Aurivillius-type ceramics, a class of high temperature piezoelectric materials: Drawbacks, advantages and trends

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Progress in Solid State Chemistry 37 (2009) 15e39 www.elsevier.com/locate/pssc

Aurivillius-type ceramics, a class of high temperature piezoelectric materials: Drawbacks, advantages and trends Alberto Moure a,*, Alicia Castro b, Lorena Pardo b b

a Instituto de Cera´mica y Vidrio, CSIC, C/Kelsen, 5, 28049 Madrid, Spain Instituto de Ciencia de Materiales de Madrid, CSIC, c/Sor Juana Ine´s de la Cruz, 3 Cantoblanco, 28049 Madrid, Spain

Abstract The obtention of reliable and high performance piezoelectric ceramics for uses at high temperatures is still an open issue in the field of electroceramics. The materials used nowadays for such applications present limitations due to different causes: low piezoelectric coefficients, difficulties in processing that lead to the necessary use of single crystals, high cost of raw materials and more. In this sense, an increasing interest in materials with the so-called Aurivillius-type structure has occurred during recent years, due to their relatively high piezoelectric coefficients and high ferroeparaelectric phase transition temperature. However, some difficulties must be overcome, such as processing for obtaining highly dense ceramics and determining their real piezoelectric behaviour at high temperature. In this work, a review of the processing and properties of ceramics with this structure is shown. Effects of the use of precursors obtained by an alternative route mechanical activation on the microstructure are explained. A complete piezoelectric characterization at working temperatures (>300  C), barely found in the literature, is also shown. The effects of trapped charges in the dielectric permittivity and in the piezoelectric radial resonance are also discussed. Ó 2009 Elsevier Ltd. All rights reserved. Keywords: Functional ceramics; Mechanochemical activation; Sintering; Hot pressing; Recrystallization; Microstructure; Electrical properties; Piezoelectric properties; Bismuth layered structure ceramics

Contents 1.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

* Corresponding author. Fax: þ34 91 735 5843. E-mail address: [email protected] (A. Moure). 0079-6786/$ - see front matter Ó 2009 Elsevier Ltd. All rights reserved. doi:10.1016/j.progsolidstchem.2009.06.001

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2.

Processing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Processing of ceramic precursors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. Processing of ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

19 19 21

3.

Electrical studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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4.

Piezoelectric studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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5.

General conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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1. Introduction The study of piezoelectric materials for uses as actuators, sensors and transducers at high temperature (>300  C) is a continuous issue in the field of electroceramics. The most commonly used piezoelectric material operating at high temperature is quartz (SiO2), in the form of single crystals [1]. It presents a high resistivity and its properties are practically independent of the temperature. However, the piezoelectric coefficients are low (d11 ¼ 2.3 pC/ N), and the maximum working temperature is limited by a structural phase transition at 573  C. Besides, as is well known, the growth of defect-free single crystals and their further orientation are difficult and expensive processes. Other single crystals that have been tested lately, such as lithium tetraborate, Li2B4O7, or gallium phosphate, GaPO4 [1] share these drawbacks. Some ferroelectric materials, such as LiNbO3 and LiTaO3, are proposed for these applications. They have a relatively high ferroeparaelectric phase transition but the very difficult processing of ceramics for such compositions [2] also limits their applicability. Similar problems are found for relaxor materials such as those with Pb(Zn1/3Nb2/3)O3e PbTiO3 (PZNePT), Pb(Mg1/3Nb2/3)O3ePbTiO3 (PMNePT) or Pb(Yb1/2Nb1/2)O3ePbTiO3, for which their processing as ceramics is difficult as well, although nanostrucutred ceramics have been recently achieved by spark plama sintering [3]. They have effective d33 coefficients as high as 2500 pC/N along the h001i direction in single crystals with composition close to their MPBs [4,5]. To achieve a high piezoelectric activity, it is necessary to texture the ceramics, for example by template grain growth [6,7]. Values of d33 in textured ceramics can be 1.8 times higher than for randomly oriented ones. A piezoelectric coefficient d33 higher than 1600 pC/N has been measured [6] for 0.675Pb(M1/3Nb1/3)O3e0.325PbTiO3. Furthermore, although their TC can reach values up to 350  C (for Pb(Yb1/2Nb1/2)O3ePbTiO3), their piezoelectric activity vanishes at temperatures lower than 200  C [8,9]. This makes their use as piezoelectric at T > 300  C impossible. Thus, poled polycrystalline ferroelectrics for uses as high temperature piezoceramics must fulfil three conditions: a high ferroeparaelectric phase transition temperature, high piezoelectric coefficients, and an easy processing of the ceramics from relatively low-cost raw materials. Moreover, other issues must be solved, such as the existence of other phase transitions, the pyroelectric effect in some resonance modes and the high conductivity produced by carrier transport at temperatures that interfere with the charge induced by piezoelectricity. Recent investigations have been carried out on various (1  x)PbTiO3e(x)Bi(Me0 Me00 )O3 solid solutions, where the Me0 and Me00 cations occupy the octahedral sites of the perovskite structure in ratios that give a þ3 average valence. Materials with TC in the range of 500  C can

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be obtained for a given composition [10]. Piezoelectric coefficients are optimized for MPB compositions [11e15], where phases with different distortions of the perovskite structure (such as rhombohedral and tetragonal) co-exist. Values of d33 ¼ 460 pC/N, and a TC of 450  C have been obtained in the system (1  x)BiScO3exPbTiO3 for x ¼ 0.64 [11]. Limitations to these solid solutions come from the processing. It has been reported that further increase of TC in these systems can be achieved at the expense of reducing the tolerance factor and thus the stability of the perovskite phase [16]. This makes the processing more difficult as single phase ceramics, which can only be obtained by pressure-assisted methods. A solution can be the synthesis by mechanical activation, which allows ceramics with low tolerance factor to be obtained by conventional sintering as recently proved [17]. However, although they have high TC, it is not clear to date that piezoelectric activity is maintained up to this temperature. It has been proved that piezoelectric coefficients drop to zero at temperatures of TC/2 for Bi(Mg1/2Ti1/ 2)O3ePbTiO3 [18] ceramics. Another family of ferroelectric materials with even higher ferroeparaelectric transition temperature for some of their components, and, which is therefore, a good candidate for its use as high temperature piezoelectrics (even higher than 500  C), has the Aurivillius-type structure. These materials have as general formula (Bi2O2)(An1BnO3nþ1) and the crystalline structure is built by alternating layers of [Bi2O2]2þ and pseudo-perovskite blocks, with n octahedral layers in thickness [19,20]. The major component of the spontaneous polarization lies in the ae b plane of the perovskite-like layers for even n values, while a component in the c-axis is also found for odd n values [21,22]. The main difficulty for their practical applications is to obtain mechanically stable and dense ceramics. Its lamellar crystal growth makes the compaction of the precursor ceramic powder more difficult, and the reduced mass transport during sintering leads to porous ceramics. Hot pressing and hot forging of Aurivillius compounds have been traditionally used to produce highly dense ceramics [23e26] with anisotropic properties because of the developed textured microstructure, in which the lamellar grains align with the c-axis (with lower or null ferroelectric activity) parallel to the applied pressure (Fig. 1). Dense ceramics with high degrees of texture have also been obtained lately by spark plasma sintering (SPS) [27e29] at lower processing temperatures and times than by hot pressing. The enhancement of the piezoelectric activity is clear. The d33 values are increased by a factor higher than 2 for textured ceramics with respect to the ordinary fired ones [27]. Recently, a pressure-less route by slip casting in a magnetic field (higher than 10 T) have also been employed to obtain textured Aurivillius ceramics [30,31]. The advantage with respect to pressure-assisted methods is that the ferroelectric aeb plane can be oriented in more favourable directions, for example, perpendicular to the major faces of a disk, avoiding further mechanization of the samples for piezoelectric applications. High degrees of orientation were obtained (Lotgering factor up to 0.88). Again the piezoelectric activity achieved gives d33 values which are more than double for textured ceramics [30] than the conventionally sintered ones. However, the texture obtained by any of these methods is detrimental for the simultaneous appearance of good mechanical and ferroelectric properties. In the case of pressure-assisted methods, the highest ferroelectric and piezoelectric activities occur in samples cut in parallel to the applied pressure (Fig. 1), with the lamellar grains, and the aeb polar plane, aligned in parallel to the thickness, t, of the sample. Nevertheless, the samples that are cut in parallel (Fig. 1) have poor mechanical properties as shown in the literature for textured Aurivillius ceramics: bend strength in PbBi2Nb2O9 [32], Pb4Bi4Ti7O21 and Bi4Ti3O12 [33] or the fracture toughness in PbBi2Nb2O9 [25], and SrBi2Nb2O9 [34]. When grains are aligned in the way

