Effect of sintering temperature on the structural, dielectric and ferroelectric properties of tungsten substituted SBT ceramics

Physica B 406 (2011) 374–381

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Effect of sintering temperature on the structural, dielectric and ferroelectric properties of tungsten substituted SBT ceramics Indrani Coondoo a,n, Neeraj Panwar b, A.K. Jha a a b

Thin film and Materials Science Laboratory, Department of Applied Physics, Delhi Technological University (formerly Delhi College of Engineering), Delhi 110042, India Department of Physics, University of Puerto Rico, San Juan, PR 00931, USA

a r t i c l e i n f o

a b s t r a c t

Article history: Received 23 September 2010 Accepted 28 October 2010

Structural, dielectric and ferroelectric properties of tungsten (W) substituted SrBi2(Ta1 xWx)2O9 (SBTW) [x¼ 0.0, 0.025, 0.05, 0.075, 0.1 and 0.2] have been studied as a function of sintering temperature (1100–1250 1C). X-ray diffraction patterns confirm the single-phase layered perovskite structure formation up to x¼0.05 at all sintering temperatures. The present study reveals an optimum sintering temperature of 1200 1C for the best properties of SBTW samples. Maximum Tc of  390 1C is observed for x¼ 0.20 sample sintered at 1200 1C. Peak-dielectric constant (er) increases from  270 to  700 on increasing x from 0.0 to 0.20 at 1200 1C sintering temperature. DC conductivity of the SBTW samples is nearly two to three orders lower than that of the pristine sample. Remnant polarization (Pr) increases with the W content up to xr0.075. A maximum 2Pr ( 25 mC/cm2) is obtained with x¼ 0.075 sample sintered at 1200 1C. The observed behavior is explained in terms of improved microstructural features, contribution from the oxygen and cationic vacancies in SBTW. Such tungsten substituted samples sintered at 1200 1C exhibiting enhanced dielectric and ferroelectric properties should be useful for memory applications. & 2010 Elsevier B.V. All rights reserved.

Keywords: Ceramics X-ray diffraction Dielectric Ferroelectric

1. Introduction The discovery of ferroelectricity was made in 1921 by Valasek [1] in Rochelle salt (sodium potassium tartrate tetrahydrate; NaKC4H4O6  4H2O). The applications of such ferroelectric materials exploit the unique properties of ferroelectricity such as high dielectric constant, polarization hysteresis, piezoelectricity, electro-optic and pyroelectric properties. These materials are utilized in various devices such as high-permittivity capacitors, pyroelectric sensors, actuators, ultrasonic transducers, electro-optic devices, positive temperature coefficient of resistance (PTCR) components, ferroelectric memories (FeRAMs), etc. [2,3]. The potential utilization of these properties in new generation devices has motivated intensive studies. For example, the high dielectric permittivity of perovskite-type materials, like SrxBa1 xTiO3, can be advantageously used in dynamic random access memories, while the large values of switchable remanent polarization of ferroelectric materials are suitable for nonvolatile FeRAMs. The most popular ferroelectric material for nonvolatile memory applications is PbZrxTi1 xO3 (PZT); however, in addition to Pb toxicity issue, this material has a serious fatigue degradation problem that limits its utility [4,5]. In recent years, SrBi2Ta2O9 (SBT) has emerged as a potential material for nonvolatile ferroelectric memories. Besides being non-toxic, it exhibits many desirable properties viz. almost no

n

Corresponding author. E-mail address: [email protected] (I. Coondoo).

0921-4526/$ - see front matter & 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.physb.2010.10.074

