Thermal and electrical properties of rare-earth co-doped ceria ceramics

Materials Chemistry and Physics 112 (2008) 711–718

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Thermal and electrical properties of rare-earth co-doped ceria ceramics V. Prashanth Kumar, Y.S. Reddy, P. Kistaiah, G. Prasad, C. Vishnuvardhan Reddy ∗ Department of Physics, Osmania University, Hyderabad 500007, India

a r t i c l e

i n f o

Article history: Received 7 August 2007 Received in revised form 1 February 2008 Accepted 15 June 2008 Keywords: Rare-earth-doped ceria Sol–gel method Powder X-ray diffraction Thermal expansion Impedance spectroscopy Electrolyte

a b s t r a c t Sol–gel method was used to prepare rare-earth co-doped ceria ceramics with composition of Ce0.8 Gd0.2−x Yx O2 (0 ≤ x ≤ 0.1). The phase identification, morphology, thermal expansion and electrical properties of samples were studied by XRD, SEM, Dilatometry and Impedance spectroscopy. All com¯ space group). The lattice parameters were determined by X-ray pounds have the cubic structure (Fm3m powder diffraction. The thermal expansion is linear for all the samples. The thermal expansion coefficient (TEC), determined in the temperature range from room temperature to 1000 ◦ C, increases with x. The results showed that Ce0.8 Gd0.15 Y0.05 O2 and Ce0.8 Gd0.1 Y0.1 O2 posses higher conductivity and lower activation energy than pure gadolinium-doped ceria (GDC). Impedance measurements indicate the presence of unassociated defect which relax with an activation energy in the range 1–1.2 eV. The Cole–Cole plots show semicircles in the temperature range 250–350 ◦ C. © 2008 Elsevier B.V. All rights reserved.

1. Introduction Doped ceria oxides have been extensively studied as one of the most promising electrolyte materials for intermediate temperature solid oxide fuel cells (ITSOFCs) [1–4]. These materials demonstrate much higher ionic conductivity at relatively lower temperatures in comparison to that of the traditional electrolyte yttria-stabilized zirconia (YSZ). The substitution of Ce4+ by suitable trivalent cations such as Gd3+ , Sm3+ , Y3+ or La3+ enhances the chemical stability, increases the ionic conductivity and suppresses the reducibility of ceria-based materials. The most effective substitutes are Gd2 O3 and Sm2 O3 possibly due to the fact that they minimize the changes in lattice parameter. So far, many studies have been carried out on doped ceria [5–14]. Some singly rare-earth-doped electrolytes, such as Ce1−x Smx O2−y , Ce1−x Gdx O2−y , Ce1−x Yx O2−y , etc. and few codoped ceria-based electrolytes have been investigated [3,4,15–19]. Sm and Gd doping is known to improve the electrical conductivity and compatibility with electrodes of SOFCs [20,21]. It is desirable to optimize the preparation protocols of the material at lower temperatures. Zhang et al. reported that a maximum conductivity is observed for Ce0.8 Gd0.2 O2 in Ce1−x Gdx O2 compositions [22]. It is interesting to see how the activation energy and absolute conductivity of ceria (r = 0.97 Å) change when it is co-doped with yttrium (r = 1.019 Å) and gadolinium (r = 1.053 Å) rare-earth ions of different radii. Also to the best of our knowledge, no studies are reported on

∗ Corresponding author. Tel.: +91 40 27682242; fax: +91 40 27090020. E-mail address: [email protected] (C. Vishnuvardhan Reddy). 0254-0584/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.matchemphys.2008.06.030