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Fig. 1. Ideal orientations of the lamellar grains under hot pressing of an Aurivillius-type structure ceramic and scheme of parallel (to the applied pressure) cut sample.

shown in Fig. 1, the paths of crack propagation are easier, along grain boundaries, in parallel to the aeb planes. Propagation is more difficult in the perpendicular direction (c-axis), where there is not a grain alignment. The fracture through the grains due to the crystalline structure [32] also contributes to the damage of the mechanical properties in the polar direction, inside aeb planes, of parallel cut samples. An additional difficulty is that the direction with the largest piezoelectric activity is the same as the maximal conductivity, which makes the unavoidable poling process and the application at high temperature more difficult. Alternative routes must be used to obtain the ideal microstructure of the ceramics: simultaneously non-textured, so isotropic, and dense ceramics with an appropriate grain size to make the poling process easier. The processing of dense ceramics can be improved by the use of alternative synthesis methods, such as co-precipitation [35,36], hydrothermal synthesis [36,37], molten salt synthesis [38,39], solegel [40,41], synthesis from metaleorganic polymeric precursors [42] or recently the so-called citrate-gel method [43,44]. However, these chemistry-based routes require high purity inorganic or organometallic reactants that are sensitive to light or humidity and are more expensive than the widely available oxides and carbonates. A cleaner and economic alternative is the use of mechanical activation. In this method, part of the energy necessary to produce the chemical reaction is supplied by mechanical means. During the process, the particle size of the crystals is reduced and the homogeneity of the mixture is increased. It improves the reactivity of the precursors [45] and allows the compaction of the green pellet to be optimized. The processing of highly dense ceramics [46e50] is possible in a single thermal process, in which synthesis, grain growth and sintering take place. A thorough review of the application of this route to ferroelectric materials in general has been recently published [51]. Aurivillius-type ceramics studied and reviewed in this work were processed by this method. The study of electric properties of these compounds is a relatively well-known field. Most of the work has been centred on Bi4Ti3O12-based compounds (n ¼ 3) [52e55], or on other systems with n ¼ 2, such as SrBi2Ta2O9 [56]. However, the study of the piezoelectric properties at working temperatures (>300  C) has been usually omitted in most of the published literature. Normally, the piezoelectric characterization for these materials is based on the values of the coefficients at room temperature. This is not really representative of the performance of these

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materials at working conditions, since these materials are not competitive for piezoelectric applications at low temperature ranges, but at higher ones. As the temperature increases, the piezoelectric activity can decrease with respect to that measured at room temperature [57]. This will be discussed in this review. A few works have been recently published which take into account these features. Piezoelectric activity as a function of temperature has been reported for structures with n ¼ 2 [58,59] and n ¼ 4 [60]. Some data have also been given by measuring piezoelectricity at room temperature after annealing at different temperatures [61]. This only gives an idea about the depoling temperature, but not about the piezoelectric activity at those temperatures. If the ceramic is not depoled, the piezoelectricity as a function of temperature on heating and cooling is the same, but the values of the piezoelectric coefficients at room temperature are different than at higher temperatures [62]. A well-known characteristic of these compounds is that they present an anomaly in the dielectric permittivity and dielectric loss factor at temperatures below the ferroeparaelectric phase transition, especially at low frequencies. Firstly, some authors attributed it to ferroe ferroelectric phase transitions [63] or spaceecharge relaxation [64]. It seems to be well established that a contribution to the origin of this anomaly is related with oxygen ion-jump mechanisms that occur preferentially in the aeb plane [65,66]. The dielectric anomaly below the ferroeparaelectric transition could limit the working temperature of the ceramics, even when the material is still piezoelectric. However, to the best knowledge of the authors, current literature ignores the influence of these anomalies in the piezoelectric resonances at high temperature. Our studies on the effect of the mechanical activation in the processing of Aurivillius-type ceramics, applied to structures with n ¼ 2 based on (Bi3TiNbO9)1xe(SrBi2Nb2O9)x system with x ¼ 0.65 and 1.00 (TC between 760 and 910  C, approximately) is shown in this work, along with a critical review of the electrical, piezoelectric and resonance characteristics of Aurivillius-type structure ceramics at the expected working temperature in comparison with recently published literature. 2. Processing An alternative route based on mechanical activation was used to obtain the precursors of ceramics based on Bi3TiNbO9eSrBi2Nb2O9. The results will be shown in the following section. For the processing of the ceramics, three routes were employed: conventional sintering, hot pressing and, as an alternative, a thermal treatment called here recrystallization after hot pressing. 2.1. Processing of ceramic precursors Ceramics with Bi3TiNbO9 (BTN) and (SrBi2Nb2O9)0.35(Bi3TiNbO9)0.65 (SBN/BTN 35/65) compositions were obtained from mechanically activated precursor powders. These precursors were prepared by energetic milling in a vibrating mill of a stoichiometric mixture of analytical grade Bi2O3, Nb2O5, TiO2, and SrCO3 for solid solution SBN/BTN 35/65 members. The mixture was placed in a stainless-steel pot with a 5 cm steel ball in a vibrating-type mill (Fritsch Pulverisette 0). Milling was carried out during 3 weeks for BTN and 5 weeks for SBN/BTN 35/ 65. Fig. 2 shows the effect of the milling on the mixtures. Complete amorphous powders, according to the X-ray diffraction (XRD), were obtained for BTN precursors [48], while remaining peaks corresponding to oxides and carbonates of the initial mixture are present for

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20

a

b

336h

*

840h 168h 672h

Intensity (a.u.)

24h

504h

+

3h

0h

10

20

30

2θ (°)

40

50

60

10

20

30

40

*

+

**

0h

*

Intensity (a.u.)

