The effects of Co and Cr on the electrical conductivity of cerium gadolinium oxide

Solid State Ionics 282 (2015) 54–62

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The effects of Co and Cr on the electrical conductivity of cerium gadolinium oxide Samuel Taub a, Kerstin Neuhaus b, Hans-Dieter Wiemhöfer b, Na Ni a, John A. Kilner a, Alan Atkinson a,⁎ a b

Department of Materials, Imperial College, London SW7 2AZ, UK Institute for Inorganic and Analytical Chemistry, University of Münster, 48149 Münster, Germany

a r t i c l e

i n f o

Article history: Received 10 June 2015 Received in revised form 25 September 2015 Accepted 25 September 2015 Available online xxxx Keywords: Doped ceria Ionic conductivity Electronic conductivity Transition metal Chromium Cobalt

a b s t r a c t The effects of transition metal oxide dopants (Co and Cr) on the electrical conductivity of Ce0.9Gd0.1O1.95 (CGO) have been examined. Co is an effective sintering aid for CGO whereas Cr inhibits densification. After sintering both transition metals are found segregated at grain boundaries and as second phase precipitates at triple grain junctions. 2% Co addition has negligible effect on bulk conductivity, but increases the grain boundary intrinsic conductivity, whereas Cr addition in the range 0.01–4% also has negligible effect on bulk conductivity, but decreases the grain boundary intrinsic conductivity even at very low concentrations. For both Co and Cr these changes in conductivity are ionic and electronic conductivity remains negligible. The changes can be accounted for by changes to the boundary core charge (electrical potential) and surrounding space charge regions induced by the Co and Cr. Both Co and Cr increase the electronic conductivity mainly due to electron hopping along the network of transition metal ions segregated in the boundary cores. © 2015 Elsevier B.V. All rights reserved.

1. Introduction Doped cerium oxides have higher ionic conductivities than those of conventional zirconia solid solutions [1,2] and are therefore of interest as electrolytes in solid oxide fuel cells (SOFCs) operating in the intermediate temperature range (500–700 °C). In this temperature range relatively cheap stainless steels can be used as cell supports and interconnectors. Doping of the ceria with Gd or Sm, at a level of 10 to 20% of the cation sites (abbreviated to cat.%), gives the highest ionic conductivities [3,4]. In this paper we focus on Gd-doped ceria Ce(1 − x)GdxO(2 – x / 2), which is often shortened to a CGO or GDC. In this shorthand CGO10 refers to x = 0.1 and CGO20 to x = 0.2. Ceramic powders of these materials typically require sintering temperatures in excess of 1300 °C for full densification which makes the co-firing of the electrolyte and stainless steel supports difficult. As a result, a sintering aid is usually added in order to reduce the sintering temperature to an acceptable value (e.g. below 1000 °C). Co-doping ceria with the transition metal oxides (TMOs) of Co, Cu, Mn, Ni and Fe has been shown to improve densification when they are added at low concentrations (typically 1–2 cat.%) [5–9]. In particular, low concentration CoO-doping has proved to be an effective sintering aid without having a deleterious effect on the conductivity [10,11], with 2 cat.% being the optimum doping level [7,9]. As well as deliberately adding TMOs as sintering aids, the use of stainless steel could ⁎ Corresponding author. E-mail address: [email protected] (A. Atkinson).

http://dx.doi.org/10.1016/j.ssi.2015.09.024 0167-2738/© 2015 Elsevier B.V. All rights reserved.

potentially lead to elements from the steel (particularly Cr) entering the electrolyte during manufacture or during long term operation. In an earlier publication [12] we reported the effects on TMOs on the densification and microstructure of CGO10. That work confirmed that Co is a very effective sintering aid, but showed that Cr is a strong inhibitor of densification. Nevertheless, the effects of both TMOs on the microstructure were similar in terms of segregation at grain boundaries and second phase formation. It was concluded that the different effects on densification were due to differences in the electrical charge distribution at the grain boundaries which led to different effects on cation diffusion in the grain boundary cores. There are still uncertainties as to the effect of Co on the resulting electrical properties. Kleinlogel et al [10] reported that the addition of less than 2 cat.% Co had no influence on the total conductivity of CGO20 (Ce0.8Gd0.2O1.90), in agreement with Jud et al [13] for slowcooled Co-doped CGO. However, these authors did also report an increase in grain boundary electronic conductivity for samples that had been quenched after sintering. Lewis et al [14], on the other hand, reported that low level Co-doping increased the specific grain boundary ionic conductivity in CGO at lower temperatures, which was in agreement with the findings of Avila-Paredes et al [15] for Co-doped CGO10. Variable valence sintering aids and nano-sized grains have both also been shown to increase the electronic charge carrier concentrations in CGO [16,17]. In materials that are predominant oxygen ion conductors, common conductivity measurement techniques, such as electrochemical impedance spectroscopy (EIS) and 4-point DC measurements, are unsuitable for determining minority charge carrier concentrations,

