Effect of Nb substitution for Ti on the electrical properties of Yb2Ti2O7-based oxygen ion conductors

SOSI-13264; No of Pages 10 Solid State Ionics xxx (2014) xxx–xxx

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Effect of Nb substitution for Ti on the electrical properties of Yb2Ti2O7-based oxygen ion conductors L.G. Shcherbakova a, J.C.C. Abrantes b,c, D.A. Belov a,d, E.A. Nesterova d, O.K. Karyagina e, A.V. Shlyakhtina a,⁎ a

Semenov Institute of Chemical Physics, Russian Academy of Sciences, ul. Kosygina 4, Moscow 119991, Russia UIDM, ESTG, Instituto Politécnico de Viana do Сastelo, Apartado 574, 4901-908 Viana do Castelo, Portugal Ceramics Dept., CICECO, University of Aveiro, 3810 Aveiro, Portugal d Faculty of Chemistry, Moscow State University, Leninskie gory, Moscow 119991, Russia e Emanuel Institute of Biochemical Physics, Russian Academy of Sciences, ul. Kosygina 4, Moscow 119991, Russia b c

a r t i c l e

i n f o

Article history: Received 27 April 2012 Received in revised form 27 December 2013 Accepted 15 January 2014 Available online xxxx Keywords: Oxide ion conductivity Pyrochlore Acceptor doping Donor doping Impedance spectroscopy

a b s t r a c t We have studied the effect of niobium doping on the electrical conductivity of Yb2Ti2O7-based oxygen ion conductors. Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) pyrochlore solid solutions were synthesized through coprecipitation followed by firing at 1550 °C for 4 h. The materials were examined by XPS, XRD, scanning electron microscopy and impedance spectroscopy. Yb2(Ti0.99Nb0.01)2O7 was shown to have the highest oxygen ion conductivity in air (2.3 × 10−3 S/cm at 750 °C), which is however markedly lower than that of undoped Yb2Ti2O7. In the (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) system, the highest conductivity is offered by (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (σ = 4.44 × 10−3 S/cm at 650 °C). Additional oxygen vacancies created by Ca doping in pyrochlore structure reduce the detrimental effect of Nb4+ doping on the oxide ion transport up to 5% Nb. The conductivity of the Yb2(Ti0.99Nb0.01)2O7 and (Yb0.8Tb0.1Ca0.1)2 [Ti0.95Nb0.05]2O6.9 solid solutions was measured both in air and under reducing conditions (5% H2 in N2 and CO2 atmospheres). A comparative study of both these compositions under 5% H2 in N2 atmosphere showed that the transport mechanism was not affected by complex doping of the lanthanide and titanium sublattices in the Yb2Ti2O7-based materials and was related to oxygen vacancies. Conductivity measurements in CO2 were done to ensure correct evaluation of the ionic conductivity of (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9, because in air it seems to be a mixed p-type and ionic conductor. © 2014 Elsevier B.V. All rights reserved.

1. Introduction The conduction mechanism in the A2B2O7 (3 +/4 +) pyrochlores was analyzed in detail in the literature [1–7]. The pyrochlore structure typically contains oxygen vacancies and oxygen is therefore transported in pyrochlores through the vacancies. The highest oxygen ion conductivity in air among rare-earth pyrochlores, exceeding the conductivity of Y2O3-stabilized ZrO2, a well-known solid electrolyte with the fluorite structure, was first achieved in the Ln2 Ti2 O 7 (Ln = Dy–Yb) titanates acceptor-doped with calcium and magnesium on the Ln site: (Ln1 − xMx)2Ti2O7 − δ (Ln = Dy–Yb; M = Ca, Mg; x = 0.1) [8–11]. Among these solid solutions, the highest ionic conductivity (4 × 10− 2 S/cm at 800 °C) is offered by (Yb 0.9 Ca 0.1 ) 2 Ti 2 O 6.9 [8,9]. Deng et al. [12] obtained high oxygen ion conductivity at a lower calcium content, in (Yb1.96Ca0.04) Ti2O7 − δ, which correlates with the results reported in [8–11]. Under reducing conditions, this solid solution is chemically stable and also has high ionic conductivity (~1 × 10−2 S/cm at 750 °C) [12].

⁎ Corresponding author. Tel.: +7 499 1378303; fax: +7 499 2420253. E-mail addresses: [email protected], [email protected] (A.V. Shlyakhtina).

The formation of oxygen interstitial compounds based on the pyrochlore structure is more difficult. For example, Thompson et al. [13] have reported that oxygen interstitial species are introduced in Ce2Zr2O7. Chemical intercalation of oxygen into Ce2Zr2O7 was carried out in the hard conditions (Ar atmosphere, solution of NaOBr in H2O). Donor doping (with cations having a higher valence than that of titanium) on the Ti site of the R2Ti2O7 (R = Gd, Y) pyrochlores considerably increases their electronic conductivity under reducing conditions and turns them to the promising SOFC anode materials [14–16]. In this context, the Gd2(MxTi1 − x)2O7 (M = Mo, Mn, Ru) solid-solution systems [14–16] and Yb0.96Ca1.4TiNbO6.98 [12] have been investigated in sufficient detail. O. Sakai et al. were the first to synthesize Yb2Nb2O7 pyrochlore [17] and to investigate its magnetic susceptibility. Donor doping with niobium on the Zr site in Gd2Zr2O7 was studied by Xia et al. [18] for the Gd2(Zr1 − xNbx)2O7 + δ (x = 0, 0.1, 0.2) solid solutions. The bulk conductivity of the zirconate solid solutions under reducing conditions is only slightly higher than that in air, in contrast to the titanate pyrochlores. In view of this, and given the results reported in [12,14–16], the study of donor substitutions in the titanate pyrochlores is of most interest. In particular, doping of (Yb1.9Ca0.1)Ti2O6.9 with molybdenum and magnesium ((Yb1.9Ca0.1)(TiMo0.5 Mg0.5)O7 − δ) led to a marked increase in electronic conductivity, from 7 × 10−4 S/cm in air