fatigue after 1012 switching cycles, good retention characteristics, low switching fields and low leakage current [6,7]. Bismuth Layered Structure Ferroelectric (BLSF), SrBi2Ta2O9, is a member of the Aurivillius family with a general formula (Bi2O2)2 + (An 1BnO3n + 1)2 , where A¼Ca2 + , Ba2 + , Sr2 + , Bi2 + , etc.; B¼Fe3 + , Ti4 + , Nb5 + , Ta5 + , W6 + , Mo6 + , etc. and n denotes the number of corner sharing octahedra forming the perovskite like slabs [8]. The composition SBT with n¼2 is orthorhombic with space group A21  am and shows ferroelectric behavior at room temperature [9]. There are numerous studies on the modification of composition in SBT for improving the dielectric and ferroelectric properties [10–12]. The revelation that improved dielectric and ferroelectric properties can be obtained by the substitution of the A or B sites in SBT with a smaller cation having higher valency that has increased the scope for other substitutions of that kind. Recently, tungsten (W6 + ) has been investigated as a substituent for bismuth titanates and lanthanum doped bismuth titanates in which the remanent polarization is reported to enhance when a small amount of Ti4 + was replaced by W6 + [13,14]. This has been explained by considering the fact that the substitution introduces cation vacancies that suppress oxygen vacancies in the structure. This is especially relevant for SBT, where the degradation in ferroelectric properties and fatigue is widely believed to originate from the existence of oxygen vacancies [15]. In the present work, the hexavalent ion W6 + has been chosen as a donor cation for substituting the pentavalent Ta5 + sites in SBT and its influence on the microstructure; dielectric and ferroelectric properties have been investigated as a function of sintering temperature.

I. Coondoo et al. / Physica B 406 (2011) 374–381

2. Experimental Samples of series SrBi2(Ta1  xWx)2O9 (SBTW), with x ¼0.0, 0.025, 0.050, 0.075, 0.10 and 0.20 were synthesized by conventional solidstate reaction method from the starting materials like SrCO3, Bi2O3, Ta2O5 and WO3 (from Aldrich) in their stoichiometric ratios. 2–3 wt% more Bi in the form of Bi2O3 was also added to compensate the Bi loss at high temperature. The powder mixtures were thoroughly ground and passed through a sieve of appropriate size and then calcined at 900 1C in air for 2 h. The calcined mixtures

375

were ground and admixed with about 1–1.5 wt% polyvinyl alcohol (Aldrich) as a binder and then pressed at  300 MPa into disk shaped pellets. The pellets were air-sintered at 1100, 1150, 1200 and 1250 1C for 2 h. X-ray diffractograms of the sintered samples were recorded using Bruker diffractometer in the range 101r2y r701 with CuKa radiation. The sintered pellets were polished to a thickness of 1 mm and coated with a silver paste on both sides for its use as electrodes and cured at 550 1C for half an hour. The dielectric measurements were carried out using Solartron1260 Gain-phase analyzer operating at

Fig. 1. XRD patterns of SrBi2(WxTa1  x)2O9 samples sintered at (a) 1100 1C, (b) 1150 1C, (c) 1200 1C and (d) 1250 1C.

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oscillation amplitude of 50 mV. The Polarization-Electric field (P–E) hysteresis measurements were recorded at room temperature using an automatic PE loop tracer based on Sawyer–Tower circuit.

3. Results and discussion 3.1. X-ray diffraction XRD patterns of the samples having different concentrations of tungsten and sintered at varying temperatures are shown in Fig. 1(a)–(d) where all samples exhibit the characteristic peaks of SBT. The peaks have been indexed using the observed interplanar spacing d and the obtained lattice parameters were refined with the help of a computer program—POWDIN [16]. It is observed that the single-phase layered perovskite structure is maintained in the range 0.0rx r0.05 at all sintering temperatures. In samples with x40.05 content a very low intensity unidentified peak is observed (indicated by n) at all sintering temperatures deciding the solubility limit of W concentration in SBT. In all the diffraction patterns a marginal shift of peaks towards higher diffraction angle with increase in W concentration is observed implying a decrease in lattice parameters. The decrease in lattice parameters confirms the occupancy of W6 + at Ta sites. This can be understood from the fact ˚ is smaller in comparison to that the ionic radius of W6 + (0.60 A) ˚ [17]. The in-plane lattice parameters a and b decrease Ta5 + (0.64 A) significantly with higher W concentration (x 40.05), while much change is not observed in the c parameter (Table 1). 3.2. Microstructural studies Fig. 2 shows the scanning electron micrographs of the samples sintered at different temperatures. A systematic study of the micrographs reveals porous and loosely packed grains for pristine SBT samples sintered at 1100–1200 1C (Fig. 2a). However, when sintered at 1250 1C, densely packed microstructure is observed. Also, an increase in the grain size with increase in sintering temperature is noticed. In all the W-substituted samples, randomly oriented and anisotropic plate-like grains are observed and the average grain size also increases gradually with increase in the sintering temperature from 1100 to 1200 1C. However, the average grain size increases sharply in SBTW samples sintered at 1250 1C. For example, in the sample with x¼ 0.075, the average grain size increases from 2–3 to 8–9 mm on varying the sintering temperature from 1100 to 1250 1C. On the other hand, the average grain size also enhances with W content for the same sintering temperature. 3.3. Dielectric studies It is well known that in most cases the dielectric constant of the ferroelectric materials depends upon the composition, grain size, secondary phases etc. [18]. Fig. 3(a–d) exhibits the temperature

