Synthesis-phase–composition relationship and high electric-field-induced electromechanical behavior of samarium-modified BiFeO3 ceramics

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ScienceDirect Acta Materialia 83 (2015) 149–159 www.elsevier.com/locate/actamat

Synthesis-phase–composition relationship and high electric-field-induced electromechanical behavior of samarium-modified BiFeO3 ceramics ⇑

Julian Walker,a,c, Peter Bryant,b Valsala Kurusingal,b Charles Sorrell,a Danjela Kuscer,c Goran Drazic,d Andreja Bencan,c Valanoor Nagarajana and Tadej Rojacc a

The School of Materials Science and Engineering, University of New South Wales, Sydney, Australia b Thales Australia, Rydalmere, Australia c Electronic Ceramics Department, Jozef Stefan Institute, Jamova Cesta, 39, Ljubljana SI-1000, Slovenia d Laboratory for Materials Chemistry, National Institute of Chemistry, Ljubljana, Slovenia Received 25 June 2014; revised 25 September 2014; accepted 30 September 2014

Abstract—Solid-state (non-activated) and mechanochemical activation (activated) synthesis methods were used to produce Sm-modified BiFeO3 ceramics of composition Bi0.88Sm0.12FeO3. The first part shows that the formation of Bi0.88Sm0.12FeO3 using the two synthesis methods followed a different reaction pathway on annealing the powders. The non-activated ceramics reacted by forming two intermediate phases, isostructural to BiFeO3 and SmFeO3, and then inter-diffusing, forming the final Bi0.88Sm0.12FeO3 solid solution. Unlike the non-activated samples, the activated ceramic powders formed Bi0.88Sm0.12FeO3 phase on annealing the powders, without apparent intermediate phases. As revealed by transmission electron microscopy, the non-activated reaction pathway caused the Pbam phase to form as chemical inhomogeneous (Sm-rich) isolated nano-sized grain inclusions in the final ceramics. Conversely, the activated reaction pathway caused the Pbam phase to form chemically homogeneous nano-regions within the R3c phase grains. The results demonstrate the important role of processing in the appearance of the frequently discussed anti-polar Pbam phase in this system. In the second part, the high electric-field-induced polarization and strain behaviors of these ceramics were studied by means of polarization–electric (P–E) and strain–electric field (S–E) hysteresis loops, and the S–E loops were compared with those of unmodified BiFeO3. Bipolar S–E loops of Bi0.88Sm0.12FeO3 had a distinctive butterfly shape with less frequency dependence relative to BiFeO3 at driving-field frequencies of 0.1–100 Hz. BiFeO3 ceramics exhibite strong driving electric-field-frequency-dependent domain switching, the origins of which were previously attributed to a domain-wall pinning mechanism and “hardening” behavior. This study shows that Sm-modification induces a “hardening–softening” transition in BiFeO3 ceramics. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Bismuth ferrite; Rare-earth; Synthesis–structure relationship; Electromechanical behavior

1. Introduction The discovery of a compositionally induced polar-tonon-polar phase transition [1] in samarium (Sm)-modified bismuth ferrite (BiFeO3) has stimulated considerable interest in these materials as a potential lead-free piezoelectric [2]. Epitaxial pulsed-laser-deposited (PLD) thin-films of Sm-modified BiFeO3 exhibited a reduced coercive field (Ec) and increased converse longitudinal piezoelectric coefficient (d33) relative to unmodified BiFeO3, attributed to the presence of the phase boundary. Similar phase transitions [3] and property enhancements [4] were later observed in BiFeO3 modified with other rare-earth (RE) species, such as Nd, Gd and Dy. In addition, a phase with a space group Pbam and an anti-polar cation ordering, isostructural to

⇑ Corresponding

author at: Electronic Ceramics Department, Jozef Stefan Institute, Jamova Cesta, 39, Ljubljana SI-1000, Slovenia. Tel.: +386 51 700 972.; e-mail: [email protected]

orthorhombic PbZrO3, was observed in the vicinity of the phase transition [5]. While thin films are important for micro- and nanoelectronic applications, a significant proportion of existing piezoelectric devices, such as sensors, actuators and transducers, rely on piezoelectric materials in the form of bulk ceramics [6]. Despite considerable efforts, however, the ferroelectric and electromechanical properties of RE-modified BiFeO3 ceramics [7,8] are not yet able to match those of thin-films [4,9,10]. Similarly to unmodified BiFeO3 [11–13], it is believed that the main problem is a poor understanding of the processing of these new RE-modified BiFeO3 compositions, and thus an inability to reproduce their functional properties effectively. The synthesis techniques used to produce RE-modified BiFeO3 ceramics have varied greatly, with the literature detailing numerous conventional solid-state [3,5], rapid liquid-phase sintering [8,14,7] and various wet chemical [15,16] methods. These processing techniques have been used in studies reporting the functional responses of the

http://dx.doi.org/10.1016/j.actamat.2014.09.058 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.

