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Grain Growth Anomaly and Dielectric Response in Ti-rich Strontium Titanate Ceramics Luís Amaral,†,§ Manuela Fernandes,† Ian M. Reaney,⊥ Martin P. Harmer,‡ Ana M. R. Senos,†,* and Paula M. Vilarinho† †

Department of Materials and Ceramic Engineering, CICECO, University of Aveiro, 3810-193 Aveiro, Portugal Center for Advanced Materials and Nanotechnology, Lehigh University, Bethlehem, Pennsylvania 18015, United States ⊥ Department of Engineering Materials, University of Sheffield, Sheffield, S1 3JD, UK ‡

ABSTRACT: A grain growth anomaly in Ti-rich strontium titanate ceramics is reported. Here we show that three discontinuities on the temperature dependence of grain growth take place with drops in the grain size at temperatures around 1500, 1550, and 1605 °C. We also show that similar discontinuities can be observed in the dependence of the grain boundary activation energy for conductivity and in the grain boundary thickness, assessed by impedance spectroscopy (IS). These notable coincidences are reported for the first time and strongly support the formation of different grain boundary complexions in polycrystalline oxides with transitions in between the observed grain growth regimens, which may be correlated to different grain boundary mobility and dielectric properties. These results call into question the discussion on the role of nonstoichiometry of SrTiO3 and complexions on the microstructure development and open opportunities to design properties of functional materials.

1. INTRODUCTION A grain growth anomaly in strontium titanate (SrTiO3) was recently reported.1 Two drops in the grain boundary mobility with increasing temperature from 1200 to 1600 °C in oxygen atmosphere were observed, independently of the Sr/Ti ratio. In this temperature range, grain growth kinetics did not follow the classical Arrhenius-type temperature dependence. Changes in the faceting behavior of the grain boundaries at high temperatures were pointed as the possible explanation for the anomalous grain growth behavior observed in SrTiO3.1,2 There is a clear relevance in these findings but the overall scenario behind the grain growth anomaly is not yet known. Indeed, the correlation with grain boundary complexions3,4 is unclear and more research on this topic is missing. Moreover, no relation with the dielectric properties of the material has been established. More work on the topic is clearly needed, since the understanding of this anomaly may offer an alternative and powerful way of controlling the microstructure and tailoring the dielectric response of SrTiO3 based materials just by tuning the sintering temperature. Strontium titanate-based materials have a great interest for a wide range of applications in microelectronics, namely, in tunable microwave devices due to the high electric-field dependence of the permittivity5 and more recently as possible thermoelectric materials.6 Additionally, using SrTiO3 as a model system, the understanding of these findings may be extended to other technologically important systems and key steps can be taken in the direction of using grain boundary engineering for materials design. © 2013 American Chemical Society

The importance of the sintering cycle on the microstructure development during sintering of materials has long been recognized.7−9 For example, the activation of grain boundary or volume diffusion mechanisms was shown to be influenced by the heating rate9 or temperature.10 Moreover, it is known that the formation of interfacial or grain boundary phases and the occurrence of grain boundary “phase” transitions drastically change materials transport and physical properties.4,11 Grain boundary transitions and kinetics are often affected by interface anisotropy.12 Grain boundaries can be faceted (atomically ordered) or rough (atomically disordered) and faceting or roughening transitions may occur.13 Grain boundary faceting behavior is an important factor in determining the grain boundary mobility and, consequently, the grain growth behavior of a material.13 In correlation with faceting-roughening transitions of the grain boundaries, different regimens from abnormal to stagnant grain growth were found to occur.13,14 Grain growth studies in alumina enabled the observation of six grain boundary phases or complexions with different levels of atomic disorder, corresponding to different grain boundary mobility.4 The grain boundary phases consisted in intrinsic or clean boundaries, submonolayers, bilayers, trilayers, nanoscale intergranular films of an equilibrium thickness, and complete wetting films of an arbitrary thickness.4,15,16 The formation of these complexions has been interpreted from the interplay of Received: April 10, 2013 Revised: October 15, 2013 Published: October 30, 2013 24787

dx.doi.org/10.1021/jp4035437 | J. Phys. Chem. C 2013, 117, 24787−24795

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a better understanding of the role of grain boundaries in SrTiO3.