72h

50

60

2θ (°)

Fig. 2. XRD patterns of the stoichiometric mixture: a) Bi2O3, TiO2, Nb2O5 and b) the same oxides with SrCO3, precursors of BTN and SBN/BTN 35/65 ceramics, after different milling times. [(Symbols: TiO2 (þ), SrCO3 (*), Nb2O5 (C), Bi2O3 (-).)]

composition SBN/BTN 35/65. The presence of SrCO3 limits the action of the mechanical activation. In any case, the reactivity of the mixture is increased by the prolonged milling. Fig. 3 shows the XRD patterns of the mechanically activated mixtures after several thermal treatments. The crystallization temperature is reached at 700  C (a reduction of 350  C with respect to traditional solid-state method) for BTN and at 800  C (reduction of 250  C) for SBN/ BTN 35/65. Furthermore, the particle size of the mechanically activated precursors even after thermal treatment is much lower than that obtained by traditional solid-state reaction [67], which favours the compact of the green pellet and the densification of the ceramics, as will be shown below. This crystallization temperature is lower than that generally obtained for Aurivillius ceramics. A further decrease in the synthesis temperature has been recently published. For example, Du et al. have isolated pure Bi4Ti3O12 phase at 550  C by one step aqueous sole gel method [41], although the effect of this route in the sintering is not shown. Synthesis at temperatures as low as 500  C for CaBi4Ti4O15 have been recently obtained by Deshmukh et al. from a mixture of Ca(OH)2, Bi(OH)3 and Ti(OH)4 [68]. This means a reduction of 300  C with respect to traditional solid-state reaction. Unfortunately, no data of sintered ceramics are given. The effect of the loss of (OH) groups during sintering may be a problem for good densification of the ceramic. Du et al. obtained the same synthesis temperatures for Pr and V doped Bi4Ti3O12 by the citrate-gel method [44], allowing a decrease in sintering temperatures to only 900  C-4 h. Good ferroelectric properties (2Pr and 2Ec values were 35 mC/cm2 and 148 kV/cm, respectively, measured at 300 kV/cm) are shown in his work. All of these routes are chemistry based, with the disadvantages pointed out in the introduction of this review, but these examples show the effort in the study of Aurivillius ceramics to reduce the processing temperatures and improve the sinterabilty of precursors. As shown in the following section, combined mechanical activation and processing of recrystallization after hot pressing fulfils these conditions.

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39 P

FRT

P PPP P

P

P

21

P P P P

P

P

800ºC

**

Intensity (a.u.)

700ºC

600ºC

500ºC F F

390ºC

F

F

330ºC 250ºC 100ºC RT

10

20

30

40

50

60

70

2θ (°) Fig. 3. XRD patterns at increasing temperatures of the stoichiometric mixture: a) Bi2O3, TiO2, Nb2O5 and b) the same oxides with SrCO3 milled during 5 weeks. [Aurivillius phase (P), Fluorite phase (F), SrCO3 (*), Nb2O5 (C), Bi2O3 (-); RT: room temperature; FRT: final room temperature.]

2.2. Processing of ceramics To obtain the ceramics, the precursor powder was shaped by uniaxial pressing as disks of approximately 10 mm diameter and 2 mm thickness, which were then isostatically pressed at 2000 kg cm2. Fig. 4 shows the XRD patterns for BTN (a) conventionally sintered, and (b) hotpressed ceramics. The peak intensities of the former are the same as for the powder, indicating that sintering produces isotropic ceramics. Fig. 4b reveals the development of (00l ) textured ceramics at certain processing temperatures. The XRD patterns of the hot-pressed ceramics at temperatures lower than 1050  C are similar to the sintered ones, indicating that the ceramic grains grow without any preferential direction. At 1050  C it is observed that the (00 10) peak has a higher intensity than the others, indicating that at this temperature the grains are arranged with the c-axis parallel to the applied pressure, which is typical of Aurivillius-type ceramics. As said, this texture is detrimental to the simultaneous appearance of good mechanical and ferroelectric properties. The limit for the application of BTN ceramics comes from their low piezoelectric coefficients (d33 ¼ 6 pC/N) [69], related with the high ferroeparaelectric phase transition temperature. Solid solutions of Bi3TiNbO9 with SrBi2Nb2O9 are expected to reduce the temperature and make the poling process easier. Fig. 5 shows the XRD patterns of the sintered and hot-pressed ceramics at different temperatures for SBN/BTN 35/65 compositions. It also shows the results of the ceramics firstly hot pressed and then recrystallized to higher temperatures, with the aim of increasing the grain size while maintaining the high densification. Three features must be pointed out: the target

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22

a

b

115

0010

1100ºC-2h

006

020 214 0010 018 119 026

111 008 113

220 2010

135

1050 ºC - 1h

Intensity (a.u)

Intensity (a.u)

006

115

1100ºC-1h

1050ºC-2h

111

008

020

214

018

119

113

026

2010 135

220

1000 ºC - 1h 900 ºC - 1h

1050ºC-1h 850 ºC - 1h 1000ºC-1h

10

20

700 ºC - 1h

30

40

50

60

10

20

30

2θ (º)

40

50

60

2θ (º)

Fig. 4. XRD patterns of BTN ceramics: a) conventionally sintered and b) hot pressed, from mechanically activated precursors.

phase is crystallized in only one step (hot pressing or sintering), the hot-pressed ceramics are not textured at low temperatures (up to 1000  C), and that when firstly hot-pressed isotropic ceramics are thermally treated at higher temperatures without external pressure, the resulting ceramics are also isotropic. Table 1 shows the results of the quantitative microstructure analysis of sintered, hot-pressed and recrystallized ceramics after hot pressing of the SBN/BTN 35/65 composition. Several features can be remarked: a) sintered ceramics show a limit to achieve porosity lower than 5%, since a minimum of porosity of 7% is obtained at 1200  C-2 h. Higher sintering temperatures

020 018 00 10

115

214 026 119

1100 ºC - 1h 220

2010

135

1050 ºC - 1h 1000 ºC-HP + 1100 ºC - 1h

Intensity (a. u.)

1000 ºC-HP + 1050 ºC - 1h

*

1000 ºC-HP 900 ºC-HP + 1150 ºC - 1h 900 ºC-HP + 1100 ºC - 1h 900 ºC-HP + 1050 ºC - 1h 900 ºC-HP + 1000 ºC - 2h 900 ºC-HP

20

30

40

50

60

2θ θ (º) Fig. 5. XRD patterns of sintered, hot pressed (HP) and recrystallized after hot-pressing SBN/BTN 35/65 ceramics from mechanically activated precursors. [Peak marked with (*) corresponds to the alumina powder used during hot pressing.]