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ceramics, including dielectric permittivity [17], pyroelectric properties [14], magnetic and multiferroic properties [18,19], ferroelectric domain switching behavior and piezoelectricity [8,7]. The very different processing methods, however, seemingly result in processing-dependent properties, and few studies provide detailed discussion on the specific influence that the methods of synthesis may have on the functional response of the RE-modified BiFeO3 ceramics. In complex materials systems, such as the Pb(Zr,Ti)O3based [20] and (K,Na)NbO3-based compositions [21,22], it has been shown that the method of synthesis can have a considerable impact on the ceramic’s compositional homogeneity. The processing-related inhomogeneity may then have a crucial influence on the piezoelectric properties, particularly when compositions are at or near a phase boundary [20,22]. The issue of the compositional homogeneity in Sm-modified BiFeO3 ceramics has not yet been addressed, so the inconsistency in the functional responses reported in the literature is not surprising and calls for systematic investigations. This study presents a detailed analysis of the interrelationships between the method of synthesis, the reaction and the resulting phase composition of Bi0.88Sm0.12FeO3 (BSFO) ceramics. This composition was chosen because, previously, Fujino et al. [1] discovered a sharp rise in the electromechanical properties at this composition, just before the phase boundary at 14 mol.% Sm. Secondly, the composition with 12 mol.% Sm is located far enough away from the phase boundary to avoid the particularly complex phase coexistence found for the 14 mol.% composition [23], yet close enough to have a ceramic with functional properties of interest. For example, the first report on bulk Smmodified BiFeO3 found promising multiferroic properties close to this composition [24]. The ceramics were prepared by two processing methods: a solid-state method and a mechanochemical activationassisted method. These methods, of which neither is novel, but both are established ceramic processing methods, were compared with the purpose of understanding the processing–properties relationship. Different reaction pathways were observed with each synthesis technique, the consequences of which were variations between the chemical homogeneities and phase compositions of the two ceramics. In particular, the phase coexistence between the polar, rhombohedral R3c matrix phase and anti-polar, orthorhombic Pbam phase, was effected by the synthesis method. The polarization and strain responses of BSFO ceramics under applied electric-fields above the coercive field were also examined. In particular, the strain–electric-field (S–E) hysteresis loops are discussed by comparison with those of unmodified BiFeO3 (BFO) ceramics. The comparison between BSFO and BFO revealed that the electric-fieldinduced strain in BSFO showed more limited frequency dependence relative to BFO, with open butterfly-shaped S–E loops across the driving electric-field frequency range 0.1–100 Hz. This was in contrast to the behavior of BFO ceramics, which exhibited strong frequency-dependent electric-field-induced strain in the same frequency range. This indicates that Sm modification of BiFeO3 results in a so-called “softening” effect, similar to that observed, for example, in donor doping of lead zirconate titanate (PZT) [25–27], which is technologically one of the most important methods for controlling properties in piezoelectric ceramics [27].

2. Experimental procedure Two types of ceramics with nominal composition BSFO were prepared. According to the synthesis procedure, the two ceramics are denoted as non-activated and activated, where non-activated refers to a method of solid-state synthesis, and activated refers to a method of mechanochemical activation-assisted synthesis. Both methods of synthesis began with ball-milling of the individual ceramic powders: Bi2O3 (99.998%, Alfa Aesar), Fe2O3 (99.999%, Alfa Aesar) and Sm2O3 (99.99%, MetallÒ rare earths). Each powder was milled in two stages for a total of 6 h in a polyurethane (PE) jar, using 3-mm and 10-mm yttrium stabilized zirconia (YSZ) milling balls. The milling was conducted using a planetary ball mill (Retsch, PM400) at rotational frequency of 200 min1. Surfactant ammonium polyacrylate (PAA) (MW = 4000 g mol1, 50 wt.% solution with H2O, Sigma Aldrich) was added to the powders [28] as a steric stabilizer [29]. The initial powders were milled to mean particle sizes <1 lm, with the following d50 values: d50(Fe2O3) = 0.48 lm, d50(Bi2O3) = 0.34 lm, d50(Sm2O3) = 0.78 lm. After milling, the powders were dried to 50 °C for 12 h and sieved with <45 lm mesh. A stoichiometric mixture of as-milled Bi2O3, Fe2O3 and Sm2O3 was mixed to a nominal composition of BSFO, by redispersing the three powders in ethanol and by 1 h of milling (planetary mill, Retsch PM400) at 200 min1 of rotational frequency, using 3-mm YSZ milling balls. The mixture was then dried and sieved as previously described. To produce the non-activated samples, the powder mixture, prepared as described above, was uniaxially pressed into 8-mm-diameter pellets at 150 MPa and reactively sintered, i.e. directly reacted and sintered in a single heating process, at temperatures between 650 °C and 900 °C with a heating rate of 10 °C min1 and held at the maximum temperature for 4 h. The activated BSFO powders were prepared by the use of mechanochemical activation, as reported for BFO [30,31]. The powder mixtures were first prepared by the same method as those mixtures of the non-activated method, as described above. Following this preparation, a heating cycle was applied to remove PAA from the powder. The heating cycle was as follows: the powder was placed in an alumina crucible, heated at 2 °C min1 to 350 °C in air, held for 30 min and cooled to room temperature at 10 °C min1. Mechanochemical activation was then conducted for a total of 40 h, on batches of 25–30 g of powder in a 180 ml tungsten carbide–cobalt (WC–Co) container with WC–Co milling balls with diameter 10 mm and a mill rotational frequency of 300 min1. Following mechanochemical activation, the powder was milled again in a PE container with 3 mm YSZ milling balls and ethanol for 4 h at rotational frequency of 200 min1. Following the final milling step, the powders were dried, sieved, isostatically pressed and reactively sintered. These procedures were performed as described above for the non-activated samples, and reaction sintering was conducted within the same temperature range and with the same holding time. The bulk densities of samples were calculated geometrically, i.e. from sample dimensions and mass. Geometrical densities of non-activated and activated samples were 8050 ± 100 kg m3 and 8190 ± 50 kg m3, respectively. The relative densities were obtained by scanning electron microscopy (SEM) image analysis of the porosity of the