grain boundary premelting, prewetting, and multilayer adsorption and are dependent on the temperature, atmosphere, impurities, and stresses.4,17 The coexistence of two or more grain boundary complexions with different mobility was proposed by Harmer and coauthors as the general cause of abnormal grain growth in alumina.4 Reverse transitions of complexions with the temperature, that is, decreasing disorder of grain boundaries with increasing temperature, have been reported in some metallic systems but not in oxides.18,19 In a much lesser extent, grain boundaries of normal and abnormal grains of slightly Sr- and Ti-rich SrTiO3, sintered in oxygen between 1425 and 1600 °C, were also investigated.3 Four types of grain boundaries were observed: straight and atomically flat, straight and atomically rough, stepped, and curved, which were correlated with the boundary complexions defined in alumina.4 The formation of the different grain boundary types was attributed to nonstoichiometry at the grain boundaries.3 Ti-rich, neutral, and Sr-rich grain boundaries were detected, independently of the bulk composition. Ti-rich boundaries showed a retarding effect on grain boundary mobility, attributed to a solute drag mechanism, whereas neutral or Sr-rich (100)-facets showed the highest mobility.3 No amorphous films were reported at the grain boundaries for either normal or abnormal grains and, accordingly, the higher mobility complexions observed in alumina4 were not found. However, considering the SrO-TiO2 phase diagram,20 a liquid phase may exist above the eutectic temperature of 1440 °C on the TiO2-rich side. The presence of such a phase and its distribution during sintering are key factors determining grain boundary properties and, consequently, the overall microstructure evolution. In general and as well-known, the microstructure change during liquid phase sintering is faster than in solid state sintering because of the accelerating role of the liquid phase on the mass transport.21 In Ti-excess SrTiO3 with low donor concentration, a continuous liquid film wets the grain boundaries and faster grain growth occurs.22 It was reported as well that in the presence of a thin liquid film between grains, dislocations promote grain growth of SrTiO3.23,24 In another work, 1 mol % TiO2 was added to the SrTiO3 powders to create a TiO2-excess liquid above the eutectic temperature.14 Transmission electron microscopy (TEM) showed liquid phase in the triple junction of the grains and the interface between the SrTiO3 grains and the liquid was faceted.14 Even below eutectic temperatures, nanometer-thick impuritybased intergranular ‘‘glassy” films (IGFs) have been widely observed at grain boundaries in ceramics.11 Enhanced densification related to these IGFs has been considered the origin of activated sintering below the eutectic temperatures.25 Impurities have also been reported to change equilibrium dihedral angles and increase significantly the penetration length of liquids in polycrystalline materials.26 Here we report a systematic grain growth study performed over a wide range of temperature on a 0.5 mol % Ti-rich SrTiO3 composition (ST 0.995), and bulk and grain boundary electrical evolution with the sintering temperature is assessed by impedance spectroscopy. This technique allows to address separately the grain boundary impedance response and therefore to detect changes in the grain boundary dielectric properties.27 In this way, changes in the grain growth behavior are correlated to changes in the grain boundary properties noticed from alterations in their electrical response. The present work brings new data to the discussion, contributing to

2. EXPERIMENTAL METHODS Raw powders of SrCO3 and TiO2 (pro analysis (purity ≥99%), Merck) were ball milled in alcohol in a planetary mill, using Teflon pots with zirconia balls for 5 h. After calcination at 1100 °C for 2 h, the powders are monophasic. Powders were then milled again under the same conditions. Laser diffraction and SEM characterization showed particle size ranging from ∼0.03 to ∼1 μm with average particle size determined by Coulter of 0.59 μm.28 Powders were uniaxially pressed at 50 MPa and then isostatically pressed at 200 MPa. In order to define the temperature for the grain growth experiments, a dilatometric analysis was performed in which green pellets were heated up in air at 15 °C min−1 and the linear shrinkage was recorded by a computer assisted dilatometer Linseis, model 4 L70-2000. Figure 1 shows the relative density dependence on the

