Enhanced ionic conductivity of Ag addition in acceptor-doped Bi_0.5Na_0.5TiO₃ ferroelectrics

Abstract

A promising lead-free ferroelectric material, Bi_0.5Na_0.5TiO₃ (BNT), was recently proposed as an oxide ion conductor. In a survey of BNT perovskite-structured oxides that exhibit ionic conductive grain, it was found to be highly sensitive to stoichiometry and the synthesis processes. The conduction mechanism and defect properties of these BNT therefore could be greatly influenced. Herein, we show composites made by a ferroelectric matrix of acceptor-doped BNT with a slight amount of Ag₂O in which the defect and transport properties could be tailored through different sintering atmospheres. The bulk conductivity increased dramatically by ∼2800 times to 1.36 mS cm⁻¹ at 502 °C.

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Materials that exhibit high oxygen ion conductivity are of considerable interest for a range of environmentally friendly electrochemical applications, including solid oxide fuel cells, gas sensors and active catalyst supports.^1–4 In the previous decades, study was devoted to improving performance of such existing materials and exploring new compositions. Materials with excellent oxide ion conducting properties range from fluorite-type,⁵ perovskite-structured^6,7 and new structured oxides such as La₂Mo₂O₉,^8–10 SrSiO₃,^11–13 and apatite oxides.^14,15 They exhibit different conduction mechanisms with the vacancy/interstitial-sites mediated transport resulted from extrinsic, intrinsic and structural disorders involved in them. Compared with the commercialized yttria-stabilized zirconia (YSZ), the perovskite-structured oxide ion conductors are promising alternatives because of their relatively low thermal expansion and better crystal lattice matching, such as lanthanum gallate (LSGM) with oxide ion conductivity higher than that of YSZ.¹⁶ Moreover, the perovskite structure with a general formula of ABO₃ displays diverse chemical compositions and properties, which can be optimized.¹⁷ The majority of these perovskite oxides are acceptor doped with low-valent cations, obtaining the formation of extrinsic oxygen vacancies and/or electronic species as charge compensating defects.^18–20

As a typical ferroelectric candidate to replace toxic lead zirconate titanate (PZT), perovskite-structured Bi_0.5Na_0.5TiO₃ (BNT) was recently proposed to be a potentially good oxide ionic conductor in the intermediate temperature range in an oxygen-deficient composition form.^20–24 In a previous study, we indicated a lone pair substitution theory as well as extrinsic disorders to clarify the significant structural evolution and the origin of their oxide ionic conductive performance in A/B sites con-substituted BNT perovskites.²⁵ The oxygen vacancies are generated by the acceptor doping and the A-sites disorder provides the weak Bi–O bonds, which facilitate delocalized conduction of the oxygen ions. However, the Na/Bi nonstoichiometry in BNT was also verified to display a dramatic effect in its electrical properties and conductive mechanism.²⁶ In fact, BNT structure with grain conductive properties is particularly sensitive to stoichiometry and synthesis processes. Based on extensive experiments, most of the potential contents with acceptor doping exhibit abnormal giant dielectric permittivity and rarely perform high conductive grain as expected. Moreover, the sintering temperature with a difference of only 10 °C will bring a noticeable change in the complex impedance characteristics. Thus, it raises a problem of the nature of defect species caused by acceptor doping that will significantly influence their oxide ionic conductive properties.

In a previous study, Ag was introduced to A-site of BNT and was supposed to form a morphotropic phase boundary to improve their piezoelectric properties. In this communication, slight amounts of Ag₂O were introduced in acceptor doped BNT to modify the defect species that originated from the percolative capacitor conception. With the addition of metallic oxide in ceramics, the same electric field can induce higher polarization and produce much greater conductivity.²⁷ Moreover, a preferential oxidation–reduction reaction is supposed at or along the grain boundaries, which will be helpful to stabilize the oxygen vacancies of grain structure in a certain oxygen partial pressure range. High permittivity values on these oxygen deficient materials can be ascribed to a Maxwell–Wagner interfacial polarization effect due to the presence of interfacial boundaries consisting of crystal defects such as screw dislocations.²⁸ The condition for the uncertainty of oxide ion conductive properties in such a BNT can be greatly improved through different sintering atmospheres.

