Structural and electrical properties of strontium substituted Y2BaNiO5

Journal of Alloys and Compounds 620 (2015) 91–96

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Structural and electrical properties of strontium substituted Y2BaNiO5 Narendar Nasani a, Devaraj Ramasamy a, Isabel Antunes b, Budhendra Singh a, Duncan P. Fagg a,⇑ a b

Nanotechnology Research Division, Centre for Mechanical Technology and Automation, Department of Mechanical Engineering, University of Aveiro, 3810-193 Aveiro, Portugal Department of Materials and Ceramic Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal

a r t i c l e

i n f o

Article history: Received 31 July 2014 Received in revised form 12 September 2014 Accepted 15 September 2014 Available online 22 September 2014 Keywords: Haldane energy gap Barium yttrium nickelate Electrical conductivity Solid solution Protonic ceramics

a b s t r a c t The Y2xSrxBaNiO5 (x = 0, 0.1, 0.2 and 0.3) acceptor substituted system has been synthesized by solid state reaction. Structural and microstructural properties have been characterized by X-ray diffraction (XRD) and scanning electron microscopy (SEM), respectively. Lattice volume is shown to decrease linearly with increasing Sr content until composition x = 0.2, highlighting the limit of the solid solution. The electrical response in the temperature range (700–100 °C) was assessed by A.C. impedance spectroscopy in wet and dry O2 and N2 atmospheres. Conductivity measurements as a function of oxygen partial pressure (pO2) were also performed. The data reveal that the conductivity Y2BaNiO5 can be increased by one and half orders magnitude by Sr-doping and is independent of both water vapour and oxygen partial pressures (pH2O and pO2). The low activation energy for electrical conduction (0.216–0.240 eV) suggests a thermally activated electron hopping mechanism, while the observed pO2 and pH2O independence of conductivity suggests that charge compensation for Sr is predominantly by formation of Ni3+ rather than formation of oxygen vacancies. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction The current article assesses the structural and electrical properties of materials based on the base composition Y2BaNiO5; a material that is most commonly studied as Haldane energy gap material. Since the discovery of the Haldane energy gap in one dimensional anti-ferromagnetic materials (1D AF), several materials types have been shown to offer this property [1,2]. Haldane materials demonstrate an energy gap between singlet ground state and the first excited triplet state when the spin quantum number is an integer. Such materials have generated interest as they can facilitate the study of ground and excited states of quantum models as well as offering the possibility of switchable devices. Rare earth doped R2BaNiO5 (R = Y, rare earth) compounds [3]. Tm2BaNiO5 and Yb2BaNiO5 were reported to be dimorphic with Ni ions in pyramidal (Pnma) or octahedral (Immm) structural coordination [4,5]. The magnetic susceptibility [6,7] and spin thermal conductivity [8] behaviour in these compounds have been extensively explored, while electronic structures have been probed by using various techniques viz. extended X-ray absorption fine structure (EXAFS), X-ray absorption near-edge structure (XANES) [9], electron energy loss spectroscopy (EELS) [10], nuclear magnetic resonance spectroscopy (NMR) [11], X-ray and/or neutron diffraction techniques [3] and photoemission spectroscopy [12]. ⇑ Corresponding author. Tel.: +351 234 370830; fax: +351 234 370953. E-mail address: [email protected] (D.P. Fagg). http://dx.doi.org/10.1016/j.jallcom.2014.09.127 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

Doped materials have also been studied, Novák et al. studied the electronic structure of Y2xMxBaNiO5 (M = Sr, Ca) by density functional theory (DFT) [13] while, Capsoni et al. studied the effect of different ion substitutions at the Ni-site (Y2BaNi1xMxO5, M = Mg, Zn) using Raman spectroscopy [14]. In the case of Mg substitutions, structural symmetry (orthorhombic with Immm space group) with NiO6 octahedra was preserved. On the contrary, with the addition of Zn2+ ions, both Immm and Pnma symmetries were found to coexist in samples with x = 0.13 while, at higher concentrations of Zn2+ ions, single structural symmetry (Pnma) having NiO5 pyramidal coordination was observed, with the assimilated Zn composition showing the existence of isolated pyramid structure. Structural changes and magnetic susceptibility of Y2BaNi1xZnxO5 oxides were also studied by Saez–Puche et al. noting similar structural transformations and shift to Curie–Weiss behaviour on inclusion of Zn [15]. Y2BaNiO5 (BYN) is one of the typical Haldane gap oxide materials having 1D AF nature [16] with orthorhombic symmetry (Immm space group) having isolated 1D chains of vertex-sharing NiO6 octahedra along the z-axis. The NiO6 octahedra are compressed along the z direction which leads to the presence of two short apical Ni–O bond distances (1.88 Å). Nevertheless, the ground state has insulating properties with a large optical gap of about 2.3 eV [16]. The most important structural feature in these compounds is their chemical bonding which leads to 1D magnetic behaviour, due to the absence of Ni–O–Ni interaction between neighboring chains, allowing Y2BaNiO5 to be an ideal 1D anti-ferromagnet with