Thermal treatment

1050  C-1 h sintered 1100  C-1 h sintered 1150  C-2 h sintered 1200  C-2 h sintered HP at 900  C1 hþ As-HP 1000  C-2 h 1050  C-1 h 1100  C-1 h 1150  C-1 h HP at 1000  C1 hþ As-HP 1050  C-1 h 1100  C-1 h

Pore area distributions Area (mm2)

sA (mm2)

Porosity (%)

Number of Grain size distributions measured pores Area sA (mm2) (mm2)

2.6  0.2 3.7  0.9 3.8  0.5 4.4  0.9 1.8  0.1 2.2  0.1 2.5  0.2 2.9  0.2 6.2  1.1 2.5  0.1 2.8  0.2 2.8  0.2

3.3  0.4 3.3  1.2 3.1  0.7 7.3  3.3 1.0  0.1 1.7  0.1 1.7  0.1 1.9  0.2 5.8  1.3 1.6  0.1 1.6  0.2 3.5  0.3

39.1  3.5 18.4  1.1 13.4  0.1 7.1  0.4 0.20  0.01 1.4  0.2 3.7  0.6 3.3  0.2 6.3  0.4 0.51  0.09 1.0  0.1 3.5  0.6

1191 1607 1205 1133 172 924 1703 1327 1147 271 519 1363

0.24  0.01 3.7  0.3 6.8  0.8 8.1  1.1 0.22  0.01 2.04  0.04 5.1  0.9 11.3  0.8 18.1  3.4 0.53  0.01 4.1  0.3 7.5  0.6

0.42  0.04 5.6  0.7 10.6  1.8 13.0  2.5 0.28  0.02 2.67  0.07 10.1  2.4 21.0  1.9 33.2  7.9 0.73  0.02 4.9  0.5 10.9  1.2

Dmax =Dmin

sR

1.51  0.07 1.60  0.08 1.92  0.08 2.22  0.1 1.58  0.06 1.61  0.07 1.8  0.08 1.94  0.07 1.88  0.13 1.75  0.13 1.47  0.06 1.7  0.1

0.36  0.03 0.47  0.04 0.59  0.04 0.89  0.04 0.38  0.03 0.46  0.04 0.79  0.06 0.74  0.03 0.65  0.07 0.76  0.09 0.35  0.03 0.47  0.05

Number of d33 (pC/N) measured grains

252 1108 746 965 831 895 1396 1118 772 808 1161 855

e e 20 18 7 e e e 13 7 13 12

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

Table 1 Microstructural parameters of sintered, hot-pressed and recrystallized SrBi2Nb2O9eBi3TiNbO9 35/65 ceramics from mechanically activated precursors.

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cannot be used because they lead to abnormal grain growth [69]; b) hot pressing at moderate temperatures produces almost fully dense ceramics (porosity lower than 1%). However, lower values of d33 are achieved at room temperature (7 pC/N) with respect to sintered ceramics (up to 20 pC/N, higher than for BTN), due to the small grain size which is obtained (<1 mm2); c) recrystallization allows ceramics with lower porosity than the sintered ones and larger grain size than the hot-pressed ones to be obtained. That is, the ideal microstructure is obtained through control of the processing parameters: initial hot pressing and final recrystallization temperatures. The final microstructure depends on the initial conditions. Recrystallized ceramics after hot pressing at 900  C-1 h develop grains with a larger area and aspect ratio than those recrystallized after hot pressing at 1000  C-1 h. Movement of grain boundaries by mass diffusion across them during recrystallization is favoured in the microstructure of the ceramic hot pressed at 900  C-1 h, which has higher grain boundary density per volume. The increase in the porosity content after recrystallization is a consequence of the deterioration of the microstructure after hot pressing, as well as in the final stage of sintering when high temperatures or long times are used. It has been proposed in other materials, such as calcium-modified lead titanate [70] or in alumina [71], that deterioration is related to the ratio between the rates of grain growth and densification. The higher the ratio, the more the microstructure is deteriorated. The densification rate during the recrystallization is very low. After hot pressing, a fully dense ceramic is obtained, and the energy supplied during the recrystallization is mainly employed in grain growth. Thus, the ratio increases, and deterioration takes place. As the recrystallization temperature increases with respect to the hot-pressing temperature, the grain growth rate increases exponentially with temperature [72], and the deterioration is more important. The porosity increases due to the appearance of new pores and to the increase in the pore area as the grain grows [73]. Fig. 6 shows more graphically this behaviour, for the ceramics initially hot pressed at 1000  C. Similar features are observed for the ceramics recrystallized after hot pressing at 900  C [73]. The obtained microstructure produces an increase of the d33 coefficient with respect to the initial hot-pressed ceramics. Values of 13 pC/N at room temperature are obtained for recrystallization after hot pressing of the ceramics. The evolution of electromechanical and piezoelectric coefficients with increasing temperature are shown in Section 4 of this work. A significant improvement in elastic properties is also produced, which will be clearly shown in the same section. As a summary, mechanical activation allows precursors of Bi3NbTiO9 and Bi3NbTiO9e SrBi2Nb2O9 ceramics with high reactivity to be obtained. Ceramics can be processed by a single thermal treatment at relatively low temperatures. However, a limit of temperature and time exists for obtaining dense ceramics with homogenous microstructure. Hot pressing from amorphous precursors provides highly densified BTN ceramics with controlled texture: isotropic and fine grained at low temperatures, 700e1000  C, and textured and coarse-grained at 1050  C for 1 h. Textured ceramics have poor mechanical properties, while non-textured have low piezoelectric coefficients. Isotropic ceramics of nominal composition SrBi2Nb2O9eBi3TiNbO9 35/65 with controlled microstructure are obtained in a process called recrystallization after hot pressing. Porosity increases as the difference between recrystallization and initial hot-pressing temperatures increases, but it is lower than that of the sintered samples at the same or higher temperatures. An increase in the recrystallization temperature produces the increase in grain area and aspect ratio that also depends on the initial hot pressing and final recrystallization temperature.

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25

Fig. 6. Optical micrographs of the polished surfaces of (1) hot-pressed ceramic at 1000  C-1 h and the corresponding recrystallized ceramics at (2) 1050  C-1 h and (3) 1100  C-1 h; SEM and optical micrographs of polished and thermally etched surfaces of (4) hot-pressed ceramic at 1000  C-1 h and the corresponding recrystallized at (5) 1050  C-1 h and (6) 1100  C-1 h.

3. Electrical studies Impedance measurements between 100 Hz and 5 MHz as a function of temperature (from 200 to 950  C) were performed with a HP-4194A analyser on parallel-faced ceramics with painted Pt electrodes sintered at 700  C, in an experimental set-up for the temperature control described elsewhere [26]. The heating and cooling rates were 2  C/min, stabilizing the temperature for 1 min. From these experimental data and the sample geometry, complex permittivity, 3* ¼ 30  i300 , and dielectric loss factor, tan d, were obtained. Fig. 7 shows the real part of the dielectric permittivity and the loss factor for sintered Bi3TiNbO9 ceramics at high frequency (1 MHz), while Fig. 8 shows the same features at low frequency (1 kHz). With the aim of comparing the effect of the microstructure on the permittivity, Fig. 9 shows the same parameters for hot-pressed isotropic ceramics, that is, those processed at 700, 900 and 1000  C, at 1 MHz. Fig. 7 and 9 clearly show the ferroeparaelectric phase transition temperature, with a maximum in permittivity located at T w 905e910  C. Microstructure has a twofold influence: porosity and grain size and shape. A decrease of the porosity results in an increase of permittivity, if sintered and hot-pressed ceramics are compared. The sample with higher porosity (sintered at 1000  C-1 h) has the lower permittivity at lower temperature (30 ¼ 88 at 300  C). The value of 30 at the same temperature is 258 for the ceramic hot pressed at 700  C1 h. In general, all the hot-pressed ceramics (with porosity lower than 1%) have higher permittivity than the sintered ones.

