A novel high-k ‘Y5V’ barium titanate ceramics co-doped with lanthanum and cerium

ARTICLE IN PRESS

Journal of Physics and Chemistry of Solids 68 (2007) 650–664 www.elsevier.com/locate/jpcs

A novel high-k ‘Y5V’ barium titanate ceramics co-doped with lanthanum and cerium Da-Yong Lua,b,, Xiu-Yun Sunb, Masayuki Todaa a

Department of Chemistry and Chemical Engineering, Faculty of Engineering, Yamagata University, 4-3-16 Jonan, Yonezawa 992-8510, Japan b Department of Applied Chemistry & Research Center of Materials Science and Engineering, Jilin Institute of Chemical Technology, Chengde Street 45, Jilin 132022, China Received 29 May 2006; received in revised form 7 February 2007; accepted 9 February 2007

Abstract Structural, dielectric, and ferroelectric properties of a novel high-k ‘Y5V’ (Ba1xLax)(Ti1x/4yCey)O3 ceramics (where x ¼ 0.03 and y ¼ 0.05, denoted by BL3TC5) with the highest ‘Y5V’ dielectric response (e0 410 000) among rare-earth-doped BaTiO3 ceramics to date are investigated in detail using SEM, TEM, XRD, DSC, EPR, Raman spectroscopy (RS), temperature and frequency, electric field dependences of dielectric permittivity (e0 ), and temperature and electric field dependences of ferroelectric hysteresis loops. The BL3TC5 diffusion of ferroelectric phase transition occurs around dielectric peak temperatures (Tm) near a room temperature characteristic of dielectric thermal relaxation. Powder XRD data and defect complex model were given. ‘‘Relaxor’’ behavior associated with an order/ disorder model and formation of a solid solution were discussed. The EPR results provided the evidence of Ti vacancies as compensating for lattice defects. High-k relaxor nature of BL3TC5 is characterized by an average cubic structure with long-range lattice disordering and local polar ordering; a slow change of the e0 (T) and Pr(T) curves around Tm; no phase transition observed by DSC; and a broad, red-shifted A1 (TO2) Raman phonon mode at 251 cm1 accompanying the disappearance of the ‘‘silent’’ mode at 305 cm1 and a clear anti-resonance effect at 126 cm1 at room temperature. r 2007 Elsevier Ltd. All rights reserved. Keywords: A. Ceramics; C. Raman spectroscopy; D. Dielectric properties; D. Electron paramagnetic resonance (EPR); D. Ferroelectricity

1. Introduction The continuous trend of miniaturization in the field of dielectrics requires higher and higher volumetric efficiencies, which can be realized in two ways: (1) raising dielectric permittivity (e0 ), and (2) reducing dielectric layer thickness (i.e., fine-grained microstructure). As the most important dielectric material, barium titanate (BaTiO3) is the simple compound with the highest room-temperature (RT) dielectric permittivity (e0 RT) of 1600 [1,2]. With technological progress, e0 RT of BaTiO3 has been improved to a limit value of about 5000 when the grain size is controlled over an optimum size range of 0.7–1.0 mm [3–6]. The dielectric Corresponding author. Department of Chemistry and Chemical Engineering, Faculty of Engineering, Yamagata University, 4-3-16 Jonan, Yonezawa 992-8510, Japan. Tel.: +81 238 26 3167; fax: +81 238 26 3165. E-mail address: [email protected] (D.-Y. Lu).

0022-3697/$ - see front matter r 2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.jpcs.2007.02.018

response of BaTiO3 shows first-order phase transition (FPT) behavior at the cubic–tetragonal phase transition (i.e., ferroelectric phase transition) point of 125 1C (i.e., Curie temperature TC), accompanied by a very strong but sharp dielectric peak (e0 m ¼ 10 000) [1–6]. Because BaTiO3 is the most easily compatible host compound for different dopants, and the occupation by different types of ions at equivalent sites in BaTiO3 can lead to a shift of the Curie peak towards RT and a so-called ‘‘diffusion of ferroelectric phase transition’’ (DPT) [7] around Tm (Tm representing the temperature at which maximum permittivity (e0 m) occurs), it is feasible to raise permittivity by utilizing the diffused Curie peak around Tm near RT when modifying BaTiO3 chemically. [Note: In ‘‘relaxor’’-type materials, Tm does no longer correspond to the phase transition of material with DPT in the substance. In this paper, we still use the classic terminology ‘‘diffuse phase transition’’ (DPT) as contrasted to FPT.]