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Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

x

OO ↔Oi ″ þ VO

::

::

x

OO ↔1=2O2 þ VO þ 2e′

a

0 Yb2[Ti1-xNbx]2O7

-2 -3 -4 -5

ð1Þ

-6

ð2Þ

-7

With these equilibria, we expect an ionic domain at high oxygen partial pressures, where the intrinsic oxygen vacancies formed in Eq. (1) are main charge carriers. At very low oxygen partial pressures, due to equilibrium (2) n-type conduction is expected, and the material becomes an electronic conductor. Niobium (Nb5 +) was expected to act as a donor dopant, which would increase the electronic conductivity of the material, which can be interesting for the application as material for SOFC anodes. On the other hand, Ca additions can increase the oxygen ion conductivity, creating oxygen vacancies to compensate the formation of negatively charged defects. Tb additions cannot affect the defect equilibrium, but under oxidizing conditions the oxidation state Tb4+ can exist, with suppression of the p-type conductivity. The expected compensation mechanisms for dopant additions can be summarized by the following equations (using Kroger–Vink notation):

VO.. x=0 x = 0.01 x = 0.04 -20

•

x

x

-15

n

-10

-5

0

5

log(pO2/Pa)

b

0 (Yb0.8Tb0.1Ca0.1)2[Ti1-xNbx]2O7±δ -1

Oi''

-2 -3 -4 -5 VO.. x=0 x = 0.05 x = 0.1

-6

Yb2 O3 þ Nb2 O5 →2YbYb þ 2NbTi þ 7OO þ Oi ″;

Oi''

-1

log([ ])

to 0.3 S/cm, after reduction in 5% H2/N2 at 850 °C for 24 h [12]. At the same time, when doped with a considerable amount of variable valence ions, most donor-substituted materials based on rare-earth titanates show not only a marked increase in electronic conductivity under reducing conditions but also a drop in ionic conductivity. Recent work [19] has shown that, even when Ln2Ti2O7 (Ln = Dy, Yb) materials are doped with variable valence cations, the resultant (Ln0.8Ca0.1 Tb′0.1)2Ti2O7 − δ (Ln = Dy, Yb) solid solutions are highly stable under reducing conditions, but their electronic conductivity is not high. The main defect equilibria in this type of undoped pyrochlore involve anti-Frenkel defects (Eq. (1)) and defects promoted by the redox reaction represented by Eq. (2) [14]:

log([ ])

2

ð3Þ

n

-7 ••

x

2CaO þ 2TiO2 →2CaYb þ 2TiTi þ 6OO þ VO ; x

x

x

Tb2 O3 þ 2TiO2 →2TbYb þ 2TiTi þ 7OO ; ●

x

TbYb þ e′ →TbYb :

-20

ð4Þ

-15

-10

-5

0

5

log(pO2/Pa) ð5Þ ð6Þ

Using thermodynamic data found in the literature [20,21] and converting to mole fractions, we find that, at 1000 °C, Kred and KaF are approximately 10−13 Pa1/2 and 10−6, respectively. Combining the equilibria described by Eqs. (1) and (2) with this thermodynamic data and using the DefChem software package [22], we can calculate a hypothetical defect diagram of this material. Fig. 1 illustrates the effect of Nb doping on the Yb2Ti2O7-based material and this material modified with Tb and Ca additions. Niobium and terbium are aliovalent ions: they can assume, respectively, the 5 + or 4 + and 4 + or 3 + states, but if their charge is the same as that of the host ion in the pyrochlore structure, no charge compensation occurs, and the defect diagram remains unchanged. So, to predict eventual changes due to charge compensation, the defect diagrams (Fig. 1) were calculated under the assumption that all of the Nb, Tb and Ca are, respectively, in the 5+, 4+ and 2+ states. In Fig. 1a and b, it is evident that the electron concentration increases, which will result in an n-type conductor throughout the oxygen partial pressure range, at least for compositions with higher Nb content, since the electron mobility can be several orders of magnitude higher than the oxygen vacancy mobility. Tb and Ca additions (Fig. 1b) have no effect on the defect diagram because we assume, for now, that all of the Tb is in the 4 + state, which is compensated by Ca additions, which have the same concentration as Tb. However, if we predict Tb reduction to the 3+ state (assuming a constant equilibrium value of 104 in order to get the reduction within the pO2 range), the Ca will be uncompensated and the Ca2 + will act as acceptor in the pyrochlore structure, with

Fig. 1. Expected defect diagram for (a) Yb2[Ti1 − xNbx]2O7 ± δ (x = 0, 0.01, 0.04 and 0.1) and (b) (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O7 ± δ (x = 0, 0.05 and 0.1). Arrows represents the increase of Nb content and the effect on each defect concentration.

evident changes in the defect diagram (Fig. 2). Due to the differences in carrier mobilities, we will still expect to obtain an n-type conductor throughout the oxygen partial pressure range. If Nb assumes the 0 (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O7±δ

Tb4+

-1

Oi''

-2

log([ ])

x

′

-3

-4

VO..

-5 n -6 -20

-15

-10

-5

0

5

log(pO2/Pa) Fig. 2. Expected defect diagram for (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O7 ± δ assuming the coexistence of Tb 3+ and Tb 4+.

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

4 + state, the behavior in both systems will tend to be similar to that of the undoped material, since there is no charge compensation for Ti4 + substitution. In this paper, we report the synthesis and oxygen ion conductivity of Yb2Ti2O7-based solid solutions with a low (≤ 10%) degree of Nb substitution on the Ti site: Yb2(Ti1 − xNbx)2O7 (x = 0.01, 0.04, 0.1). In addition, to obtain higher oxygen ion conductivity we examined the (Yb0.8Tb0.1Ca0.1)2(Ti1 − xNbx)2O6.9 (x = 0, 0.05, 0.1) solid solutions, with substitutions on the lanthanide site. The materials that had high oxygen ion conductivity in air, Yb 2 (Ti0.99 Nb 0.01 ) 2 O 7 and (Yb0.8 Tb0.1 Ca0.1)2[Ti0.95 Nb0.05]2 O6.9, were also studied in reducing atmospheres:5% H2 in N2 and CO2.