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samples, and were 94% and 97% for non-activated and activated BSFO, respectively. The X-ray diffraction (XRD) patterns were recorded using a Bruker D8 Discover diffractometer, in a 10–90° 2h range with a step of 0.04° and an acquisition speed of 1.5° min1. The analysis was conducted ex situ, meaning that the samples were heat treated at four selected temperatures between 650 °C and 900 °C, held for 4 h, cooled to room temperature and then analyzed by XRD. Rietveld refinement [32] was conducted using the TOPAZ R software package (Version 2.1, 2003, Coelho software). The peak shape was refined with a Voigt function, and the background with a linear function. Peak intensities, peak shapes and scale factors were fitted before refinement of unit cell parameters, atomic positions and then atomic occupancies. For R3c phase, ICSD#15299 [32] was used for structural fitting. The orthorhombic space group Pbnm was used for structural fitting of the non-polar phase ICSD#162895, as previously used in studies of Sm-rich (Bi1xSmx)FeO3 [33] and SmFeO3 compositions [34]. Pbam phase was fitted using structural data from ICSD#160460, and Bi25FeO39 was fitted using ICSD #41937. SEM was performed using a Philips XLF-30 SEM instrument equipped with a field emission gun and an energy-dispersive X-ray spectroscopy (EDXS) analyzer. Analytical transmission electron microscopy (TEM) (JEOL JEM 2010 and Cs-probe corrected JEM-ARM200CF both equipped with a JEOL EDXS detector) was employed for the study of structure properties at the nanometer and atomic scales. The specimens were prepared by mechanical grinding, dimpling and final Ar-ion milling. Electrical properties were measured from samples sintered at 820 °C (non-activated) and 800 °C (activated) for 4 h. The sintered samples were sliced using a diamond wire saw and ground to 0.2-mm thickness. Electrodes were applied by sputtering, first a 30-nm chrome (Cr) layer, followed by a 100-nm gold (Au) layer. Polarization-field (P–E) and strain–field (S–E) hysteresis loops were measured using an aixACCT TF 2000 analyzer with a SIOS Mebtechnik GmbH laser interferometer and a TREK model 609E-6 (5 kV) high-voltage amplifier.

3. Results and discussion 3.1. Synthesis and structure The non-activated and activated powder mixtures of Bi2O3–Fe2O3–Sm2O3, mixed according to the target stoichiometry BSFO, were reacted to four different temperatures in the range 650–900 °C. XRD analysis was then used to investigate the reaction and phase formation as a function of the reaction temperature. These XRD results are shown in Fig. 1a and b for non-activated and activated samples, respectively. In both non-activated and activated samples, the main phase identified was the rhombohedral R3c phase [35,36], which is isostructural with BiFeO3. In addition, Pbam [3], Pbnm (isostructural to SmFeO3 [34]) and parasitic, non-perovskite phase Bi25FeO39, were also identified at various temperatures (Fig. 1). To investigate the evolution of phases during the reaction, selected portions of the XRD patterns in Fig. 1 were analyzed more closely. The XRD spectra in Fig. 2a show the 2h region 31.4–32.6° of non-activated reacted powder,

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highlighting the major peaks of rhombohedral BiFeO3 with space group R3c [35,36]. Fig. 2b also displays a 2h region of the non-activated reacted powder’s spectra, this time in the 2h range 32.25–32.5°, where the major SmFeO3 non-polar, orthorhombic Pbnm [34] phase peaks occur, as well as one major peak of parasitic Bi25FeO39 phase. Fig. 2c and d shows the same 2h regions as Fig. 2a and b, but for the activated reacted powders. These two 2h regions were selected to demonstrate the difference in the reaction pathway of the two ceramics. In the non-activated reacted powders, the reaction was observed to progress by way of two intermediate phases identified as rhombohedral R3c (polar) phase, isostructural with BiFeO3 (Fig. 2a) and orthorhombic Pbnm (non-polar) phase, isostructural with SmFeO3 (Fig. 2b). The peaks of rhombohedral R3c phase shift incrementally to higher 2h values as the reaction temperature is increased (see arrows in Fig. 2a). Conversely, the peaks of orthorhombic Pbnm phase shift to lower 2h values (see arrows in Fig. 2b), and reduce in intensity before disappearing completely within the test temperature range. The presence of Pbam phase is also detected at annealing temperatures P820 °C using Rietveld refinement (refer to Fig. 5). With respect the Bi25FeO39 parasitic phase (Fig. 2b), in unmodified BiFeO3 ceramics, the Bi25FeO39 phase is known to occur as part of the diffusion-controlled reaction between Bi2O3 and Fe2O3 [11], which explains why it is seen in the present case at temperatures <820 °C. The rhombohedral R3c peak shift to higher 2h values (Fig. 2a, arrows) indicates that the corresponding R3c unit cell parameters reduce with increasing temperature. Quantitative analysis of this effect was carried out with Rietveld refinement. The unit cell volumes, at different annealing temperatures for both non-activated and activated reacted powders, are shown in Fig. 3. The non-activated reacted powder’s unit cell volumes show an initial decrease from ˚ 3 between reaction temperatures of 374 to 370 A 650 °C and 820 °C. Between 820 °C and 900 °C, the unit cell volume reaches an approximate plateau, showing that, beyond 820 °C, the reduction in the unit cell parameters was minor. The behavior of the non-activated reacted powders is probably related to interdiffusion of Bi3+ and Sm3+ from the initially formed BiFeO3 and SmFeO3 isostructural phases, and will be discussed in detail in Section 3.3. The activated reacted powder XRD profiles show a small 2h shift of the R3c peaks to lower 2h values, as a function of temperature (Fig. 2c). This indicates a small appar˚ 3 at ent increase in the unit cell volume, i.e. from 369 A 3 ˚ 650 °C to 370 A at 900 °C, as determined by Rietveld refinement (Fig. 3). Additionally, activated reacted powders had no detectable Pbnm peaks (Figs. 1b and 2d). The absence of the Pbnm phase and the behavior of the R3c unit cell volume as a function of temperature, for the activated reacted powders, strongly suggest that activated samples exhibit a reaction pathway contrasting to that of the nonactivated samples. The temperature dependence of the R3c unit cell volume of the activated powders (Fig. 3) can be explained by considering the physical properties of these powders. This is demonstrated by analysis of the phase evolution of powder mixtures during the mechanochemical activation process (Fig. 4). Fig. 4 shows the XRD patterns from powder mixtures before mechanochemical activation and at 10 h intervals of mechanochemical activation, up to 40 h. The powder mixture before mechanochemical activation