dependence er (at 100 kHz) of the samples. All the SBTW samples exhibit sharp transition at their respective Tc whereas pristine SBT shows broadened transition. The effect of sintering temperature on the peak er and Tc is shown in Fig. 4(a and b). At a particular sintering temperature, peak dielectric constant value increases with tungsten concentration whereas with increase in sintering temperature it increases up to 1200 1C and beyond that it decreases in all substituted samples. As far as Tc is concerned it first decreases for samples with x¼0.025 and then shows an increasing trend up to x¼0.20 at all sintering temperatures. These results indicate that the optimum sintering temperature for maximum dielectric constant and Tc is 1200 1C with tungsten substitution. Generally, in isotropic perovskite ferroelectrics substitution at B-site (located inside an oxygen octahedron) with smaller ions results in a larger polarization due to the availability of enlarged ‘‘rattling space’’ for smaller B-site ions, which eventually shifts the Curie point to a higher temperature [19]. However, in the anisotropic layered perovskites, the crystal structure may not change as free as that in the isotropic perovskites with substitution due to the structural constraint imposed by the (Bi2O2)2 + interlayer [20,21]. Moreover, since the valency of the substituted cation (W6 + ) is higher than that of Ta5 + , the substitution creates cationic vacancies at Sr-site (VSr ) to maintain electrical neutrality of the lattice structure [20–22]. The corresponding defect representation can be expressed as  Null ¼ WTa þ

1 00 Vsr 2

ð1Þ

 implies W occupying the Ta-site. When the W content is where WTa low, the lattice structure under the constraint of Bi–O layer and the presence of cation vacancies at the A-site possibly have resulted in an increased stress value. Given that the ionic radius of W6 + is smaller than Ta5 + there would be a competition between shrinking tendency of the lattice and structural constraint (that would oppose the shrinking tendency) imposed by the Bi–O interlayer. This is corroborated by a negligible change in the in-plane lattice parameters for lower concentration (Table 1). In such a scenario the perovskite structure would be less stable and may cause a decrease in Tc [23] as observed for samples with x ¼0.025 (Fig. 4b). On comparing the variation of in-plane lattice parameters (Table 1), one observes that for tungsten concentrations x 40.05, the decrease in the parameters a and b is much greater than that for concentrations x r0.05. This indicates that at higher concentrations the structural constraint imposed by the Bi–O interlayer was overcome leading to structural distortion. It is this enhancement of the ferroelectric structural distortion along with the introduction of cation vacancies at the A-site that leads to an eventual increase in Tc value [20,24]. Also, the high Tc is indicative of enhanced polarizability [20] that explains the increase in peak-er with increase in W concentration. Moreover, it has been reported that the cation vacancies make the domain motion easier and increase the dielectric constant [25,26]. Thus an increase in er with W content is observed. The peak-er value

Table 1 Lattice parameters a, b and c of SrBi2(WxTa1  x)2O9 samples sintered at 1100, 1150, 1200 and 1250 1C. x

0 0.025 0.05 0.075 0.1 0.2

1100 1C

1150 1C

1200 1C

1250 1C

a

b

c

a

b

c

a

b

c

a

b

c

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

˚ (A)