Figure 1. Relative density evolution with the temperature (heating rate of 15 °C min−1). At 1400 °C, the ST 0.995 samples show nearly full densification.

temperature of as-pressed ST 0.995 samples. It is clear that the onset of densification is around 1050 °C with a maximum of shrinkage rate at approximately 1260 °C and that at 1400 °C the samples are already densified, showing a relative density of nearly 99%. Therefore, for all the samples in the entire temperature range of the grain growth experiments of this work (1400−1650 °C), the matter transport is dominated by the grain growth process and not by the competing densification phenomena. For the grain growth experiments, the samples were heated up at 15 °C min−1 to temperatures between 1400 and 1650 °C in air for 2 h of holding time and then cooled at the same rate of 15 °C min−1. All the samples reached a relative density >98%, as measured by the Archimedes method. Polished and thermally etched sections of sintered samples were observed by scanning electron microscopy/energy dispersive spectroscopy (SEM/EDS) (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. 24788



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Figure 2. SEM microstructures of samples sintered at several temperatures (indicated on the images). Despite the continuous increase of the sintering temperature, the decrease of the grain size is obvious. Arrows signalize the regimen transitions with decreasing of the grain size.

For TEM (Hitachi 9000) observation of the liquid phase in grain boundaries and triple points, the samples were polished using a Gatan disc grinder in order to reduce their thickness to approximately 30 μm. The samples were then glued to a copper ring and ion beam milled using a Gatan Precision Ion Polishing System (Model 691). The dielectric properties were characterized by impedance spectroscopy. Silver paste electrodes were painted on the sample bottom and top surfaces and impedance measurements (HP 4284A Precision LCR Meter) were carried out between 100 Hz and 1 MHz in a selected temperature range of 200−700 °C. The collected impedance data were normalized by multiplying by the geometric factor A/d (A is the area of the electrode and d is the thickness of the sample) and analyzed with the software ZView (Scribner Associates Inc.).

Figure 3. Average grain size as a function of the sintering temperature defining four grain growth regimens (regimen A, from 1400 to 1500 °C; regimen B, from 1500 to 1550 °C; regimen C, from 1550 to 1605 °C; and regimen D, from 1605 to 1650 °C). Lines are only to guide the eyes.

3. RESULTS AND DISCUSSION Figure 2 presents the SEM microstructures of the samples sintered for 2 h at increasing temperatures from 1400 to 1650 °C. Similarly to the reported grain growth anomaly,1 it can be seen that the grain size does not increase continuously with the increase of the sintering temperature, as would be expected since grain growth transport is a thermally activated process. In fact, and differently from the previous report,1 four grain growth regimens can be defined with transitions at temperatures around 1500, 1550, and 1605 °C where the grain size decreases. Similar observations can be drawn from the dependence of the average grain size and grains equivalent diameter distribution on the temperature depicted in Figures 3 and 4, respectively. The represented distributions correspond to the temperatures at the beginning and end of the grain growth regimens identified in Figure 2. It is indeed clear that there is an increasing of the average grain size and an enlargement of the

grain size distribution within every regimen. But, this “normal” variation with the sintering temperature is followed by a drop in the average grain size accompanied by a narrower grain size distribution with a higher number of small grains, which corresponds to the beginning of each new regimen. A careful analysis of the sintered samples by SEM (Figure 5) revealed the presence of an amorphous phase at the grain boundaries and triple points. Some crystalline material could be observed inside the amorphous phase in certain triple points as well, probably originated during cooling. Figure 5a presents a SEM micrograph of a ST 0.995 sample sintered at 1450 °C showing in detail a triple pocket with liquid phase and some crystalline material inside. 24789



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composition under study. Additionally, the Sr-rich crystalline material in the triple pocket may contain Zr-rich impurities, as observed in one of our previous works.28 These impurities may come from the milling media and are mainly located in those small crystals. Their concentration should be less than 270 ppm as previously measured by inductively coupled plasma spectroscopy (ICPS).28 From the TEM analysis of selected samples, no Ti precipitates could be detected for the samples sintered at 1450 and 1500 °C (Figure 6). However, in the triple points the Figure 4. Distribution of the grain equivalent diameter for sintering temperatures corresponding to the beginning and end of the grain growth regimens indicated as A, B, C, and D.