S. Taub et al. / Solid State Ionics 282 (2015) 54–62

even when studied as a function of oxygen partial pressure. In principle an ion-blocking technique can measure the electronic conductivity, but nevertheless contradictory results have been presented in the literature. For example, whilst Fagg et al [11] reported a decrease in n-type electronic conductivity with addition of Co to CGO20, Schmale et al [18] reported notable increases in the electronic conductivity of (Co + Pr)doped CGO20. It is thus apparent that many discrepancies still exist in the literature relating to the effects of low concentration Co-doping in CGO and the mechanisms responsible for these changes in conductivity. In contrast to Co additions, there are no studies in the literature regarding the effects of Cr doping on the electrical properties of CGO. In this current work, we present a comprehensive study of the effects of two commonly found TMO dopants (Co and Cr) on the electrical conductivity of CGO using on a combination of electrochemical impedance spectroscopy and ion blocking measurements, based on the Hebb–Wagner polarisation technique. These results, combined with microstructural analyses have then been used as a basis to propose a mechanism responsible for the modified electrical properties that result from low level TMO doping in CGO. 2. Experimental methods Commercially available Ce0.9Gd0.1O1.95 (CGO) powder (NexTech Materials, Ohio, USA), with a crystallite size of approximately 36 nm, was doped with transition metals in the form of nitrate (Co(NO3)2·6H2O and Cr(NO3)3·9H2O, Alfa Aesar, Lancashire, UK) solutions in ethanol. The details are given in [12] and the procedure was repeated for the as-received CGO, but without adding the Co or Cr nitrates, to provide a dopant-free reference. In the case of Co-doped CGO the nominal Co concentration was 2 cat.% (designated 2CoCGO in this paper) and was chosen because previous studies had shown this to be a suitable concentration as a sintering aid. In the case of Cr-doped CGO, samples with Cr levels in the range 0.01 to 4 cat.% were prepared. These are designated 0.01CrCGO and so on. The powders were uniaxially pressed at 75 MPa for 15 s in a cylindrical stainless steel die before being isostatically pressed at 300 MPa for 30 s. The resulting green pellets were surrounded by powder of the same composition, to prevent loss of dopant by evaporation, and sintered as summarised in Table 1. The sintering conditions (heating rate, dwell temperature and dwell time) were designed to produce specimens of close to full density guided by the results of previous work [12] which also confirmed that there was no significant loss of Co or Cr during sintering. Densities were measured by the Archimedes method using theoretical densities determined by Rietveld refinement of x-ray diffractograms. The x-ray analysis showed no evidence of second phase formation in any samples, but previous high resolution transmission electron microscopy revealed the presence of second phase precipitates (below the detection limit of x-ray powder diffraction) at three-grain junctions in 2CoCGO and 2CrCGO samples [12]. The microstructures of the CGO, 2CoCGO and 2CrCGO are described in detail in [12]. The bulk and specific grain boundary contributions to the total conductivity were evaluated using AC impedance spectroscopy. Discs, approximately 10 mm in diameter and 1 mm in thickness, were pressed and sintered to high density in accordance with Table 1. All samples were subsequently ground flat using 800 and 1200 grit SiC paper and polished to a 0.25 μm finish using a succession of diamond suspensions. Pt electrodes were applied as a paste and fired at 900 °C for 30 min.

Table 1 Sintering conditions at heating rates of 5 °C min−1. Material

Dwell temperature (°C)

Dwell time (min)

CGO Co-doped CGO Cr-doped CGO

1350 1020 1500

300 45 300

55

Impedance spectra were obtained using a Solartron 1260 Frequency Response Analyser over the range 1 to 107 Hz with an excitation voltage of 50 mV. Measurements were recorded as a function of temperature between 120–700 °C in air. All data were recorded using ZPlot and equivalent circuit modelling was performed in the associated ZView software package (Scribner Associates Inc., North Carolina, USA). The impedance response was modelled using a series resistor and inductor for the contact leads and a resistor in parallel with a constant phase element (CPE) for each of the bulk and grain boundary responses, which were connected in series. The impedance of a CPE is given by Z = Q− 1(jω)− n and the equivalent capacitance by C = (R1 − nQ)1/n, where R is the parallel resistance in the CPE. The electronic conductivity was measured using the Hebb–Wagner technique using ion-blocking electrodes. Sintered pellets, 8 mm diameter and 3 mm thickness, were ground flat using 800 and 1200 grit SiC papers and then polished to 0.25 μm finish using diamond suspensions. The experimental arrangement for conductivity measurement was as described in [18]. One electrode was a spring-loaded Pt wire with 0.5 mm tip and shielded from the surrounding atmosphere by encapsulating the contact in a glass bead. The counter electrode was porous Pt paste. The samples were heat treated for 12 h at 800 °C before assembling the setup to ensure a complete equilibration of the samples in air, which prevents cracking after the glass encapsulation. After assembly, the samples were briefly heated to 850 °C to sinter the Pt paste and glass bead. Before any measurements were made the sample was again equilibrated in the measuring rig at 600 °C for 48 h under no applied potential. The oxygen pressure at the counter electrode was controlled using flowing gas mixtures. For each test temperature the applied voltage at the point electrode was reduced from 0 to −0.6 V (with respect to the counter electrode) and then increased back to 0 V with a step size of 0.02 V. At each step the sample was held for 30 min to ensure that a steady state concentration gradient was attained at each applied voltage, and then the steady state electronic current was measured. At each test temperature the sample was equilibrated for 1 h before starting the polarisation experiment at the new temperature. In order to deduce an accurate value for the local oxygen activity at the point electrode it is necessary to make a correction for the contribution of the contact resistance to the potential drop across the sample by a current interruption method [18]. This could only be done with confidence at 700 °C and above because of the high sample resistance at lower temperatures. Therefore measurements made at temperatures below 700 °C could have errors in the estimated oxygen activity, at low oxygen activities, that varies from sample to sample and so detailed comparison between samples at low oxygen activities is restricted to temperatures of 700 °C and above. (Scanning) transmission electron microscopy ((S)TEM) crosssections were prepared from sintered, bulk specimens using a focused ion beam (FIB) milling system (Helios NanoLab, FEI Company, The Netherlands, 2 − 30 keV Ga + incident beam energy with currents of 16 pA–21 nA). A ~ 20 nm gold coating was deposited onto the surface of the sintered pellets by sputter deposition prior to milling in order to prevent charging of the samples. To reduce the damage caused by the high energy Ga + beam and improve the quality of the specimens for subsequent TEM analysis, the specimens were polished at the last stage with 2 keV Ga + . (S)TEM work was carried out on a FEI Titan 80–300 S/TEM operated at 300 kV equipped with a monochromator, a Cs aberration image corrector and Gatan Tridiem 866 imaging filter. Focal series micrographs of grain boundaries were acquired at different objective lens focus values (using a spherical aberration coefficient Cs ~− 4 μm) and exit-wave reconstruction was performed using TrueImage software (FEI). Electron energy loss spectroscopy (EELS) was performed in STEM mode, and the incident angle α and collection angle β for EELS acquisitions were ~ 10 and ~ 14 mrad, respectively. The energy resolution of EELS under experimental conditions was ~ 0.7 eV, as defined by the full width at half-maximum of the zero-loss peak (ZLP).

