Manipulating the grain boundary properties of BaCeO3-based ceramic materials through sintering additives introduction 38 D O I: 1 0. 15 82 6/ ch im te ch .2 01 9. 6. 2. 01 Vdovin G. K., Rudenko A. O., Antonov B. D., Malkov V. B., Demin A. K., Medvedev D. A. Chimica Techno Acta. 2019. Vol. 6, No. 2. P. 38–45. ISSN 2409–5613 G. K. Vdovina, A. O. Rudenkoa, B. D. Antonova, V. B. Malkova, A. K. Deminab, D. A. Medvedevab a Institute of High Temperature Electrochemistry, Yekaterinburg 620990, Russia b Graduate School of Economics and Management, Ural Federal University, Yekaterinburg 620002, Russia e-mail: dmitrymedv@mail.ru Manipulating the grain boundary properties of BaCeO3‑based ceramic materials through sintering additives introduction BaCeO3-based materials represent a well-known family of proton-conducting electrolytes, which can be used in different solid oxide electrochemical devices. An effective operation of the latter across an intermediate-temperature range requires improved transport of PCEs, including their grain (G) and grain boundary (GB) components. In the present work, some 3d-elements in a small amount were used as sintering additives to verify the possibility of improving the GB conduc- tivity of BaCe0.9Gd0.1O3–δ. It is shown that copper oxide (CuO) can be considered as one of the most effective sintering agents, since its use enables decreasing the GB density of the BCG ceramic material at the reduced sintering temperatures. The obtained results form a new tactic for designing new protonic electrolytes, whose conductivity might be prevail over ones containing Ni-based modifiers. Keywords: BaCeO3; SOFCs; SOECs; sintering additives; impedance spectroscopy; proton- conducting electrolytes. Received: 11.06.2019. Accepted: 28.06.2019. Published: 05.08.2019. © Vdovin G. K., Rudenko A. O., Antonov B. D., Malkov V. B., Demin A. K., Medvedev D. A., 2019 Introduction P r o t o n - c o n d u c t i n g o x i d e (PCO) materials occupy a  special place in the high-temperature electrochemistry due to its unique features consisting in pro- ton transportation in an oxide matrix [1–5]. These features allows PCOs to be utilized as electrolytes for various types of electro- chemical devices (EDs) such as solid oxide fuel cells (SOFCs), solid oxide electrolysis cells (SOECs), pumps and sensors [6–9]. As a result of high realizable conductivity levels of PCOs, the mentioned devices can operate at  reduced temperatures (below 600 °C) compared with the conventional systems based on oxygen-ionic electrolytes [10, 11]. In  order to  further improve the  EDs’ performance and efficiency, different strategies aimed at the electrolyte modi- fications can be exploited [1–3, 12–14]: the decrease of their thickness, designing the structures with a higher ionic mobil- ity, purposeful modification of  ceramics morphology. The latter is a highly promis- ing strategy, since the overall conductivity of  polycrystalline PCOs is  known to  be determined by a high resistance of grain boundaries [14]. One of the most obvious 39 ways to modify the grain boundaries rests on  adding the  low-melting phases [15], which intensify mass transport and pro- motes grain growth. Nickel oxide (NiO) is often used as such an additive. However, no improvement (or even deterioration) of grain boundary transport is observed [16–20]; this is  due to  two undesirable effects: very low solubility of  Ni2+-ions in  a  Ce-sublattice of  BaCeO3-based ma- terials and, consequently, the sedimenta- tion of  proton-blocking Ni-containing phases (NiO, BaY2NiO5) along the grain boundaries. Therefore, a  rational search of possible alternative to NiO is a matter of fundamental and applied interests as- sociated, respectively, with design of PCOs with optimized properties and their suc- cessful application in EDs. In the present work, on a well-known example of  PCOs, BaCe0.9Gd0.1O3–δ, pos- sibility