New Insights Into the Effect of Nonstoichiometry on the Electric

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C: Plasmonics; Optical, Magnetic, and Hybrid Materials

New Insights Into the Effect of Nonstoichiometry on the Electric Response of Strontium Titanate Ceramics Luís Amaral, Alexander Tkach, Paula M. Vilarinho, and Ana Maria R. Senos J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.8b10438 • Publication Date (Web): 18 Dec 2018 Downloaded from http://pubs.acs.org on December 23, 2018

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The Journal of Physical Chemistry

New Insights into the Effect of Nonstoichiometry on the Electric Response of Strontium Titanate Ceramics

Luís Amarala,b, Alexander Tkacha, Paula M. Vilarinhoa*, Ana M. R. Senosa a

Department of Materials & Ceramic Engineering, CICECO – Aveiro Institute of

Materials, University of Aveiro, 3810-193 Aveiro, Portugal b Center

of Physics and Engineering of Advanced Materials (CeFEMA), Instituto Superior

Técnico, Universidade de Lisboa, 1049–001 Lisbon, Portugal (current address)

* Corresponding author: E-mail: [email protected] Tel: +351 234 370 354/259; Fax: +351 234 370 204

Abstract The effect of nonstoichiometry (Sr/Ti ratio from 0.995 to 1.02) on the bulk and grain boundary contributions to the electrical response of strontium titanate (STO) ceramics is investigated. Nonstoichiometric STO exhibits lower electrical resistivity than its stoichiometric counterpart. This decrease is systematic for both Ti and Sr-excess and with a greater effect on the grain boundary resistivity compared to bulk. Moreover, systematic variations with the degree of nonstoichiometry are observed for bulk and grain boundary conductivity, activation energy and capacitance. These changes are correlated to a high concentration of point defects and induced by the nonstoichiometry, increasing the charge carrier concentration. As a result, nonstoichiometry can be used to tailor the microstructure and properties of strontium titanate, particularly grain boundary properties.

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1. Introduction Strontium titanate (SrTiO3, STO) is an important material from both the technological and scientific perspective. STO is an incipient ferroelectric (or quantum paraelectric),1 in which the real part of the dielectric permittivity, ε', follows the Curie-Weiss law in the paraelectric regime but saturates at cryogenic temperatures rather than exhibiting the expected ferroelectric discontinuity. When the Curie temperature, Tc, is close to 0 K, the ferroelectric order is suppressed by quantum fluctuations. Characterized by high dielectric permittivity, high tunability of ε' by dc electric field and low dielectric losses,2 STO is particularly interesting for capacitor and tunable microwave devices.3 Additionally, the electrical properties of the ceramic grain boundaries (GBs), rather than those of grain interiors, are of major importance in barrier-layer devices based on STO.4 On the other hand, such blocking GBs are highly disadvantageous if fast transport in the electroceramic material is required, as for example in STO-based thermoelectric applications.5,6 Indeed, increasing attention has been paid to the GB structure of STO ceramics and its relation to the microstructure development7 and electrical properties.8,9 Our previous studies revealed that defects in ceramic bulk and GBs related to the stoichiometry play an important role in the sintering kinetics of STO for both densification and grain growth.10,11 Correlation between grain growth, GB mobility and defect concentration was recently highlighted by sintering STO under weak electric fields.12 Grain growth acceleration due to defect redistribution under weak electric fields was reported. Enhanced boundary mobility near the negative electrode was attributed to increased concentration of migrating oxygen vacancies, while strontium vacancies accumulate at the positive electrode. According to the authors the electric field itself does not change the grain growth in STO, but rather affects the local defect concentration.12 The broad range of GB properties has been recognized for a long time and new