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Fig. 1. XRD patterns for (a) non-activated and (b) activated Bi2O3–Fe2O3–Sm2O3 reacted powders after annealing at different temperatures between 650 °C and 900 °C. The major peaks are indicated as follows: R3c, indexed; Bi25FeO39, d; SmFeO3 Pbnm, O; Pbam, +.

Fig. 2. XRD patterns for (a, b) non-activated and (c, d) activated Bi2O3–Fe2O3–Sm2O3 reacted powders after annealing at different temperatures between 650 °C and 900 °C: (a, c) 2-theta region where major R3c peaks appear; (b, d) 2-theta region where major non-polar Pbnm phase peaks appear.

contains peaks of all three oxide constituents: Bi2O3 (B), Fe2O3 (F) and Sm2O3 (S). With increasing activation time, the individual oxide peaks broadened and reduced in

intensity. A broad “hump” between 26 and 32° 2h, appears in the XRD pattern of the 10 h activated mixture and persists up to 40 h of activation. In addition to this

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Fig. 3. Unit cell volumes of R3c phase in reacted Bi2O3–Fe2O3–Sm2O3 non-activated (circles) and activated (squares) powders after annealing at different temperatures between 650 °C and 900 °C. Error bars show calculated error from Rietveld refinement. The lines are drawn as a visual guide.

Fig. 4. XRD profiles of Bi2O3–Fe2O3–Sm2O3 powder mixture, before (no activation) and after 10, 20, 30 and 40 h of mechanochemical activation. Major peaks of Bi2O3, Fe2O3 and Sm2O3 are labelled B, F and S, respectively. All peaks that are not labelled belong to the pattern of Bi2O3.

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broadened widths, relative to the “no activation” mixture. Peak broadening, increased background intensity relative to the peak intensity, together with the observed “hump”, indicate a progressive reduction in the crystallite size of constituent oxides and the formation of amorphous phases with increase activation time. This behavior is often observed as a result of mechanochemical activation [37]. The minor increase in the R3c unit cell volumes of the activated reacted powders, annealed at increasing temperatures (Fig. 3), is probably related to increased crystallization of the amorphous phase. Thus, the contrasting behaviors of the non-activated and activated powders during annealing (see Figs. 1–3) may be attributed to the different crystalline nature of the powders, i.e. the nonactivated powder is a physical mixture of three crystalline oxides (see Fig. 4, “No activation”) and the activated powder is composed of a mixture of both nanocrystalline and amorphous phases (see Fig. 4, “40 h”). Fig. 5a and b presents the phase composition (in wt.%) determined by Rietveld refinement for non-activated and activated reacted powders, respectively, as a function of the reaction temperature over the range 650–900 °C. In Fig. 5a, three phases are clearly identified in non-activated reacted powders at 650 °C, i.e. R3c, Pbnm and the parasitic phase Bi25FeO39. The increase in the reaction temperature brings about the reduction and disappearance of the Bi25FeO39 first, followed by disappearance of Pbnm, supporting the qualitative observations made from Figs. 1a and 2a, b. The anti-polar, Pbam phase first appears at 740 °C and increases in content up to 900 °C. In this same temperature range, the Pbnm phase disappears, and the concentration of the R3c phase decreases. The phase composition of the activated reacted powders in Fig. 5b shows that, over the temperature range 650– 900 °C, neither Bi25FeO39 nor Pbnm phases were detected by XRD. In comparison to the non-activated samples, the activated samples show a more stable phase composition as a function of temperature. The largest change occurs between 650 °C and 740 °C, with R3c shifting from 75 to 85 wt.% at the expense of the Pbam phase. The results of the phase composition presented in Fig. 5 were confirmed with SEM and energy dispersive spectroscopy analysis (not shown). 3.2. Anti-polar Pbam phase formation

“hump”, in the 40 h activated pattern, some peaks are visible at 32.8°, 35.2°, 49.3° and 59.8° 2h. These peaks belong to Fe2O3 and exhibit reduced intensities and

The phase diagrams for the Sm-modified BFO system [1,3,4,38] show that the two binary compositions of the diagram are the polar rhombohedral R3c (BiFeO3) and the

Fig. 5. Phase composition as a function of the annealing temperature for (a) non-activated and (b) activated Bi2O3–Fe2O3–Sm2O3 powders. Error bars express typical ±5% error of XRD technique.