of  improving its grain bound- ary transport was checked via an addition of CuO, Co3O4 and NiO as second dopants. To launch bridges between structural, mi- crostructural and transport properties of the materials obtained, the X-ray diffrac- tion (XRD), scanning electron microscopy (SEM) and electrochemical impedance spectroscopy (EIS) analyses were used. Experimental BaCe0.9Gd0.1O3–δ (BCG) material and its doping derivatives (BaCe0.89Gd0.1M0.01O3–δ, BCGM, where M = Cu, Co and Ni) were prepared using traditional solid state syn- thesis method. BaCO3, CeO2, Gd2O3, CuO, Co3O4 and NiO powders (with purity not less than 99.5 %) were taken in stoichio- metric amounts and thoroughly mixed via a mortar and pestle. The obtained mixtures were first pre-synthesized at 1100 °C for 5 h. The resulting powders were again mechani- cally activated, uniaxally pressed into discs and then sintered at 1450 °C for 3 h. One part of  the  obtained ceramic samples was crushed and characterized by XRD (diffractometer Rigaku D / MAX- 2200VL / PC, Japan), while another part was studied by SEM (JEOL JSM-5900 LV, Japan). Electrochemical characterization was performed for  the  Ag|BCG|Ag or Ag|BCGM|Ag symmetrical cells by uti- lizing an  Amel 2550 potentiostat / galva- nostat (Italy) and a  MaterialsM 520 fre- quency response analyser (Italy). These cells were fabricated in the following sequence: polishing the discs, an Ag paste painting and its sintering at 800 °C for 1 h. The im- pedance spectra were obtained for  wet (pH2O = 0.03 atm) air atmosphere in a fre- quency range of 10–2–106 Hz with an ampli- tude of 30 mV. The spectra were analysed by  an  equivalent circuit method, using a ZView software for model processing. Results and discussion The  XRD pattern of  the  sintered BCG material (Fig.  1) shows the  forma- tion of a single-phase product, the crys- tal structure of  which can be indexed as an orthorhombically-distorted perovs- kite with lattice parameters of a = 6.221 Å, b = 8.770 Å, c = 6.244 Å and space group of  Pmcn. The  doping with 3d-elements does not change symmetry of  the  per- ovskite structure, while the  lattice pa- rameters of BCGM are virtually constant (a = 6.220 ± 0.003 Å, b = 8.766 ± 0.005 Å, c = 6.237 ± 0.007 Å). These results can be explained by the extremely small amount of  sintering additives introduced into the BaCeO3 structure, 0.005 mol.%. 40 Despite of small concentrations used, the M-doping affects considerably the mi- crostructural parameters of  the  sintered materials that can be seen from qualita- tive (Fig.  2) and quantitative (Table  1) analyses. In  detail, the  processes associ- ated with grains growth, their close pack- age, pores disappearance and densifica- tion occur in all the cases. Nevertheless, a  degree of  these processes is  different and increases in the sequence of BCGCo– BCGCu–BCGNi, indicating dissimilar na- ture of the sintering additives. To reveal grain (σg.) and grain boundary (σg.b.) contributions of the total conductiv- ity (σtotal), the EIS analysis was successfully performed. The  electrochemical charac- terization was carried out in different tem- Fig. 1. XRD patterns of the sintered BCG and BCGM materials (a) and example of refinement for BCG (b) Fig. 2. Surface morphology of the sintered ceramic materials: BCG (a), BCGCo (b), BCGCu (c) and BCGNi (d) 41 perature ranges, boundaries of which were determined by  accuracy of  the  analysis. The  obtained spectra (Fig.  3) were ana- lyzed using an equivalent circuit scheme of  Ro–(R1Q1)–(R2Q2)–(R3Q3), where R is the resistance, Q is the constant phase element,  indexes of  1, 2, 3 correspond to  grain, grain boundary and electrode processes, respectively. Correlation of ho- dographs’ elements with these processes is performed analyzing the characteristic capacitance (C) and frequency (f) values calculated as follows: ( )1/ 1nC R Q R−= ⋅ ⋅ (1) ( ) ( )1/ 