gadolinia and yttria co-doped ceria. Hence, in this paper we report the results of our investigations on electrical properties and thermal expansion of Y, Gd co-doped ceria synthesized through sol–gel method. 2. Experimental The sample with the general formula of Ce0.8 Gd0.2−x Yx O2 (0 < x < 0.1) was synthesized by the sol–gel method. Cerium nitrate, yttrium nitrate, and gadolinium oxide were used as starting materials. Cerium nitrate and yttrium nitrate were dissolved in water and the desired amount of gadolinium oxide was dissolved in nitrate solution. The individual solutions were mixed together and an aqueous solution of citric acid corresponding to every mole of metal atom was added. The amount of citric acid used was necessary to bind all the metals present in the solution. The pH value of the mixed solution was adjusted to ≈7 with ammonia solution under continuous stirring at 80 ◦ C and homogeneous sol was formed. With the evaporation of water, sponge-like gel was obtained, which was calcined at 600 ◦ C to get the final composition powder. Then the dried powders were ground in an agate mortar and then pallets were made using hydraulic press at 5 MPa. The pellets were sintered finally at 1300 ◦ C in air for 4 h. The circular pellets were used for electrical measurements. The crystal structure of the samples was determined by X-ray diffraction (XRD) analysis using a Pananalytical X’Pert Pro X-ray diffractometer using Cu K␣ radiation (1.5418 Å; operated at 40 kV, 30 mA) at room temperature in the Bragg angle region of 20◦ < 2 < 90◦ with a scan rate of 2◦ min−1 . The crystallite size, D, of the samples was estimated using the Scherrer formula [23]: D=

0.9 ˇ cos 

Scanning electron microscopy (SEM) micrographs were taken on the intersection of the specimens using a Hitachi-S 3000N Scanning electron microscope. Thermal expansion measurements were performed using Netzsch 402PC dilatometer in air from room temperature to 1000 ◦ C. The rectangular samples used for these measurements were of dimension 25 mm × 6 mm × 6 mm with polished frontal faces. The rectangular pellets were pressed in an hydraulic press and sintered

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Fig. 2. Linear thermal expansion curves for Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1). Table 2 The average thermal expansion coefficients of Ce0.8 Gd(0.2−x) Yx O2 Fig. 1. X-ray diffraction patterns of Ce0.8 Gd0.2−x Yx O2 powders sintered at 900 ◦ C.

at 1300 ◦ C. Alumina standard was used as a reference. Heating and cooling rates are 3 ◦ C min−1 , with annealing time half an hour at maximum temperature. The dc electrical conductivity of the samples was measured in air by two-probe method using a conventional set-up on sintered circular pellets. Measurements were performed from RT to 600 ◦ C with a heating rate of 5 ◦ C min−1 . The conductivity measurements were taken using Keithley 236 source/measure unit. The data acquisition is automated using a PC and Test Point software. Silver paste was used as electrodes on large area faces of the pellet for all electrical measurements. The impedance spectra of the samples were recorded in air using AUTOLAB PGSTAT 30 Frequency Response Analyzer (FRA) in the frequency range 100 Hz to 1 MHz and in the temperature range RT (≈30 ◦ C)–475 ◦ C.

3. Results and discussions From Fig. 1, it can be seen that the prepared co-doped ceria samples are formed in single phase with cubic fluorite structure with ¯ (ICDD 075-0162). This is same as ceria struca space group Fm3m ture. The average crystallite sizes of the sample powders calculated by the Scherrer formula are between 20 and 25 nm. This indicates that the fine powders synthesized by the sol–gel method can be successfully sintered into electrolytes with single phase [24,25]. Unit cell parameters are calculated Using XLAT program by least square refinement technique from the X-ray diffractograms. The dependence of unit cell parameter on dopant concentration is shown in Table 1. As can be seen, the unit cell parameter decreases with increasing yttrium content. The radius of Y3+ (120 pm) is smaller than the radius of Gd3+ (180 pm) [26], hence it is expected that the lattice parameters of the co-doped samples will have smaller values as observed in the present study. The densities of pellets sintered at 1300 ◦ C are measured by the Archimedes method. The measured sintered density (d), the theoretical density (dth ), and relative density d/dth (%) are summarized in Table 1. Measured density of all the samples is more than 95% of

Composition (x)