5.5285 5.5103 5.5102 5.5080 5.5104 5.5020

5.5128 5.5005 5.4964 5.4985 5.4979 5.4817

25.0131 24.9882 25.0253 25.0349 25.0426 25.0603

5.5243 5.5215 5.5172 5.5005 5.4822 5.4852

5.5337 5.5090 5.5062 5.4807 5.4689 5.4628

25.1001 24.9767 25.0046 25.0156 25.0354 25.0701

5.5212 5.5314 5.5270 5.5251 5.5242 5.5233

5.5139 5.5202 5.5199 5.5045 5.5060 5.4939

24.9223 25.1079 25.0585 25.0567 25.0850 25.0861

5.5246 5.5323 5.5361 5.5059 5.4972 5.4955

5.5143 5.5237 5.5219 5.4854 5.4782 5.4786

24.9017 25.0678 25.0366 24.8736 25.0867 25.1515

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Fig. 2. SEM micrographs of fractured surfaces of (a) x¼ 0.0, (b) x¼ 0.025, (c) x ¼ 0.075 and (d) x¼ 0.10 in SrBi2(Ta1  xWx)2O9 sintered at 1100, 1150, 1200 and 1250 1C.

of the pristine SBT increases with sintering temperature. However, in the SBTW compositions, peak-er value increases with the sintering temperature up to 1200 1C and exhibits a decrease at 1250 1C (Fig. 4a).

3.4. DC conductivity The electrical conductivity of ceramic materials encompasses a wide range of values. In insulators, the defects w.r.t. the perfect crystalline structure act as charge carriers and the consideration of charge transport leads necessarily to the consideration of point defects and their migration [27]. In the low temperature region (ferroelectric phase), extrinsic conduction dominates, whereas intrinsic ionic conduction occurs in the high temperature paraelectric phase [28,29]. Intrinsic conductivity results from the movement of the component ions, whereas the conduction resulting from impurity ions present in the lattice is known as extrinsic conductivity. Oxygen ion conduction in solids generally occurs via a hopping mechanism [30]. The compositional variation of DC conductivity (sDC) for the samples prepared at different sintering temperatures is shown in Fig. 5(a–d). The corresponding activation energy (EAC) in the high temperature region is shown in Table 2. It is observed that throughout the temperature range, sDC of the SBTW samples is

nearly two to three orders lower than that of the pure samples. The two predominant conduction mechanisms indicated by slope changes in the two different temperature regions are observed in Fig. 5. The conduction in the high temperature paraelectric phase with such high activation energies is normally associated with the motion of oxygen vacancies. It is well known that electrical properties are controlled by the inherent defects or charge carriers, which are produced during processing [31]. SBT sample contains a certain amount of inherent defects (oxygen vacancies, Vo ), resulting from the loss of Bi2O3 during sintering process [32] and exhibits n-type conductivity due to the creation of oxygen vacancies that leave behind two trapped electrons [32] 1 Oo - O2 m þ Vo þ2e0 2

ð2Þ

where Oo is an oxygen ion on an oxygen site, Vo is an oxygen vacant site and e0 represents electron. The conductivity in the perovskites can be described as an ordered diffusion of oxygen vacancies [33]. The oxygen vacancy (Vo ) is effectively double positively charged with respect to the neutral lattice and is considered to be the most mobile intrinsic ionic defect in the perovskite oxides [34,35]. Its motion is manifested by the enhanced ionic conductivity associated with an activation energy value of  1 eV [36]. Chen et al. [37] reported that the ionic conductivity increases rapidly

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Fig. 3. Variation of er with temperature at 100 kHz in W-doped samples sintered at (a)1100 1C, (b) 1150 1C, (c) 1200 1C and (d) 1250 1C.