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S. Taub et al. / Solid State Ionics 282 (2015) 54–62

3. Results

Table 2 Activation energy, Ea (kJ mol−1), for total bulk and grain boundary conductivity.

3.1. Ionic conductivity At temperatures between 150–300 °C, two arcs were observed in the Nyquist plots (not including the electrode response) with a CPE circuit element index, n, typically in the range 0.85–0.90. The corresponding capacitances were approximately 5 × 10−11 F cm−2 for the bulk and 2 × 10−8 F cm−2 for the grain boundary processes, which are typical for doped ceria [19]. The grain boundary capacitances tended to be slightly higher for CrCGO than for CGO and slightly lower for CoCGO than CGO. At temperatures above ~ 550 °C, the Nyquist plots only showed frequency dependence due to the electrode impedance. Between these temperatures, one arc was observed, with a capacitance value typical of a grain boundary process. The low frequency intercept of this arc with the real axis gave only the overall resistance of the electrolyte and the bulk and grain boundary contributions could not be separated in this temperature range. Thus it was possible to calculate the total resistivity, ρtot, and (for temperatures below 350 °C) the individual contributions from bulk, ρbulk, and grain boundaries, ρgb. It has become common in the literature to take the reciprocals to give σbulk and σgb. However, this can be misleading as the bulk and grain boundaries are in series in the model and therefore the conductivities cannot be added to compute the total. Furthermore, the grain boundary contribution to the total resistivity depends on the number of grain boundaries crossed by the current and thus depends on microstructure. We therefore choose to discuss the grain boundary contribution in terms of its specific conductivity, σ*gb, defined as the effective conductivity of a slab of material of width δ at each grain boundary which has the same effect on the total conductivity, σ
gb ¼

δ σ D gb

Ebulk a Egb a

CGO

2CoCGO

2CrCGO

76 ± 1 93 ± 3

68 ± 2 88 ± 1

77 ± 2 102 ± 3

with the bulk lattice conduction and Cgb that associated with the grain boundary conduction. The bulk conductivity and specific grain boundary ionic conductivity (in the temperature range in which they could be separated) for CGO without TMO doping are shown in Fig. 1. There is good agreement between measurements taken during heating and cooling indicating that the samples were in a stable condition and that no significant permanent changes were induced by heating. The activation energy (Table 2) for bulk ionic conductivity in CGO was measured to be 76 ± 1 kJ mol− 1 and that for grain boundary conductivity was 93 ± 3 kJ mol−1 over this temperature range. It can be seen in Figs. 1, 3 and 4 that the Arrhenius plots for grain boundary conductivity all show slight upward curvature at the lower temperatures. This could be due to small changes in boundary segregation with temperature. The values of activation energy in Table 2 are averages over the temperature ranges shown in the corresponding figures.

a)

ð1Þ

where D is grain size. Since the capacitance of a parallel plate capacitor is inversely proportional to the distance separating the plates, it is possible to rewrite Eq. (1) as σ
gb ≃

C bulk σ C gb gb

ð2Þ

1 nm

assuming that the relative permittivities of the bulk and grain boundaries are the same [20]. In Eq. (2) Cbulk is the capacitance associated

Counts (a.u.)

b)

at GB 515

525

535

545

Grain interior 555

565

Energy Loss (eV)

Fig. 1. Arrhenius plot of conductivity of CGO without transition metal doping. Measurements were recorded upon heating and cooling to/from 700 °C.