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developments in their understanding are still being debated. Dillon et al.13 introduced the concept of interface complexion, in which different interface properties in the same material were associated with chemical/structural differences in the interface regions. Different complexions depend for example on material processing conditions and can largely control the behavior of interfaces such as GBs,14,15 with a consequently strong impact on the materials properties. Four different GB complexions have been identified in STO,16 correlated with different mobilities. Changes in the GB faceting behavior at high temperatures were considered to be responsible for sudden drops in the GB mobility of stoichiometric STO with increasing temperature.17 Later, we established for Ti-rich STO for the first time the relation between these GB mobility discontinuities and electrical properties, by correlating the grain growth regimen transitions to abrupt changes in the activation energy for GB conductivity.18 GB electrical response was much more affected by the sintering temperature than that of the bulk counterpart, strongly reinforcing the idea of the key role played by the GBs in the grain growth anomaly observed for Ti-rich STO.18 The anomaly was found to consist of four grain growth regimens separated by sharp grain size decreases despite the increasing sintering temperature.18 The anomalous behavior was related to a first order reverse change of complexions driven by the temperature effect on the GB wettability and electrostatic potential.18 Moreover, after annealing these ceramics in oxygen atmosphere, an increase and levelling-off in the GB conductivity activation energy and a four-fold decrease in the GB permittivity were observed, thus, confirming a key role of oxygen vacancies in the relation between the grain growth and conductivity anomalies of as-sintered Ti-rich ST ceramics.19 In spite of the changes observed mainly at the GB, our previous results demonstrated that the variation of the Sr/Ti ratio has only a weak effect on intrinsic quantum paraelectric



behavior and no dielectric anomaly is observed at low temperatures.11 However, the relation between nonstoichiometry, GB electrical properties and microstructure development has not yet been addressed, although it was found to be very important for instance in Na0.5Bi0.5TiO3 ceramics.20 Within this context, the premise of this work is to establish the effect of nonstoichiometry on the electrical properties of STO at elevated temperatures. The electrical response of stoichiometric and nonstoichiometric STO is separated into the contributions from the intrinsic bulk and extrinsic grain boundaries using impedance spectroscopy (IS)21,22 and correlated to the microstructure and GB properties.

2. Experimental Ceramics of STO compositions with Sr/Ti ratios of 0.995, 0.997, 1.000, 1.005, 1.010 and 1.020 (hereafter named as ST 0.995, ST 0.997, ST, ST 1.005, ST 1.01 and ST 1.02, respectively) were prepared by conventional mixed oxide method. Raw powders of SrCO3 and TiO2 (pro analysis Merck, purity ≥ 99%) were ball milled in alcohol in a planetary mill for 5 h, using Teflon pots with zirconia balls. After calcination, at 1100 ˚C for 2 h, the powders were monophasic. The crystallographic structure and phase content of STO powders were analyzed by X-ray diffraction (XRD, PANalytical X’Pert Pro diffractometer, Cu-Kα radiation, 45 kV and 40 mA), used from 10° to 50° 2θ with a step size of 0.025°. The particle size distribution of the calcined powders was obtained using laser diffraction system (Coulter LS230). Powders were milled again under the same conditions, uniaxially pressed at 50 MPa and further isostatically pressed at 200 MPa. The pressed samples of all the compositions were then heated at 5 ˚C min-1 up to 1500 ˚C, kept at that temperature for 5 h, and cooled at 5 ˚C min-1 to room temperature inside


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the furnace. After the sintering cycle, all the samples attained relative densities above 96%, as measured by the Archimedes method. Polished and thermally etched sections of sintered samples were observed by scanning electron microscopy (SEM, Hitachi S-4100). Using ImageJ software, the grain size distribution of the sintered samples was determined, taking more than 600 grains in at least three SEM micrographs. The area of the section of the grains was measured and then its circular equivalent diameter calculated. The average grain size, G, was determined from the average equivalent diameter, by using a multiplying factor of 1.22. The error bars were derived from the standard deviation of the values measured in the different images. For transmission electron microscopy (TEM, Hitachi 9000) observations, the two faces of the bulk samples were polished using a Gatan disc grinder in order to reduce its thickness to ~30 μm. The sample was then glued to a copper ring and ion beam polished using a precision ion polishing system (Gatan 691). Bottom and top surfaces of the samples were then polished and silver paste electrodes were painted on both surfaces. IS measurements were carried out using precision LCR meter (HP 4284A) between 100 Hz and 1 MHz in a temperature range from 200 to 700 ˚C. The collected impedance data were normalized by multiplying by the geometric factor A/d (A is the area of the electrode and d the thickness of the sample) and analyzed with the software ZView (Scribner Associates Inc.).