Fig. 4. (a) Variation of peak dielectric constant with sintering temperature in SrBi2(Ta1  xWx)2O9 samples observed at 100 kHz and (b) variation of Curie temperature with W concentration at different sintering temperatures.

for temperatures above 300 1C having an activation energy of nearly 0.81 eV. In the present case also, Ea in the high temperature region for pure SBT lies in the range 0.77–0.82 eV. It is known that the oxygen vacancies can be suppressed by the addition of

donors because the donor oxide contains more oxygen per cation than the host oxide it replaces [38]. Other studies on layered perovskites have also reported a decrease in conductivity with addition of donors [28,39]. In the present study, the Ta5 + -site

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-12

-15 ln σdc (Ωcm)-1

ln σdc (Ω cm)-1

-20 -24

-20 -25

-28

-30

-32

-35 1.0

1.5

2.0

2.5

0.0 0.025 0.050 0.075 0.10 0.20

-10

0.0 0.025 0.050 0.075 0.10 0.20

-16

379

3.0

3.5

1.5

1.0

1000/T (K)-1

0.0 0.025 0.050 0.075 0.10 0.20

ln σdc (Ω cm)-1

-16

2.5

3.0

-20 -24

-12

3.5

0.0 0.025 0.050 0.075 0.10 0.20

-16 ln σdc (Ω cm)-1

-12

2.0

1000/ T (K)-1

-20 -24 -28 -32

-28

-36 1.2

1.4

1.6

1.8

2.0

1000/T (K)

2.2

2.4

2.6

-1

1.0

1.5

2.0

2.5

1000/T (K)

3.0

3.5

-1

Fig. 5. Variation of DC conductivity with temperature in SrBi2(Ta1  xWx)2O9 samples prepared at sintering temperature (a) 1100 1C, (b) 1150 1C, (c) 1200 1C and (d) 1250 1C.

Table 2 Activation energy of SrBi2(WxTa1  x)2O9 samples sintered at 1100, 1150, 1200 and 1250 1C. x in SBTW

1100 1C Ea (eV)

1150 1C Ea (eV)

1200 1C Ea (eV)

1250 1C Ea (eV)

0 0.025 0.05 0.075 0.1 0.2

0.72 1.25 1.28 1.26 1.26 1.27

0.69 1.87 1.85 1.89 1.83 1.79

0.76 1.92 1.96 1.98 1.86 1.74

0.77 1.72 1.75 1.75 1.73 1.69

substitution by W6 + in SBT can be formulated using a defect chemistry expression as Ta2 O5

WO3 þ Vo 2

1  W þ 3Oo 2 Ta

ð3Þ

It shows that the oxygen vacancies are reduced upon the substitution of the donor W6 + ions for Ta5 + ions. Hence, it is reasonable to believe that the conductivity in W-substituted SBT is suppressed by donor addition. It is well known that DC conductivity depends on both the concentration and mobility of charge carriers given as [40]

sDC ¼ nem

ð4Þ

where n is the concentration of charge carriers, e the charge of the carrier and m the mobility of the charge carrier. Therefore reduced

sDC indicates a decrease in concentration and/or mobility of charge carriers [41]. As per the above discussion, the high sDC observed in pure SBT samples (Fig. 5) can be attributed to the high concentration of Vo and their motion. Whereas in W-substituted samples reduced sDC is observed because the transport phenomena involving oxygen vacancies is greatly reduced. As the concentration of W increases, more oxygen vacancies are compensated and thereby the energy needed for the charge carriers to migrate (by hopping) also increases, leading to higher activation energy. As a result, the conductivity decreases because there are fewer oxygen vacancies available with sufficient energy to move around and the energy barrier between the scarcer vacancies increases [28]. The high Ea value of  1.75–2 eV in the high temperature region in the SBTW ceramics is consistent with the fact that the oxide ionic conduction reduces in these samples [41]. Also, it is observed that in the W-substituted samples with x40.05, sDC increases, which is reflected as a decrease in Ea value (Table 2). By comparing with Eq. (4), the decrease in the Ea (or increase in sDC) for samples with x40.05 could be associated with an increase in the concentration of mobile charge carriers in the form of Vo [42]. This correlation can be ascribed to the existence of the multiple valence states of tungsten. Since tungsten is a transition element, the valence state of W ions in a solid solution most likely varies from W6 + to W4 + depending on the surrounding chemical environment [43,44]. When W4 + are substituted for Ta5 + sites, oxygen vacancies would be created, i.e. one oxygen vacancy would be created for every two W4 + ions entering the lattice leading to an increase in mobile charge carriers concentration. This increase in oxygen vacancy concentration for SBTW samples with x40.05 would result in