Fig. 2. TEM characterisation of grain boundaries in CGO without TMO dopants sintered at 1350 °C for 5 h. (a) a focal series exit-wave reconstructed phase image of the grain boundary: (b) EELS oxygen k edge spectra from the boundary core and the grain interior.

S. Taub et al. / Solid State Ionics 282 (2015) 54–62

Analysis of the grain boundaries by TEM [12] showed grain boundary Gd enrichment (5.8 mass % above bulk composition) over a region no more than 2 nm in width, which indicates Gd segregation to the boundary cores. A similar analysis for Si revealed no Si above the limit of detection (0.5 mass %). The lattice image of a boundary in Fig. 2 shows no indication of amorphous material at the boundary core. Also shown in Fig. 2 are EELS spectra of the oxygen K edge at the boundary core and in the grain interior. The similarity in the spectra indicates that the oxygen environment in the cores has a similar degree of order to the oxygen in the grain interiors. It is therefore concluded that the grain boundary conductivity, which is 2–3 orders of magnitude lower than the bulk conductivity, is representative of “clean” boundaries in CGO. The conductivity results for 2CoCGO are summarised in Fig. 3. The addition of 2 cat.% Co is shown to increase the total conductivity at the lower temperatures below approximately 400 °C. The results for the different contributions to the total conductivity are shown in Fig. 3a. These reveal that the Co has had negligible effect on the bulk conductivity and that the increase in total conductivity is due to a significant increase in the specific grain boundary conductivity. The Co addition has reduced the activation energy (Table 2) for grain boundary conduction slightly from 93 ± 3 to 88 ± 1 kJ mol−1. These findings are in general agreement with earlier studies [9,11,16]. The effects of low concentration Cr doping on the bulk and specific grain boundary conductivity of CGO are shown in Fig. 4. Cr concentrations between 0.01 and 4 cat.% are seen to have negligible effect on

Fig. 3. Conductivity as a function of temperature in 2CoCGO (a) bulk and specific grain boundary conductivity measured upon heating (b) total conductivity measured upon heating and cooling.

57

Fig. 4. Arrhenius plot of conductivity for Cr-doped CGO (0.01–4 cat.% Cr) measured upon heating. The dashed lines represent the bulk and specific grain boundary conductivities of CGO without transition metal addition.

the bulk conductivity, which is due to the low solubility of the dopant within the host lattice. The specific grain boundary conductivity, on the other hand, shows a dramatic decrease of over two orders of magnitude even for Cr concentrations as low as 0.01 cat.% (100 ppm). Furthermore, the reduction of intrinsic grain boundary conductivity is approximately independent of the total Cr doping level within the range studied. These observations are in direct contrast to the results of Zajac et al [21], who studied Cr-doping of CGO15. In that study, Cr concentrations between 0.5–2 cat.% were observed to increase the grain boundary conductivity for samples sintered at 1500 °C. This difference might be due partly to volatilisation of Cr species during sintering in their samples, which were not analysed chemically after sintering. TEM analysis of a grain boundary in a Cr-doped specimen is shown in Fig. 5. The Si concentrations were again below the detection limit of the analysis method in agreement with the analysis of the Cr-free grain boundary. Hence it can be concluded that the effect of Cr on the grain boundary intrinsic conductivity is not due to an attraction of Si to Cr in the grain boundary core. Fig. 5 shows a lattice image of a grain boundary in 2CrCGO. There is no evidence of any amorphous grain boundary phase, although the boundary core appears to be more disordered than in the Cr-free case (Fig. 2). The increased level of disorder at the grain boundary in the 2CrCGO sample was confirmed by the change of the energy loss near-edge structure (ELNES) of the oxygen K-edge (Fig. 5b) compared with the Cr-free sample shown in Fig. 2b. ELNES probes the local density of unoccupied electronic states, which is influenced by the local atomic arrangement. Three characteristic peaks can be clearly distinguished in the ELNES spectrum from the grain interior of the Cr-free sample as a result of a well-defined local symmetry in the arrangement of oxygen atoms. In order of increasing energy loss, the first peak is attributed to a transition O 1s → O 2p hybridised with Ce 4f [22]. The second and third peaks are related to the transition O 1s → O 2p hybridised with Ce 5d and are a signature of tetrahedrally coordinated oxygen atoms in the fluorite structure [23,24]. In the ELNES spectrum from the boundary core of the Cr-free specimen (Fig. 2b) all three peaks are clearly still present indicating preservation of the local order of the fluorite lattice even in the boundary core. In the case of CrCGO sample (Fig. 5b) the 3 peaks are again displayed in the ELNES spectrum from the grain interior, but the peaks are not visible in the ELNES spectrum from the boundary core. This is a result of the presence of many different local atomic arrangements, and essentially the loss of local symmetry, consistent with a highly disordered core structure. The activation energy for bulk conductivity in the Cr-doped specimens (Table 2) was measured to be 77 ± 2 kJ mol−1, in good agreement with the Cr-free experiments. However, the activation energy for the

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S. Taub et al. / Solid State Ionics 282 (2015) 54–62

a)

2 nm

Fig. 6. Examples of steady-state I–V curves for CGO, 2CoCGO and 2CrCGO in the temperature range 750–800 °C. The sign of the applied voltage refers to the polarity of the blocking electrode with respect to the reference electrode.