3. Results and discussion 3.1. The electrical response The impedance spectra of the studied ceramics show two main contributions: a high frequency low capacitance contribution of the order of 10-11 F/cm assigned to the bulk; and a low frequency high capacitance contribution of the order of 10-9 - 10-8 F/cm, which



is often found in well sintered ceramics with narrow intergranular regions and attributed to the GBs.21 Additionally, at high temperatures, e.g. 700 ˚C, the onset of the electrode interface relaxation process could be seen in some cases, with an associated capacitance of the order of 10-6 F/cm. Thus, the impedance spectra obtained in this work are modeled with an equivalent circuit consisting of a series of three blocks of a resistor and a constantphase element (CPE) in parallel, related to bulk, GB and sample-electrode interface contributions, respectively.21 However, we should note here that the electrode-sample interface contribution is generally residual and should be more pronounced at even higher temperature or lower frequency. As presented in Figure 1, semicircles regarding the two main contributions of bulk and GBs are resolved in the complex impedance spectra of all the ceramics under study. In the measured frequency range, the semicircles regarding the bulk relaxation are observed between approximately 200 and 500 ˚C, whereas the semicircles concerning the GBs are present between approximately 450 and 700 ˚C. Figure 1a shows the Nyquist plots (imaginary Z'' vs. real part Z' of the complex specific impedance) at 300 ˚C for all the ceramics under study. At this temperature and in the measured frequency range (102-106 Hz), the semicircles related to the GB contribution are not completely resolved and only a small low frequency peak appears. An important observation is that the well resolved bulk semicircles change perceptibly with small stoichiometric variations between the samples. A decrease of the bulk resistivity (smaller semicircles) is evident for all the nonstoichiometric ceramics relative to stoichiometric.


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Figure 1. Impedance spectra at (a) 300 ˚C and (b and c) 600 ˚C for STO ceramics with Sr/Ti ratios from 0.995 to 1.02 sintered at 1500 ˚C for 5 h. (c) shows an enlargement of (b) near the origin. The points regarding the frequencies of 102, 103, 104, 105 and 106 Hz are marked with full symbols for all the curves. It is clear that the impedance spectra are strongly affected by the nonstoichiometry of STO.



GB semicircles are well resolved and dominate the impedance spectra at 600 ˚C, as can be seen from Figure 1b,c, allowing assessment of the variation in GB with induced nonstoichiometry. Only for ST 1.02, the observed semicircle is poorly resolved at this temperature, although the GB semicircles are detected in the measured frequency range and fitted at lower than ~550 ˚C. Compared to stoichiometric ST, smaller semicircles are evident for nonstoichiometric ceramics in Figure 1b and more strikingly in Figure 1c, which correspond to the data in Figure 1b zoomed near the origin. Only by enlarging the region near the origin, can semicircles for higher degrees of nonstoichiometry be made apparent, highlighting the major impact of the small stoichiometric variations on the GB properties. Moreover, this effect is systematic; the size of the semicircles decreases proportionally to the degree of nonstoichiometry for both types of excess. Furthermore, from the similarity of the semicircles for ST 0.995 and ST 1.01, one can conclude that the effect of Ti-excess on decreasing GB resistivity is about twice that of Sr-excess. A quantitative analysis of the variation of the bulk (at 300 ˚C) and GB (at 600 ˚C) resistivity with the Sr/Ti ratio is provided in Figure 2, which presents the values obtained by fitting the complex impedance spectra of Figure 1. The decrease in bulk and GB resistivity with the nonstoichiometry is evident. The magnitude of the decrease of GB resistivity is much higher, however, thus requiring the use of a log scale to be fully displayed. Regarding the bulk resistivity, shown in Figure 2a for Sr/Ti ratios between 0.995 and 1.010, the decrease is nearly proportional to the degree of nonstoichiometry for both Ti and Sr excess. The bulk resistivity of Ti-rich ST 0.995 is 7 times lower than that of the stoichiometric ceramics, whereas a decrease by a factor of 9 is observed for Sr-rich ST 1.01 (which has a degree of nonstoichiometry twice that of ST 0.995). On the other hand, the same degree of nonstoichiometry (0.5 mol.%)