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

26

a

1000

BTN ceramics sintered at: 1000 °C - 1 h. P>25% 1050 °C - 1 h. P=10.9% 1050 °C - 2 h. P=13.9%

ε'

1100 °C - 2 h. P=8.5%

100

b

1 MHz

10

tan δ

1

0.1

0.01

200

300

400

500

600

700

800

900

Temperature (ºC) Fig. 7. Real part of dielectric permittivity and dielectric loss factor at 1 MHz as a function of temperature for Bi3TiNbO9 sintered ceramics (solid symbols: heating; hollow symbols: cooling; P: porosity).

The grain size has an influence in the values of permittivity and also in the shape of the phase transition. It can be observed if isotropic ceramics hot pressed at 700, 900 and 1000  C (Figs. 7e9) are compared during the heating run measurements, because porosity is similar and its effect can be neglected. The permittivity decreases as the hot-pressing temperature increases. The values of 30 at 240  C are 250, 201 and 134 for the ceramics hot pressed at 700, 900 and 1000  C, respectively. The reduction in the aspect ratio of the grains influences the dielectric permittivity. Thus, the increase in 30 in fine grained hot pressed, Aurivillius-type ceramics is related with the anisotropic growth behaviour of these structures [74], which hot pressing inhibits. On the other hand, the grain size is also responsible for the broadening in the transition. A diffuse phase transition is produced due to an inhomogeneous tension distribution originated by a reduced grain area [75]. There is a slight variation of the temperature for the maximum of permittivity from one grain to another. The observed transition would be the convolution of the different transition anomalies. The dependence of the permittivity and dielectric loss factor with frequency is observed when Figs. 7 and 8 are compared. The characteristic peak corresponding to the phase transition

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

a

100000

27

BTN ceramics sintered at: 1000 °C - 1 h. P>25% 1050 °C - 1 h. P=10.9% 1050 °C - 2 h. P=13.9%

10000

ε'

1100 °C - 2 h. P=8.5%

1000

b

100

1 KHz

100

tan δ

10

1

0.1

0.01 200

300

400

500

600

700

800

900

Temperature (°C) Fig. 8. Real part of dielectric permittivity and dielectric loss factor at 1 kHz as a function of temperature for Bi3TiNbO9 sintered ceramics (solid symbols: heating; hollow symbols: cooling; P: porosity).

cannot be distinguished at 1 kHz. At that frequency, a shoulder can be discerned between 400 and 600  C, in both heating and cooling runs. In the ceramic sintered at 1000  C-1 h an additional anomaly located at higher temperature (650e750  C) is observed. The anomalies observed here seem to have a principal explanation of a ferroeferroelectric phase transition, reported by other authors [63,76] in the same range of temperature for this composition, determining most probably the appearance of the reversible anomalies (in the range of 400e 600  C) in the dielectric properties of BTN ceramics. The ferroeparaelectric phase transition temperature and the dependence of the permittivity with temperature in SBN/BTN 35/65 ceramics are shown in Fig. 10. Recrystallized ceramics with controlled microstructure were chosen to make the comparison. The ceramics to be compared are named as: type A (hot pressed at 900  C and then recrystallized at 1150  C), type B (1000 þ 1050  C, with the highest conductivity) and type C (1000 þ 1100  C, with the lowest conductivity). Measurements at 1 MHz and 1 kHz are presented. The maximum in permittivity is located around 700e750  C, and increases as the difference between the recrystallization and

28

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

a

BTN ceramics hot pressed at: 700 °C - 1 h. P=0.5% 900 °C - 1 h. P=0.4% 1000 °C - 1 h. P=0.5%

ε'

1000

b

1 MHz 100

tan δ

1

0,1

0,01 200

300

400

500

600

700

800

900

Temperature (ºC) Fig. 9. Real part of dielectric permittivity and dielectric loss factor at 1 MHz as a function of temperature for Bi3TiNbO9 hot-pressed ceramics (solid symbols: heating; hollow symbols: cooling; P: porosity).

initial hot-pressing temperatures increases. It has been proposed that an increase of internal stress decreases the TC value [75]. The increase in the grain size and the porosity development explained above allows the internal stress to be relaxed, and as consequence, TC increases. As for BTN, reversible anomalies are observed in the low frequency measurement, but with a more pronounced hysteresis. A double contribution to the anomalies is found. One (reversible) seems to come from ferroeferroelectric phase transition as for BTN. The other one, responsible also for the hysteresis has another origin. It has been proposed [77] in similar solid solutions that oxygen vacancies and the associated defect dipoles are accumulated at grain boundaries and domain walls, and they are frozen at room temperature. In Aurivillius-type ceramics, the generation of oxygen vacancies is assumed to be related with the evaporation of Bi [78]: Bi2 O3 ¼ 2Bi þ 3=2O2 þ 2V000 € Bi þ 3VO

ð1Þ

The defect dipoles are constituted as: VO¨eV000 Bi or VO¨eO00 ad, where O00 ad are absorbed oxygen ions. At high temperatures, or when the domains disappear in the paraelectric phase, those

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

29

ε'

1000

100

1 MHz SBN/BTN 35/65 Type A (P=6.3%. = 18.1 μm2) Type B (P=1.0%. = 4.1 μm2) Type C (P=3.5%. = 7.5 μm2)

ε'

1000

100

1 KHz 200

300

400

500

600

700

800

Temperature (ºC) Fig. 10. Real part of dielectric permittivity at: a) 1 MHz, and b) 1 kHz as a function of temperature for SrBi2Nb2O9e Bi3TiNbO9 35/65 recrystallized after hot pressing. Type A) 900  C hot-pressed ceramics and recrystallized at 1150  C1 h; Type B) 1000  C hot-pressed ceramics recrystallized at 1050  C-1 h; and Type C) 1000  C hot-pressed ceramics recrystallized at 1100  C-1 h (solid symbols: heating; open symbols: cooling; P: porosity; hAGi: mean grain area).

defect dipoles are de-iced, and they contribute to an increase of the dielectric permittivity. The defects need longer times (days) to return to their initial iced state. This mechanism explains the thermal hysteresis and is a contribution to the observed anomalies of the dielectric permittivity on heating. Data of the dc conductivity were obtained by fitting impedance arcs, Z0 eZ00 , using the EQUIVCRT program [79]. This program uses a fitting of non-linear least squares that allows the simultaneous determination of all the parameters of the equivalent circuit, as it is described in detail in [80]. These equivalent circuits comprise two or three (depending on the temperature) RQC circuits in series, each one consisting of a resistance (R) in parallel to a capacitor (C ) and to a constant phase element (Q). Fig. 11 shows the Arrhenius plot of the dc conductivity of the same three ceramics as Fig. 10. The scheme of the equivalent RQC circuit has been added. The control of the microstructure explained above allows the effect of grain size and shape and porosity in the conductivity of the ceramics to be studied. For the three ceramics, two steps in the conduction

SBN/BTN 35/65. Type A ceramics P=6.3 %. G=18.1 μm2

1E-4 1E-5

σdc (S/cm)

R Ea=1.63 e.V.

1E-6 1E-7

C

Ea=1.80 e.V. Q

1E-8

Ea=0.39 e.V.