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The P–E loops of SrBi2(Ta1  xWx)2O9 samples synthesized with different sintering temperatures and measured at 50 Hz at room temperature are shown in Fig. 6. It is observed that W-substitution results in the formation of well-defined hysteresis loops as compared to pure SBT. Fig. 7 shows the compositional dependence of remanent polarization (Pr) of SrBi2(Ta1  xWx)2O9 samples at different sintering temperatures. As can be observed, Pr depends on W concentration as well as sintering temperature. For the same sintering temperature Pr first increases with x and then decreases (Fig. 7). The optimum tungsten content for maximum 2Pr ( 25 mC/cm2) is x ¼0.075. It is also worthwhile to note that for the same tungsten content, the optimum sintering temperatures for maximum 2Pr are 1250 1C for pure SBT and 1200 1C for SBTW samples. It is known that ferroelectric properties are affected by compositional modification, microstructure and lattice defects like intrinsic oxygen vacancies [40,45]. In soft ferroelectrics with higher-valent substituents, the defects are cation vacancies whose mobility is extremely low below Tc. Thus the interaction between cation vacancies and domain walls is much weaker than that in hard ferroelectrics (with lower-valent substituents), wherein the associated mobile oxygen vacancies are likely to assemble in the vicinity of domain walls thereby pinning them and making their polarization switching difficult, leading to a decrease in Pr values [46–48]. Based on the obtained results and aforementioned discussion, it can be understood that in pure SBT oxygen vacancies assemble at sites like domain boundaries/ walls leading to strong domain pinning. Hence the well-saturated P–E loops for pure SBT are not obtained. Whereas in SBTW samples, the

Fig. 7. Compositional variation of Pr of SrBi2(Ta1  xWx)2O9 samples prepared at different sintering temperatures (1100, 1150, 1200 and 1250 1C). Inset shows the compositional dependence of Ec at different sintering temperatures.

1150 °C 1200 °C 1250 °C

15

P (μC / cm2)

1100 °C

x = 0.0

1100 °C

10

1150 °C 1200 °C

5

1250 °C

8 4

10

20

30

40

-80 -60 -40 -20

0 -4

E (kV / cm)

x = 0.075

15

1100 °C

10

1200 °C

1150 °C 1250 °C

5

-80 -60 -40 -20

0 -5

0

20 40 60 E (kV / cm)

80

P (μC / cm2)

1250 °C

P (μC / cm2)

20

1200 °C

40

60

80

E (kV / cm)

-12

-15

1150 °C

20

-8

-10

1100 °C

x = 0.025

12

0

C

0 -40 -30 -20 -10 0 -5

P (μC / cm2)

3.5. Ferroelectric studies

associated cation vacancy formation due to the substitution of Ta5 + by W6 + suppresses the concentration of oxygen vacancies. A reduction in the number of oxygen vacancies reduces the pinning effect on the domain walls, leading to an enhanced remnant polarization and decrease in Ec (inset of Fig. 7). Also, it is known that domain walls are relatively free in larger grains that cause an increase in Pr values

an increase in sDC (using Eq. (4)) and consequently a decrease in the Ea. Such an explanation has also been reported by Shannigrahi and Yao [44].

20

x = 0.20

15 10 5

-80 -60 -40 -20

0 -5

-10

-10

-15

-15

-20

-20

0

20

40

60

80

E (kV / cm)

Fig. 6. Variation in P–E loops at room temperature as a function of sintering temperature for SrBi2(Ta1  xWx)2O9 samples with x ¼(a) 0.0, (b) 0.025, (c) 0.075 and (d) 0.20.