2. Experimental Yb 2 [Ti 1 − x Nb x ] 2 O 7 (x = 0.01, 0.04, 0.1) и (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti 1 − x Nb x ] 2 O 6.9 (x = 0; 0.05; 0.1) ceramics were prepared using coprecipitation. The starting chemicals used were Yb 2 O 3 , Tb 4 O 7 and NbCl 5 powders, which were dissolved in HCl, and a TiCl4 solution in HCl. The niobium chloride solution was unstable and was stabilized by C 2 H 5 OH. The H 2 O:C 2 H 5 OH ratio in the water–alcohol solution of NbCl 5 was 3:1. The concentrations of the rare-earth chloride solutions were 0.178 M (Yb) and 0.124 M (Tb), and those of the TiCl 4 and NbCl5 solutions were 1.44 M and 0.194 M. The solution of metal chlorides contained 1.9 M HCl. The Yb 2 [Ti 1 − x Nb x ] 2 O 7 (x = 0.01, 0.04, 0.1) solid solutions were synthesized via reverse precipitation (pH 10.8–11), by adding titrated ytterbium chloride and titanium chloride solutions and an ethanolic NbCl5 solution to aqueous ammonia. The precipitant used was aqueous ammonia with NH 4 OH:H 2 O = 1:1. The (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) solid solutions were prepared through precipitation of ytterbium, terbium, titanium, and niobium hydroxides with ammonia (pH 10.8–11). Calcium was precipitated separately, from a 0.34 M Ca(NO3)2 solution, using ammonium oxalate as a precipitant. After the precipitation, the mother liquor was decanted and the suspensions were poured together, thoroughly mixed, and centrifuged. The resultant precipitate was washed with warm water four times to remove Cl− and HOOCCOO−, dried in air at 140 °C for 24 h, and calcined at 650 °C for 2 h. The resultant powders were reground and pressed at 249–332 MPa into pellets, which were then fired at 1550 °C for 4 h. As shown earlier [23], the synthesis temperature can be raised when a small amount of Tb is added. In this way, homogeneous, dense samples can be obtained at higher temperatures, with no grain-boundary segregation of the calcium-containing phase. The density of our samples was determined by measuring their mass and dimensions. The densities of the samples thus prepared are listed in Table 1. X-ray diffraction (XRD) patterns of polycrystalline samples were collected at room temperature on a DRON-3 M automatic diffractometer (Cu Kα radiation, λ = 1.5418 Å, Bragg reflection geometry, 35 kV, 28 mA) in the 2θ range 13° to 65° (scan step of 0.02° or 0.1°). Phase compositions were determined using PCPDFWIN v. 2.3. To determine lattice parameters and analyze the samples for phase purity, WinXPow software and the PDF database were used.

3

The valence of the niobium in the solid solutions was determined by X-ray photoelectron spectroscopy (XPS) on a Kratos Axis Ultra DLD spectrometer using a monochromatic Al Kα X-ray source and 150-W electron gun. The pass energy was 160 eV for survey scans and 40 eV for individual lines. The analysis area was about 300 × 700 μm. The binding energy scale was preliminary calibrated using the following lines of standards (cleaned by ion sputtering): Au 4f5/2 at 83.96 eV, Cu 2p3/2 at 932.62 eV, and Ag 3d5/2 at 368.21 eV. A charge neutralizer was used to minimize the surface charging effect. As the binding-energyscale reference, we used the C 1s level (285.0 eV) arising from the carbon adsorbed on the sample surface. The microstructure of the sintered ceramics was examined using scanning electron microscopy (JEOL JSM-6390LA). Before SEM observation the ceramic samples were manually ground, polished with diamond paste, and finally thermally etched at 1400 °C for 0.5 h. For electrical measurements, disk-shaped polycrystalline samples (diameter 6–9 mm and thickness 1–3 mm) were prepared. Contacts to the sample faces were made by firing ChemPur C3605 paste, containing colloidal platinum, at 950–1000 °C. The measurements were made by a two-probe method in air in the frequency range 10 mHz to 3 MHz at 13 fixed temperatures from 300 to 1000 °C and an applied sinusoidal voltage of 0.5 V peak, using a Novocontrol Beta-N impedance analyzer and a NorECs ProboStat ceramic cell fitted with platinum electrodes and a Pt/Pt–Rh thermocouple. In data processing, we used ZView software [24]. Total conductivity versus oxygen partial pressure data were obtained between 800 and 1000 °C, by the reoxidation method, using a mixture of 95% N2 and 5% H2 to reduce the oxygen partial pressure (measured by a ZrO2-based sensor, closed to the sample) to 10−17–10−13 Pa. Subsequently, the gas flow was shut off and the impedance spectra were measured at 10-min intervals until the oxygen partial pressure increased to 105 Pa, which took more than 6 h. In this temperature range, the spectra show only one arc, which is related to the electrode reaction process, so the intersection of each spectrum with the real axis of the Nyquist plot was used to evaluate the total conductivity. The ac measurements in a reducing atmosphere were made by impedance spectroscopy in the frequency range 20 Hz to 1 MHz and at temperatures from 300 to 750 °C using a Hewlett-Packard 4284A precision LCR bridge.

