a LiAlO2 Single Crystal Studied by 7Li NMR Spectroscopy and Conductivity Measurements. J. Phys. Chem. C 2012, 116, 14243− 14247. (54) Fischer, D. M.; Duwe, P.; Indris, S.; Heitjans, P. Tracer Diffusion Measurements in Solid Lithium: A Test Case for the Comparison Between NMR in Static and Pulsed Magnetic Field Gradients after Upgrading a Standard Solid State NMR Spectrometer. Solid State Nucl. Magn. Reson. 2004, 26, 74−83. (55) Blanc, F.; Sperrin, L.; Jefferson, D. A.; Pawsey, S.; Rosay, M.; Grey, C. P. Dynamic Nuclear Polarization Enhanced Natural Abundance 17O Spectroscopy. J. Am. Chem. Soc. 2013, 135, 2975− 2978. (56) Fuda, K.; Kishio, K.; Yamauchi, S.; Fueki, K.; Onoda, Y. 17O NMR Study of Y2O3-Doped CeO2. J. Phys. Chem. Solids 1984, 45, 1253−1257. (57) Fuda, K.; Kishio, K.; Yamauchi, S.; Fueki, K. 17O NMR Relaxation Study of δ-Bi2O3. J. Phys. Chem. Solids 1985, 46, 1141− 1146. (58) Kim, N.; Stebbins, J. F. Vacancy and Cation Distribution in Yttria-Doped Ceria: An 89Y and 17O MAS NMR Study. Chem. Mater. 2007, 19, 5742−5747. (59) Ando, M.; Oikawa, I.; Noda, Y.; Ohki, S.; Tansho, M.; Shimizu, T.; Kiyono, H.; Maekawa, H. High Field 17O NMR Study of Defects in Doped Zirconia and Ceria. Solid State Ionics 2011, 192, 576−579. (60) Heinmaa, I.; Joon, T.; Kooskora, H.; Pahapill, J.; Subbi, J. Local Structure and Oxygen Ion Dynamics in La Doped Ceria: 17O NMR Study. Solid State Ionics 2010, 181, 1309−1315. (61) Subbi, J.; Heinmaa, I.; Poder, R.; Kooskora, H. Solid State NMR Spin-Lattice Relaxation Investigation of Oxygen Dynamics in Scandium Doped Ceria from 60 to 1073 K. Solid State Ionics 2013, 239, 15−20. (62) Assmus, W.; Whippey, N. Skull Melting. Chem. Ing. Tech. 1983, 55, 716−717. (63) Ionic Conductivity in Oxides. Heyne, L. In Mass Transport in Oxides, Proceedings of a Symposium Held at Gaithersburg, MD, October 22−25, 1967; Wachtman, J. B., Jr., Franklin, A. D., Eds.; 1967; p 149. (64) Lee, D.-K.; Fischer, C. C.; Valov, I.; Reinacher, J.; Stork, A.; Lerch, M.; Janek, J. An EMF Cell with a Nitrogen Solid Electrolyte − On the Transference of Nitrogen Ions in Yttria-Stabilized Zirconia. Phys. Chem. Chem. Phys. 2011, 13, 1239−1242. (65) Wagner, C. Beitrag zur Theorie des Anlaufvorgangs (Theory of the Tarnishing Process). Z. Phys. Chem. 1933, B21, 25−41. (66) Banerji, A.; Grover, V.; Sathe, V.; Deb, S. K.; Tyagi, A. K. CeO2Gd2O3 System: Unraveling of Microscopic Features by Raman Spectroscopy. Solid State Commun. 2009, 149, 1689−1692. (67) Weber, W. H.; Hass, K. C.; McBride, J. R. Raman-Study of CeO2 − 2nd-Order Scattering, Lattice-Dynamics, and Particle-Size Effects. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 48, 178−185. (68) Keramidas, V. G.; White, W. B. Raman-Spectra of Oxides with Fluorite Structure. J. Chem. Phys. 1973, 59, 1561−1562. (69) Nakajima, A.; Yoshihara, A.; Ishigame, M. Defect-Induced Raman-Spectra in Doped CeO2. Phys. Rev. B: Condens. Matter Mater. Phys. 1994, 50, 13297−13307. (70) Bastow, T. J.; Stuart, S. N. 17O NMR in Simple Oxides. Chem. Phys. 1990, 143, 459−467. (71) Balazs, G. B.; Glass, R. S. AC-Impedance Studies of Rare-EarthOxide Doped Ceria. Solid State Ionics 1995, 76, 155−162.

(72) Wang, S.; Kobayashi, T.; Dokiya, M.; Hashimoto, T. Electrical and Ionic Conductivity of Gd-Doped Ceria. J. Electrochem. Soc. 2000, 147, 3606−3609. (73) Göbel, M. C.; Gregori, G.; Maier, J. Electronically Blocking Grain Boundaries in Donor Doped Cerium Dioxide. Solid State Ionics 2012, 215, 45−51. (74) Heitjans, P.; Schirmer, A.; Indris, S. In Diffusion in Condensed Matter − Methods, Materials, Models; Heitjans, P., Kärger, J., Eds.; Springer: Berlin, 2005.

8577


Observing Local Oxygen Interstitial Diffusion in Donor-Doped Ceria by

Recommend Documents