Figure 6. TEM micrographs of ST 0.995 samples sintered at (a) 1450 and (b) 1500 °C. The inset is the electron diffraction pattern from the triple pocket. Amorphous phase was identified in the triple junction and along the grain boundaries.

Figure 5. (a) SEM micrograph of a sample sintered at 1450 °C showing a triple point. (b,c) EDS line profile along lines 1 and 2 in (a), respectively. A Ti-rich amorphous phase and Sr-rich crystalline material are seen inside the triple pocket.

presence of the liquid phase is evident, confirming the SEM observations. No diffraction contrast emerged as the sample was tilted in the TEM, pointing to the amorphous nature of this phase. Indeed, Ti-rich SrTiO3 compositions have an eutectic point at 1440 °C,20 so traces of liquid phase can be expected in these samples. The presence of the amorphous phase and crystalline material in the triple pocket is illustrated in the

Figure 5b,c shows EDS line profiles across the triple point. It is shown that the amorphous phase is Ti-rich whereas the crystalline material inside is Sr-rich, despite the Ti-excess of the 24790



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diffraction pattern of the triple point in the inset of Figure 6a. The diffraction halo is originated by the glassy phase, whereas the small spots are from the small crystals inside. Additionally, in Figure 6b traces of liquid phase can be clearly seen along the boundaries, as lighter regions in the micrograph. The solubility limits of SrO and TiO2 in SrTiO3 are quite low and have been discussed for more than forty years.29−32 According to Witek et al.29 the solubility limit of TiO2 is smaller than 0.1 mol % at 1000 °C and the amount held in solution is higher for faster cooling rates. However, there is no second phase detectable by XRD in SrTiO3 with Sr/Ti of 0.997.28,33 For a slightly higher excess, Sr/Ti = 0.995, the composition used in this work, a second-phase seen by SEM was reported.29 Baürer et al.,34 for Sr/Ti = 0.996, reported rutile after sintering at 1400 °C, but the detected diffraction pattern peak became less pronounced after quenching the sample from high temperature. Additionally, after sintering at 1500 °C no such phase was evident.3,34 These observations indicate that the solubility of excess TiO2 is higher at high temperatures34 and is in agreement with a solubility of more than 2 mol % TiO2 at elevated temperatures as reported by Eror and Balachandran.32 Impedance spectroscopy provides the ability of analyzing separately bulk and grain boundary contributions to the dielectric response. This ability turns out to be very useful for the purpose of detecting changes in the grain boundary dielectric response and in this way to correlate different grain boundary properties to different grain boundary complexions, as in the present case. The resistivity of grains and grain boundaries in Ti-rich strontium titanate ceramics was assessed by impedance spectroscopy and new data for the anomaly of the grain growth observed in strontium titanate are brought into the discussion. Complex impedance spectra were collected between room temperature and 700 °C (from 100 Hz to 1 MHz) for all the samples under investigation, sintered between 1400 and 1650 °C. An equivalent circuit was used to fit the impedance data.27 The equivalent circuit is constituted of a series of three blocks of a resistor and a Constant Phase Element (CPE) in parallel, modeling bulk, grain boundaries, and sample−electrode interface contributions. The complex specific impedance spectra are presented in Figure 7a,b. At both measurement temperatures (300 and 600 °C) well-resolved semicircles are observed. The assignment of the semicircles to bulk and grain boundary contributions is supported by the magnitude of the obtained capacitance values.27 At 300 °C (Figure 7a) the spectra are mainly dominated by the bulk response whereas the grain boundary contribution is dominating at 600 °C (Figure 7b). The capacitance of bulk and grain boundaries is nearly independent of the measurement temperature. The average of the capacitance values obtained from 200 to 500 °C for the bulk and from 500 to 700 °C for the grain boundary contribution regarding the several sintering temperatures is presented in Figure 8. The error bars represent the standard deviation of the several values. It can be seen that the bulk capacitance is approximately 2 × 10−11 F/cm, almost independently of the sintering temperature, whereas the grain boundaries show capacitance values of around 4 × 10−8 F/cm. In the case of grain boundary capacitance, a slight tendency to increase with increasing sintering temperature may be observed. The magnitude of these values is in the expected range for bulk and grain boundary contributions.27 One important remark regarding the impedance spectra in Figure 7 is that the sintering temperature has a strong effect on