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S = 1. Lannuzel et al. successfully incorporated acceptor dopants (Sr2+ and Ca2+) for Y3+ into Y2BaNiO5 compounds and their studies revealed that the magnetic and transport properties, measured at low temperatures, could be modified by such doping, with a decrease in resistivity of BYN of two orders magnitude measured at 300 K and the generation of new magnetic states with spin greater than S = 1/2 inside the Haldane gap [10]. Recently, Chen et al. studied the dielectric properties of Y2BaNiO5 in the frequency range of 20 Hz to 1 MHz between 20 K and room temperature and demonstrated high dielectric permittivity (e0  104) at room temperature below 100 kHz [17]. As highlighted, most previous studies of BYN have been devoted towards the understanding of basic structural, magnetic and electrical properties of Haldane gap systems measured at low temperatures. Nevertheless, Y2BaNiO5 has also been identified as an impurity phase in proton conducting ceramic electrolyte materials, particularly yttrium doped barium cerates and zirconates (BaCe(1xy)ZryYxO3d), when NiO is used as sintering additive [18–21]. It has also been reported to be present in nickel cermet anodes Ni–BaCe(1xy)ZryYxO3d prepared by typical synthesis techniques [22,23]. Tong et al. mentioned that the Y2BaNiO5 impurity phase can play vital role in the grain growth and densification of Ba(Zr,Y)O3d materials when one uses NiO as sintering aid to reduce the sintering temperature below 1600 °C [18,19]. Thus, the presence of BYN as an impurity phase in typical components of protonic ceramic fuel cells (PCFCs) may contribute to their subsequent electrochemical performance. It is, therefore, necessary to understand the electrical properties of BYN also at temperatures that are relevant for operation of PCFC devices (500–700 °C). This factor, therefore, led the authors to study the conductivity behaviour of Y2BaNiO5 in controlled atmospheres as a function of temperature to assess its subsequent properties when incorporated in protonic ceramic fuel cells or in further electrochemical applications that operate at elevated temperatures. To the best of author’s knowledge, there are no conductivity studies previously available on BYN under such conditions. In the current article, we also investigate the effect of strontium doping on structural and electrical properties of Y2xSrxBaNiO5 for different doping concentrations (x = 0, 0.1, 0.2 and 0.3).

2. Experimental section 2.1. Synthesis of Y2xSrxBaNiO5 (x = 0, 0.1, 0.2 and 0.3) Y2xSrxBaNiO5 (x = 0, 0.1, 0.2 and 0.3) was prepared by a combination of mechanical activation and solid state synthesis [24] using Y2O3 (Hermann C. Starck), SrO2 (Sigma Aldrich), BaO2 (Sigma Aldrich) and NiO (99% Strem Chemicals ABCR GmbH&Co.). Y2O3 was dried overnight at 900 °C and kept at 250 °C prior to weighing. Stoichiometric quantities of all the reactant powders were mixed and ball milled in zirconia vials under high energy at 650 rpm for 420 min using zirconia balls to mechanically activate the precursor powders. For simplicity the Y2xSrxBaNiO5 compositions are given the nomenclature, SBYN0, SBYN1, SBYN2 and SBYN3, for x = 0, 0.1, 0.2, and 0.3, respectively. The obtained powder was dark brown in colour. The activated powder was uniaxially pressed into pellets with 1 cm diameter and sintered at 1400 °C for 6 h with a heating and cooling rate of 4 °C/min in ambient air. Then the sintered pellets were crushed into powder in agate mortar using pestle and analyzed for phase purity. For electrical measurements, bar shaped pellets (1.3  0.3  0.4 cm3) were prepared by cold isostatic pressing of the activated powder at 300 MPa for 15 min followed by similar sintering conditions as above. The geometric density of sintered samples was found to be 85–90% that of theoretical density, with the higher density samples being those with the higher Sr-content.