12nf R Q − −= ⋅ ⋅ π (2) Here, n is the power index in a frequen- cy dependence part of the constant phase element. Depending on the temperature, the C value varies between 20 and 80 pF for  the  first arc and between 3 and 8  nF for the second one, while the f values amount hundreds kHz and hundreds Hz, respec- tively. Both levels of these parameters relate with grain and grain boundary properties. It should be noted that Ro  = 0.001  Ω was purposefully introduced in the equiv- alent circuit scheme; it imitates the origin of the coordinates, providing a correct fitting. As  can be seen from Fig.  4a, Cu- and Ni-doping of  BCG results in  an  in- crease of  σg., while Co-doping has an opposite (but minor) effect. For ex- ample, the  σg. value at  200  °C reaches 5.01·10–5, 3.27·10–5 and 2.61·10–4  S  cm–1 for BCG, BCGCo and BCGCu, respectively. This parameter cannot be precisely deter- mined for the BCGNi material at 200 °C, but its σg. level is by ~2 times lower than that of BCGCu at the lower temperatures (100–150 °C). Table 1 Microstructural parameters of the BCG and BCGM materials sintered at 1450 °C for 3 h: ρ is the relative density, L is the total shrinkage, D is the average grain size, γ is the grain boundary density a Composition ρ, % L, % D ± 5 %, µm γ ± 5 %, µm–1 BCG 86 8.2 0.8 4.66 BCGCo 91 12.3 3.4 1.10 BCGCu 94 17.5 6.9 0.54 BCGNi 97 22.3 9.6 0.39 a estimated on the base of the following equation: γ = 3.722·D–1, see ref. [21]. Fig. 3. Impedance spectra obtained for the Ag|BCG|Ag symmetrical cell at different temperatures in wet air atmosphere (a) and example of the fitting procedure for the spectra obtained at 250 °C (b) 42 Discussing the grain boundary trans- port (Fig. 4b), the σg.b. improvement is ob- served for  all the  BCG-modified mate- rials and can be related with decreasing the grain boundary density (Table 1). Com- parison of the BCGCu and BCGNi samples allows formulation of the assumption that nickel is only partially dissolved in the Ce- sublattice of BCG, while another part local- izes onto grain boundaries. This proposal is based on the fact that BCGNi exhibits the lowest grain boundary density, which should provide the highest σg.b.; however, this is not confirmed experimentally. The resulting conductivity (σtotal, Fig. 4c) of the BCGM samples is higher than that of BCG, showing that the grain boundary transport determines the overall proper- ties even for BCGCo (at least in the entire studied temperature range). The BCGCu material exhibits the maximal achievable σtotal values, ranging from 6.33·10 –6 S cm–1 at 100 °C to 1.86·10–4 S cm–1 at 200 °C. As shown in Fig. 5, the apparent acti- vation energies (Ea) of σg. fall in the range of  0.45–0.49  eV, being in  close agree- ment with a characteristic value of 0.5 eV for proton transportation of PCOs [22–24]. According to  these data, the  M-doping does not affect the grain transport proper- ties of BCG. Another scenario is observed for the grain boundary transport proper- ties, when Ea of σg.b. decreases by 25–40 % comparatively 0.92 eV reaching for the ba- sic BCG oxide. It might be also associated with the  meaningful decrease of  grain boundary density serving as  a  barrier to ionic charge transfer. Fig. 6 displays the most interesting re- sult — a ratio between the grain boundary (Rg.b.) and total (Rtotal) resistances of BCG and BCGM. This ratio decreases signifi- cantly for  the  latter samples as  a  result of weakening the effect of grain boundaries on the overall transport of BCGM. A viv- id example can be seen when Rg.b. / Rtotal = = 0.5: for  BCG this level is  reached at ~250 °C, whereas for BCGM — at 100– 125 °C. In terms of real operation of PCOs (400–600 °C), the introduction of 3d-elem- etns in small amounts can improve the out- Fig. 4. Grain (a), grain boundary (b) and total (c) conductivities of the BCG and BCGM ceramic materials Fig. 5. Apparent