Figure 7. Complex specific impedance spectra at (a) 300 °C and (b) 600 °C regarding samples sintered between 1400 and 1650 °C. A strong effect of the sintering temperature on the impedance response of bulk and grain boundaries is evident.

Figure 8. Average capacitance of bulk and grain boundaries as a function of the sintering temperature. Both contributions are nearly independent of the sintering temperature.

the resistivity of both bulk and grain boundaries. Regarding both contributions, the variation of the size of the semicircles (and consequently of the resistivity) is not systematic with the sintering temperature. Concerning the bulk contribution, the impedance spectra at 300 °C (Figure 7a) may be roughly divided in three groups. In the first group, samples sintered at 1425, 1450, 1525, 1535, 1610, or 1625 °C show bulk resistivity lower than 0.4 MΩ·cm (with the exception of the sample sintered at 1625 °C with a slightly higher resistivity). On the other hand, the second group that consists in samples sintered at 1400, 1500, 1550, 1600, and 1605 °C exhibits a bulk resistivity around 1.6 MΩ·cm. After that, samples sintered at 1475, 1575, 1590, and 1650 °C show a bulk resistivity close to 24791



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or higher than 2 MΩ·cm. It is interesting to note that these samples, which show high bulk resistivity, are located near the end of the respective grain growth regimens and, with the exception of the sample sintered at 1600 °C, the samples showing intermediate bulk resistivity correspond to the beginning of the grain growth regimens. The samples with the lowest bulk resistivity are located in the middle of the grain growth regimens. An exception is the sample sintered at 1535 °C, which corresponds to the end of regimen B. All samples showed relative densities above 98% and therefore changes in the bulk resistivity cannot be explained by differences in densification. The observation of a nonsystematic variation of the bulk resistivity with the sintering temperature and its relation with the grain growth regimens clearly suggest a relation with the incorporation of Ti-excess and its dependence on the sintering temperature. Indeed, different levels of solubility of the Ti-excess in the bulk affect the bulk defect concentration and, consequently, the bulk resistivity. Concerning the grain boundaries impedance response at 600 °C (Figure 7b) the dependence on the sintering temperature is also not systematic. In addition, one important observation is that the magnitude of the variation of the grain boundary resistivity is much higher than that of the bulk. In fact, the grain boundary resistivity changes dramatically from around 2 kΩ·cm for samples sintered at 1475 and 1610 °C to approximately 12 kΩ·cm for samples sintered at 1550 °C. The Arrhenius dependence of the bulk and grain boundary conductivity on the measurement temperature is shown in Figure 9 for samples sintered at the several temperatures. Regarding the bulk conductivity (Figure 9a), all lines are nearly parallel, revealing similar sintering temperature dependence of the conductivity. On the other hand, the slope of the lines concerning the dependence of the grain boundary conductivity

on the measurement temperature (Figure 9b) shows a marked variation with the sintering temperature. The dispersion of the conductivity values at 500 °C is much larger than that of the higher temperature of 700 °C. This shows that the differences in the grain boundary characteristics have a lower effect on the grain boundary conductivity at higher temperatures, where the mobility of the species involved in the conduction process is higher. These observations are clearly illustrated in Figure 10 presenting the activation energy for conductivity of bulk and

Figure 10. Activation energy for bulk and grain boundary conductivity as a function of the sintering temperature. A discontinuous variation defining four regimens is also observed, more obvious for the grain boundary conductivity.