Counts (a.u.)

b)

Grain interior 515

525

535

545

at GB 555

565

Energy Loss (eV) Fig. 5. TEM characterisation of a grain boundary region in a 2CrCGO specimen. (a) A focal series exit-wave reconstructed phase image of a typical grain boundary: (b) EELS oxygen k edge spectra from the boundary core and the grain interior.

intrinsic grain boundary conductivity was increased to 105 ± 5 kJ mol−1 averaged over all the Cr-doped samples. 3.2. Electronic conductivity

where a is the contact radius of the micro-contact and the oxygen pressure at the blocking electrode is given by application of the Nernst equation. Fig. 7 shows the resulting electronic conductivity of 2CoCGO as a function of oxygen partial pressure between 350–650 °C. (Note that at these temperatures there could be errors in the estimated oxygen pressure due to no reliable correction for the contact resistance being possible at these low temperatures and low oxygen pressures as explained in the experimental section.) Increasing the temperature is shown to increase the electronic conductivity and shift the conductivity minimum towards higher oxygen partial pressures. Similar findings have been reported elsewhere in the literature for Ce0.8Gd0.2 − xPrxO2− δ [18]. At lower temperatures, the n-type conductivity of 2CoCGO is approximately proportional to pO2− 1/6, in agreement with the findings of Fagg et al [11]. A comparison between the electronic conductivity of 2CoCGO and undoped CGO, as a function of oxygen partial pressure at the higher temperatures studied, is presented in Fig. 8 and clearly shows that Co has increased the electronic conductivity across the entire range of oxygen pressure studied and particularly in the p-type region. The comparison can be made reliably at these temperatures because it was possible to make a reliable correction for the contact resistance at these higher temperatures. These results are in very good agreement with those of Fagg et al [11,16] for 2CoCGO20. At these higher temperatures the

Fig. 6 shows selected results from polarisation measurements (i.e. the steady-state I–V curves) of CGO, 2CoCGO and 2CrCGO, where the negative sign of the applied voltage refers to the polarity of the blocking electrode with respect to the reference electrode. There is detectable hysteresis in the I–V curves of 2CoCGO which is attributed to a redox process involving the Co ions. Hysteresis is also observed for the CGO sample, although to a much lesser extent, and is probably caused by the partial reduction of Ce4 + to Ce3 + (and re-oxidation) at higher temperatures [25]. The increasing gradient at the higher voltages (lower average oxygen activities in the sample) corresponds to an increase in n-type conductivity [26]. The addition of Co-oxide and an increase in temperature are both shown to displace the onset of the enhanced n-type conductivity to higher oxygen partial pressures. According to the Hebb–Wagner model, the electronic conductivity, σe, can be calculated as a function of oxygen partial pressure at the blocking electrode, pO2, from the slope of the steady-state I–V curve;

σe ¼


 1 ∂I 2πa ∂V

ð3Þ

Fig. 7. Electronic conductivity of 2CoCGO as a function of oxygen partial pressure measured between 350 and 650 °C. The dashed line indicates a −1/6 power proportionality.

S. Taub et al. / Solid State Ionics 282 (2015) 54–62

Fig. 8. Electronic conductivity of 2CoCGO as a function of oxygen partial pressure at 750 and 800 °C compared with data for CGO.

electronic conductivity of CGO is approximately proportional to pO2 to the power −1/4 in the n-type regime and +1/4 in the p-type regime in agreement with previously reported results for CGO20 [26]. The electronic conductivity of 2CoCGO does not show a simple dependence on oxygen activity at these higher temperatures and this is attributed to changes in valency of the Co ions. The apparent activation energy for n-type electronic transport in 2CoCGO at constant oxygen activity was measured to be 138 ± 4 kJ mol−1. Table 3 summarises values of the total conductivity of CGO and 2CoCGO and values of the electronic conductivity measured at pO2 = 0.21 atm. Also shown is the electronic transference number te. These values indicate that the enhanced electronic conductivity due to the Co cannot account for the majority of the increased total conductivity observed in 2CoCGO and that the increase is mainly ionic. Furthermore, although there is an increase in electronic conductivity with Co doping, the electronic transference number remains below 2 × 10−3. The electronic conductivity of 2CrCGO (2 cat.% Cr-doped CGO) is shown in Fig. 9. Unlike 2CoCGO, 2CrCGO exhibits a clear dependency on pO2− 1/4 in the n-type regime and pO21/4 in the p-type regime as found for CGO. A comparison between the electronic conductivity of 2CrCGO and CGO, as a function of oxygen partial pressure, is presented in Fig. 10 for 750 and 800 °C. The addition of 2 cat.% Cr is shown to enhance both n-type and p-type conductivity by approximately equal factors. The apparent activation energy for electronic conductivity on 2CrCGO was measured to be 132 ± 4 kJ mol−1 for p-type conduction and 221 ± 4 kJ mol−1 for n-type conduction. 4. Discussion 4.1. Ionic conductivity The results for electronic conductivity in CGO, 2CoCGO and 2CrCGO show that the electronic transfer number in air is always sufficiently low that the impedance results can be interpreted as measurements of ionic conduction. The effect of Co on bulk ionic conduction is small (Fig. 3), but noticeably positive at the lower temperatures (b200 °C). This