1E-9

Ea=0.37 e.V. 0,9

1,1

1,3

1,5

1,7

1,9

1000/T (1/K)

SBN/BTN 35/65. Type B ceramics P=1.0 %. G=4.1 μm2

1E-4

σdc (S/cm)

1E-5 1E-6

Ea=1.50 e.V.

1E-7

Ea=1.59 e.V.

1E-8

Ea=0.39 e.V.

1E-9

Ea=0.85 e.V. 0,9

1,1

1,3

1,5

1,7

1,9

1000/T (1/K) SBN/BTN 35/65. Type C ceramics P=3.5 %. G=7.5 μm2

1E-4

σdc (S/cm)

1E-5 1E-6

Ea=1.79 e.V. Ea=1.82 e.V.

1E-7

Ea=1.00 e.V.

1E-8 Ea=0.51 e.V.

1E-9 0,9

1,1

1,3

1,5

1,7

1,9

1000/T (1/K) Fig. 11. Arrenhius plot of the d.c. conductivity for SrBi2Nb2O9eBi3TiNbO9 35/65 ceramics recrystallized after hot pressing. Type A) 900  C hot-pressed ceramics and recrystallized at 1150  C-1 h; Type B) 1000  C hot-pressed ceramics recrystallized at 1050  C-1 h; and Type C) 1000  C hot-pressed ceramics recrystallized at 1100  C-1 h (solid symbols: heating; open symbols: cooling; P: porosity; hAGi: mean grain area).

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

31

can be observed. In both steps, the conductivity has a linear relationship with 1/T, with different slopes. A slight thermal hysteresis occurs for the three ceramics, and the conductivity decreases during the cooling run. Fig. 11 also shows the activation energy of the conductivity in each step of the Arrhenius plot, which presents some thermal hysteresis, being smaller during the cooling run in the low temperature region and higher in the high temperature range. The dependence of the activation energy with the microstructure has the same tendency as the conductivity. The steps shown in the Arrhenius plot are common in Aurivillius-type ceramics [81e83]. At low temperature, the conduction is dominated by defects such as extrinsic unintentionally introduced impurities [83], which in this case can be small amounts of Fe coming from the prolonged milling. At high temperature, it is dominated by intrinsic defects (migration of oxygen vacancies, with activation energy values [82], similar (w1.4 eV) to those shown in Fig. 11). The decrease in its concentration results in an increase in the transition temperature between intrinsic (high temperature) and extrinsic (low temperature) conduction [81]. The recrystallization at higher temperatures produces an increase in grain size (Table 1), and therefore, the reduction of the concentration of defects [84]. This would produce the increase in the temperature of the intrinsiceextrinsic dominated transition, as is confirmed by the results for the type A and B ceramics (Fig. 11). Summarizing this section, sintered isotropic ceramics have a lower dielectric permittivity than the hot-pressed ceramics, due to their higher porosity. The permittivity increases as the area and the aspect ratio of the grains is reduced, due to the anisotropy of the dielectric properties related with the crystalline structure. The observed anomalies in their dielectric permittivity and losses below the ferroeparaelectric phase transition temperature are due to a structural phase transition (reversible) and to the de-icing of defect dipoles located close to domain walls (irreversible). The Arrhenius plot shows a change between the domination of intrinsic defects and extrinsic impurities at temperatures that depend on the characteristics of the ceramic microstructure, mainly the grain size. 4. Piezoelectric studies As was pointed in the introduction, there is a lack of papers reporting the piezoelectric activity at the supposed working temperatures (>300  C). For the piezoelectric studies at high temperature, two issues must be taken into account: the values of the piezoelectric, electromechanical and elastic coefficients, and the resonance at the temperature range where electric anomalies are found. The works referred to [58e60] in the first section of this paper apply the resonanceeantiresonance method according to IEEE standard [85] to calculate coefficients at the different temperatures. It is widely accepted that Standards have limitations for low sensitivity or highloss materials [86]. To overcome these limitations, the results shown in this review have been calculated with a different procedure. The method for the piezoelectric, elastic and dielectric complex characterization of the ceramics is summarized below. Firstly, the ceramics were poled in a silicon oil bath at 200  C with fields as high as 120 kV/ cm, up to saturation. Dielectric, elastic and electromechanical coupling factors corresponding to the radial vibration mode were measured by an automatic iterative method described elsewhere [87], from complex admittance measurement. The coefficients are calculated by solving the set of non-linear equations that results when experimental impedance data at the resonance and antiresonance frequencies, along with some auxiliary frequencies, are introduced into the

32

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

analytical solution of the wave equation. For measurements at high temperature, the data acquisition and calculation must be separated, and the displacement of the resonance towards lower frequencies with the increase of temperature has to be taken into account. To do this, after each measuring temperature has been stabilized, an acquisition of an admittance-frequency file with 200 points is first collected in a wide frequency interval, from which the resonance is located, and then a new file with 400 data points in a narrower interval around the resonance is collected. The main modification of this measuring strategy with respect to the original [87] online method used at room temperature is that values of jYj, and its phase angle q, and auxiliary frequencies, are taken by linear interpolation of the measured data in the suitable range of frequencies. The data acquisition was also obtained with an impedance analyser HP4194, in the same experimental set-up used for the dielectric characterization. To illustrate the possibility of calculating the piezoelectric and elastic complex coefficients, Table 2 shows the full characterization at room temperature of the radial resonance of the ceramics which were sintered and recrystallized after hot pressing of SBN/BTN 35/65. The method allows the shown values at each measurement temperature to be obtained. Complex values of all the parameters are given in Table 2, together with the corresponding quality factor 0 00 0 00 of the compliance Qm (s11E) ¼ s11E /s11E , for s11E* ¼ s11E  is11E . It is the unique shown, although it is possible to list the quality factors for each coefficient. The recrystallized ceramics have larger values of Q than the sintered ones, which seems to have a direct relation with porosity. The higher level of densification obtained by recrystallization after hot pressing brings about these results. The improvement of Q by the control of microstructure in Aurivillius-type structure ceramics is thus suggested, as it was demonstrated for modified lead ceramics [88]. The iterative method used thus allows not only the piezoelectric and elastic parameters to be determined in materials with high losses (as when measurement temperature increases), but it is also a powerful tool to quantify and relate them with the microstructure of the ceramics. Fig. 12 shows the evolution with temperature of the real parts of the piezoelectric coefficient d31, the electromechanical coupling factor kp, the elastic compliance coefficient sE11, and the planar frequency number Np. It is shown for both sintered and recrystallized ceramics at the conditions shown in the figure. Table 2 Complex piezoelectric, elastic and dielectric characterization at room temperature of the radial resonance of thin disks of ceramics sintered and recrystallized after hot pressing. HP 1000  C-1 hþ HP 1000  C-1 hþ HP 900  C-1 hþ

Processing conditions Sintered at 

d33 (pC/N) kp (%) sp cP11 (1010 N m2) sE11 (1012 m2 N1) Qm (sE11) sE12 (1012 m2 N1) 3T33 (RT) 3T33 (500  C) d31 (pC/N) 12 m2 N1) sD 11 (10 D 12 m2 N1) s12 (10 Np (kHz mm)