was analysed by structural and profile fitting model [25]. The microstructure of sintered pellets were observed using Scanning Electron Microscopy (SEM) (model Hitachi SU-70). Electrical measurements were performed on bar shaped samples (4 probe configuration) with an Autolab (Electrochemie) PGSTAT302N frequency response analyzer using 50 mV amplitude in 1 MHz–0.01 Hz frequency range in the direction of decreasing temperature in the range 700–100 °C with 50 °C intervals. Prior to the electrical measurements, contacts were painted on the dense bars with Pt paste and fired at 900 °C for 30 min with a heating/cooling rate of 4 °C/min. All samples were measured under dry/wet N2 and O2 atmospheres with a flow rate of 50 ml/ min. Humidification was achieved by bubbling gases through distilled water in a saturated KCl solution, producing approximately 86% relative humidity at ambient conditions. Drying was performed using Varian gas driers. Humidity levels were measured using a humidity meter (Jumo). The oxygen partial pressure (pO2) was controlled by mixing O2 with N2 with 2 h stabilization times at each measurement during the pO2 experiment. Stability of the data was confirmed by performing repeated impedance measurements at each condition after a further 30 min. The obtained impedance data was analyzed by using ZView software (Scribner Associates).

3. Results and discussion 3.1. Phase formation and structural analysis For the composition x = 0 (SBYN0) the desired BYN (JCPDS-0810002, Y2BaNiO5) phase was identified in the XRD patterns, along with small traces of Y2O3 (JCPDS-155173) and Ba3Y4O9 (JCPDS087118), Fig. 1. For SBYN1, SBYN2 and SBYN3, the BYN phase was identified with minor traces of the Ba3Y4O9 impurity peak, but without any Y2O3 impurities. Capsoni et al. also found the impure phases Y2O3 and BaY2O4 in Y2BaNi1xMxO5 (M = Mg, Zn) polymorphs prepared by solid state reaction under synthesis conditions involving multiple firing, grinding steps with a final maximum firing temperature of 1200 °C [14]. The impurity phases observed in the current work and in the work of Capsoni are in accordance with the phase equilibrium study of Lopato et al. of the Ba–Y–O system, who showed that BaY2O4 is the stable phase for temperatures <1400 °C while the Ba3Y4O9 phase is stable at >1400 °C. The phase diagram of the Y2O3–BaO–NiO system constructed by Buttrey et al. in the 1000–1350 °C temperature range, suggests that BaY2NiO5 is a point composition [26]. Hence, any slight deviation of stoichiometry would lead to the presence of impurity phases. From the phase diagram, the observed impurity, Ba3Y4O9, would suggest Ba-excess and entrance into the BaY2NiO5 + Ba3NiO4+x + Ba3Y4O9 region. Nonetheless, the composition of each compound was confirmed by energy dispersive X-ray analysis (SEM-EDAX) and was shown to be stoichiometrically accurate at the resolution of this technique. Attempts to remove

2.2. Materials characterization The powder X-ray diffraction (XRD) data and phase purity were analyzed by using a Philips X’Pert diffractometer in the 2h = 20–80° angular range, Detector X’Celerator, active length 2.5460°, step width 0.02° and counting time 30 s/step, operated at 45 kV and 40 mA with Cu Ka radiation. The obtained diffraction pattern

Fig. 1. X-ray diffraction patterns of Y2xSrxBaNiO5 (X = 0, 0.1, 0.2 and 0.3) calcined at 1400 °C for 6 h in ambient air.