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in Ti-rich ST 0.995 and Sr-rich ST 1.005 leads to a more significant decrease of the bulk resistivity in the case of Ti excess: factors of 7 and 2 were observed, respectively. The bulk resistivity of Ti-rich ST 0.997 ceramics decreases by factor of 1.5 comparing with stoichiometric ST. Regarding Sr/Ti ratio of 1.020 (ST 1.02), the bulk resistivity is also considerably decreased (by a factor of 4). However, this decrease is smaller than that of ST 1.01, which has half the degree of nonstoichiometry.

Figure 2. (a) Bulk and (b) GB resistivity dependence on the Sr/Ti ratio at 300 and 600 ˚C, respectively, showing a strong decrease for nonstoichiometric STO ceramics. Note the use of a log scale in (b) due to the much higher decrease of the GB resistivity.

Similarly to the effect on the bulk resistivity but with higher magnitude, a very significant impact of the small nonstoichiometric variations is clearly observed in the GB resistivity (see Figure 2b). Indeed, a decrease by factors of 4 and 17 is found in the GB resistivity ACS Paragon Plus Environment


of ST 0.997 and ST 1.005 compositions, correspondingly. A more dramatic change occurs in the GB resistivity of the Ti-rich composition ST 0.995, which diminishes 55 times in relation to the stoichiometric one, while a factor of 59 is observed in the case of the Srrich ST 1.01.

Figure 3. (a) Bulk and (b) GB conductivity dependence on the temperature. Grain boundaries are more affected than bulk by nonstoichiometry.

The dependence of the bulk and GB conductivity in logarithmic scale on inverse temperature is found to be linear, as presented in Figure 3, expanding the observations remarked at 300 and 600 ˚C to a wider temperature range. It can be seen that both bulk and GB conductivity increase with nonstoichiometry in a nearly systematic way, in


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agreement with the decrease of the resistivity previously observed. Additionally, the temperature, at which the GB contribution is detected in our measurement frequency range, is lowered by the nonstoichiometry. Furthermore, the lines for bulk conductivity of all the samples in Figure 3a are almost parallel, showing similar variation with inverse temperature. On the other hand, strong changes in the line slope are observed for the GB conductivity in Figure 3b. This variation in the slope is reflected in Figure 4, which presents the activation energy for bulk and GB conductivity obtained for all the Sr/Ti ratios.

Figure 4. Activation energy for bulk and GB conductivities as a function of the Sr/Ti ratio in STO ceramics. Bulk activation energy is nearly independent of the composition; grain boundaries are markedly dependent on the stoichiometry, revealing different dependences for Sr and Ti-excess.

Despite the fact that the resistivity of both bulk and GBs strongly decreases for nonstoichiometric ceramics, the activation energy for conduction of bulk and GBs reveals very different dependences on the composition. The values for the bulk conductivity are practically independent of the composition, showing a very slight tendency to increase with the Sr/Ti ratio. All the activation energy values for the bulk conductivity are in the range 0.87 ≤ Ea,B ≤ 1.00 eV, which is in agreement with previously reported values for