3. Results and discussion 3.1. XPS determination of the Nb valence state The valence of the niobium (4 + or 5 +) in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) solid solutions synthesized through coprecipitation can be accurately determined by XPS. Xia et al. [18] suggested, without presenting experimental evidence, that Gd2(Zr1 − xNbx)2O7 + δ (x = 0, 0.1, 0.2) solid solutions prepared by solid-state reactions in air contained pentavalent niobium. In this study, the solid solutions with the lowest and highest niobium contents in the two systems were characterized by XPS. Fig. 3a, b shows the Nb 3d and O 1s XPS spectra of the Yb2 [Ti1 − xNbx]2O7 (x = 0.01, 0.1) and (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0.05, 0.1) solid solutions, and Table 2 lists the Nb 3d5/2 and O 1s

Table 1 Characteristics of the solid solutions. Sample no

Composition

Heat treatment

Color

Average grain size, μm

Relative density, %

а, Å

1 2 3 4 5 6

Yb2[Ti0.99Nb0.01]2O7 Yb2[Ti0.96Nb0.04]2O7 Yb2[Ti0.9Nb0.1]2O7 (Yb0.8Tb0.1Ca0.1)2Ti2O6.9 (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9

1550 1550 1550 1550 1550 1550

Beige Light yellow Light yellow Beige Beige Beige

5.53(±0.17) 4.01(±0.09) 6.52(±0.03) 11.99(±0.09) 3.39(±0.05) 2.15(±0.03)

89.5 87.9 87.0 90.2 88.8 89.4

10.032(2) 10.039(2) 10.053(2) 10.026(2) 10.052(2) 10.070(2)

°С–4 h °С–4 h °С–4 h °С–4 h °С–4 h °С–4 h

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

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L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

Fig. 3. (a) Nb 3d and (b) O 1s XPS spectra of (1) Yb2(Ti0.99Nb0.01)2O7, (2) Yb2(Ti0.9Nb0.1)2O7, (3) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 and (4) (Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9.

binding energies. The chemical shift ΔEb of XPS lines is known to correlate with the effective charge on the atoms involved in bonding. Accordingly, the oxidation state of niobium is a quasi-linear function of the energy position of the Nb 3d level. This can be used to assess the chemical state of niobium in its compounds with oxygen. Kuznetsov [25] analyzed the Nb 3d and O 1s spectra of Nb metal, NbO, NbO2, and Nb2O5

and presented the Nb 3d5/2 binding energies in these materials: 202.2 eV in Nb metal, 204.4 eV in NbO, 206.4 eV in NbO2, and 207.9 eV in Nb2O5. In this study, the Nb 3d5/2 binding energy in Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.1) and (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0.05, 0.1) ranges from 206.5 to 206.8 eV (Table 2), which corresponds to the valence state Nb4+. The O 1s binding energy in NbO2 is 530.7 eV [26]. The

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx Table 2 Oxidation state of the niobium in the solid solutions. Sample no

1 3 5 6

Composition

Yb2[Ti0.99Nb0.01]2O7 Yb2[Ti0.9Nb0.1]2O7 (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9

Binding energy Eb, Ev Nb 3d5/2

O 1s

206.5 206.8 206.5 206.7

530.2 530.7 530.4 530.6

multicomponent oxides under consideration range in O 1s binding energy from 530.2 to 530.7 eV (Table 2). Thus, the present XPS data demonstrate that, when niobium is incorporated into the solid solutions through coprecipitation, the single-phase materials contain only tetravalent niobium, and their formulas should have the form Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb0.1Ca0.1)2 [Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1). Atuchin et al. [27] assume that Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE varies much less than the Nb 3d5/2 BE. In connection with this, they propose that various niobium compounds can be compared using Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE in order to find out whether they contain Nb5+, Nb4+, or a mixture of these. The Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE values obtained in this study for compounds prepared by the same procedure are listed in Table 3. Since the peak position for the first composition, containing 1% Nb, was found with a rather large uncertainty, this value can be omitted, and the average of Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE is the same, confirming Atuchin et al.'s assumption. According to Atuchin et al. [27], the maximum value of Δ(O–Nb5+) is ~323.5 eV. In this study, this parameter is larger (323.9–324 eV). Therefore, the synthesized solid solutions most likely contain Nb4+ with trace levels of Nb5+. We analyzed reports that had mentioned the presence of Nb4+ in NbxTi1 − xO2 powders synthesized from solution-derived precursors and then annealed in air [28,29]. In both studies, Nb-doped TiO2 powders were prepared by a liquid-phase method (sol–gel process). XPS data in the former report demonstrate that low-temperature (450 and 650 °C) decomposition of organic precursors leads to the formation of NbxTi1 − xO2 containing only Nb4+, whereas higher temperature (1050 °C) annealing yields NbxTi1 − xO2 containing Nb4+ with Nb5+ impurities. In the latter study, where Nb-doped TiO2 was also synthesized through low-temperature decomposition of an organic precursor (sol–gel process, annealing at 700 °C), EPR data showed that only Nb4+ was present. Thus, when low-temperature precursors containing titanium, niobium, and oxygen are decomposed at T ≤ 700 °C, as in our study, Nb4+ stabilization is possible. Coprecipitation ensures almost complete reaction between Ln2O3 and TiO2 even at room temperature. For example, in the case of Nd2Ti2O7 the coprecipitation product was a hydroxo complex of composition Nd 2[TiO 2(OH) 2] 3 × nH2 O [30]. The hydroxo complex is dried at 100 °C for 24 h and then rare-earth titanates are synthesized from the coprecipitated precursors in two steps in our work. Firing at 650 °C for 2 h ensures water removal and further synthesis of the pyrochlore phase (which reaches completion at ~ 850 °C according to previous data [30,31]). In the second step, the precursor decomposition product is reground, pressed, and annealed at