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the minor impurity phase of Ba3Y4O9 by further mechanical activation were unsuccessful. Nonetheless, Rietveld refinement highlighted that the content of impurity phases in all compositions was 2 wt%, Fig. 2; a very minor level (1 mol%) that can be considered negligible. Moreover, the absence of any strontium related impurity phases suggests the incorporation of the Sr2+ dopant in the base composition BYN, in agreement with the article aim. All the Y2xSrxBaNiO5 (x = 0, 0.1, 0.2 and 0.3) compositions exhibit an orthorhombic system with space group Immm. The lattice parameters and the atomic positions were refined using the Rietveld method with profile and structural model using Fullprof software, Table 1. Satisfactory global agreements for Rietveld refinement were obtained. The unit cell parameters are shown to decrease with increasing strontium content up to the composition x = 0.2, Table 1, and a corresponding decrease in unit cell volume is also noted, Fig. 3. As the ionic size of Sr is greater than that of the Y host, Sr2+ (1.13 Å) cf. Y3+ (0.93 Å) [10,27], the observed lattice volume contraction can only be explained by consideration of the potential charge compensation mechanisms for the acceptor dopant. For instance, due to a decreasing unit cell size related to the formation of oxygen vacancies or due to the oxidation of Ni2+ to Ni3+ with an associated reduction in ionic size. The most likely of these charge compensation mechanisms will be discussed in the later text based on the electrochemical results. The inversion of the lattice parameters and associated lattice volume at the composition x = 0.3, suggests that this composition exceeds the solid solubility of Sr in the BYN structure. The solid solution limit of x = 0.2 corresponds well with that previously noted by Lannuzel et al. [10], while being smaller than that reported for the Ca-substituted system, Y2xCaxBaNiO5, where solubilities up to x = 0.5 have been documented [28]. The lower solubility in the current case is likely

Table 1 Rietveld refinement results for pure phase of Y2xSrxBaNiO5 materials synthesized by mechano-synthesis.

a

(a) Basic structural details Structure

Space group

Orthorhombic

Immm

(b) lattice parameters (in Å) and angle (in °) Sample a b c

a

b

c

Vol. (Å3)

SBYN0 SBYN1 SBYN2 SBYN3

90 90 90 90

90 90 90 90

90 90 90 90

245.093 244.844 244.496 244.651

3.75579 3.75444 3.75108 3.75068

5.76000 5.75655 5.75465 5.75604

11.32939 11.32874 11.32654 11.33216

(c) Goodness of fita Sample Rb Rf

v2

c/a

S (goodness of fit) = Rwp/Rexp

SBYN0 SBYN1 SBYN2 SBYN3

4.63 5.05 3.46 7.28

3.0165 3.0174 3.0195 3.0214

0.8333 0.8165 0.8041 0.9001

6.07 8.75 4.18 6.84

4.94 7.66 3.78 5.15

Rb: Bragg R-factor, Rf: Crystallographic Rf – factor and v2: Goodness of fit.

SBYN0

Fig. 3. Variation of lattice volume as a function of the strontium content for Y2xSrxBaNiO5.

to be due to the larger ionic size mismatch between Sr and the Y host than that for Ca. Lattice volume decreases with increasing acceptor substitution have also been documented in the BaY2xCaxNiO5 system, in agreement with the current work [28]. 3.2. Microstructure

SBYN2

The microstructure of Y2BaNiO5 (BYN) was observed by scanning electron microscopy (SEM). Fig. 4 represents the fractured cross sectional SEM images of series of compositional pellets sintered at 1400 °C for 6 h. The SBYN0 pellet can be observed to have not attained full densification at 1400 °C and to contain pores, corresponding to its lower measured density of 85%. In contrast, the Sr substituted BYN analogues are all observed to be dense, Fig. 4, and to exhibit larger grain sizes. An example Energy Dispersive X-ray Analysis image (EDAX) for the sample x = 0.1 is included in Fig. 4e and highlights that the composition is homogenous throughout the material. Note also that no impurities can be observed by this technique, due to their aforementioned low concentration (<1 mol%). 3.3. Electrical conductivity of Y2xSrxBaNiO5

Fig. 2. Observed, calculated, and difference X-ray powder diffraction profile from the Rietveld refinement for SBYN0 and SBYN2 calcined at 1400 °C for 6 h.