Percolative Bi_0.5Na_0.5Ti_1−xGa_xO_3−δ@yAg₂O composites were prepared by mixing different amounts of Bi_0.5Na_0.5Ti_1−xGa_xO_3−δ and commercial Ag₂O (99.99%) powders, which were ball milled with ethyl alcohol for 6 h. The obtained powders were compacted into pellets by cold isostatic pressing. In case of bismuth volatility, they were dealt with by a burying sintering process at 1050–1100 °C for 4 h at a 3 °C min⁻¹ heating rate in air and nitrogen, respectively. Attempts to obtain samples with higher density resulted in pellets experiencing a stay at 300 °C for 2 h due to the exothermic chemical reactions of Ag₂O decomposition. Polycrystalline Bi_0.5Na_0.5Ti_1−xGa_xO_3−δ were fabricated by the conventional solid state ceramic route with mixing an appropriate amount of high purity Bi₂O₃ (99.9%, Sinopharm), Na₂CO₃ (99.8%, Sinopharm), TiO₂ (98%, Sinopharm), and Ga₂O₃ (99.999%, Sinopharm) as raw powders. Then, they were mixed and ball milled with ethyl alcohol for 6 h, then calcined at 950 °C for 4 h.

Phase structures of the powders and the sintered pellets were investigated by X-ray diffraction (XRD) with Cu Kα radiation in the Bragg angle (2θ) range from 10° to 90° (XRD-7000, Shimadzu, Kyoto, Japan). The Rietveld program FullProf was used for full-pattern matching and structural refinements. The Archimedes method was employed to measure the specimen density. The theoretical density was derived from the lattice parameters obtained from the XRD diffraction pattern. X-ray photoelectron spectroscopy (XPS) measurements were obtained using a spectrometer (VG ESCALAB220i-XL, Thermo Scientific, MA, USA).

The sintered pellets were polished for property investigation. Platinum electrodes were coated on both polished surfaces and fired at 950 °C for 3 h to ensure maximum conductivity and adherence. A frequency response analyzer (CH660, ChenHua, Shanghai, China) was used to obtain the impedance spectra in the frequency range of 0.1 Hz to 1 MHz from 150 °C to 750 °C. After 30 min of thermal stabilization, each spectrum was obtained at 50 mV AC voltage and the impedance data was collected by Z-viewer software. The weak-field dielectric response with a 50 mV signal level was measured using a precision impedance analyzer (4294A, Agilent, CA, USA) associated with a temperature controller (TP94, Linkam, Surrey, UK) at a 3 °C min⁻¹ heating rate.

Fig. 1 presents the dielectric permittivity (ε′) and loss tangent (tan δ) of Bi_0.5Na_0.5Ti_0.9Ga_0.1O_3−δ@1 wt%Ag₂O (BNTG) sintering in air and nitrogen as a function of temperature. Both the ε′ in lower frequency increase by several orders of magnitude and show much frequency dependence in the high temperature range wherein it begins at about 200 °C for air sintered BNTG (A-BNTG) and 150 °C for nitrogen sintered BNTG (N-BNTG). For the high frequency, the maximum ε′ at the temperature (T_m) is about 2600 that can be ascribed to the intrinsic contribution that originates from cation displacement in BNT perovskite. The sudden increase in ε′ is due to the localized charge carriers hopping between spatially fluctuating lattice potentials. Both the loss of A-BNTG and N-BNTG are high indicating the potential high conductive grain involved in them. By comparison, N-BNTG displays a plateau in ε′ of ∼10⁵ with loss peaks present in the tan δ–T curves.

Fig. 1 Temperature dependences of dielectric permittivity (ε′) and loss tangent (tan δ) of Bi_0.5Na_0.5Ti_0.9Ga_0.1O_3−δ@1 wt%Ag₂O composites sintered in air and nitrogen, respectively.