5

1550 °C for 4 h. Thus, it is quite possible that most of the niobium in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb0.1Ca0.1)2 [Ti 1 − x Nb x ] 2 O 6.9 (x = 0, 0.05, 0.1) ceramics is in the tetravalent state. Thus, the donor effect of Nb5 +predicted by Eq. (3), which would increase the electronic conductivity of the solid solutions, is missing. This can be understood in terms of the specifics of the lowtemperature synthesis of the rare-earth titanates from solution and the ability of the pyrochlore structure to stabilize B cations in a tetravalent state as the synthesis temperature is raised. XPS data for the Yb 2[Ti1 − xNbx] 2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb 0.1Ca 0.1)2 [Ti 1 − x Nb x ] 2 O 6.9 (x = 0, 0.05, 0.1) solid-solution series indicate isovalent substitution of niobium on the titanium site, with no charge compensation. Subsequent electrochemical measurements in a wide range of oxygen partial pressures (Section 3.4) confirmed this result for Yb2(Ti0.99Nb0.01)2O7 and (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9. 3.2. Effect of Nb substitution for Ti on the conductivity of Yb2Ti2O7 Fig. 4 shows the XRD patterns of the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) materials. The three samples are seen to have the pyrochlore structure, with the typical superlattice reflections (relative to the fluorite subcell) 111, 311, 331, 511, and 531. The unit-cell parameters of the solid solutions are listed in Table 1. It can be seen that, with increasing niobium content, the cubic cell parameter a of the Yb2 [Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) pyrochlores increases, which correlates with the six-coordinate ionic radii of Ti 4 + and Nb 4 + : 0.605 and 0.68 Å, respectively. Indeed, isovalent doping with niobium, which has a greater ionic radius, should lead to an increase in unit-cell parameter in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) series. The relative densities of the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) ceramics are listed in Table 1. Fig. 5 illustrates the microstructure of thermally etched Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) surfaces. The average grain size was found to be 5.53 μm at x = 0.01, 4.01 μm at x = 0.04, and 6.52 μm at x = 0.1 (Table 1). As seen in Fig. 5, the percentage of small grains is rather large at x = 0.01 and 0.04, whereas at x = 0.1 there are very few small grains. Fig. 6a, b, shows the impedance spectra of Yb2[Ti0.99Nb0.01]2O7 (x = 0.01) in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) solid-solution series at 400 and 500 °C in air. The spectra consist of semicircles that represent the bulk and grain-boundary contributions at high and medium frequencies, respectively, and the electrode polarization region at low frequencies. Classic impedance spectra, with all three semicircles resolved, typical of oxygen ion conductors, were only obtained for Yb2(Ti0.99Nb0.01)2O7 (Fig. 6a, b). In those spectra, we were able to identify the bulk and grain-boundary contributions (Table 4) [32]. Fig. 7 shows Arrhenius plots of bulk conductivity for Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1). Bulk conductivity exhibits Arrhenius behavior only at x = 0.01 and 0.04, whereas at the higher niobium concentration there is a distinct break at ~ 750 °C. Nb-vacancy clusters have significant binding energies, suggesting greater vacancy trapping. This is consistent with the observed increase in activation energy for ion migration at higher Nb doping level in Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1), and would lead to two distinct regions or non-linear portions in the conductivity Arrhenius plot. At high Nb content we observed a change in the slope of the Arrhenius plot (Fig. 7, curve 3),

Table 3 Analysis of Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE in penta- and tetravalent niobium compounds (NIST, literature data [24,26]) and comparison with the Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE values obtained in this study. Δ(O–Nb) = O 1s BE–Nb 3d5/2 BE, eV (this work)

Δ(O–Nb5+) = O 1s BE–Nb 3d5/2 BE, eV

Δ(O–Nb4+) = O 1s BE–Nb 3d5/2 BE, eV

(1) 323.7 (surface) (2) 323.9 (surface) (3) 323.9 (surface) (4) 323.9 (surface) (4) 324 (grinding) Without (1) because of an error, bEΔ (O–Nb)this work N= 323.9

Atuchin et al. [26] 322.8–323.5 Kuznetsov M.V. [24] 323.1 bEΔ (O–Nb5+) N= 323.13

Atuchin et al. [26] 324.3–325.03 Kuznetsov M.V. [24] 325.0 bEΔ (O–Nb4+) N= 324.78

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

6

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

Intensity (a.u.)

222

440 622

400 111

311

331 511

444

531 533

10

20

30

40

2θ,

50

60

1 2 3 70

o

Fig. 4. XRD patterns of Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1): (1) Yb2(Ti0.99Nb0.01)2O7, (2) Yb2(Ti0.96Nb0.04)2O7, (3) Yb2(Ti0.9Nb0.1)2O7.

which suggests a change of the activation energy for conduction and consequently a change in conductivity mechanism. Note that the activation energy for electronic defect formation is usually much higher than that for ionic defects, so we can have a transition mechanism between low and high temperatures. The activation energies for bulk conduction in the solid solutions are presented in Table 6. The activation energy is seen to increase with niobium concentration: from Ea = 0.91 eV in Yb2(Ti0.99Nb0.01)2O7 (x = 0.01) to 1.16 eV in Yb2(Ti0.96Nb0.04)2O7 (x = 0.04) and to 2.18 eV (T N 700 °C) in Yb2[Ti0.9Nb0.1]2O7 (x = 0.1). Table 5 presents temperature-dependent bulk conductivity data for Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) in air. The ionic conductivity of Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) gradually decreases with increasing x and is lower than that of the undoped Yb2Ti2O7 [33]. We suppose that Nb4+–O2− may produce a detrimental effect on the oxide ion transport behavior of the pyrochlore structure, because the Nb4+–O2− bond is stronger than the Ti4+–O2 − bond, producing trapping effect [34] (See Table 6.). 3.3. Effect of Nb substitution for Ti on the conductivity of Yb2Ti2O7 doped with terbium and calcium Fig. 8 shows the XRD patterns of (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1). The three samples are seen to have the pyrochlore structure, like in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) series. The XRD patterns contain the typical superlattice reflections 111, 311, 331, 511, and 531. An additional line at 2θ = 26.6° (in rare-earth titanates, zirconates, and hafnates) is typical of high-conductivity pyrochlores. In earlier studies [35,36], it was tentatively identified as 300. The observed splitting of the 444 line (Fig. 8, scan 1) is due to the α1–α2 doublet (the calculated and observed dα1 values are 1.4477 Å; the calculated and observed dα2 values are 1.45107 and 1.4510 Å, respectively). The unit-cell parameters of the solid solutions are listed in Table 1. Like in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) series, the lattice parameter increases with niobium concentration, which is due to the larger octahedral ionic radius of niobium in comparison with titanium. The relative densities of the (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) ceramics are given in Table 1. Fig. 9 illustrates the microstructure of thermally etched (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) surfaces. The average grain size was found to markedly decrease with increasing niobium content: r = 11.99 μm at x = 0, r = 3.39 μm at x = 0.05, and r = 2.15 μm at x = 0.1 (Table 1). As seen in Fig. 9b, inset the (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 sample contains faceted crystallites of a Ca compound at grain boundaries. No such compound was detected by XRD. The shape of the impedance spectra of the (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) solid solutions is typical of traditional solid electrolytes (Fig. 6c, d). In the spectra of the three