59

Fig. 9. Electronic conductivity as a function of oxygen partial pressure for 2CrCGO measured between 450 and 700 °C. The dashed line indicates conductivity proportional to pO2−1/4.

negligible effect reflects the low solubility of the transition metal ions in the ceria lattice. The small increase in conductivity at the lower temperatures in 2CoCGO could be due to a slight increase in oxygen vacancy concentration following some Co2+ substituting for Ce4+ [14]. However, no similar increase is seen in 2CrCGO (Fig. 4) and this could reflect an even lower solubility for Cr in the ceria lattice. On the other hand, the effects of Co and Cr on the intrinsic grain boundary conductivity are large and different. Co produces an increase in specific boundary conductivity by approximately a factor of 2, whilst Cr produces a large decrease, by over an order of magnitude, that is not sensitive to the Cr content for overall Cr contents between 0.01 and 4 cat.%. The grain boundary compositions for 2CoCGO and 2CrCGO have been studied using high resolution analytical TEM [12] and the results are summarised in Table 4. As discussed in [12] the compositions quoted are averages over approximately 2 nm either side of the boundary centre line due to electron beam broadening by the sample degrading the obtainable resolution. Therefore the concentration differences in the boundary cores could be larger than those given in Table 4. The observed segregation of Gd to the grain boundaries in CGO is consistent with reports in the literature [27–29]. Both Co and Cr dopants have a similar effect on boundary composition in that Gd is segregated (as in CGO), Ce is depleted and the transition metal is segregated at the boundary. For both Co and Cr most of the transition metal is present as second phase precipitates at triple grain junctions. In the case of Co the precipitates were identified as being similar to Co3O4 and for Cr the precipitates are a mixed oxide of Cr and Gd. Therefore, from analysis of the grain boundary compositions there is no obvious reason why Co and Cr should have such markedly different effects on the intrinsic grain boundary ionic conductivity. In a fast ion-conducting ceramic the grain boundaries have lower intrinsic conductivity than the bulk grains and the oxygen ion (in CGO) pathway is across the grain boundaries rather than along them. This is often approximated by the so-called “brick layer” model of regular cube-shape bulk grains (bricks) separated by continuous layers of grain boundary. Grain boundaries in fast ion conductors are known to

Table 3 Values of the total conductivity and electronic conductivity of CGO and 2CoCGO measured between 600 and 700 °C at 0.21 atm. All conductivity values are given in S cm−1. T (°C)

600 650 700

CGO

2CoCGO

σtotal

σe

te

σtotal

σe

te

1.33 × 10−2 2.21 × 10−3 3.00 × 10−2

3.76 × 10−7 1.81 × 10−6 5.12 × 10−6

2.83 × 10−5 8.19 × 10−4 1.71 × 10−4

1.51 × 10−2 2.68 × 10−2 3.77 × 10−2

6.12 × 10−6 1.25 × 10−5 6.45 × 10−5

4.05 × 10−4 4.66 × 10−4 1.71 × 10−3

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S. Taub et al. / Solid State Ionics 282 (2015) 54–62

Fig. 10. Electronic conductivity as a function of oxygen partial pressure for CGO and 2CrCGO at temperatures between 750 and 800 °C.

have a complicated structure with an electrically charged core and narrow space charge regions of opposite charge extending a short distance on either side of the core. The resistance to the motion of the oxygen ions across the boundary therefore contains contributions from both the space charge layers and the boundary cores acting in series. The space charge region has a lower conductivity than the bulk because it is depleted in oxygen vacancies and the core also has lower conductivity because it is disordered (despite being oxygen deficient). In fluoritestructured oxygen ion conductors there is strong evidence that the core carries a net positive charge and that the space charge is negative [30]. The core positive charge is due to an excess of cation charge over the stoichiometric value and could be regarded as either cation excess or oxygen deficiency due to the preferred segregation of oxygen vacancies to the core. The compensating negative space charge is due to a depletion of oxygen lattice vacancies which have effective positive charge. Guo and Waser [30] give the following expression relating the boundary specific conductivity to the electrical potential difference, Δϕ, at the boundary core relative to the bulk: σ b expð2eΔϕ=kB T Þ ≅ σ
gb 4eΔϕ=kB T

ð4Þ

where e is the magnitude of the electronic charge. Applying this equation to the present results at 500 K, the boundary core potential for CGO without the transition metals is estimated to be 0.22 V. This is very close to the value of approximately 0.25 V deduced in a similar way by Guo et al. [31] for ceria doped with 10 mol% Y2O3. In addition, Guo and Waser derive the following expression for the difference in activation energies for bulk and grain boundary conduction:   1 dðΔϕÞ bulk ¼ ð2eΔϕ−kB T Þ 1 þ Egb : a −Ea TΔϕ dð1=T Þ

ð5Þ

The effect of the transition metals on grain boundary intrinsic ionic conductivity could therefore come from their modification of either

Table 4 Grain boundary composition of samples as the difference between the measured grain boundary concentration and that of the bulk [12]. Material

Ce (mass %)

Gd (mass %)

Co or Cr (mass %)