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non-polar orthorhombic Pnma or Pbnm (SmFeO3) with the polar-to-non-polar phase boundary in the proximity of 14– 15 mol.% Sm. The anti-polar ordered phase orthorhombic Pbam exists in the compositional range 10–15 mol.% Sm and increases in concentration towards the polar-tonon-polar phase boundary [3]. The Pbam phase is thought to occur to relieve stresses resulting from the substitution of Bi3+ with the smaller Sm3+ cation into the BiFeO3 perovskite [39]. As a result, a phase formation model has been proposed, where Pbam forms as nano-cluster-regions within the R3c matrix, and the number of regions grows with increasing Sm concentration [9]. Despite comprehensive and detailed literature on the appearance of Pbam phase, however, no systematic

study exists that considers the role of methods of synthesis on influencing the Pbam phase formation. The issue of the anti-polar ordered Pbam phase formation in relation to the processing method was addressed here by employing high-resolution TEM (HR-TEM). The results of the analysis of the anti-polar Pbam phase in the non-activated and activated BSFO ceramics reacted to 820 °C and 800 °C, respectively, are shown in Fig. 6. With this technique, regions of Pbam phase were first identified locally in the ceramics and, once identified, the size, chemical composition and coexistence of the Pbam and R3c phases were assessed. Fig. 6a shows a bright-field (BF) TEM image of grains of non-activated BSFO, revealing the presence of

Fig. 6. (a) BF TEM image of the non-activated BSFO ceramics annealed at 820 °C. The insets show SAED patterns from the matrix (upper) and the inclusion (lower) in [0 1 0] zone axis. The 1=4 (0 0 1) reflections belonging to the Pbam phase are circled in lower SAED pattern. (b) BF TEM image of activated BSFO ceramics annealed at 800 °C. (c) HR-TEM image of activated BSFO ceramic in [1 1 0]p zone axis direction. Upper inset shows SAED pattern, and bottom insets show FFT of selected areas; 1=4 (0 0 1) reflections belonging to Pbam are circled. (d) Atomic resolution HAADF STEM image of the boundary (dashed line) between R3c and Pbam phase in activated BSFO ceramics in the [0 0 1]p zone axis. (e) Enlarged view of the Pbam and R3c phase with the corresponding polyhedral models illustrating the Fe diagonal displacement (marked with arrows).

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Fig. 7. Proposed reaction pathways for materials prepared by (a) solid-state (non-activated) and (b) mechanochemically activated (activated) synthesis methods. Depictions of materials with compositions located in the polar region of the polar-to-non-polar phase boundary system (<15 mol.% Sm). (a) Non-activated pathway in three stages, homogeneous oxide mixture, reaction of intermediate phases BiFeO3 and SmFeO3 with subsequent inter-diffusion of Bi3+ and Sm3+ during annealing. (b) Activated ceramic reaction in two stages, first with amorphous oxide mixture, then formation of final phase composition without intermediate phases.

ferroelectric domains with dimensions ranging from tens of nanometers up to 100 nm, and 200 nm large grain inclusion. By comparison of the EDXS spectra performed on the inclusion and on the matrix (see supplementary material), it was concluded that a higher amount of Sm and a lower amount of Bi was detected in the grain inclusion relative to the matrix grain. The top selected-area electron diffraction (SAED) pattern in Fig. 6a identifies the matrix phase as the polar R3c phase, while the SAED pattern associated with the Sm-rich isolated grain inclusion (lower SAED patter in Fig. 6a) shows additional superstructure 1=4 (00 l) reflections. These superstructure reflections are consistent with the quadrupled unit cell of the complex Pbam phase with octahedral tilting around the c-axis, as previously described in the literature [3,5]. These TEM results show that, in the non-activated ceramics, there are localized grain inclusions of nonhomogeneous Sm distribution within the BSFO matrix, identified by an orthorhombic Pbam crystal structure. This is in agreement with XRD phase analysis (Fig. 5a), which identified that the Pbam phase contributed to 20 wt.% of the ceramics phase composition. The TEM data from the activated BSFO ceramics is presented in Fig. 6b–e. Notably, the activated ceramics show a fine domain structure over a similar size range to nonactivated ceramics, but without any identifiable grain inclusions present (Fig. 6b). Fig. 6c shows a HR-TEM image of the activated ceramic and two different phases where identified in the activated sample. The first phase is the Pbam phase with quadrupled periodicity, which is seen as 1=4 (00 l) superstructure reflections circled in the SAED (upper right image in Fig. 6c) and in the fast Fourier transform (FFT) pattern (right FFT image in Fig. 6c). The second phase is the R3c phase without the superstructure reflections (left FFT in Fig. 6c). In contrast to the non-activated ceramic, the activated BSFO was chemically homogeneous on the level of EDXS analysis, i.e. no segregation of Sm was detected in the Pbam phase regions, and Pbam did not appear as isolated grain inclusions between R3c grains, instead appearing within R3c grains. The appearance the Pbam phase is consistent with the phase compositions determined from XRD spectra (Fig. 5b) where the Pbam phase accounts for 10 wt.%. In order to further understand the chemical composition of the Pbam phase of the activated ceramics, an atomically resolved Cs-probe aberration corrected TEM was employed. High angular dark-field scanning transmission electron microscopy images (HAADF-STEM) were fitted with the R3c (ICSD#15299) and Pbam (parameters from