Fig. 5. SEM images of Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1): (a) Yb2(Ti0.99Nb0.01)2O7, (b) Yb2(Ti0.96Nb0.04)2O7, (c) Yb2(Ti0.9Nb0.1)2O7.

samples, the bulk and grain-boundary contributions and the electrode polarization region can be identified. Fig. 6c, d shows impedance spectra of (Yb0.8 Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 at 400 and 500 °C, respectively. The bulk and grain-boundary conductivities of all the samples studied can be evaluated using the EQUIVCRT program [32] (Table 4). Fig. 10 shows Arrhenius plots of bulk and grain-boundary (Fig. 10, inset) conductivity for (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1).

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

7

a -Z" (kΩ•cm)

50

Data (400oC) Fit

25

0 0

50

100

150

200

250

Z' (kΩ•cm)

-Z" (kΩ•cm)

b

10

5

Data (400oC) Fit

0 0

10

20

30

40

Z' (kΩ•cm)

-Z" (kΩ•cm)

c

10

5 Data (400oC) Fit

0 0

10

20

30

40

50

Z' (kΩ•cm)

-Z" (kΩ•cm)

d

1.0

Data (400oC) Fit

0.5

0.0 0

1

2

3

4

5

6

Z' (kΩ•cm) Fig. 6. Impedance spectra of (a, b) Yb2(Ti0.99Nb0.01)2O7 and (c, d) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 at (a, c) 400 and (b, d) 500 °C.

The activation energies for bulk and grain-boundary conduction in the solid solutions are listed in Table 5. The activation energy slightly decreases on doping to x = 0.05, from Ea = 0.87 eV in (Yb0.8Tb0.1Ca0.1)2Ti2O6.9 to 0.86 eV in (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9, and increases to Ea = 0.95 eV in (Yb0.8 Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9. Table 6 presents

temperature-dependent bulk conductivity data for (Yb0.8Tb0.1Ca0.1)2 [Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) in air. The highest 650 °C conductivity among the (Yb0.8Tb0.1Ca 0.1) 2 [Ti1 − xNb x] 2 O 6.9 (x = 0; 0.05; 0.1) solid solutions is offered by (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (x = 0.05). So additional oxygen

Table 4 Parameters of equivalent electrical circuit, used for fitting.

R1 (Ohm) T1 (F) P1 R2 (Ohm) T2 (F) P2 R3 (Ohm) T3 (F) P3

Yb2(Ti0.99Nb0.01)2O7 (d = 7.27 mm, h = 2.00 mm) 400 °C 29,300 ± 100 (2.42 ± 0.09) × 10−11 0.953 ± 0.002 28,900 ± 700 (6.13 ± 0.54) × 10−8 0.832 ± 0.012 76,400 ± 3700 (1.94 ± 0.05) × 10−5 0.449 ± 0.018



500 °C 4330 ± 35 (3.15 ± 0.69) × 10−11 0.941 ± 0.013 2440 ± 90 (1.07 ± 0.32) × 10−7 0.820 ± 0.030 18,500 ± 860 (8.60 ± 0.23) × 10−5 0.463 ± 0.014

Z CPE ¼

1 T ðiωÞP



(Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (d = 7.77 mm, h = 1.96 mm) 400 °C 3790 ± 24 (8.50 ± 1.20) × 10−11 0.908 ± 0.009 5320 ± 190 (2.49 ± 0.34) × 10−7 0.752 ± 0.016 15,700 ± 870 (7.09 ± 0.24×)10−5 0.425 ± 0.021

1 — bulk contribution; 2 — grain-boundary contribution; 3 — electrode–sample interface.

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

500 °C 665 ± 3.0 (7.06 ± 0.40) × 10−10 0.792 ± 0.003 384 ± 16 (8.30 ± 2.1) × 10−8 0.894 ± 0.022 1370 ± 39 (2.21 ± 0.09) × 10−4 0.379 ± 0.014

8

L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

1

Table 6 Temperature-dependent bulk conductivity of the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8 Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1) solid solutions in air. Composition

-1

Bulk conductivity (S/cm) 800 °С

-2

Yb2[Ti0.99Nb0.01]2O7 Yb2[Ti0.96Nb0.04]2O7 Yb2[Ti0.9Nb0.1]2O7 (Yb0.8Tb0.1Ca0.1)2Ti2O6.9 (Yb0.8Tb0.1Ca0.1)2 [Ti0.95Nb0.05]2O6.9 (Yb0.8Tb0.1Ca0.1)2 [Ti0.9Nb0.1]2O6.9

-3 -4 -5

1 1(bulk+g.b.) 2 3

-6 0.7

0.9

1.1

1.3

1000/T,

1.5

1.7

Fig. 7. Arrhenius plots of bulk conductivity for Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1): (1) Yb2(Ti0.99Nb0.01)2O7, (2) Yb2(Ti0.96Nb0.04)2O7, (3) Yb2(Ti0.9Nb0.1)2O7.