activation energy values calculated for different types of conductivities 43 put properties (power density, current den- sity) of PCO-based electrochemical devices by tens of percent. Considering transport properties of BCGM and data on low-melting phases in corresponding systems [25–27], it can be concluded that the studied dopants behave differently respectively each other. For ex- ample, cobalt is assumed to act as a dopant, fully incorporated in the Ce-site of BaCeO3; nickel is mostly localizes at grain boundary region due to the mentioned low solubil- ity in the Ce-sublattice, although a certain amount can be nonetheless incorporated; finally, copper demonstrates dual nature: it has a high solubility (at least, more than 5  mol.% [21]), but can be also formed as a sediment at grain boundaries because of very low melting temperatures detected for a Ba–Cu–O system [15, 27]. Conclusions This work shows that the doping strate- gy of barium cerate with transition elements is one of the simplest and most effective methods aimed at  fabricating the  gas- tight ceramic samples at reduced sintering temperatures. This effect is achieved due to the intensification of diffusion processes caused by the appearance of a liquid phase. The  latter leads not only to  an  increase in the relative density of the ceramic mate- rials, but also to grain size growth, which may be favorable for designing new poly- crystalline materials and rational engineer- ing their grain boundary parameters. Acknowledgements This study was performed within the framework of the Russian Science Foundation [grant no. 18-73-00001]. The characterization of powder and ceramic materials was carried out at the Shared Access Centre “Composition of  Compounds” of  the  Institute of  High Temperature Electrochemistry [28]. References 1. Meng Y, Gao J, Zhao Z, Amoroso J, Tong J, Brinkman KS. Review: recent progress in low-temperature proton-conducting ceramics. J Mater Sci. 2019;54(13):9291–312. DOI: 10.1007 / s10853-019-03559-9. 2. Kim J, Sengodan S, Kim S, Kwon O, Bu Y, Kim G. Proton conducting oxides: A review of materials and applications for renewable energy conversion and storage. Renew Sustain Energy Rev. 2019;109:606–18 DOI: 10.1016 / j.rser.2019.04.042. Fig. 6. Temperature dependence of grain boundary resistance contribution of the total resistance of the BCG and BCGM ceramic materials 44 3. Mohd Rashid NLR, Samat AA, Jais AA, Somalu MR, Muchtar A, Baharuddin NA, Wan Isahak WNR. Review on zirconate-cerate-based electrolytes for proton-con- ducting solid oxide fuel cell. Cerm Int. 2019;45(6)6605–15. DOI: 10.1016 / j.ceramint.2019.01.045. 4. Putilov LP, Tsidilkovski VI. Impact of bound ionic defects on the hydration of acceptor- doped proton-conducting perovskites. Phys Chem Chem Phys. 2019;21(12):6391–406. DOI: 10.1039 / C8CP07745B. 5. Wang W, Medvedev D, Shao Z. Gas humidification impact on the properties and performance of perovskite‐type functional materials in proton‐conducting solid oxide cell. Adv Funct Mater. 2018;28(48):1802592. DOI: 10.1002 / adfm.201802592. 6. Putilov LP, Demin AK, Tsidilkovski VI, Tsiakaras P. Theoretical modeling of the gas humidification effect on the characteristics of proton ceramic fuel cells. Appl Energy. 2019;242:1448–59. DOI: 10.1016 / j.apenergy.2019.03.096. 7. Tarutin A, Lyagaeva J, Farlenkov A, Plaksin S, Vdovin G, Demin A, Medvedev D. A reversible protonic ceramic cell with symmetrically designed Pr2NiO4+δ-based electrodes: fabrication and electrochemical features. Materials. 2019;12(1):118. DOI: 10.3390 / ma12010118. 8. Danilov N, Lyagaeva J, Vdovin G, Medvedev D. Multifactor performance analysis of reversible solid oxide cells based on proton-conducting electrolytes. Appl Energy. 2019;237:924–34. DOI: 10.1016 / j.apenergy.2019.01.054. 