Khomchenko et al. [40], and ICSD#160460) structural models. Fig. 6d shows the HADDF-STEM image of the activated sample, where different Fe–Bi diagonal distances were measured. The off–centering of B-site (Fe) atoms was correlated with the structural models (Fig. 6e) and, consequently, two phases, Pbam and R3c, were identified. A coherent boundary between R3c and Pbam phase was identified and is marked on Fig. 6e. The intensities of A and B atomic sites in both phases vary randomly, and no segregation of Sm or Bi ions on the A sites was detected. The TEM data confirm the coexistence of the R3c and Pbam phases in both the non-activated and activated ceramics, consistent with the results of the XRD analysis (Figs. 1, 2 and 5). They also show that the precise chemical composition and microstructural nature of the Pbam phase varies between the two materials (see supplementary material). In non-activated ceramics, the Pbam phase was found to exist as isolated grain inclusions in some instances, which disrupted the domain structure (Fig. 6a) and was associated with an increase in the concentration of Sm relative to the matrix. In contrast, the activated ceramic TEM data (Fig. 6b–e) shows the Pbam phase as nano-regions of varying size within the grains of the R3c phase, with no difference detected between the chemical composition of the Pbam and R3c regions, probed with an atomic-level resolution. The TEM study indicates that the chemical and phase coexistence of anti-polar Pbam phase in activated ceramics is consistent with the nano-cluster and growth model presented for PLD thin-film materials [9]. In contrast, the chemical and microstructural properties of the anti-polar Pbam phase in non-activated ceramics do not correlate with the same nano-cluster growth model, existing instead as Sm-rich isolated grain inclusions. The difference between the two ceramics is explained as a function of their different reaction pathways, which are presented in the following section. 3.3. Reaction pathways Fig. 7a shows schematically the reaction pathway of the formation of non-activated BSFO, which can be subdivided into three stages. The initial stage is the powder precursor (Fig. 7a, left schematic), a homogeneous mixture of the three starting oxides, Bi2O3, Fe2O3 and Sm2O3, confirmed with XRD (Fig. 4, “no activation”). The second stage is characterized by two key mechanisms occurring simultaneously: (i) the formation of intermediate phases, i.e. the Bi-rich phase that is isostructural with R3c BiFeO3 and

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the Sm-rich phase that is isostructural with Pbnm SmFeO3, and (ii) inter-diffusion of Bi3+ and Sm3+ between these two intermediate phases at elevated temperatures (Fig. 7a, middle schematic). The intermediate phase formation in non-activated BSFO is identified by the XRD analysis of powders annealed to temperatures <820 °C (Figs. 1a, 2b and 5a). The inter-diffusion of the Bi3+ and Sm3+ cations between the intermediate phases R3c and Pbnm is identified by two observations. First, the peak shift (Fig. 2a) and the change in the unit cell volume of the R3c phase (Fig. 3) and, secondly, the simultaneous peak shift and intensity reduction of the Pbnm phase (Fig. 2b), both as a function of increasing annealing temperature. Reduction in the R3c unit cell volume is known to occur as the smaller Sm3+ ions (128 pm) [41] progressively substitute for the larger Bi3+ ions (136 pm) [42] at the A-sites of the perovskite structure [9,43]. An opposite trend is also expected and, in fact, observed for Pbnm phase (see peak shift in Fig. 2b) as larger Bi3+ ions substitute for smaller Sm3+ ions. The reduction in the Pbnm peak intensities (Fig. 2b) with increasing temperature indicates that the amount of this phase reduces (Fig. 5a) as the Bi3+ content in this phase increases through diffusion. Above 85 mol.% of Bi3+ diffused into Pbnm, the composition should cross the polar-to-non-polar phase boundary [1], resulting in a transition from Pbnm to the coexisting R3c and Pbam phases. Of the R3c and Pbam phases, the Pbam phase is more prevalent at compositions in close proximity to the polar-to-non-polar phase boundary [3]. As a result, it is likely that, when the composition of Pbnm non-polar phases reach the point of compositionally induced phase transition, they in fact transition initially to Pbam phase. Thus, if residual chemical gradients remain in the ceramic as a result of incomplete inter-diffusion, it is likely that Sm-rich grains may retain a Pbam structure, as observed by TEM (Fig. 6a). The third and final stage of the non-activated reaction pathway is characterized by the formation of R3c–Pbam coexisting BSFO solid solution (Fig. 7a, right schematic), which is expected for this composition from the BiFeO3– SmFeO3 phase diagram [3,23]. Fig. 7b represents the activated BSFO reaction pathway occurring in two stages. The initial powder oxide precursor consists of nanocrystallites and an amorphous phase with presumably intimately mixed Bi, Fe and Sm oxides (Fig. 7b, left schematic). The amorphous phase was the result of the mechanochemical activation, confirmed by the observation of peak intensity reduction, peak broadening and the appearance of a broad “hump” in the XRD patters of activated powders (Fig. 4). The second stage of the reaction proceeds by the reaction of the amorphous phase, occurring by the apparent direct nucleation and crystallization of a relatively homogeneous Sm3+ A-site substituted BSFO perovskite (Fig. 7b, right schematic). This scenario is consistent with the absence of intermediate phases in the activated ceramics (Figs. 1, 2 and 5), and with proposed mechanochemical reaction mechanisms [44]. The direct reaction pathway without intermediate phase formation is supported by similar observations in studies on the mechanochemical processing of PZT [45,46], lead manganese niobate [47] and KNN-based materials [21]. Note that, while the exact origin of the reaction mechanism in the activated powder is not known, different reaction pathways between the