grain boundaries for the several sintering temperatures, obtained from the respective conductivity temperature dependences. The activation energy regarding the bulk contribution is almost independent of the sintering temperature with values around 0.9 eV, which is in agreement with previously reported values.35,36 Only a small tendency to increase with increasing sintering temperature is detectable. On the other hand, the activation energy for grain boundary conductivity denotes a very strong variation with the sintering temperature with values oscillating from 0.46 to 1.76 eV. Typical reported values of activation energy for the grain boundary contribution are within the range of 1.4−1.6 eV.35 For some of the sintering temperatures, much lower values have been found in this work, suggesting in those cases a lower temperature sensitivity of the grain boundary conductivity. It is again important to consider that microstructural nonideality, such as a nonuniform grain size distribution, may originate deviations from the ideality of the brick layer model and therefore may affect the calculated values of the activation energy of grain boundary conductivity obtained from the impedance spectra.36,37 Nonetheless, the results clearly show that grain boundaries are much more affected than the bulk by the sintering temperature. Also, grain size may play a role, once that higher grain boundary resistivity is expected for smaller grain size (or higher grain boundary surface area).35,38 However, the differences in grain size are constant along the whole measurement temperature and, therefore, they alone may not explain the different dependences of the grain boundary conductivity on the temperature illustrated by the activation energy values in Figure 10. Moreover, the strong variation in the activation energy for grain boundary conductivity is not continuous with the increasing sintering

Figure 9. Arrhenius-type temperature dependence of (a) bulk and (b) grain boundary conductivity. Grain boundaries are much more affected by the sintering temperature than the bulk counterpart. 24792



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where δGB stands for the grain boundary thickness, CB and CGB are the bulk and grain boundary capacitance, respectively, and G is the average grain size.46 An assumption is made that the dielectric constants of bulk and grain boundaries are similar. The grain boundary thickness values of Figure 11 are in good agreement with previously reported values.43,47,48 Discontinuities are also present in the evolution of the grain boundary thickness with the sintering temperature. In fact, Figure 11 shows a striking similarity to the dependence of activation energy for grain boundary conductivity on the sintering temperature and to the grain growth regimens previously identified. Only the grain boundary thickness of the sample sintered at 1475 °C does not fit exactly to the observed grain growth regimens; that sample is positioned in the grain growth regimen A whereas the grain boundary thickness is lower than that of the sample sintered at 1450 °C. On the other hand, a coincidence is verified in the relative position of that sample regarding the grain boundary thickness and the activation energy for grain boundary conductivity. At this point, one important aspect must be highlighted. There is a remarkable coincidence in the behavior of the several parameters investigated in this work. The variation with the sintering temperature observed for the average grain size, grain boundary activation energy for conductivity, and grain boundary thickness is strikingly similar as illustrated in Figure 12. In fact, all the parameters seem to be related and to define

temperature. In fact, four regimens, as observed for grain growth can be defined (Figure 10). Glassy films wetting the grain boundaries are known to constitute a blocking layer for oxygen ion conduction.38 For example, Si-rich second phases spreading from the triple points were found to markedly enhance the grain boundary electrical barrier effect on zirconia samples.39,40 Additionally, Badwal38 attributed the inflections or peaks in the grain boundary resistivity of ZrO2-based electrolytes variation as a function of the sintering temperature to the dynamic nature of the grain boundary phase (composition, location, and wetting properties) and the amount and type of impurities in the starting powders. In the present case, the amount and type of impurities is obviously the same in all samples because they were prepared from the same powder and under the same conditions. In this work, the coincidence of grain growth behavior and grain boundary dielectric response strongly suggests the predominance of different grain boundary complexions with first order reverse transitions in between the several regimens. The nature of the grain boundaries of strontium titanate is rather complex and the electrical response of these regions is dependent not only on the dynamic nature of the grain boundary phase but also on space charge effects.41−45 Therefore, the strong effect of the sintering temperature observed on the grain boundary conductivity is suggested to be related to the wettability of amorphous film and variation of the electrostatic interaction, inducing reverse complexion transitions. These complexions may have different mobility and electrical properties and therefore influence the sintering behavior and the dielectric response of the material. Indeed, the discontinuities in grain size and activation energy for grain boundary conductivity are strong indications of different grain boundary properties, which are observed in this work by two different techniques. Another strong indication of the presence of different complexions changing with the sintering temperature is the grain boundary thickness variation with the sintering temperature, presented in Figure 11. These values were estimated from the bulk and grain boundary capacitance values in Figure 8 and the average grain size in Figure 4 applying the relation δGB =

CB G CGB

Figure 12. Illustration of the similar trends observed for the dependence of the average grain size, grain boundary activation energy for conductivity, and grain boundary thickness on the sintering temperature.