The conductivity of pure and strontium modified BYN samples were evaluated as a function of temperature under a range of

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(a) SBYN0

(b) SBYN1

(c) SBYN2

(d) SBYN3

(e) SBYN1-EDS

Fig. 4. Cross sectional SEM image of (a) undoped and Sr doped BYN (b), (c), (d) and EDAX elemental mapping of Sr doped BYN (e) pellet sintered at 1400 °C for 6 h in ambient air.

controlled atmospheres. Fig. 5a depicts the conductivity of SBYN0, SBYN1 and SBYN2 compounds as a function of temperature in N2 gas atmospheres. The conductivities of all samples exhibit positive temperature dependences, while the conductivities in dry and wet N2 atmospheres are observed to be effectively equal. With the incorporation of Sr2+ in the BYN lattice, a significant increase in conductivity can be observed, yielding an increase of approximately one and half orders of magnitude over that of the base composition [17,28]. However, the measured conductivities for samples SBYN1 and SBYN2 are shown to be similar across most of the temperature range, only revealing slight improvements with increased Sr substitution at the highest temperatures. To clarify this feature, Fig. 5b plots this data on a linear conductivity scale. The observed non-linear increase in conductivity with increasing acceptor dopant content is in agreement with that previously observed in the system Er2xCaxBaNiO5, where compositions x = 0.34 and x = 0.19 showed similar levels of conduction, which were an order of magnitude higher than the base composition [28]. Fig. 6 shows the variation of the conductivity for SBYN0, SBYN1 and SBYN2 compositions as a function of temperature measured in wet and dry O2 atmospheres. The resultant trends are shown to mirror those measured under N2, in all aspects. Previously, Lannuzel et al. proposed that Sr/Ca doping can strongly affect the magnetic and transport properties of BYN by alterations in its electronic structure [10], a suggestion supported by DFT calculations by Novak et al. [13]. For the current BYN

system, the conductivities measured in Figs. 5 and 6 are shown to be thermally activated, with activation energies, determined using the Arrhenius relation in the temperature range 500– 100 °C, outlined in Table 2. Activation energy values lie in the range 0.216–0.240 eV for all samples and can be noted to be effectively independent of atmosphere. A significant feature of the present work is that the activation energy 0.216–0.240 eV of BYN is significantly lower than the expected energy gap 2.3 eV of the undoped BYN Haldane system, in agreement with that observed by DiTusa [29] for single crystal samples measured in a sub-ambient temperature range. Recently Chen et al. documented dielectric and resistivity behaviours of polycrystalline samples of the undoped BYN system and also reported activation energies of Ea  0.12 eV at temperatures below 25 °C [17]. Thus, the current activation energies of Table 2, measured at much higher temperatures, concur with those at lower temperature, which have previously been suggested to be due to a small polaronic electron hopping conduction mechanism where electrons are localized due to the presence of local distortions in the crystal lattice sites [28]. Such values of activation energy are significantly lower than those which could be related oxide-ion conduction. This concurs with the absence of direct O–O bond linkages between NiO6 chains in the BYN crystallographic structure; a factor which would be expected to hinder potential oxide-ion conduction. Protonic conductivity can also be negated due to the water vapour partial pressure independence of the conductivities demonstrated in Figs. 5 and 6.

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Table 2 Activation energy of Y2xSrxBaNiO5 (X = 0, 0.1 and 0.2) system in O2 and N2 atmospheres. X

Activation Energy, Ea (eV) (500–100 °C) Oxygen

0 0.1 0.2

Nitrogen

Dry

Wet

Dry

Wet

0.216 0.221 0.233

0.216 0.227 0.232

0.223 0.223 0.234

0.224 0.229 0.240

3.3.1. Conductivity dependence on oxygen partial pressure (pO2) Fig. 7 plots the oxygen partial pressure (pO2) dependence of conductivity of compositions within the solid solution, at temperatures in the range 600–700 °C. No significant variations in conductivity can be observed with changing pO2 under these conditions, for any composition. The conductivities of SBYN1 and SBYN2 are shown to be higher than SBYN0, as expected from the results of Figs. 5 and 6. If charge compensation in the generic material, Y2xSrxBaNiO5, would exclusively determined by ionic defects, the substitution of Y for Sr would be described in Kröger-Vink notation as

BaO þ ð2  xÞYO1:5 þ xSrO þ NiO x

x ! BaxBa þ xSr=Y þ ð2  xÞYxY þ NiNi þ 0:5xV O þ ð5  0:5xÞOO

ð1Þ

However, R. Castañer et al., using X-ray absorption spectroscopy, highlighted that acceptor doping can also influence the valence state of nickel cation in BYN, leading to a mixed Ni2+ to Ni3+ valence state [28]. This alterative charge compensation mechanism can be expressed by the equilibrium of oxygen vacancies and Ni valence with the surrounding atmosphere, described by the equation K ox 1 x  x V O þ O2 ðgÞ þ 2NiNi () OO þ 2NiNi 2

ð2Þ

With mass action constant: Fig. 5. Conductivity of Y2xSrxBaNiO5 (x = 0, 0.1 and 0.2) measured in N2 gas atmosphere.