non-activated and activated powders are referred to here based on the differences observed by XRD analysis, within the investigated temperature range (650–900 °C). 3.4. High electric-field-induced polarization and strain hysteresis In order to explore the polarization and strain behavior of the as-prepared BSFO ceramics under applied electric field, particularly the possible role of the different phase content (Fig. 5) and appearance (Fig. 6) of the Pbam phase, the polarization–electric (P–E) and strain–electric-field (S– E) hysteresis was measured (Fig. 8a, b). The P–E loops of both non-activated and activated ceramics (Fig. 8a) indicate ferroelectric domain switching. A higher electrical conductivity contribution was observed in the activated ceramics, indicated by the increased “roundness” of the P–E loop relative to the non-activated ceramics. The polarization behavior is in agreement with the S–E loops (Fig. 8b), which showed a distinct butterfly shape with peak-to-peak strains of 0.2%. As the electrical conductivity can only directly affect the polarization response, the strain loop of the activated samples (Fig. 8b) did not reveal the large phase angle as was observed in the P–E loop (Fig. 8a). In ferroelectric ceramics, such strain response is commonly attributed to the switching process, particularly to the switching of ferroelectric/ferroelastic non-180° domain walls, which may also contribute significantly to the total measured strain in polycrystalline ferroelectrics under an electric-field [6]. While the observed P–E and S–E loops of BSFO are consistent with domain switching (Fig. 8a, b), it should be noted that contributions from field-induced phase transformations between coexisting Pbam and R3c phases cannot be ruled out. The possible role of the phase transition in the field-induced strain of BSFO is a subject of investigation at present and will be published separately. The origin of the increased electrical conductivity observed in the activated ceramics is presently not well understood and is most likely complex, but the literature detailing the nature of mechanochemically processed ceramics offers some possible explanations. In general, the mechanochemical activation disrupts the crystalline order of the powders, resulting in amorphous phases [37,44], as observed in the present case (Fig. 4). Owing to the large impact energies to which the powder particles are exposed during high-energy milling, the activation can induce micro-strains and crystal disorder in the treated powders. These micro-strains may persist, even after annealing at high temperatures, and may thus still be present in the final ceramics [48]. However, it was shown by theoretical analysis on CaMnO3 that tensile/compressive strains affect the formation energy of point defects, such as oxygen vacancies [49]. One can thus envision that the defect chemistry of the ceramics prepared mechanochemically might be modified owing to the presence of micro-strains. In turn, this may effect the electrical conductivity. Though possible, further focused studies would be required to confirm this scenario. Interestingly, the comparable S–E behavior of the nonactivated and activated ceramics suggests only a minor role of the different appearance of the Pbam phase (Fig. 6) and Pbam phase concentrations (Fig. 5) present in the two ceramics. Pinched P–E loops have been observed in thin film RE-modified BiFeO3 with near MPB compositions, the proposed origin consists of a field-induced Pnma–R3c (non-polar-to-polar) phase transition [4], with potential

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involvement also of a Pbam–R3c (antipolar-to-polar) phase transition [9]. However, there appears to be no evidence of this behavior in the macroscopic responses of these ceramics. In earlier studies [30,50,51], unmodified BiFeO3 (BFO) has been shown to behave as a “hard” ferroelectric material, exhibiting characteristic pinched P–E and S–E loops, aging and a strong dependence of the field-induced strain responses on the driving electric-field frequency. Fig. 8c and d compares the S–E loops of BFO ceramics (Fig. 8c) with those of BSFO (Fig. 8d) measured at driving electric-field frequencies in the range 0.1–100 Hz and at similar field amplitudes of 150 and 160 kV cm1, respectively. The BFO ceramics were prepared by a mechanochemically assisted method, the details of which are reported by Rojac et al. [30]. The S–E loops of the BFO ceramics, shown in Fig. 8c, exhibit a strong electric-field frequency dependence. This manifests as a variation in the remnant and peak-to-peak strain as a function of the electric-field frequency. The S–E loops measured at 100 Hz and 10 Hz are characterized by a strong decrease in the strain as the field is reduced from the maximum value to zero, resulting in no or small remnant strain and a small peak-to-peak strain (<0.1%). Electric-field frequencies of 1 Hz and 0.1 Hz result in remnant strains and significantly higher peak-to-peak strain (Spp 0.25% at 0.1 Hz), relative to that measured at higher frequencies (Spp 0.04% at 100 Hz). This frequency behavior was previously shown to be, in part, related to the electric-field-induced rearrangement of domain-wall pinning centers at low driving field frequencies [4]. Hence, while the domain walls are pinned on application of high field frequencies (e.g. 100 Hz), this rearrangement of pinning centers at low driving frequencies (e.g. 0.1 Hz) results in a