the p-type behavior of undoped and stoichiometric STO23 and acceptor-doped STO,24-28 being as well consistent with ionic charge transport attributed to oxygen vacancies.27,29 On the other hand, the activation energy for the GB conductivity shows a remarkable variation with the Sr/Ti ratio and very different behavior for Sr and Ti excess. In fact, nonstoichiometry lowers the GB activation energy in every case, when compared to that of the stoichiometric STO (Ea,GB = 1.73 eV). However, for Ti-rich compositions the decrease is proportional to the amount of Ti-excess, whereas in the case of Sr excess, the lowest degree of nonstoichiometry investigated (ST 1.005) induces the largest decrease in the activation energy (Ea,GB = 0.37 eV), which then increases for higher fractions of Sr excess. Furthermore, the same degree of nonstoichiometry (0.5 mol.%) leads to a much stronger decrease of the GB activation energy in the case of Sr excess (ST 1.005, Ea,GB = 0.37 eV) than in the case of the Ti-rich ceramics (ST 0.995, Ea,GB = 1.48 eV). Typical values of activation energy for the GB conduction are within the range of 1.4 – 1.6 eV, as predicted for acceptor doped STO27 and confirmed for Fe-doped bi-crystals30 as well as for undoped25 or Fe-doped STO ceramics.31 The value reported for Ni-doped STO24,27 is also close to half the bandgap, Wg/2 ≈ 1.6 eV. This activation energy is associated to a charge transport across the GB space charge layers with “W”-shaped spatial conductivity profile.27 Such profile forms when partial electron conductivity exceeds that of holes at the GB core due to enhanced concentration of electrons. For lower potential barriers and/or partial electron conductivity, the GB conductivity profile may show a “V”-shape, dominating at low temperatures, with an associated activation energy around 1.0 eV, which is combined with the ionization energy of the bulk acceptor.27 In this context, the activation energy for GB conductivity found for Ti-rich ST 0.995 and ST 0.997 (1.59 eV), as well as stoichiometric ST (1.78 eV) and Sr-rich ST 1.02 (1.51 eV), is in agreement with the previously reported values for STO and consistent with a “W”-type


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GB conductivity profile, despite the fact that the activation energy of the stoichiometric composition is slightly higher. On the other hand, the activation energies for ST 1.005 and ST 1.01 (1.11 eV) are lower, especially that of 0.37 eV only for ST 1.005, which suggest a “V”-shaped profile. It is important to consider here that some microstructural features, such as a non-uniform grain size distribution, may originate deviations from the ideality of the brick layer model and influence the values for the activation energy of GB conductivity obtained from the impedance spectra.25,32 Nevertheless, the systematic variation of the activation energy concerning both Sr and Ti excess strongly suggests an effective direct relation between nonstoichiometry and activation energy for GB conductivity. Concerning the values of bulk capacitance, they are nearly independent of the temperature, as expected for STO at elevated temperatures.24,25 A similar nearly temperature-independent behavior is observed for the GB capacitance. Therefore, the average capacitance values measured at the several temperatures are presented in Figure 5 for all the ceramics. It is again evident that the GBs are more affected by nonstoichiometry than the grain interiors. Systematic variation of GB capacitance with the degree of nonstoichiometry is obvious for both Sr and Ti excess. Regarding Ti-rich compositions, increasing the degree of nonstoichiometry leads to a steady increase of GB capacitance. On the other hand, Sr-rich compositions reveal the highest capacitance for ST 1.005 with further monotonous marked decrease to a value close to that of the stoichiometric composition for ST 1.02.



Figure 5. Bulk and GB capacitance as a function of the Sr/Ti ratio. Grain boundaries are more affected by nonstoichiometry of STO ceramics than grain interiors.

3.2. Microstructure and defect chemistry The results above presented are a clear indication of the role of nonstoichiometry on the electrical behavior of STO and its importance at the GB level. Their comprehension requires knowledge and understanding of the microstructure and defect chemistry of these STO compositions. In spite of the stoichiometry variations between the compositions under study, the XRD analysis of the calcined powders showed only peaks consistent with the cubic crystallographic structure of STO, as presented in Figure 6a. No second phases were detected in any of the compositions, within the detection limits of the XRD technique. The particle size distribution of the calcined powders (see Figure 6b), analyzed by laser diffraction (Coulter), was found to be equivalent for all the compositions as well. A bimodal size distribution could be observed in all the cases, with peaks around 0.2 and 2-3 µm. The first peak of the distributions is always below 1 µm whereas the second one corresponds mainly to agglomerates in the powder suspensions that are destroyed after pressing.10 On this ground, we do not expect differences in the microstructure


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development of STO nonstichiometric ceramics triggered by the initial morphology of the STO powders.

Figure 6. (a) XRD patterns of calcined powders of stoichiometric and nonstoichiometric STO compositions, presenting however only the peaks corresponding to the STO phase. (PDF card #84-0444 is also shown). (b) Particle size distribution of the calcined powders with stoichiometric and nonstoichiometric STO compositions, presenting similar bimodal size distributions.