vacancies created by Ca doping in the pyrochlore structure reduce the detrimental effect of Nb doping on the oxide ion transport up to 5% Nb. The higher Nb concentration, like in (Yb 0.8 Tb0.1Ca 0.1 ) 2 [Ti0.9Nb0.1]2O6.9, decreases oxygen ion conductivity in pyrochlores in accordance with [34]. Fig. 10, inset shows Arrhenius plots of grain-boundary conductivity for the (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) ceramics. The Nb-free sample, (Yb0.8Tb0.1Ca0.1)2Ti2O6.9, and the sample with 5% of the titanium substituted with niobium, (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti0.95Nb0.05]2O6.9, are similar in conductivity but differ markedly in grain size: 11.99 and 3.39 μm, respectively. Thus, in this system, for low Nb contents, the grain size has an insignificant effect on the grain-boundary conductivity. But when Nb content attains 10% ((Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9) the grain-boundary conductivity component clearly decreases. It should be emphasized that both the bulk and grain-boundary conductivities of these materials exhibit Arrhenius behavior. 3.4. Conductivity of Yb 2 (Ti 0.99 Nb 0.01 ) 2 O 7 and (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti 0.95Nb0.05]2O6.9 as a function of oxygen partial pressure and in CO2 atmosphere The above XPS data on the valence of the niobium (4+) in the solid solutions demonstrate that these materials are ordinary anion-deficient oxygen ion conductors with isovalent niobium substitution on the titanium site. Therefore, the conduction mechanism in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) system is similar to that in undoped. Yb2Ti2O7, and the conduction mechanism in the (Yb0.8Tb0.1Ca0.1)2 [Ti1 − xNbx]2O7 (x = 0; 0.05; 0.1) system is similar to that in (Yb 0.9 Ca 0.1 ) 2 Ti 2 O 6.9 [37]. As shown by Savvin et al. [19] for (Yb0.8Tb 0.1 Ca 0.1 ) 2Ti2 O 7 − δ , terbium experiences grain-boundary segregation, as evidenced by the fact that it is the grainTable 5 Activation energy for bulk and grain boundary conduction in the solid solutions. Sample Composition no

1 2 3

Yb2[Ti0.99Nb0.01]2O7 Yb2[Ti0.96Nb0.04]2O7 Yb2[Ti0.9Nb0.1]2O7

4 5 6

(Yb0.8Tb0.1Ca0.1)2Ti2O6.9 (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 (Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9

Activation energy for bulk ionic conduction, Ea (eV)

Activation energy for grain-boundary conduction, Ea (eV)

0.91 1.16 2.18 (T N 700 °C) 0.99 (T b 700 °C) 0.87 0.86 0.94

1.18

Activation energy for bulk ionic conduction excluding sample 3. a We failed to identify the grain-boundary contribution.

750 °С

650 °С

550 °С

−3

3.24 × 10 2.33 × 10−4 1.12 × 10−5 1.26 × 10−2 1.38 × 10−2

−3

2.30 × 10 1.31 × 10−4 3.88 × 10−6 9.43 × 10−3 9.03 × 10−3

−4

8.93 × 10 3.46 × 10−5 5.17 × 10−7 3.74 × 10−3 4.44 × 10−3

2.37 × 10−4 6.44 × 10−6 9.79 × 10−8 1.12 × 10−3 1.33 × 10−3

6.50 × 10−4

3.80 × 10−4

1.70 × 10−4

5.56 × 10−5

1.9

K-1

boundary contribution that decreases by an order of magnitude under reducing conditions, whereas the bulk contribution remains essentially unchanged. This behavior originates from the considerable difference in ionic radius between terbium and ytterbium. Tb3 + + h• → Tb4 + process for oxidizing conditions and Tb4 + + e′ → Tb3 + process for reducing conditions is only valid for the grain-boundary conductivity of (Yb0.8Tb0.1Ca0.1)2Ti2O6.9. The grain bulk is rare-earthdeficient, and this reduces the bulk conductivity in comparison with (Yb0.9Ca0.1)2Ti2O6.9 [23]. The comparative study of the best high-temperature conductors (Yb2(Ti0.99Nb0.01)2O7 and (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9) in the two Nb-doped series was performed at different partial pressures. The results are presented in Fig. 11a, b. One can see a plateau with predominant ionic conductivity and a significant increase in n-type conductivity under reducing conditions, typical of Yb2Ti2O7 based materials [33]. The plateau for pure Yb2Ti2O7 [32] and Yb2(Ti0.99Nb0.01)2O7 is the same (up to ~10−10 Pa), whereas that for (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 extends up to 10−12 Pa. The weak rise in the n-type conductivity of the last sample under reducing conditions appears unusual. The complex doping of the Yb sublattice with Ca and Tb may be the reason for this strange behavior, but it is consistent with the p-type conductivity at high oxygen partial pressures. Tb as an aliovalent element can assume the 3+ and 4+ valence states. To compensate Ca additions, Tb should remain in the 4 + state, but at low partial pressures Tb is reduced to the 3 + state. Such complex equilibria can also be responsible for the slight increase in conductivity at high oxygen partial pressures, probably due to an increase in p-type conductivity. The observed behavior of the conductivity as a function of oxygen partial pressure is in conflict with the defect diagrams presented in the Introduction section. The explanation for this behavior can be the unexpected oxidation state of Nb (Nb4 +), however this confirms the result obtained by XPS. In this case, Nb 4 + and Ti4 + are isovalent, the Nb additions do not need to be charge compensated,

222

Intensity (a.u.)

log σ·T, S·K·cm-1

0

440 400

111 311

622

331 511

531 533

4441

2

a

3

a

10 1.20 1.13 1.37

20

30

40

50

60

70

2θ, o Fig. 8. XRD patterns of (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti 1 − x Nb x ] 2 O 6.9 (x = 0, 0.05, 0.1): (1) (Yb0.8Tb0.1Ca0.1)2Ti2O6.9; (2) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9; (3) (Yb0.8Tb0.1Ca0.1)2 [Ti0.9Nb0.1]2O6.9.

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L.G. Shcherbakova et al. / Solid State Ionics xxx (2014) xxx–xxx

9 2

log σ·T, S·K·cm-1

1 0 -1 -2 -3

1

-4

2

-5 -6 0.7

3

0.9

1.1

1.3

1.5

1.7

1.9

1000/T, K-1

2

log σ·T, S·K·cm-1

1 0 -1 -2 -3

1(bulk)

1(b.+g.b.)

2(bulk)

2(b.+g.b.)

3(bulk)

3(b.+g.b.)