CGO 2CoCGO 2CrCGO

−3.6 −7.5

+5.8 +2.1 +4.7

+0.7 +0.3

the space charge regions or the boundary cores as proposed earlier by Avila-Paredes and Kim [15]. In an earlier paper on the influence of transition metals on sintering of CGO we proposed that the different effect of Co and Cr on sintering was due to their different ionic charge. Although the concentrations of Co and Cr at the grain boundaries are similar (Table 4) the higher charge of Cr would lead to an increase in charge at the boundary core and Co to a decrease. Thus the Co would increase the concentration of oxygen vacancies in the space-charge region (reduce the depletion) and increase the ionic conductivity in the space charge region as observed. For 2CoCGO at 500 K, the observed increase in grain boundary conductivity relative to CGO corresponds, according to the space charge model and Eq. (4), to a reduction in the core potential to 0.19 V. Conversely, for 2CrCGO at 500 K, the decrease in grain boundary conductivity relative to CGO corresponds to an increase in the core potential to 0.28 V. Thus the grain boundary conductivity is very sensitive to small changes in the core potential as a result of the exponential dependence in Eq. (4) and hence is sensitive to relatively small changes in segregation of charged species at the boundary core. If this is the mechanism responsible for changes in the grain boundary conductivity, then the activation energy for bulk and grain boundary conductivity would be expected to be related by Eq. (5). Applying Eq. (5) to the present data for CGO predicts that the activation energy for grain boundary conduction should be 25 kJ mol−1 higher than for bulk conduction. The measured values in Table 2 show a corresponding value 17 ± 4 kJ mol−1 which is in reasonable agreement with the predicted value and further supports the validity of the space charge model in CGO without TMO additions. For 2CoCGO the application of Eq. (5) predicts that the difference in activation energies for bulk and grain boundary should be 16 kJ mol−1, slightly smaller than the 25 kJ mol−1 predicted for CGO. The corresponding measured values (Table 2) show a difference of 20 ± 3 kJ mol−1 for 2CoCGO which is the same as that for CGO to within experimental error. Although the experiments do not show the smaller difference predicted by the space charge model, the discrepancies are small and close to the experimental errors. Therefore the effects of Co on grain boundary conductivity are broadly consistent with the space charge model. The increase in core charge from the Cr ions would give the opposite effect and reduce the conductivity, also as observed. Application of Eq. (4) predicts an increase of core potential from 0.22 V to 0.28 V as a result of Cr addition. The Cr is observed to increase the activation energy for grain boundary conductivity to 102 ± 3 kJ mol− 1, which is 25 ± 4 kJ mol−1 higher than for bulk conduction. Application of Eq. (5) predicts that the difference in activation energies for grain boundary and bulk conduction in 2CrCGO should be 27 kJ mol− 1 which is in good agreement with the measured value. Therefore the effect of Cr on the grain boundary conductivity is also consistent with the space charge model. This suggests that motion across the boundary core is having negligible influence. The TEM result in Fig. 5 shows no evidence of any amorphous second phase at the boundary (as is known to occur with low levels of Si contamination), but there is some indication that the boundary core is more disordered by the Cr when compared with Fig. 2. It is also known that impurities segregated at grain boundaries can have profound effects on the core structure leading to different boundary “complexions” [32,33]. Nevertheless, the present results indicate that the main influence of the boundary core is not due to its structure, but due to changes in electrical potential caused by segregation. 4.2. Electronic conductivity The power law dependences on oxygen activity seen in Figs. 7, 8, 9 and 10, suggest oxygen activity dependent electronic carrier concentrations that are related to lattice defect populations in CGO, CoCGO and CrCGO, although CoCGO and CrCGO show differences compared to CGO. Simple defect equations for oxidation of CGO result in the

S. Taub et al. / Solid State Ionics 282 (2015) 54–62

following equilibria between defect populations (strictly speaking these should be activities, but are usually approximated as being proportional to concentrations). Ke ¼