TEC (10−6 /◦ C) RT–800 ◦ C

TEC (10−6 /◦ C) RT–1000 ◦ C

0.0 0.05 0.1

12.21 12.24 12.63

12.69 12.99 13.11

the theoretical value. It is observed that the density of the samples increases with increase of doping concentration. The SEM studies reveal the dense texture with small porosity in the samples which also confirms the high relative density of the samples. Relatively dense samples, with density greater than 90% of the theoretical value, are required for the measurement of thermal expansion coefficient (TEC) [27]. The thermal expansion characteristic (l/l) of Ce0.8 Gd0.2−x Yx O2 (x = 0, 0.05 and 0.1) obtained during the heating from RT to 1000 ◦ C in air is shown in Fig. 2. The thermal expansion depends on the electrostatic forces within the lattice, which depend on the concentration of positive and negative charges and their distances within the lattice. The thermal expansion increases if the attractive forces decrease. The thermal expansion of a lattice with a certain structure and a fixed oxygen to metal stoichiometry is characterized by a steady thermal expansion coefficient, ˛, caused by the thermal lattice vibrations. The thermal expansion coefficients (TEC) between 50–800 ◦ C and 50–1000 ◦ C are calculated from the curves and listed in Table 2. The TEC value increases with increasing Y content. The thermal expansion curves of Ce0.8 Gd0.2−x Yx O2 (x = 0, 0.05 and 0.1) are linear and have two slopes. The temperature, at which the change of slope occurs, decreases with increasing temperature and this is called onset temperature. The value of onset temperature is located at approximately 380, 350 and 290 ◦ C for x = 0, 0.05 and 0.1, respectively. The slope of the curves is slightly increased with increasing the yttrium content at high temperatures. These increases were found to be reversible in subsequent heating and cooling cycles at a rate of 3 ◦ C min−1 in air. The lattice expansion observed at high temperatures in these oxides may be attributed

Table 1 Crystallographic data of Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1) sintered at 1300 ◦ C Composition (x)

Crystal structure

a (Å)

Volume (Å3 )

Bulk density (d) (gm cm−3 )

XRPD density (dth ) (gm cm−3 )

d/dth (%)

0.0 0.05 0.1

Cubic Cubic Cubic

5.465 5.456 5.452

163.20 162.43 162.05

6.79 6.74 6.67

7.14 7.04 6.91

95.1 95.7 96.5

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Fig. 3. Variation of electrical conductivity with temperature for Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1).

to the loss of lattice oxygen and the formation of oxygen vacancies. The increasing of TEC is accompanied by the decreasing of the unit cell parameter. This is in good agreement with the previous reports [28–30]. The variations of dc conductivity ( dc ) with temperature, obtained from the bulk resistance and sample dimensions are presented in Fig. 3. The conductivity follows the Arrhenius behavior of the form:

 E  a

dc = 0 exp −

kT

where  0 is the pre-exponential factor, T is the absolute temperature, Ea is the activation energy for conduction and k is the Boltzmann’s constant. The conductivity plots of the sample show slight deviation from linearity in the lower-temperature region (below 200 ◦ C). The high temperature region (200–600 ◦ C) however has linear dependence. The activation energies of conduction for all compositions are listed in Table 3. The activation energy Ea decreases with increasing x and in the range 1.25–1.13 eV. It is observed that the conductivity of the present class of samples increases with increase in yttrium doping and changes by about five orders when the temperature is increased from 200 to 600 ◦ C. From the observed trends in the conductivity behavior, the conductivity is high at 600 ◦ C and may be very high at higher temperatures (beyond 600 ◦ C), which is ideal for high temperature solid electrolytes. The value of the conductivity at 600 ◦ C is sufficient to use these samples as electrolyte in ITSOFC. Impedance is the opposition of a material to the flow of alternating current. It is just a specific branch of the tree of electrical measurements. Electrical properties of materials can be characterized using ac impedance data. There are various formalisms by Fig. 4. Complex impedance plots of Ce0.8 Gd0.2−x Yx O2 (a) x = 0, (b) x = 0.05 and (c) x = 0.1.

Table 3 Activation energies of Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1) x

dc conduction (eV)

ac conduction (eV)

Relaxation (eV)

0 0.05 0.1

1.25 1.15 1.13

1.09 1.05 1.03

1.20 1.05 1.03

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Fig. 6. Relaxation time versus inverse of temperature for Ce0.8 Gd0.2−x Yx O2 .

which one can exploit the wealth of the data. The effects, which do not show in one formalism, show up in the other. All the four formalisms are inter-related. The different formalisms are electric modulus (M), admittance (Y), permittivity (ε) and impedance (Z) [31,32]: Z ∗ = Z  − jZ  Y ∗ = Y  + jY  = jωC0 ε∗ M ∗ = M  + jM  = jωC0 Z ∗ ε∗ = ε − jε = (M ∗ )−1 Complex impedance and modulus techniques are used to separate out the contribution of grain and grain boundary effects of the ceramics [33]. By additional use of complex electric modulus for-

Fig. 5. Complex impedance plots of Ce0.8 Gd0.2−x Yx O2 (a) x = 0, (b) x = 0.05 and (c) x = 0.1.