-4 -5 0.7

0.9

1.1

1.3

1000/T,

1.5

1.7

1.9

K-1

Fig. 10. Arrhenius plots of bulk and grain-boundary conductivity (inset) for (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0, 0.05, 0.1): (1) (Yb0.8Tb0.1Ca0.1)2Ti2O6.9; (2) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9; (3) (Yb0.8Tb0.1Ca0.1)2[Ti0.9Nb0.1]2O6.9.

Fig. 9. SEM images of (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti 1 − x Nb x ] 2 O 6.9 (x = 0, 0.05, 0.1): (a) (Yb 0.8 Tb 0.1 Ca 0.1 )2 Ti2 O 6.9 ; (b) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9; (c) (Yb0.8Tb0.1Ca0.1)2 [Ti0.9Nb0.1]2O6.9. The Fig. 9, b, inset presents (above) SEM image and (below) Ca X-ray map of (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9.

and the defect diagram of Nb-doped material is the same as that of Yb2 Ti2 O 7 (Fig. 1a, x = 0). For (Yb0.8Tb 0.1 Ca 0.1 ) 2[Ti 0.95 Nb 0.05 ] 2 O 6.9 , based on the defect chemistry predictions in Fig. 1b, the result suggests that Tb and Nb are in the 4 + state. So, Tb4 + is compensated by Ca2 + in Yb3 + position and Nb4 + does not need to be compensated at Ti4 + position, and the conductivity as a function of oxygen

Fig. 11. Total conductivity as a function of oxygen partial pressure at different temperatures for (a) Yb2(Ti0.99Nb0.01)2O7 and (b) (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9.

Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019

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partial pressure exhibits the same behavior as the undoped composition with x = 0 (Fig. 1b). From the conductivity behavior as a function of oxygen partial pressure, assuming a ¼ power law dependence, we can evaluate the ionic and n-type components of the conductivity (Fig. 12). The activation energy for ionic (0.8 eV) and electronic (2.7 eV) components are the same for Yb2(Ti0.99Nb0.01)2O7 and (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9, which means that the transport mechanism is not influenced by complex doping on the lanthanide site in Yb2Ti2O7-based materials synthesized using coprecipitation. This difference between the activation energies for electronic and ionic conduction is in agreement with the transition mechanism observed for Yb2(Ti0.9Nb0.1)2O7 (Fig. 7) and consistent with the activation energies presented in Table 5. Due to the possibility of p-type conductivity in air for (Yb0.8Tb0.1Ca0.1)2 [Ti0.95Nb0.05]2O6.9, this material was investigated in CO2 atmosphere to ensure relatively low oxygen pressure, within the electrolyte domain of this pyrochlore, and to get true ionic conductivity. Note that the bulk conductivity of the best conductor (Yb0.8Tb0.1Ca0.1)2[Ti0.95Nb0.05]2O6.9 is not affected by the CO2 atmosphere, however the grain-boundary contribution exhibits true Arrhenius behavior (better linearity) when measured in CO2 atmosphere. The deconvolution of impedance spectra is easier for a pure ionic material than for a mixed p-type and ionic conductor.

4. Conclusions We have studied two series of Yb2Ti2O7-based pyrochlore solid solutions with a low degree of Nb substitution on the Ti site. The specifics of the coprecipitation process, which enables the preparation of the rareearth titanates at rather low temperatures, and the stability of B cations in a tetravalent state in the pyrochlore structure even at a considerable increase in synthesis temperature allowed us to obtain two series of ytterbium titanates containing niobium as an isovalent substituent. XPS data for Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.1) and (Yb0.8Tb0.1Ca0.1)2 [Ti1 − xNbx]2O6.9 (x = 0.05, 0.1) indicate isovalent substitution of niobium on the titanium site. In both solid-solution systems, the cubic pyrochlore cell parameter increases with niobium content, which is consistent with the relationship between the ionic radii of Ti4+ and Nb4+. Niobium substitution in the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) system reduces the bulk conductivity in air, whereas in the (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) system the highest 650 °C conductivity is offered by the x = 0.05 material. Additional oxygen vacancies created by Ca doping in the pyrochlore structure reduce the detrimental effect of Nb doping on the oxide ion transport up to 5% Nb. The solid solutions with high oxygen ion conductivity, Yb2[Ti1 − xNbx]2O7 with x =

1 0

Log (σ.Ω.m)

-1

Ionic 0.8 eV

-2 -3 -4

Electronic 2.7 eV

-5 -6 0.75

0.8

0.85

1000/T

0.9

0.95

1

(K-1)

Fig. 12. Arrhenius plots of ionic and electronic conductivity components for the Yb 2 (Ti 0.99 Nb 0.01 ) 2 O 7 (closed symbols) and (Yb 0.8 Tb 0.1 Ca 0.1 ) 2 [Ti 0.95 Nb 0.05 ] 2 O 6.9 (open symbols).

0.01 and (Yb0.8Tb0.1Ca0.1)2[Ti1 − xNbx]2O6.9 with x = 0.05, were investigated both in air and under reducing conditions. The bulk conductivity of the Yb2[Ti1 − xNbx]2O7 (x = 0.01, 0.04, 0.1) and (Yb0.8Tb0.1Ca0.1)2 [Ti1 − xNbx]2O6.9 (x = 0; 0.05; 0.1) solid solutions has oxygen vacancy conductivity type. The transport mechanism of charged species is not affected by complex doping of the lanthanide sublattice and Nb4 + doping in the Ti sublattice in Yb2Ti2O7-based materials. Acknowledgments This work was supported by the Russian Foundation for Basic Research (grant no. 13-03-00680) and the Chemistry and Materials Science Division of the Russian Academy of Sciences (basic research program n 2: Advanced Metallic, Ceramic, Glassy, Polymeric, and Composite Materials). This work was supported in part by M.V. Lomonosov Moscow State University Program of Development. References [1] M.P. van Dijk, K.J. de Vries, A.J. Burggraaf, Solid State Ionics 16 (1985) 211. 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Please cite this article as: L.G. Shcherbakova, et al., Solid State Ionics (2014), http://dx.doi.org/10.1016/j.ssi.2014.01.019