The microstructures of the sintered STO ceramics with Sr/Ti ratios from 0.995 to 1.020 obtained by SEM are illustrated in Figure 7a. After the sintering step at 1500 ˚C, dense



ceramics with coarse microstructures and enlarged grain size distribution (see Figure 7b) were observed for Ti-rich compositions, despite the fact that an increase in the number of small grains was found in ST 0.995, corresponding to the largest Ti excess. By its turn, Sr-rich ceramics reveal finer microstructures (Figure 7a) and narrower grain size distribution (Figure 7b).

Figure 7. (a) SEM micrographs, (b) equivalent grain diameter distribution and (c) average grain size as a function of the Sr/Ti ratio of STO ceramics with Sr/Ti ratios from 0.995 to 1.02, sintered at 1500 ˚C for 5 h. Nonstoichiometry has a strong effect on the microstructure development of STO ceramics. Sr-rich samples present smaller grain size while Ti-excess leads to enlarged grain size distributions.

The average grain size variation with the Sr/Ti ratio is presented in Figure 7c. Relatively to stoichiometric ST, 0.3 mol.% excess of Ti (ST 0.997) leads to an increase of the


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average grain size, which is followed by a strong decrease for ST 0.995, with a 0.5 mol.% Ti excess. Such decrease can be related to the reported grain growth anomaly in ST 0.995, which results in GB mobility drops with increasing temperature.17,18 Conversely to ST 0.997, the average grain size observed for Sr-rich ST 1.005 and ST 1.01 ceramics is always smaller than that of the stoichiometric composition, being similar to that of Tirich ST 0.995. With the exception of the highest Ti-excess case, the general effect of the nonstoichiometry on the microstructure development is in agreement with previous reports on increased mass transport with decreasing Sr/Ti ratio.10,11,33 The strong variations found in the sintered microstructures can arise from changes in the defect chemistry induced by the several degrees of nonstoichiometry10,11 as well as from the presence of liquid phase18 traces consistent with the eutectic point at 1440 ˚C for Ti-rich compositions.34 Then the predominance of different interface complexions13,16 and, consequently, different interface kinetics is expected. TEM micrographs of ST 0.995, ST and ST 1.02 sintered at 1500 ˚C, shown in Figure 8, reveal rather rough grain boundaries in all the samples, although sharp edges suggesting GB faceting can be detected in the Tirich (Figure 8a) and stoichiometric ones (Figure 8b).33 Moreover, no Ti precipitates were observed in the Ti-rich samples but traces of an amorphous phase could be detected along the GBs and in some triple points of Ti-rich (see Figure 8a) and stoichiometric ceramics.10,18,33 No diffraction contrast emerges upon sample tilting in the TEM, in agreement with the amorphous nature of this phase, which is further confirmed by the diffraction halo in the diffraction pattern of the triple point presented in the inset in Figure 8a. By its turn, Sr-rich ST 1.02 (Figure 8c) shows a pattern of planar defects inside the grains, characteristic of Ruddlesden-Popper (RP) structures, a three-dimensional mosaic of single-layered rock-salt blocks with the formula SrO∙(SrTiO3)n.10,33,35



Figure 8. TEM micrographs of samples (a) ST 0.995, (b) ST and (c) ST 1.02 sintered at 1500 ˚C. The inset in (a) is the electron diffraction pattern from the quadrupole pocket, where amorphous phase is observed. Sharp edges at the GBs are observed for ST and ST 0.995. ST 1.02 shows Ruddlesden-Popper (RP) phases.