Fig. 7. Conductivity obtained from lower frequency intercept of the complex impedance plots versus inverse of temperature.

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malism, the inhomogeneous nature of polycrystalline ceramics can be probed and grain boundary and bulk effect may be distinguished from the impedance data. In the present study Z* and M* are used. The values of Z and Z plotted on a linear scale as complex plane diagram takes the form of semicircles known as Cole–Cole plots [34]. The impedance data are obtained for all the samples in the temperature range RT–475 ◦ C and in the frequency range 100–106 Hz. In the present paper, the impedance data of Ce0.8 Gd0.2−x Yx O2 is analyzed using two different formalisms: (a) Frequency explicit plots of imaginary part of impedance (Fig. 4) and (b) Cole–Cole plots of real (Z ) and imaginary (Z ) parts of impedance (Fig. 5). All the frequency dependent graphs of Z show a peak in the temperature range 200–400 ◦ C. Below 200 ◦ C these plots show normal decrease of Z with the increasing frequency. In those plots where peaks are observed, the peak frequency is found to be a function of temperature. The inverse of peak frequency as a function of temperature is shown as log  versus 1000/T in Fig. 6. The straight-line behavior of this graph indicates the thermally active dipoles which are responsible for peak in Z versus frequency plots. The relaxation time of dipoles is found to decrease with increase in temperature. The calculated activation energies for relaxation are in the range 1–1.2 eV. The small activation energy for relaxation indicates the presence of relatively simple unassociated mobile charged defects [35]. Normally a material contains different types of dipoles, each of which has a relaxation time different from others. Hence the Cole–Cole plots in practice are not exactly semicircular, but are distorted to varying degree. The centers of these semicircles are depressed below the real axis by an angle. The relaxation time () for the circuit is given by the inverse of the frequency at the peak of semicircle at which ω = 1. In an ideal case with Debye behavior a perfect semicircle with its center lying on the real impedance axis is obtained indicating single relaxation. In polycrystalline materials in addition to the arc for dielectric relaxation within the grains (bulk relaxation), another arc due to the partial or complete blocking of charge carriers at the grain boundary may also be formed. Generally the electrode processes relax at low frequencies, grain boundaries relax at intermediate frequencies and the relaxation due to the grains of the samples occurs at higher frequencies. The contribution of various processes such as electrode reactions at the electrode—sample interface and the effect of migration of charge carriers through grains and across grain boundaries can be separated out in the frequency domain of measurements. The complex impedance plots do not show semicircles up to 250 ◦ C but as the temperature increases the curves attain greater curvature and become perfect semicircles at and above 350 ◦ C. All these plots terminate at the origin indicating the absence of series resistance in the equivalent circuit model of the sample. All the semicircles start on the real impedance axis at the lowest frequency. This starting point is found to decrease with increase in the temperature. This behavior of Cole–Cole plots is characteristic of the conducting nature of the samples and hence it may be concluded that there is no series capacitance in the equivalent circuit representation of the sample. The radius of the semicircles is decreasing with increase of temperatures indicating the relaxation time of the relaxing species as borne out by frequency explicit plots. The capacitance of an air-filled, parallel-plate condenser is given, neglecting fringing effects by C0 = A/d ε0 , where A is the area of the condenser, and d is the distance between the plates. With an applied alternating e.m.f, V, and a dielectric material between the plates, the

Fig. 8. Frequency variation of ac conductivity for Ce0.8 Gd0.2−x Yx O2 (a) x = 0, (b) x = 0.05 and (c) x = 0.1.

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Fig. 9. (a) Plots of frequency versus Z and M at different temperatures for Ce0.8 Gd0.2 O2 . (b) Plots of frequency versus Z and M at different temperatures for Ce0.8 Gd0.15 Y0.05 O2 . (c) Plots of frequency versus Z and Z at different temperatures for Ce0.8 Gd0.1 Y0.1 O2 .