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facilitated domain switching, which manifests macroscopically in an increased remnant strain and high peak-to-peak strain. The S–E loops of the BSFO ceramics in Fig. 8d are strongly contrasted with those of BFO, exhibiting much less frequency dependence. By increasing the driving field frequency from 0.1 Hz to 100 Hz, the peak-to-peak strain of BSFO ceramic was reduced from 0.25% to 0.15% (Fig. 8d), while the loops maintained an open butterfly shape at all measured frequencies. This reduction in strain with increasing frequency is much smaller than that observed by BFO ceramics over the same frequency range (Fig. 8c). The comparison of the strain responses of BSFO with those of BFO, in particular the frequency dependence, indicates that the substitution of Sm3+ for the Bi3+ in BFO leads to a transition from a relatively “hard” behavior, characteristic of BFO, to a more “soft” behavior [52] of the BSFO. Assuming that the strain response of BSFO is largely related to the domain switching of the R3c phase, as apparently suggested by the P–E and S–E loops (Fig. 8a), this “hardening–softening” transition implies that the domain walls become more mobile by the addition of Sm in BFO. This is somewhat unexpected considering that Sm2O3 is an isovalent modifier, i.e. the Sm3+ substitutes for Bi3+ in the perovskite, both of which have the same oxidation state. In-fact, “softening” in PZT, e.g. by doping with Nb2O5, is believed to be associated with the aliovalent Nb2O5 doping, and therefore with the defects, e.g. lead vacancies (V00Pb ), that compensate for the donor site (Nbd Ti,Zr) [27,53], and/or reduces concentration of oxygen vacancies by donor doping [53]. While Sm is known to exist in both Sm3+ and Sm2+ oxidation states, for example, in iodides (SmI2 [54]), no evidence was found in the literature

Fig. 8. (a) Polarization vs. electric field hysteresis, (b) strain vs. electric field hysteresis, both at 100 Hz and 180 kV cm1 on non-activated (blue) and activated (red) ceramic BSFO. (c, d) Bipolar strain hysteresis loops measured at driving electric-field frequencies of 0.1, 1, 10 and 100 Hz: (c) for BFO ceramic at driving field strengths of 150 kV cm1; (d) for non-activated BSFO ceramic at driving field strength of 160 kV cm1. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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to suggest that it may be present as Sm2+ in BSFO or other oxide compounds and solid solutions. In addition, the Sm2+ would act as an acceptor, not a donor, and would be expected to lead to “hard” behavior, which is not consistent with the observed “softening” effect (Fig. 8c, d). While the precise reason for the observed “hardening– softening” transition induced by the Sm substitution in BFO is unknown at present, it is proposed that the proximity of the BSFO composition with the polar-to-non-polar phase boundary may provide a possible explanation. For example, the “hardening” effect in acceptor doped PZT is known to diminish when the composition approaches that of the MPB [55]. Alternatively, the domain walls of BFO have previously been observed to show increased electrical conductivity [56], and the strong frequency-dependent piezoelectric response of the BFO ceramics [57] was suggested to originate from these conductive domain walls. Thus, the improved frequency stability of the strain behavior in BSFO may relate to a possible impact that Sm has on the internal structure of the domain walls and their resulting electrical conductivity. While a macroscopic effect of Sm addition to BFO is observed here, the exact microscopic origin of the behavior is yet to be discovered.

4. Summary and conclusion The present paper has established an understanding of the relationship between the methods of synthesis used to produce BSFO ceramics and their resulting structure and phase composition. Different reaction pathways were observed by reactively sintering the non-activated and activated powder mixtures. The non-activated synthesis technique was diffusion controlled and produced ceramics by way of two intermediate phases, isostructural to R3c BiFeO3 and Pbnm SmFeO3 with subsequent inter-diffusion of Sm3+ and Bi3+ between these two phases with increasing reaction sintering temperature. The final non-activated ceramics had chemically inhomogeneous Pbam phase isolated grain inclusions within the R3c microstructure. In contrast, the activated ceramics formed apparently without the intermediate phases and consisted of chemically homogeneous R3c matrix phase coexisting with nano-clusters of Pbam phase. The primary source of the differences in the reaction pathway and the final sintered ceramics was identified as the amorphous, homogeneous and reduced crystallite size of activated powder oxide mixtures, which resulted from the mechanochemical activation. The comparison of the P–E and S–E hysteresis behavior of BSFO and BFO ceramics revealed that Sm substitution reduces the driving field frequency dependence of the electric-field-induced strain behavior at frequencies between 0.1 Hz and 100 Hz. The behavior of BSFO and BFO ceramics indicated that Sm-modification of BiFeO3 results in a “hardening–softening” transition. Acknowledgments This work was jointly funded by Australian Research Council ARC Grant LP 0991794, Thales Australia, Slovenian Research Agency Program “Electronic Ceramics, Nano, 2D and 3D Structures” (P2-0105) and Project Grant “High-Performance Piezoelectric Materials for Sensors and Actuators in High-Temperature Applications” (J2-5483).

Appendix A. Supplementary data Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/ j.actamat.2014.09.058.

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