The formation of the RP structures in ST 1.02 is consistent with the higher bulk resistivity of this composition comparing with that of ST 1.01, as shown in Figure 2a. The RP planar structures accommodate some of the Sr excess and in this way can decrease the amount


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of defects present in the lattice, leading to a smaller decrease in the bulk resistivity relatively to the stoichiometric composition than that observed for ST 1.01. Furthermore, the different microstructures observed by SEM (Figure 7) and TEM (Figure 8), along with the variations of the GB resistivity (Figure 2b), activation energy (Figure 4) and GB capacitance (Figure 5), clearly show the strong impact of nonstoichiometry on the GBs of STO ceramics. The stronger effect on the GBs as compared to the bulk response can be attributed to the formation of different GB complexions with consequent variations in mobility, as previously reported,10,11,16-18 and hence different electrical properties.18 It should also be noted that the GB resistivity is usually expected to rise with the grain size decrease due to increased number of boundaries across the sample.26 However, in the present study all the nonstoichiometric samples show a decrease of the GB resistivity (Figure 2b), irrespective of the smaller grain size for most of them (Figure 7). This observation confirms the strong impact of nonstoichiometry on the defect chemistry of the material and suggests that other factors may play an important role concerning the electrical properties of STO ceramics. These factors include the formation of liquid phase during sintering of Ti-rich samples, as shown in Figure 8a, or segregation to grain boundaries, leading to different GB properties. On the other hand, the presence of RP structures on Sr-rich compositions, as presented in Figure 8c, must also be considered. Particularly regarding GB faceting, its correlation with the GB resistance has been reported by Lee et al.8 Annealing a ST bi-crystal in air at 1600 ˚C renders the faceted boundary to become defaceted. In this case a previously absent second GB semicircle then appears in the impedance spectra, indicating a strong increase in GB resistance.8 The rough GB observed by the authors above 1600 ˚C contained more oxygen vacancies than



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the faceted boundary present bellow that temperature, which were considered to induce a double Schottky barrier and obstruct the charge-carrier transport across the GB. In the present work, stoichiometric polycrystalline material and those with excess of both types of oxides presented bulk and GB impedance responses. Additionally, irrespective of the predominance of different GB structures suggested by the TEM analysis, with rough grain boundaries being generally observed and edges pointing to GB faceting being detected in the Ti-rich and stoichiometric ceramics,10,33 both types of oxide excess originated strong decrease of GB impedance contributions in comparison with the stoichiometric composition. It is, therefore, suggested that the shape of the grain boundaries alone (faceted or defaceted) may not play a crucial role in determining the GB resistivity. Instead, the presence, concentration and distribution of defects induced by nonstoichiometry is proposed to be more determinant for the GB resistivity. In this context, the possible alterations of the defect chemistry in STO induced by nonstoichiometry must be considered. As previously mentioned, the redistribution of the charged defects is known to lead to the formation of a space charge layer.24,36,37 In strontium titanate, the GB core is positively charged whereas the regions close to the grain boundaries are depleted of mobile positive charge carriers such as oxygen vacancies and holes. On the other hand, a negative space charge forms on both sides of the GB due to accumulation of compensating negative charges.36 Moreover, as observed by Baurer et al.,16 this results in local nonstoichiometry at the GB region (GB core and space charge) that can be different from that of the bulk composition (Ti or Sr- rich). The incorporation of TiO2-excess into the STO lattice requires the formation of strontium and oxygen vacancies35,38 according to Eq. (1).

TiO2  VSr''  TiTi  2OO  VO


(1)

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On the other hand, the incorporation of Sr-excess occurs through the formation of the RP structure, but small amounts of Sr excess can be incorporated by the formation of titanium and oxygen vacancies (Eq. (2)) or by a cation place-exchange (anti-site defect) and oxygen vacancies, as in Eq. (3).35,38

SrO  SrSr  VTi''''  OO  2VO

(2)

2 SrO  SrSr  SrTi''  2OO  VO

(3)

Therefore, as a result of the alterations on the defect chemistry induced by nonstoichiometry the overall concentration of defects, such as strontium vacancy V′′𝑆𝑟, ˙˙ titanium vacancy V′′′′ 𝑇𝑖 and oxygen vacancy V𝑂 is increased. Ionization of these vacancies