1150 C-2 h

1200 C-2 h

1050  C-1 h

1100  C-1 h

1150  C-1 h

20 7.39 0.26 11.720 þ 0.004i 9.127  0.003i 3171 2.329 þ 0.001i 121.65  0.15i 188.45  4.14i 4.471 þ 0.003i 9.108  0.003i 2.347 þ 0.001i 2553

18 6.63 0.21 12.299 þ 0.004i 8.490  0.003i 3329 1.747 þ 0.001i 114.71  0.62i 184.42  3.05i 3.881 þ 0.004i 8.475  0.003i 1.762 þ 0.001i 2563

12 7.20 0.22 13.8044 þ 0.0026i 7.5999  0.0014i 5252 1.6445 þ 0.0003i 125.78  0.13i 193.42  12.52i 4.1755 þ 0.0142i 7.5843  0.0014i 1.6602 þ 0.0004i 2655

12 5.86 0.21 13.6450 þ 0.0025i 7.6669  0.0014i 5512 1.6101 þ 0.0003i 126.60  0.30i 191.31  7.37i 3.4158 þ 0.0050i 7.6564  0.0014i 1.6206 þ 0.0003i 2646

13 6.33 0.21 13.0265 þ 0.0037i 8.0139  0.0023i 3551 1.6439 þ 0.0005i 120.52  0.32i 181.37  6.41i 3.6889 þ 0.0197i 8.0011  0.0024i 1.6566 þ 0.0003i 2603

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

a

b

5

33

8 7 6

3 2

KP (%)

-d31 (pC/N)

4

1150ºC-2h. Sintered

5 4 3

1150ºC-2h. Sintered 1200ºC-2h. Sintered 1000ºC-h (HP)+ 1050ºC-1h 1000ºC-h (HP)+ 1100ºC-1h 900ºC-h (HP)+ 1150ºC-1h

1200ºC-2h. Sintered

2

1000ºC-h (HP)+ 1050ºC-1h

1

1000ºC-h (HP)+ 1100ºC-1h

1

900ºC-h (HP)+ 1150ºC-1h

0

0 50

150

250

350

450

550

650

50

150

Temperature (ºC) 1150ºC-2h. Sintered 1200ºC-2h. Sintered 1000ºC-h (HP)+ 1050ºC-1h 1000ºC-h (HP)+ 1100ºC-1h 900ºC-h (HP)+ 1150ºC-1h

10.5 10.0

d NP (KHz·mm)

S11E(·10-12 m2N-1)

c

250

350

450

550

650

Temperature (ºC)

9.5 9.0 8.5 8.0

1150ºC-2h. Sintered

2675

1200ºC-2h. Sintered

2650

1000ºC-h (HP)+ 1050ºC-1h

2625

900ºC-h (HP)+ 1150ºC-1h

1000ºC-h (HP)+ 1100ºC-1h

2600 2575 2550 2525 2500

7.5 50

150

250

350

450

Temperature (ºC)

550

650

2475 50

150

250

350

450

550

650

Temperature (ºC)

Fig. 12. Thermal evolution of: a) real part of the piezoelectric coefficient d31; b) planar electromechanical factor kp; c) real part of the compliance coefficient sE11: d) planar frequency number for ceramics recrystallized after hot pressing and sintered.

The relation with the processing conditions depends on the temperature range considered. Close to room temperature, and up to 350  C, the ceramic sintered at 1150  C-2 h has the highest piezoelectric coefficient. As the temperature approaches that of the ferroeparaelectric transition, the ceramic recrystallized at 1150  C-1 h after hot pressing at 900  C has the highest d31 values. Independently of the processing conditions, piezoelectric d31 coefficients are w2 pC/N at 500  C. The same tendency is observed for the planar electromechanical coupling factor kp. The dependence with the microstructure changes as the temperature increases in the same way as d31. Values of kp > 2.7% are found at 500  C in all cases. It seems that the porosity is not a determinant factor in piezoelectric and electromechanical activity. There is a tendency towards lower kp and d31 values as the grain size decreases (see Table 1), similar to the behaviour found for samarium-modified lead titanate ceramics [89]. It seems that the intrinsic (or volume) contributions [90] are dominant in SBN/BTN 35/65, rather than extrinsic contributions (mainly due to the domain wall movements). Pr is directly related with intrinsic contributions [91]. As Fig. 12 shows, both kp and d31 have similar dependence with temperature, and it is correlated with the evolution of Pr. It has been pointed that fine grains have a higher number of defects than coarse grains, and that these defects pin the domain walls [84]. On the other hand, as has been explained, defect dipoles that pin the domain walls are de-iced at high temperatures (>400  C) [92]. The extrinsic contributions should be more

34

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

important and the differences with microstructure should be higher at those temperatures, which is not the case. The coefficients at high temperature shown here are lower than those reported by Gai et al. [58,59] for (NaBi)0.48[]0.04Bi2Nb2O9 and [(Na0.5K0.5)Bi]0.44(LiCe)0.03[]0.03Bi2Nb2O9 (where [] represents A-site vacancies) although they are close at room temperature (kp w 7e8%). Differences are higher in the values of the d33 coefficient (28 pC/N with LiCe doping). For these materials, the piezoelectric coefficients are practically independent of temperature up to close to the depoling temperature, achieving values of kp higher than 10% at 500  C. This important result is an issue to be considered. The reason why piezoelectric stability is so much different from some structures to others is not clear. It can be related with the reduction of loss factor in the doped structure, but the authors do not give any explanation. It cannot even be ruled out that, in fact, the application of the IEEE standard is not suitable for ceramics with Aurivillius structures at high temperature [86]. In any case, the processing of an isotropic ceramic with relative high coefficients confirms that these structures are suitable for application at high temperatures. The elastic compliance sE11 increases continuously with increasing temperature, and it is higher in the sintered ceramics than in the recrystallized ones. For the former, sE11 decreases as the sintering temperature increases, and for the latter it is reduced as the difference between the recrystallization and the hot-pressing temperature decreases. The values of planar frequency number Np and stiffness elastic compliance sE11 at increasing temperature for SBN/BTN 35/65 (Fig. 12) are mainly related with the final content of porosity [93]. Np and the elastic compliance sE11 decrease with decreasing porosity. This is the reason why the values of compliance sE11 are smaller for the ceramic recrystallized than for the sintered ones. The continuous increase of the coefficient with temperature shows that the material becomes softer, although this reduction is not too high in all the temperature ranges. The planar frequency number Np decreases at 500  C to 2% of its value at room temperature for the ceramic recrystallized at 1150  C-1 h after hot pressing at 900  C-1 h. The planar resonance frequency reduces at 320  C to 1.8% of its value at room temperature for pure SrBi2Nb2O9, and to w1% at 400  C when doped with La or Y [94], similar to the behaviour found in this work. The result is also similar to that found by Gai et al. [58], who reported a reduction of 4.3% at 750  C. It is remarkable that the values of Np are higher for the ceramics studied in this review, due to the lower porosity achieved by the control of the microstructure by recrystallization after hot pressing. In the referred works which study the dependence of piezoelectric activity with temperature, the authors ignore a point that is essential for the performance of these materials, as it is shown in this work. This is the possible influence of the ‘‘de-iced’’ defects from the domain walls in the resonance modes. Fig. 13 shows the experimental data of resonance (Gmax, square symbols) and antiresonance (Rmax, circle symbols) at planar resonance of type B ceramics from 250 to 425  C. Up to that temperature (not shown), spectra have a single well-defined maximum. At 250  C, a second maximum appears at higher frequencies. When the temperature increases, the second resonance peak at higher frequencies increases, and is the highest at 325  C. As the temperature increases further, the relative height of the maxima at lower frequencies is further reduced, until it almost disappears at 450  C. Although not shown, the spectra remain so up to the depoling temperature. The appearance of the second maxima occurs close to the temperatures where the anomalies in the dielectric permittivity appear (Fig. 10) and close to those reported increase in the Young’s modulus [66,92]. This increase has been explained by the effect of the charged defects (oxygen vacancies) trapped in the domain walls. The de-icing of