I. Coondoo et al. / Physica B 406 (2011) 374–381

[49–51]. In the present study, the grain size is observed to increase with increase in W concentration; however, the Pr does not monotonously increase with increase in W concentration (Fig. 7). Such a behavior of Pr and Ec beyond x40.05 seems to be possibly affected by the presence of secondary phases (as observed in XRD patterns), which hamper the switching process of polarization [52,53]. Also, as already discussed w.r.t. DC conductivity, it was concluded that there is an increase in the number of charge carriers in the form of oxygen vacancies in the SBTW samples beyond x40.05. These oxygen vacancies would assemble at sites like domain boundaries leading to domain pinning and consequently result in a reduction in the Pr and an increase in Ec values as observed for SBTW samples with x40.05. 4. Conclusions

[13] [14] [15] [16]

[17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28]

From the present study, it can be concluded that the solubility limit of tungsten content in SBT is x r0.05 with an optimum sintering temperature of 1200 1C. For the same W content, the grain size increases with sintering temperature and a sharp increase in grain size with fully developed plate-like-grained microstructure is observed for SBTW samples sintered at 1250 1C. W-substitution is effective in enhancing the dielectric and ferroelectric properties. Maximum Tc of  390 1C is observed in the sample with x ¼0.20 as compared to  320 1C for the pristine sample sintered at 1200 1C. Peak-er increases from  270 to 700 with increase in W concentration from x¼0.0 to 0.20 at 1200 1C sintering temperature. All the tungsten substituted ceramics have higher Pr than that in pure samples. The maximum 2Pr (  25 mC/cm2) is obtained with x ¼0.075 sintered at 1200 1C. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12]

J. Valasek, Phys. Rev. 17 (4) (1921) 475. K. Uchino, Ferroelectric Devices, Marcel Dekker Inc.,, New York, 1997. Y. Wu, G.Z. Cao, Appl. Phys. Lett. 75 (17) (1999) 2650. J.F. Chang, S.B. Desu, J. Mater. Res. 9 (4) (1994) 955. K. Amanuma, T. Hase, Y. Miyasaka, Jpn. J. Appl. Phys. 33 (1994) 5211. R. Jain, V. Gupta, K. Sreenivas, Mater. Sci. Eng. B78 (2–3) (2000) 63. J. Zhang, Z. Yin, M.S. Zhang, Appl. Phys. Lett. 81 (25) (2002) 4778. R.E. Newnham, R.W. Wolfe, R.S. Horsey, F.A.D. Colon, M.I. Kay, Mater. Res. Bull. 8 (10) (1973) 1183. A.D. Rae, J.G. Thompson, R.L. Withers, Acta Crystallogr. Sect. B: Struc. Sci. 48 (1992) 418. S. Ikegami, I. Ueda, Jpn. J. Appl. Phys. 13 (1974) 1572. I. Coondoo, A.K. Jha, Solid State Commun. 142 (10) (2007) 561. A. Li, H. Ling, D. Wu, T. Yu, M. Wang, X. Yin, Z. Liu, N. Ming, Appl. Surf. Sci. 173 (3–4) (2001) 307.

[29] [30] [31] [32] [33] [34] [35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47] [48] [49] [50] [51] [52] [53]