K ox ¼

½OxO ½NiNi  x

2

ð3Þ

2

1=2 ½NiNi  ½V O pO2

Under oxidising conditions the associated electroneutrality condition can be simplified to: 

bNiNi c þ 2bV O c  bSrY c =

ð4Þ

leading to the mass action relation

Fig. 6. Conductivity of Y2xSrxBaNiO5 (x = 0, 0.1 and 0.2) measured in O2 gas atmosphere.

Fig. 7. Conductivity of Y2xSrxBaNiO5 (X = 0, 0.1 and 0.2) measured as a function of oxygen partial pressure (pO2) from 700 to 600 °C temperature range.

96

K ox ¼

N. Nasani et al. / Journal of Alloys and Compounds 620 (2015) 91–96  2 ½OxO ½NiNi  2 x  ½NiNi  12 ð½Sr=Y   ½NiNi Þp1=2 O2

ð5Þ

If all charge compensation were by the formation of oxygen vacancies, ½Sr =Y   ½V O   constant, and Eq. (5) would suggest an  oxygen partial pressure dependence of ½NiNi  of pO1/4 2 . On the other hand, if charge compensation were predominately by the formation of Ni3+, then the electroneutrality equation (Eq. (4)) would  be reduced to ½Sr =Y   ½NiNi , signifying an effective independence  of the concentration of the charge carrier ½NiNi  on changing pO2. The observed pO2 independence of conductivity in Fig. 7, therefore, best relates to the latter hypothesis, of a predominant charge com pensation mechanism of formation of Ni3+ where the ½NiNi  concen= tration would be fixed by the level of acceptor doping, ½Sr Y , rather than the alternative charge compensation mechanism of formation of oxygen vacancies. This conclusion is in agreement with that suggested by Alonso [28] and Fagot–Revurat [12] by neutron diffraction, infrared absorption spectra and photo emission experiments performed at lower temperatures. The current results support the continued presence of Ni3+ at higher temperatures that lead to the enhanced conductivity noted in the Sr-substituted compositions in Figs. 5–7 due to thermally induced small polaronic hopping between Ni2+ and Ni3+.

4. Conclusions The Y2xSrxBaNiO5 (x = 0, 0.1, 0.2 and 0.3) series was synthesized by a combination of mechanical activation and solid state synthesis. The unit cell parameters and lattice volume are shown to decrease with increasing of strontium content up to composition x = 0.2. The expansion of the lattice at composition x = 0.3 suggests this composition to be beyond the solubility of Sr in the BYN lattice. The electrical conductivity was measured in wet and dry O2 and N2 containing atmospheres in the temperature range 700–100 °C. The results demonstrate that an increased conductivity of one and half orders magnitude can be obtained by Sr2+ substitutions in Y2BaNiO5 beyond that of undoped Y2BaNiO5. The low activation energy 0.216–0.240 eV for conduction, suggests a thermally activated small polaronic hopping mechanism, while negating the presence of oxide-ion conductivity. The conductivity was shown to have no appreciable dependence on atmosphere, either with water vapour or oxygen partial pressure. This result concludes the absence of protonic conductivity, while the absence of a pO2 dependence of conductivity suggests that the principal charge compensation mechanism for Sr substitution is that of the formation of Ni3+.