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four temperature regimens separated by sudden variations. This suggests the occurrence of changes at the grain boundaries at determined temperatures, namely transitions between grain boundary complexions, leading to the discontinuous behavior of the investigated parameters. A thorough HRTEM investigation of the grain boundaries of the samples sintered at temperatures around the grain growth regimen transitions is ongoing, seeking detailed information on the grain boundary complexion transitions detected in this investigation. Moreover, further work will be performed, consisting on the characterization of microstructure and dielectric properties of ceramic samples quenched just before the sintering temperatures around the grain growth regimen transitions observed in this work, which will be correlated to the present observations. These experiments may contribute to a better understanding of the reported grain growth anomaly,

Figure 11. Grain boundary thickness dependence on the sintering temperature assessed by impedance spectroscopy. A discontinuous variation with four regimens is again observed. 24793


The Journal of Physical Chemistry C namely by providing more detailed information on the microstructure and grain boundary properties at the beginning and during the grain growth experiments described in the present paper. As a final note, after consideration of the chemical and morphological nature of ceramic powders, as well as the processing parameters,49 it is expected that the clarification of the present grain boundary phenomena may be extended to other technologically important systems, greatly contributing to the possibility of using grain boundary engineering for tailoring of materials properties.

ABBREVIATIONS

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REFERENCES

IS, impedance spectroscopy; TEM, transmission electron microscopy; IGF, intergranular ‘‘glassy” film; SEM, scanning electron microscopy; EDS, energy dispersive spectroscopy; ICPS, inductively coupled plasma spectroscopy; CPE, constant phase element