The BNT family is found to be particularly sensitive to stoichiometry. Some BNTs exhibit an intensive increase in ε′ and tan δ when T_m, is reached, whereas most BNTs report a decrease in ε′ on exceeding the T_m.^{21,26,29–31} However, no literature is devoted to these obvious changes, which may be attributed to the different mounts of cationic volatilization or oxygen deficiency in sintering process. BNT with an abnormal high loss will not be necessarily accompanied with a high conductive grain in the intermediate temperature range. Fig. 2 further illustrates this by employing the complex impedance Z* of A-BNTG and N-BNTG at selected temperatures. Completely different features of ac impedance responses were observed. A-BNTG displayed a typical big arc in nearly the complete frequency range associated with grain response, which is common to other BNT based Z* characterization. The conductivity of A-BNTG was significantly influenced by the different measured atmospheres. The A-BNTG exhibited a higher conductivity in nitrogen and dramatically decreased in air, indicating an n-type electronic conduction. However, nitrogen sintered BNT phase showed a small arc at high frequencies and a large arc at low frequencies, which could be ascribed to the grain and grain boundary contributions, respectively. Thus, the bulk conductivity of N-BNTG, deduced from the intercept of the Z′ axis, increased dramatically by almost 2800 times to 1.36 mS cm⁻¹ from 0.49 nS cm⁻¹ for A-BNTG at 502 °C. Compared with the previous studies, we confirmed that a transition occurred from an n-type electronic conduction to an oxide ionic conduction mechanism. Associated with the sharp increase of the conductivity, the resistive and conducting samples exhibited different activation energies (E_a) obtained from the Arrhenius plots of grain conductivity, as summarized in Fig. 3. The resistive A-BNTG revealed a linear relationship with E_a ∼ 1.7 eV, whereas N-BNTG showed a change in slope with E_a ∼ 0.56 eV in the low temperature range and E_a ∼ 0.38 eV in the high temperature range. The E_a evolution displayed similar characteristics with the previous studies and our study, which can be ascribed to the delocalized conduction of oxygen ions in BNT, while it is noticeable that the T inflection point in the Arrhenius plot of conductive N-BNTG did not correspond to the T_m.^21,25,26

Fig. 2 Complex impedance plots of A-BNTG and N-BNTG at selected temperatures. Filled symbols denote selected frequencies.

Fig. 3 Arrhenius plots of bulk conductivity with different activation energies for A-BNTG and N-BNTG.

In this case, a single perovskite structure with rhombohedral R3c symmetry is indicated from the XRD results for both samples. Detailed information can be observed in ESI Fig. S1.† Moreover, the phase in BNTG becomes electrically inhomogeneous after the nitrogen-sintering process. In terms of defect chemistry, the acceptor doping will require charge compensation by a species with an effective positive charge. In this case, the form of oxygen vacancy is preferred as described by the following reaction:

A low concentration of oxygen in the sintering process, therefore, will greatly promote the substitution of Ti⁴⁺ by Ga³⁺. Moreover, positive oxygen vacancies will be prone to generate because of the easier volatilization of Bi cations in a reduced atmosphere compared to volatilization in air. The A-BNTG seems to have potential in exhibiting high oxide ion conductive properties; moreover, it did not provide enough oxygen vacancies in an air-sintering process. Furthermore, the smaller cation radius of B sites in ABO₃ perovskite may obtain a fraction of Ga³⁺ entering into the interstitial site. Thus, the ceramic behaves as an n-type semiconductor in the measured oxygen partial pressure, which can be accounted as follows:

Ti⁴⁺ + e′ → Ti³⁺ (1-3)

The resulting electron would not be fixed to a specific Ti atom, where it can easily migrate from one position to another. Specially, Ti⁴⁺ would turn into Ti³⁺ when the electron is bound to the positive oxygen vacancy, is related to the nearby Ti⁴⁺. The electron can migrate to the neighboring Ti⁴⁺ under the electrical field and become electrically conductive. Note that Ag addition in BNTG is helpful to obtain its high conductivity and improve the stability, especially for property repeatability because it is really highly sensitive to stoichiometry and the synthesis processes for acceptor-doped BNT as discussed above. Cycle stability of grain conductivity are provided in Fig. S2.†