Acknowledgements The authors acknowledge kind support from the FCT, FEDER and COMPETE, PTDC/CTM/100412/2008, PTDC/CTM/105424/2008, and Narendar Nasani is grateful to FCT for doctoral research Grant SFRH/BD/80949/2011. Isabel Antunes is thankful to the FCT for doctoral research grant SFRH/BD/76738/2011. References [1] F. Haldane, Phys. Rev. Lett. 50 (1983) 1153. [2] M. Yamashita, T. Ishii, H. Matsuzaka, Coord. Chem. Rev. 198 (2000) 347. [3] E. Garcia-matres, J.L. Martinez, J. Rodriguez-Carvajal, J.A. Alonso, A. Salinassanchez, R. Saez-puche, J. Solid State Chem. 103 (1993) 322. [4] A. Salinas-Sdnchez, R. Sdez-Puche, J. Rodrfguez-Carvajal, J.L. Martinez, Solid State Commun. 78 (1991) 481. [5] E. Garcia-Matres, J. Rodriguez-Carvajal, J.L. Martinez, J.A. Alonso, A. SalinasSanchez, R. Saez-Puche, Solid State Ionics 65 (1993) 915. [6] E. Garcia-Matres, J.L. Garcia-Mufioz, J.L. Martinez, J. Rodriguez-Carvajal, Phys. B 194–196 (1994) 193. [7] K. Singh, T. Basu, S. Chowki, N. Mahapotra, K.K. Iyer, P.L. Paulose, E.V. Sampathkumaran, Phys. Rev. B 88 (2013) 094438. [8] K. Kordonis, A. Sologubenko, T. Lorenz, S.-W. Cheong, A. Freimuth, Phys. Rev. Lett. 97 (2006) 115901. [9] R. Castañer, C. Prieto, A. de Andrés, J.L. Martínez, R. Sáez-Puche, J. Alloys Comp. 210 (1994) 31. [10] F.-X. Lannuzel, E. Janod, C. Payen, G. Ouvrard, P. Moreau, O. Chauvet, P. Parent, C. Laffon, J. Alloys Comp. 317–318 (2001) 149. [11] J. Das, A. Mahajan, J. Bobroff, H. Alloul, F. Alet, E. Sørensen, Phys. Rev. B 69 (2004) 144404. [12] Y. Fagot-Revurat, D. Malterre, F.-X. Lannuzel, E. Janod, C. Payen, L. Gavioli, F. Bertran, Phys. Rev. B 67 (2003) 125118. [13] P. Novák, F. Boucher, P. Gressier, P. Blaha, K. Schwarz, Phys. Rev. B 63 (2001) 235114. [14] D. Capsoni, M. Bini, V. Massarotti, P. Galinetto, Solid State Commun. 122 (2002) 367. [15] R. Saez-puche, J.M. Coronado, C.L. Otero-diaz, J.M.M. Llorente, J. Solid State Chem. 93 (1991) 461. [16] J. Darriet, L.P. Regnault, Solid State Commun. 86 (1993) 409. [17] J.W. Chen, G. Narsinga Rao, K.W. Li, J. Appl. Phys. 111 (2012) 064111. [18] J. Tong, D. Clark, M. Hoban, R. O’Hayre, Solid State Ionics 181 (2010) 496. [19] J. Tong, D. Clark, L. Bernau, M. Sanders, R. O’Hayre, J. Mater. Chem. 20 (2010) 6333. [20] A. Magrasó, Z. Li, S. Ricote, N. Bonanos, A. Manerbino, W.G. Coors, Int. J. Hydrogen Energy 37 (2012) 7954. [21] J. Tong, D. Clark, L. Bernau, A. Subramaniyan, R. O’Hayre, Solid State Ionics 181 (2010) 1486. [22] N. Narendar, G.C. Mather, P.A.N. Dias, D.P. Fagg, RSC Adv. 3 (2013) 859. [23] L. Bi, E. Fabbri, Z. Sun, E. Traversa, Energy Environ. Sci. 4 (2011) 1352. [24] I. Antunes, A. Brandão, F.M. Figueiredo, J.R. Frade, J. Gracio, D.P. Fagg, J. Solid State Chem. 182 (2009) 2149. [25] J. Rodríguez-Carvajal, Phys. B Condens. Matter 192 (1993) 55. [26] D.J. Buttrey, J.D. Sullivan, A.L. Rheingold, J. Solid State Chem. 88 (1990) 291. [27] R.D. Shannon, Acta Crystallogr. Sect. A 32 (1976) 751. [28] J.A. Alonso, I. Rasines, J. Rodriguez-Carvajal, J.B. Torrance, J. Solid State Chem. 109 (1994) 231. [29] J. DiTusa, S. Cheong, J. Park, G. Aeppli, C. Broholm, C.T. Chen, Phys. Rev. Lett. 73 (1994) 1857.