9. Volkov A, Gorbova E, Vylkov A, Medvedev D, Demin A, Tsiakaras P. Design and applications of potentiometric sensors based on proton-conducting ceramic materi- als. A brief review. Sens Actuators B. 2017;244:1004–15. DOI: 10.1016 / j.snb.2017.01.097. 10. Dai H, Kou H, Wang H, Bi L. Electrochemical performance of protonic ceramic fuel cells with stable BaZrO3-based electrolyte: A mini-review. Electrochem Commun. 2018;96:11–5. DOI: 10.1016 / j.elecom.2018.09.00. 11. Yang L, Wang S, Blinn K, Liu M, Liu Z, Cheng Z, Liu M. Enhanced sulfur and cok- ing tolerance of a mixed ion conductor for SOFCs: BaZr0.1Ce0.7Y0.2–xYbxO3–δ. Science. 2009;326(5949):126–9. DOI: 10.1126 / science.1174811. 12. Lee YH, Chang I, Cho GY, Park J, Yu W, Tanveer WH, Cha SW. Thin film solid oxide fuel cells operating below 600°C: a review. Int J Precis Eng Manuf. 2018;5(3):441–53. DOI: 10.1007 / s40684-018-0047-0. 13. Cho GY, Lee YH, Cha SW. Thin film process for thin film solid oxide fuel cells — a review. J Korean Soc Precis Eng. 2018;35(12):1119–29. DOI: http://dx.doi.org / 10.7736 / KSPE.2018.35.12.1119. 14. Kjølseth C, Fjeld H, Prytz Ø, Dahl PI, Estournès C, Haugsrud R, Norby T. Space- charge theory applied to  the  grain boundary impedance of  proton conducting BaZr0.9Y0.1O3−δ. Solid State Ionics. 2010;181(5–7):268–75. DOI: 10.1016 / j.ssi.2010.01.014. 45 15. German RM, Suri P, Park SJ. Review: liquid phase sintering. J Mater Sci. 2009;44(1):1–39. DOI: 10.1007 / s10853-008-3008-0. 16. Tong J, Clark D, Bernau L, Subramaniyan A, O’Hayre R. Proton-conducting yttrium- doped barium cerate ceramics synthesized by a cost-effective solid-state reactive sintering method. Solid State Ionics. 2010;181(33–34):1486–98. DOI: 10.1016 / j.ssi.2010.08.022. 17. Yun DS, Kim J, Kim S-J, Lee J-H, Kim J-N, Yoon HC, Yu JH, Kwak M, Yoon H, Cho Y, Yoo C-Y. Structural and electrochemical properties of dense yttria-doped barium zirconate prepared by solid-state reactive sintering. Energies. 2018;11(11):3083. DOI: 10.3390 / en11113083. 18. Fang S, Wang S, Brinkman KS, Su Q, Wang H, Chen F. Relationship between fab- rication method and chemical stability of Ni–BaZr0.8Y0.2O3−δ membrane. J Power Sources. 2015;278:614–22. DOI: 10.1016 / j.jpowsour.2014.12.108. 19. Narendar N, Mather GC, Diasa PAN, Fagg DP. The importance of phase purity in Ni–BaZr0.85Y0.15O3−δ cermet anodes — novel nitrate-free combustion route and electrochemical study. RSC Adv. 2013;3(3):859–69. DOI: 10.1039 / C2RA22301E. 20. Han D, Otani Y, Noda Y, Onishi T, Majima M, Uda T. Strategy to improve phase compatibility between proton conductive BaZr0.8Y0.2O3−δ and nickel oxide. RSC Adv. 2016;6(23):19288–97. DOI: 10.1039 / C5RA26947D. 21. Ananyev M, Medvedev D, Gavrilyuk A, Mitri S, Demin A, Malkov V, Tsiakaras P. Cu and Gd co-doped BaCeO3 proton conductors: experimental vs SEM image algorith- mic-segmentation results. Electrochim Acta. 2014;125:371–9. DOI: 10.1016 / j.electacta.2013.12.161. 22. Kreuer KD. Proton-conducting oxides. Annu Rev Mater Res. 2003;33:333–59. DOI: 10.1146 / annurev.matsci.33.022802.091825. 23. Malavasi L, Fisher CAJ, Islam MS. Oxide-ion and proton conducting electrolyte materials for clean energy applications: structural and mechanistic features. Chem Soc Rev. 2010;39(11):4370–87. DOI: 10.1039 / B915141A. 24. Kochetova N, Animitsa I, Medvedev D, Demin A, Tsiakaras  P.  Recent activity in the development of proton-conducting oxides for high-temperature applications. RSC Adv. 2016;6(77):73222–68. DOI: 10.1039 / C6RA13347A. 25. Lander JJ. The phase system BaO–NiO. J Am Chem Soc. 1951;73(6):2450–2. DOI: 10.1021 / ja01150a012. 26. Klinkova LA, Nikolaichik VI, Barkovskii NV, Fedotov VK. New phases in the barium- rich region of the BaO–BaCuO2 system. Bull Russ Acad Sci: Phys. 2009;73(8):1104–6. DOI: 10.3103 / S1062873809080243. 27. Zhang W, Osamura K, Ochiai S. Phase diagram of the BaO–CuO binary system. J Am Ceram Soc. 1990;73(7):1958–64. DOI: 10.1111 / j.1151–2916.1990.tb05252.x 28. http://www.ihte.uran.ru / ?page_id=3142.