To further identify the origin of the anomalous electrical response between A-BNTG and N-BNTG, a series of compositional, structural and defect analysis studies were carried out. By conducting SEM and EDS on different regions of the grain and grain boundary, few differences between them were observed (see ESI Fig. S3†). Both grains of the samples were large and dense, implying the role of Ga in BNT systems as a grain growth promoter. XPS results of A-BNTG and N-BNTG at ambient temperature are given in Fig. 4. They demonstrate the position of the Ti 2p doublet with 2p3/2 and 2p1/2 binding energies with 5.9 eV spin–orbit splitting (Fig. 4a).³² Both the Ti 2p peak of N-BNTG shifted to lower binding energies, 457.82 eV for 3/2 and 463.72 eV for 1/2 spin–orbit as compared with 457.97 eV and 463.88 eV for A-BNTG. This suggests an expected increase in the electronegativity and more oxygen vacancies in the TiO₆ octahedra in N-BNTG. In addition, a noticeable Ti³⁺ signal was detected in A-BNTG, which is enlarged in the inset of Fig. 4a, giving effective evidence for the supposed reaction in a different sintering atmosphere. The shift to lower binding energies of Bi 4f peak in Fig. 4b revealed a further weakening of the Bi–O hybrid orbitals and reduction in the coordination number, which was helpful for migration of oxygen ions. The O 1s profile of as-sintered and etched BNTG are presented in Fig. 3c and d. The main peaks at ∼528.9 eV correspond to the bulk Ti–O cation-oxygen bonds. The peaks at ∼530.6 eV of the as-sintered samples (Fig. 4d) are assigned to adsorbed oxygen species or oxygen vacancies. Quantification of each oxygen species were calculated and the results suggest that the amount of lattice oxygen species of N-BNTG could be significantly decreased from 57.55% to 55.37%. Due to the imbalance of the oxidation state between Ti⁴⁺ and Ga³⁺ site, oxygen vacancies were preferentially created, resulting in an increase of surface oxygen species. An extra peak observed in N-BNTG from the inset of Fig. 4d indicates the non-equilibrium cation-oxygen bonds in ABO₃ perovskite that could be related to the Ga–O bonds, giving a further illustration of the inhomogeneity in N-BNTG. In addition, the polarization hysteresis loops illustrate the ferroelectric characteristic of A-BNTG and N-BNTG, as shown in Fig. S4.† Oxygen vacancies in N-BNTG have a stronger pinning effect for the ferroelectric domain switching, thereby inducing a continuous increase in coercive field E_c and a slight decline in remnant polarization P_r.

Fig. 4 XPS of A-BNTG and N-BNTG. (a) Ti 2p, (b) Bi 4f, (c) O 1s spectra of 1 min sputtered A-BNTG and N-BNTG. (d) O 1s spectra of as-sintered A-BNTG and N-BNTG. The corresponding fitting results are shown in dashed lines. Insets in (a), (d) are enlarged views to clearly show the shoulders in the black lined boxes.

In summary, composites made by a ferroelectric matrix of acceptor-doped BNT with metallic inclusions are helpful to achieve high conducted grains, making them potentially suitable in applications such as fuel cell electrodes/electrolytes and catalyst supports. The sintered atmospheres had a significant effect on the electrical properties and conduction mechanism of such BNT system. As such, the sensitive electrical response of nonstoichiometric compositions of conductive BNT thus could be greatly optimized by controlling the defect form of acceptor doping.

Acknowledgements

This study was supported by the National Natural Science Foundation (51172187, 51372197), the Key Innovation Team of Shaanxi Province (2014KCT-04), the Major International Joint Research Program of Shaanxi Province (2012KW-10) and the Doctoral Starting-up Foundation of Xi'an University of Science and Technology (2015QDJ046).