is known to yield such charge carriers as free electrons and holes, which concentration is thus increased, decreasing the bulk and GB resistivity. Moreover, assuming that ceramic grains are terminated by Ti4+ ions, which are the cause of the GB space charge layer depleted by holes but reach in electrons,36 and that Ti vacancies can be formed at small amount of Sr excess, one can explain the sharp variation of the activation energy comparing ST and ST 1.005 ceramics. Indeed, the formation of Ti vacancies or the lack of Ti4+ ions at GB results in strong decrease of the charge transport potential barriers for holes as well as in suppression of electron concentration at the GB. Therefore, the GB conduction profile changes from “W”-shaped to “V”-shaped one with lower potential barrier as it is concluded from Figure 4. Alfthan et al. have also suggested that an increase in the Sr/Ti ratio lowers the potential barrier at GB core, and hence decreases GB resistance.39 On further increase of the Sr excess, however, the formation of the RP structure takes place, leveling out this effect. Additionally, the unavoidable presence of acceptor impurities should also be taken into consideration. As mentioned before, even undoped STO exhibits a specific concentration of oxygen vacancies due to the presence of acceptor impurities incorporated during



processing or already present in the raw materials.40-42 However, in the current work all the compositions were prepared following the same procedures and from the same raw materials, with only slight variations of the SrO/TiO2 ratio. Therefore, the effect of impurities should be basically equivalent in all the studied compositions and not sufficient to explain the observed dramatic changes in bulk and, especially, GB properties. This is also supported by the clear systematic variations with nonstoichiometry in both Ti and Sr-excess cases regarding the several investigated parameters. Nevertheless, the combined effects of nonstoichiometry (more important) and impurities may therefore originate a high concentration of defects in the material. At the same time, the different defects originated by Ti or Sr excess may somehow lead to the different behaviors regarding, for example, the variation of the GB capacitance for the two kinds of nonstoichiometry. On the other hand, it is important to note that, despite the similar effects on the resistivity, the investigated degrees of oxide excess are higher in the case of Sr than those of Ti-rich compositions. Finally, the accommodation of Sr excess is easier due to the formation of the RP structures, which does not occur on the Tirich side, where more defects can therefore be expected.

4. Conclusions The effect of nonstoichiometry (Sr/Ti ratio from 0.995 to 1.02) on bulk and GB contributions to the electrical response of SrTiO3 ceramics was systematically investigated and compared with the microstructure development. With the exception of the highest Ti excess, the grain size distribution was enlarged and grain size increased with the decrease of Sr/Ti ratio. However, impedance spectroscopy, being a very sensitive and powerful technique for the study of the effect of small stoichiometric variations, revealed that the resistivity of bulk and grain boundaries is systematically decreased in


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both Ti-rich and Sr-rich STO as compared to stoichiometric ceramics. The effect is much stronger for the grain boundaries as compared to the bulk. Therefore, it was suggested that defect chemistry plays much more important role in electrical response than microstructural development. Indeed, very small variations in the stoichiometry induce a high amount of defects, which ionization increases the charge carrier concentration, leading to a strong reduction of the resistivity, particularly at GB. Moreover, systematic dependences on the nonstoichiometry were also observed for the conductivity activation energy and capacitance (much more affected in the case of the GB than bulk contribution). These observations are also consistent with the strong impact of nonstoichiometry induced defect chemistry and correspondent modification of GB complexions, which lead to different microstructure evolution and diverse electrical response. Thus, this work highlights the importance of considering the effect of commonly observed defects and nonstoichiometry in oxide materials on their electrical properties. Furthermore, it demonstrates that nonstoichiometry offers a tool for microstructure tailoring and GB engineering enabling the design of the electrical properties and microstructure of ceramics.

Acknowledgments This work was developed within the scope of the project CICECO-Aveiro Institute of Materials, POCI-01-0145-FEDER-007679 (FCT Ref. UID/CTM/50011/2013), financed by national funds through the FCT/MEC and when appropriate co-financed by FEDER under the PT2020 Partnership Agreement. The authors acknowledge as well the financial support from the Foundation for Luso American Development (FLAD), Portugal. FCT is also acknowledged for financial support by Luís Amaral (SFRH/BD/40927/2007 and



SFRH/BPD/97453/2013) and Alexander Tkach (IF/00602/2013). The authors thank Rainer Schmidt for the helpful discussion of the IS results and Ian Reaney for TEM data acquisition and analysis.

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New Insights Into the Effect of Nonstoichiometry on the Electric

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