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Fig. 9. (Continued ).

current is given by A i = jω(ε − jε ) ε0 V d But the field strength, E between the plates, is /d and the current density J is i/A. Substituting for conductivity (which is complex) which is the ratio of J and E: ∗ =

J = jω ε0 ε + ω ε0 ε E

Dielectric conductivity is real part of the ratio J/E in the above expression, i.e., d = ω ε0 ε Dielectric conductivity represents the sum of all the loss mechanisms in the material, and is a measure of the performance of a dielectric as an insulator. The power law frequency dependence of ac conductivity corresponds to the short range hopping of charge carriers through trap sites separated by energy barriers with different heights. As the temperature increases, the ac conductivity is frequency independent with a slight increase on the higher frequency side, which means that conductivity at high temperature was mainly activated by increasing temperature and the dc conductivity dominates at high temperatures. Low-frequency intercepts of the semicircle of Cole–Cole plots (Fig. 6) on real axis give the resistance of the sample at the temperature of the plots. The conductivity of all the samples is calculated from these intercepts and plotted as a function of inverse of tem-

perature on semi-log scale (Fig. 7). The activation energies of conductivity from these plots are in the range 1.03–1.09 eV. This is observed to be very close to the activation energy of relaxation in the range 1.03–1.2 eV within the experimental errors (Table 3). The frequency variation of conductivity calculated from real part of impedance (Z ) (Fig. 8) shows two or three slopes at different temperatures of the present measurements. The low frequency and low-temperature conductivity could be because of the intrinsic defects and charge agglomerations present in the sample. In general, ac conductivity follows the Jonscher power law: (ω) = 0 (T ) + A(T )ωn where T is the absolute temperature, ω is the angular frequency,  (ω) and  0 (T) correspond to the dc and ac conductivities, respectively. As frequency increases, two dispersion regions appear for all temperatures. If the low-frequency dispersion is associated with grain boundaries (since it is associated with the larger capacitance value) and the high-frequency one with grains (smaller capacitance value) and the data of the present measurements can be fit to the following modified power law [36]: (ω) = 0 + Aωn1 + Bωn2 + Cωn3 The values of exponents n1, n2 and n3 are evaluated from the slopes of log  versus frequency plots (Fig. 8). It can be seen that there are three slopes for the curves indicating existence of three different ac conductivity mechanisms in the samples. The values of the power law exponents are less than unity. The values of n1, n2 and

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n3 obtained in the present measurements are less than unity and vary with concentration. These values are related to the entropy or disorder present in the samples. To understand the electrical microstructure of the pellets and to find out whether the overall pellet resistances represented the bulk resistance of the grains or a contribution of grain boundary and other external parameters, the experimental data are replotted as the imaginary parts of the impedance, Z , and electric modulus, M , against log frequency, as shown in Fig. 9(a–c). The data shows single peaks in both Z and M . The peaks in Z , M versus frequency curves are separated by more than 1.14 decade of frequency. This indicates non-Debye nature of the relaxation of the charge species in the sample.

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4. Conclusions [14] [15]

Nanosize rare-earth co-doped Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1) materials were prepared by the sol–gel method. It was found that the doped ceria powders are single phase with cubic fluorite structure. It demonstrated that rare-earth co-doped Ce0.8 Gd0.2−x Yx O2 (x = 0–0.1) material prepared by the sol–gel method could be easy to sinter at lower temperature and obtain dense ceramic electrolyte with better electrical property. The unit cell parameter decreases with increasing yttrium content. The slope of the thermal expansion curves is slightly increased with increasing the yttrium content at high temperatures. The temperature, at which the change of slope occurs, decreases with increasing temperature. The TEC values of all the compositions are in the range 12–13 × 10−6 ◦ C−1 and x = 0.1 sample posses high conductivity at 600 ◦ C make it useful in IT-SOFC. Impedance measurements indicate presence of thermally active unassociated defects, which show non-Debye type relaxations. The conductivity follows modified power law.

[20] [21] [22] [23]

Acknowledgements

[31] [32]

The authors thank the reviewers for their critical comments and fruitful suggestions. The authors gratefully acknowledge the financial support of the Defence Research and Development Organization (DRDO), New Delhi in the form of a research project.

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