References

1. L. Malavasi, C. A. Fisher and M. S. Islam, Chem. Soc. Rev., 2010, 39, 4370–4387.
2. J. W. Fergus, Electrochemical Sensors: Fundamentals, Key Materials and Applications, in Handbook of Solid State Electrochemistry, ed. V. V. Kharton and J. W. Fergus, Wiley-VCH, New York, 2009.
3. P. Vernoux, L. Lizarraga, M. N. Tsampas, F. M. Sapountzi, A. De Lucas-Consuegra, J. L. Valverde, S. Souentie, C. G. Vayenas, D. Tsiplakides, S. Balomenou and E. A. Baranova, Chem. Rev., 2013, 113, 8192–8260.
4. A. J. Jacobson, Chem. Mater., 2010, 22, 660–674.
5. D. J. Brett, A. Atkinson, N. P. Brandon and S. J. Skinner, Chem. Soc. Rev., 2008, 37, 1568–1578.
6. S. W. Tao, J. T. S. Irvine and J. A. Kilner, Adv. Mater., 2005, 17, 1734–1737.
7. C. Tealdi and P. Mustarelli, J. Phys. Chem. C, 2014, 118, 29574–29582.
8. P. Lacorre, F. Goutenoire, O. Bohnke, R. Retoux and Y. Laligant, Nature, 2000, 404, 856–858.
9. G. l. Corbel, E. Suard and P. Lacorre, Chem. Mater., 2011, 23, 1288–1298.
10. X. Liu, H. Fan, J. Shi, G. Dong and Q. Li, Int. J. Hydrogen Energy, 2014, 39, 17819–17827.
11. C. Tealdi, L. Malavasi, I. Uda, C. Ferrara, V. Berbenni and P. Mustarelli, Chem. Commun., 2014, 50, 14732–14735.
12. P. Singh and J. B. Goodenough, Energy Environ. Sci., 2012, 5, 9626–9631.
13. P. Singh and J. B. Goodenough, J. Am. Chem. Soc., 2013, 135, 10149–10154.
14. E. Kendrick, M. S. Islam and P. R. Slater, J. Mater. Chem., 2007, 17, 3104–3111.
15. M. S. Islam, J. R. Tolchard and P. R. Slater, Chem. Commun., 2003, 1486–1487.
16. T. Ishihara, H. Matsuda and Y. Takita, J. Am. Chem. Soc., 1994, 116, 3801–3803.
17. T. Ishihara, Perovskite Oxide for Solid Oxide Fuel Cells, Springer Verlag, New York and London, 2009.
18. M. Cherry, M. S. Islam and C. R. A. Catlow, J. Solid State Chem., 1995, 118, 125–132.
19. M. S. Islam, J. Mater. Chem., 2000, 10, 1027–1038.
20. X. He and Y. Mo, Phys. Chem. Chem. Phys., 2015, 17, 18035–18044.
21. M. Li, M. J. Pietrowski, R. A. De Souza, H. Zhang, I. M. Reaney, S. N. Cook, J. A. Kilner and D. C. Sinclair, Nat. Mater., 2014, 13, 31–35.
22. J. Shi, H. Fan, X. Liu and Q. Li, J. Eur. Ceram. Soc., 2014, 34, 3675–3683.
23. J. A. Dawson, H. Chen and I. Tanaka, J. Mater. Chem. A, 2015, 3, 16574–16582.
24. B. N. Rao, R. Datta, S. Chandrashekaran, D. K. Mishra, V. Sathe, A. Senyshyn and R. Ranjan, Phys. Rev. B: Condens. Matter Mater. Phys., 2013, 88, 224103.
25. X. Liu, H. Fan, J. Shi and Q. Li, Sci. Rep., 2015, 5, 12699.
26. M. Li, H. Zhang, S. N. Cook, L. Li, J. A. Kilner, I. M. Reaney and D. C. Sinclair, Chem. Mater., 2015, 27, 629–634.
27. F. E.-B. C. Pecharromán, J. F. Bartolomé, S. López-Esteban and J. S. Moya, Adv. Mater., 2001, 13, 1541–1544.
28. C. E. Ciomaga, M. T. Buscaglia, V. Buscaglia and L. Mitoseriu, J. Appl. Phys., 2011, 110, 114110.
29. M. Spreitzer, M. Valant and D. Suvorov, J. Mater. Chem., 2007, 17, 185–192.
30. Y. S. Sung, J. M. Kim, J. H. Cho, T. K. Song, M. H. Kim and T. G. Park, Appl. Phys. Lett., 2011, 98, 012902.
31. A. O'Brien, D. I. Woodward, K. Sardar, R. I. Walton and P. A. Thomas, Appl. Phys. Lett., 2012, 101, 142902.
32. W. Hu, Y. Liu, R. L. Withers, T. J. Frankcombe, L. Norén, A. Snashall, M. Kitchin, P. Smith, B. Gong, H. Chen, J. Schiemer, F. Brink and J. Wong-Leung, Nat. Mater., 2013, 12, 821–826.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6ra00682e

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