1 h i €O p ½e0 2 V O2 1 2

h i2 h Kh ¼ h i €O p V O2

ð6Þ



1 2

ð7Þ

In bulk CGO local electrical neutrality prevails and (except in very reducing conditions) the oxygen vacancy concentration is fixed by the Gd concentration. Thus Eqs. (6) and (7) give the electron and hole concentrations proportional to oxygen pressure to the power of −1/4 and + 1/4 respectively. This can be seen in Figs. 8 (for CGO) and 9 (for CrCGO) for temperatures above 450 °C. The reason for the lower power of −1/6 seen at lower temperatures for CoCGO (Fig. 7) is unclear. But as CrCGO shows the regular − 1/4 and + 1/4 slopes in Fig. 9, the different power law for CoCGO indicates a different influence of the added Co on the effective electronic conductivity. A possible explanation may be that the CoCGO samples exhibit an enhanced electronic conduction along the grain boundaries. When the electronic carriers are trapped (polarons) the electrons can be regarded as Ce3+ ions with an electron energy just below the conduction band, and the holes as O− ions with an electron energy just above the valence band. Furthermore, and in particular with CoCGO, there is the possibility that the grain boundaries also influence electronic conduction in CGO. In the core plus space charge model outlined earlier, the positive core potential lowers the concentration of oxygen vacancies in the space charge regions (local electrical neutrality is no longer obeyed in the space charge regions, but the space charge plus the core charge is neutral). From Eqs. (6) and (7) the lower concentration of oxygen vacancies in the space charge region will lead to an increase in the electron concentration and a reduction in the hole concentration relative to the bulk, but the concentrations will follow the same power law dependences on oxygen pressure. (As a further complication, the positive core charge will attract electrons to the core and holes to the space charge regions. The concentrations of these minority defects are sufficiently low that they will not lead to any major change to the core and space charges). Thus it is not possible to determine from the power laws alone how much electronic conduction is within the space charge regions and how much is in the bulk. For example it has been shown that in nano-crystalline ceria [25] electronic conductivity is mainly by electron transport along the grain boundaries. Bearing all these complications in mind, we now discuss the influence of the transition metal dopants on the electronic conductivity. In the CGO lattice, Co is likely to be an acceptor and Cr either an acceptor or inactive (depending on whether its charge is +3 or +4). If the acceptors are ionised then they should increase the p-type conduction in the bulk and suppress the n-type conduction. This would simply mean a shift of the minimum of the electronic conductivity to lower oxygen activity. From Fig. 10 this is clearly not the case for Cr and it is likely that its solubility in the bulk lattice is so low that its influence on the bulk electronic conductivity is small. However, if the Cr increases the grain boundary core charge, as suggested to explain its effect on ionic conduction then the oxygen vacancy concentration will fall in the space charge regions and, from Eqs. (6) and (7), the electron concentration will rise and the hole concentration will fall. This is also not consistent with the results in Fig. 10, but they do show a smaller increase in the p-type region than in the n-type. Assuming unchanged mobilities of electrons and holes, the upwards shift of the conductivity minimum in Fig. 10 from CGO to CrCGO, without change of the position with respect to the partial pressure, would imply a decrease of the thermal band gap for electron–hole pair formation which, in view of the small Cr

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concentration, is only imaginable by enhanced Cr concentration at the grain boundary. Therefore it is likely that in the case of Cr there is an increase of electron conductivity in the space charge layer (hopping between Ce3+ and Ce4 +) and an increase in hole conductivity in the cores (hopping between Cr3+ and Cr4+). At the minimum in electronic conductivity we speculate that most of the Cr is Cr3 + and that the proportion of Cr4 + rises at higher oxygen activity. At lower oxygen activities the Cr remains as Cr3+ and therefore there is no n-type conduction through the Cr network in the boundary cores. The behaviour of Co is different in that the simple power law dependences on oxygen activity are not preserved. In addition, the electronic conductivity is distinctly higher over the entire oxygen activity range as compared to that of CGO and also higher compared to CrCGO. In this case it appears that the Co ion charges are undergoing larger changes in populations with oxygen activity and there is no clear minimum in conductivity associated with a p-type to n-type transition (Fig. 8) except at low temperatures, but with slopes of − 1/6 and + 1/6 (Fig. 7). We speculate that the reason for this is that at the oxygen activity of the conductivity minimum in CGO (approximately 10−4) the Co ions are in a mixed valence state with comparable numbers of Co2+ and Co3+ ions. This gives rise to a maximum probability for hopping along the Co network in the boundary cores. At higher oxygen activity Co3+ dominates and at lower oxygen activity Co2 + dominates and conduction along the boundary cores is lowered in both extremes. This can be seen in Fig. 8 for oxygen activity below approximately 10−10. This explanation is consistent with the conclusions reached by Schmale et al. [18] regarding the influence of Co on Ce0.8Gd0.2 − xPrxO2− δ. 5. Conclusions Co is an effective sintering aid for CGO whereas Cr inhibits densification. After sintering both transition metals are found segregated at grain boundaries and as second phase precipitates at triple grain junctions. 2 cat.% Co addition has negligible effect on bulk conductivity, but increases the grain boundary intrinsic conductivity. Cr addition in the range 0.01–4 cat.% also has negligible effect on bulk conductivity, but decreases the grain boundary intrinsic conductivity. For both Co and Cr these changes in intrinsic grain boundary conductivity are ionic and electronic conductivity remains negligible (at oxygen activities studied between 0.2 and 10−13). In the case of Co, the increase can be explained by Co reducing the boundary core charge (electric potential) and increasing oxygen vacancy concentrations in the space charge regions surrounding the boundaries. Cr has the opposite effect in increasing the boundary core charge and decreasing the oxygen vacancy concentrations in the space charge regions. These effects, and associated changes in activation energies, are consistent with the space charge model for grain boundary conductivity in fast ion conductors and the explanation proposed previously to account for the effect of Co and Cr on the sintering of CGO. The direct influence of the core structure of the boundary on oxygen ion transport across the boundary is lower than that of the space charge. Both Co and Cr increase the electronic conductivity and this is concluded to be due to electronic conduction along the grain boundaries (as opposed to across them for ionic conduction). Part of their effect can be attributed to their influence on the space charge regions, but most comes from electron hopping along the network of transition metal ions in the boundary cores. It is speculated that Cr is mainly as Cr3 + ions, but some are oxidised to Cr4+ at higher oxygen activities. Co, on the other hand, is probably in a mixed state of Co2+ and Co3+ ions leading to a maximum contribution to electronic conduction at intermediate oxygen activities. Acknowledgements The authors would like to thank the EPSRC and Ceres Power Limited who supported this work in the form of a CASE award.

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