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39 22000

0.0014

G (μS)

13200

0.0006

8800 4400

8800

11000

G (μS)

T=325ºC 8800

0.0005 6600

0.0004 0.0003

0.0004

4400

0.0002 2200

4400

0.0002

0.0001

11000

0.0007

0.0006

8800

8800

0.0005

G (μS)

4400

T=375ºC

6600

0.0004 0.0003

4400

R (Ohms)

0.0003

R (Ohms)

6600

0.0004

11000

0.0007

T=350ºC

0.0005

G (μS)

4400

R (Ohms)

13200

0.0008

R (Ohms)

G (μS)

0.0006

17600

0.0006

0.0002

0.0002

2200

2200

0.0001

0.0001

11000

0.0007

0.0003

4400

G (μS)

6600

0.0004

0.0006

T=425ºC 8800

0.0005 6600

0.0004 0.0003

R (Ohms)

8800

0.0005

11000

0.0007

T=400ºC

R (Ohms)

G (μS)

8800

0.0007

T=300ºC

0.0010

4400

0.0002

0.0002 2200

0.0001 0.0000 251

0.0006

0.0002

22000

0.0014

0.0006

13200

0.0008

0.0004

0.0002

0.0006

17600

0.0010

0.0004

0.0012

T=275ºC

R (Ohms)

G R

R (Ohms)

0.0008

0.0012

17600

0.0010

22000

0.0014

T=250ºC

G (μS)

0.0012

35

252

253

254

Frequency (kHz)

255

0 256

2200

0.0001 0.0000 251

252

253

254

255

0 256

Frequency (kHz)

Fig. 13. Experimental data of resonance (Gmax, square symbols) and antiresonance (Rmax, circle symbols) at planar resonance of ceramics recrystallized at 1050  C-1 h after hot pressing at 1000  C-1 h from 250 to 425  C. (Note the change of the scale from the graphic at 325  C.)

the defects from the domain wall produces the stiffening of the material. The secondary resonance that appears corresponds to the stiffened sample regions where vacancies begin to be de-iced. As the temperature increases, the material becomes homogeneous, as more vacancies are de-iced, until virtually only one resonance can be again observed. The increase of the stiffness is consistent with the observed new resonance peak at higher frequencies. It must be remarked that the temperature at which the second resonance appears depends on the microstructure [95]. For example, for type A ceramics the second antiresonance peak at higher frequencies appears at 350  C (Fig. 14). It is related to the same factors affecting the electric

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

36

20000

0.0014

15000

10000

0.0006 0.0004

20000

0.0006 15000

G (μS)

6000

0.0004 0.0003

4000

0.0002

5000

2000

0.0001 10000

0.0007

0.0006

8000

4000

8000

0.0005

G (μS)

0.0003

T=475ºC

6000

0.0004 0.0003

4000

R (Ohms)

6000

0.0004

R (Ohms)

0.0005

10000

0.0007

T=450ºC

0.0002

0.0002 2000

0.0001

2000

0.0001

0

0.0003

4000

R (Ohms)

6000

0.0004

5 7.

0 7.

24

5

24

0

6.

6.

24

5 5.

5.

24

8000

0.0005

24

24

4.

5

T=500ºC

0

10000

0.0007

24

G (μS)

0.0006

8000

R (Ohms)

10000

R (Ohms)

0.0008

10000 T=425ºC

0.0005

0.0002

G (μS)

5000

0.0007

T=400ºC

0.0004

G (μS)

10000

0.0006

0.0002

0.0010

0.0006

0.0008

0.0004

5000

0.0014

0.0006

15000

0.0010

0.0002

0.0012

T=375ºC

R (Ohms)

0.0008

R (Ohms)

G (μS)

0.0012

G R

0.0010

20000

0.0014

T=350ºC

G (μS)

0.0012

Frequency (kHz)

0.0002 2000

0.0001

7. 5 24

7. 0

6. 5

24

24

6. 0

5. 5

24

5. 0

24

24

24

4. 5

0

Frequency (kHz)

Fig. 14. Experimental data of resonance (Gmax, square symbols) and antirresonance (Rmax, circle symbols) at planar resonance of ceramics recrystallized at 1150  C-1 h after hot pressing at 900  C-1 h from 350 to 500  C. (Note the change of the scale from the graphic at 325  C.)

anomalies. Although the ceramic is still piezoelectric, the interval in which the resonance is double establishes a practical limit to the performance of these ceramics in high temperature applications. The control of the microstructure is thus essential. The processing by recrystallization after hot pressing allows the grain growth to be controlled while retaining a low porosity. Thus, this process seems to be a good procedure to improve the properties of the ceramics and to control their working temperatures. As a summary of this section, the d31 piezoelectric coefficient and the electromechanical coupling factors kp for the best conditions of sintering and recrystallization after hot pressing do not show a significant dependence with the microstructure in the high temperature working

A. Moure et al. / Progress in Solid State Chemistry 37 (2009) 15e39

37

range. Values of d31 > 1.8 pC/N and kp > 2.7% at 500  C are found in all ceramics. The elastic compliance sE11 increases continuously with increasing temperature, and it is higher in the sintered ceramics than in the recrystallized ones due to its higher porosity. The resonance spectra of the planar mode of poled thin ceramic disks are affected by the deicing of defect dipoles from domain walls. A double resonance exists in a given temperature range, indicating that part of the sample becomes stiffer, and the corresponding resonance appears at higher frequency. The temperature at which this double resonance disappears is higher as the grain size increases, which is related with the concentration of defect dipoles. 5. General conclusions Aurivillius-type ceramics are good candidates for uses as high temperature piezoelectrics due to their high ferroeparaelectric phase transitions. Piezoelectric activity at temperatures higher than 500  C has been shown. Solutions can be achieved to the problems for its applications from different strategies: - Its low sinterability makes obtaining ceramics with high density difficult. Precursors prepared by mechanical activation allow isotropic ceramics to be obtained with porosity lower than 1% at moderate processing temperatures. Control of the microstructure (high density and controlled grain size) can be achieved by a thermal process of recrystallization after hot pressing. - De-icing of the defect dipoles produces anomalies in the dielectric permittivity and the appearance of double resonances at the supposed working temperature. This can affect the performance of these ceramics as high temperature piezoelectrics. The control of the microstructure by recrystallization after hot pressing can tailor the range of temperatures where this phenomenon occurs. - Piezoelectric activity can be further improved by suitable doping of the structures that can reduce the conductivity of the ceramics and achieve a compromise between good polarizability and a high ferroeparaelectric phase transition temperature. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]

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