381

J.K. Kim, T.K. Song, S.S. Kim, J. Kim, Mater. Lett. 57 (4) (2002) 964. W.T. Lin, T.W. Chiu, H.H. Yu, J.L. Lin, S. Lin, J. Vac. Sci. Technol. A 21 (2003) 787. A.R. James, S. Balaji, S.B. Krupanidhi, Mater. Sci. Eng. B64 (3) (1999) 149. E. Wu POWD, an interactive powder diffraction data interpretation and indexing program Version 2.1 (School of Physical Science, Flinders University of South Australia, Bedford Park S.A., JO42AU). R.D. Shannon, C.T. Prewitt, Acta Crystallogr. B25 (1965) 925. R.R. Das, P. Bhattacharya, W. Perez, R.S. Katiyar, Ceram. Int. 30 (7) (2004) 1175. E.C. Subbarao, Integr. Ferr. 12 (1996) 33. Y. Wu, C. Nguyen, S. Seraji, M.J. Forbess, S.J. Limmer, T. Chou, G.Z. Cao, J. Am. Ceram. Soc. 84 (12) (2001) 2882. P.D. Martin, A. Castro, P. Millan, B. Jimenez, J. Mater. Res. 13 (9) (1998) 2565. Y. Noguchi, M. Miyayama, T. Kudo, Phys. Rev. B 63 (21) (2001) 214102. I. Coondoo, A.K. Jha, S.K. Agarwal, J. Eur. Ceram. Soc. 27 (1) (2007) 253. H.T. Martirena, J.C. Burfoot, J. Phys. C: Solid State Phys. 7 (17) (1974) 3182. Y. Noguchi, M. Miyayama, T. Kudo, J. Appl. Phys. 88 (4) (2000) 2146. K. Singh, D.K. Bopardik, D.V. Atkare, Ferroelectrics 82 (1988) 55. R.C. Buchanan, Ceramic Materials for Electronics: Processing, Properties and Applications, Marcel Dekker Inc., New York, 1998. H.S. Shulman, M. Testorf, D. Damjanovic, N. Setter, J. Am. Ceram. Soc. 79 (12) (1996) 3124. Y. Wu, G.Z. Cao, J. Mater. Res. 15 (2000) 1583. I. Riess, Science & Technology of Fast Ion Conductors, Plenum Press, New York, 1987. B. Angadi, P. Victor, V.M. Jali, M.T. Lagare, R. Kumar, S.B. Krupanidhi, Mater. Sci. Eng. B100 (1) (2003) 93. C.A. Palanduz, D.M. Smyth, J. Eur. Ceram. Soc. 19 (6–7) (1999) 731. C.R.A. Catlow, Superionic Solids & Solid Electrolytes, Academic Press, New York, 1989. D.M. Smyth, Ferroelectrics 116 (1991) 117. W.L. Warren, K. Vanheusden, D. Dimos, G.E. Pike, B.A. Tuttle, J. Am. Ceram. Soc. 79 (1995) 536. B.H. Park, S.J. Hyun, S.D. Bu, T.W. Noh, J. Lee, H.D. Kim, T.H. Kim, W. Jo, Appl. Phys. Lett. 74 (13) (1999) 1907. T.C. Chen, C.L. Thio, S.B. Desu, J. Mater. Res. 12 (10) (1997) 2628. M.V. Raymond, D.M. Smyth, J. Phys. Chem. Solids 57 (10) (1996) 1507. M.M. Kumar, Z.G. Ye, J. Appl. Phys. 90 (2) (2001) 934. M. Yamaguchi, T. Nagamoto, O. Omoto, Thin Solid Films 300 (1–2) (1997) 299–304. Y. Wu, G.Z. Cao, J. Mater. Sci. Lett. 19 (4) (2000) 267. B.H. Venkataraman, K.B.R. Varma, J. Phys. Chem. Solids 66 (10) (2005) 1640. C.D. Wagner, W.M. Riggs, L.E. Davis, F.J. Moulder, Handbook of X-ray Photoelectron Spectroscopy, Perkin Elmer Corp., Chapman & Hall, 1990. S. Shannigrahi, K. Yao, Appl. Phys. Lett. 86 (9) (2005) 092901. H. Watanabe, T. Mihara, H. Yoshimori, C.A.P. De Araujo, Jpn. J. Appl. Phys. 34 (1995) 5240. T. Friessnegg, S. Aggarwal, R. Ramesh, B. Nielsen, E.H. Poindexter, D.H. Keeble, Appl. Phys. Lett. 77 (1) (2000) 127. Y.H. Xu, Ferroelectric Materials, Elsevier Science Publishers, Amsterdam, 1991. Q. Tan, J. Li, D. Viehland, Appl. Phys. Lett. 75 (3) (1999) 418. Y. Wu, S.J. Limmer, T.P. Chou, C. Nguyen, J. Mater. Sci. Lett. 21 (12) (2002) 947. S.B. Desu, P.C. Joshi, X. Zhang, S.O. Ryu, Appl. Phys. Lett. 71 (8) (1997) 1041. M. Nagata, D.P. Vijay, X. Zhang, S.B. Desu, Phys. Status Solidi A 157 (1) (1996) 75. J.J. Shyu, C.C. Lee, J. Eur. Ceram. Soc. 23 (2003) 1167; I. Coondoo, A.K. Jha, S.K. Agarwal, Ferroelectrics 356 (2007) 31. T. Sakai, T. Watanabe, M. Osada, M. Kakihana, Y. Noguchi, M. Miyayama, H. Funakubo, Jpn. J. Appl. Phys. 42 (2003) 2850.