(1) Baurer, M.; Weygand, D.; Gumbsch, P.; Hoffmann, M. J. Grain Growth Anomaly in Strontium Titanate. Scr. Mater. 2009, 61, 584− 587. (2) Baurer, M.; Stormer, H.; Gerthsen, D.; Hoffmann, M. J. Linking Grain Boundaries and Grain Growth in Ceramics. Adv. Eng. Mater. 2010, 12, 1230−1234. (3) Baurer, M.; Shih, S. J.; Bishop, C.; Harmer, M. P.; Cockayne, D.; Hoffmann, M. J. Abnormal Grain Growth in Undoped Strontium and Barium Titanate. Acta Mater. 2010, 58, 290−300. (4) Dillon, S. J.; Tang, M.; Carter, W. C.; Harmer, M. P. Complexion: A New Concept for Kinetic Engineering in Materials Science. Acta Mater. 2007, 55, 6208−6218. (5) Tagantsev, A. K.; Sherman, V. O.; Astafiev, K. F.; Venkatesh, J.; Setter, N. Ferroelectric Materials for Microwave Tunable Applications. J. Electroceram. 2003, 11, 5−66. (6) Koumoto, K.; Wang, Y. F.; Zhang, R. Z.; Kosuga, A.; Funahashi, R. Oxide Thermoelectric Materials: A Nanostructuring Approach. Annu. Rev. Mater. Res. 2010, 40, 363−394. (7) Palmour I. I. I., H.; Huckabee, M. L.; Hare, T. M. Rate Controlled Sintering, Principles and Practice. In Sintering - New developments; Ristic, N. N., Ed.; Proceedings of 4th Round Table, Conference on Sintering, Dubrownik, Yugoslavia, September 5−10, 1977; pp 46−66. (8) Hare, T. M.; Palmour III, H. Effect of Process Optimization Properties of Alumina Sintered Under Rate Control. In Ceramic Processing Before Firing; Onoda G. Y., Jr., Hench, L. L., Eds.; John Wiley & Sons Inc.: New York, 1978; pp 307−320. (9) Genuist, C.; Haussonne, F. J. M. Sintering of BaTiO3: Dilatometric Analyses of Diffusion Models and Microstructure Control. Ceram. Int. 1988, 14, 169−179. (10) Baurer, M.; Kungl, H.; Hoffmann, M. J. Influence of Sr/Ti Stoichiometry on the Densification Behavior of Strontium Titanate. J. Am. Ceram. Soc. 2009, 92, 601−606. (11) Luo, J. Stabilization of Nanoscale Quasi-Liquid Interfacial Films in Inorganic Materials: A Review and Critical Assessment. Crit. Rev. Solid. State 2007, 32, 67−109. (12) Rohrer, G. S. Influence of Interface Anisotropy on Grain Growth and Coarsening. Annu. Rev. Mater. Res. 2005, 35, 99−126. (13) Choi, S. Y.; Kang, S. J. L. Sintering Kinetics by Structural Transition at Grain Boundaries in Barium Titanate. Acta Mater. 2004, 52, 2937−2943. (14) Chung, S. Y.; Yoon, D. Y.; Kang, S. J. L. Effects of Donor Concentration and Oxygen Partial Pressure on Interface Morphology and Grain Growth Behavior in SrTiO3. Acta Mater. 2002, 50, 3361− 3371. (15) Dillon, S. J.; Harmer, M. P. Relating Grain-Boundary Complexion to Grain-Boundary Kinetics I: Calcia-Doped Alumina. J. Am. Ceram. Soc. 2008, 91, 2304−2313. (16) Dillon, S. J.; Harmer, M. P. Relating Grain-Boundary Complexion to Grain-Boundary Kinetics II: Silica-Doped Alumina. J. Am. Ceram. Soc. 2008, 91, 2314−2320. (17) Luo, J. Grain Boundary Complexions: The Interplay of Premelting, Prewetting, and Multilayer Adsorption. Appl. Phys. Lett. 2009, 95, 071911−1−071911−3. (18) Shi, X. M.; Luo, J. Decreasing the Grain Boundary Diffusivity in Binary Alloys With Increasing Temperature. Phys. Rev. Lett. 2010, 105, 236102−1−236102−4. (19) Luo, J. Developing Interfacial Phase Diagrams for Applications in Activated Sintering and Beyond: Current Status and Future Directions. J. Am. Ceram. Soc. 2012, 95, 2358−2371.

4. CONCLUSIONS The anomaly in the dependence of grain growth kinetics on the temperature of strontium titanate is revisited. Discontinuities on the evolution of grain growth with the sintering temperature were observed, defining four grain growth regimens with transitions at temperatures around 1500, 1550, and 1605 °C. These transitions correspond to grain size decreases despite the increasing sintering temperature. Interestingly, new related effects on the dielectric properties were also verified by impedance spectroscopy. Grain boundary dielectric 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 grain boundaries in the anomaly in the grain growth observed for Ti-rich SrTiO3. Furthermore, changes in the activation energy for grain boundary conductivity and grain boundary thickness were detected that may be well correlated to the defined grain growth regimens. We postulate that this anomalous behavior is directly related to a first order reverse change of complexions driven by the temperature effect on the grain boundary wettability and electrostatic potential. These new insights have great scientific and technological relevance in tailoring the microstructure and dielectric response of SrTiO3based materials and related systems and using grain boundary complexion behavior for materials properties design.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +351 234 370354. Fax: +351 234 425300. Present Address §

(L.A.) Materials Electrochemistry Group, ICEMS, Instituto Superior Técnico, Av. Rovisco Pais, 1049−001 Lisbon, Portugal. Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank Jian Luo and Animesh Kundu for the helpful discussion of the results. The authors acknowledge the financial support from FEDER, QREN, COMPETE, FCT, and the Foundation for Luso American Development (FLAD), Portugal. Luiś Amaral acknowledges FCT for financial support (SFRH/BD/40927/2007). 24794



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