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Rare earth (RE) is often referred to as ‘‘vitamin of the modern chemical industry’’. The dielectric response in REmodified BaTiO3 has been widely studied [8–23]. Lu has investigated the effect of RE dopants on dielectric properties of BaTiO3 [8], such as La [9,10], Ce [11–13], Nd [14,15], Sm [16,17], Eu [18], Gd [19,20], Dy [21–23], Ho [24], Er [25], Yb [26], Lu [27], and Y [28]. It was found that with lanthanide contraction, the site occupation of RE in BaTiO3 gradually transfers from Ba sites to amphoteric behavior of both Ba and Ti sites to Ti sites; solid solution limit declines; the dielectric response gradually transfers from FPT to DPT, accompanying the decrease of e0 m [9–35]. It is obvious that the occupation of RE ions at Ti sites is responsible for stronger DPT behavior due to the fact that BO6 octahedra constitute the skeleton of ABO3 perovskite and ionic size difference between 6-coordinate (6-CN) RE ions and Ti4+ ions [36]. Although the substitution of Ce4+ at Ti site results in a strong broadening of the Curie peak and further ‘‘relaxor’’-type ‘Y5V’ behavior (82%p(e0 e0 RT)/e0 RTp+22% in a temperature range 30 to 85 1C), e0 m decreases to 3000 for Ba(Ti1yCey)O3 ceramics with Tm near RT [37,38]. Individual RE doping in BaTiO3 is therefore difficult to realize both high k (e0 RT45000) and strong ‘Y5V’ DPT behaviors in the vicinity of RT. In order to exploit novel dielectrics in the RE-modified BaTiO3 family, Lu has predicted from the structure and developed La and Ce codoped barium titanate ceramics (BLTC) with high-k ‘DPT’ [39]. Among BLTC series, the (Ba1xLax)(Ti1x/4yCey)O3 ceramic (x ¼ 0.03, y ¼ 0.05, denoted by BL3TC5) with e0 RT410 000 that meets the EIA specification ‘Y5V’ is the most promising dielectric for future capacitor or DRAM applications. In this work, the novel high-k ‘Y5V’ BL3TC5 ceramic was investigated in detail by XRD, DSC, SEM, TEM, EPR, RS, temperature and frequency, electric field dependencies of dielectric permittivity, and temperature and electric field dependencies of ferroelectric hysteresis loop (FHL). Structural, dielectric, and ferroelectric properties and their interconnection were studied. 2. Experimental 2.1. Sample preparation BL3TC5 ceramics was prepared using conventional ceramic processing technique according to the formula (Ba1xLax)(Ti1x/4yCey)O3 (x ¼ 0.03, y ¼ 0.05) based on 0000 Ti-vacancy mechanism (4LaBa ! VTi ). The starting materials were reagent-grade BaCO3 (99%), TiO2 (99%), CeO2 (99.5%) (Kanto Chemical Co., Inc.) and La2O3 (99.5%) (Wako Pure Chemical Industry) powders. The stoichiometric mixture in accordance with the above formula was carefully mixed. The schedules of calcination, forming (PVA aqueous solution was added as a binder for forming process), and sintering were at 1100 1C for 5 h in air, 200 MPa for 2 min, and at 1480 1C for 24 h in air,

651

Fig. 1. BL3TC5 ceramic pellets: (a) sintered sample; (b) polished disk for dielectric measurements.

respectively. As shown in Fig. 1a, a crack-free ceramic pellet (10.3 mm in diameter  2.5 mm in thickness, resistivity r4108 O cm) with a density of 93% of theory density was obtained. A 5% Ce-doped BaTiO3 (BTC5: x ¼ 0, y ¼ 0.05) ceramics was prepared under the same conditions as BL3TC5. Both an undoped BaTiO3 (BT: x ¼ 0, y ¼ 0) and a 3% La-doped BaTiO3 ceramics (BL3T: x ¼ 0.03, y ¼ 0) with a Ti-vacancy solid solution mode (Ba10.03La0.03)Ti10.03/4O3 were prepared at 1300 1C for 3 h under the same conditions as the calcinations and forming of BL3TC5 for evolution analyses. Several BL3T powders were also sintered in order to discuss the solid solution of La in the BaTiO3 lattice. 2.2. Microstructure observations A micrograph of the polished BL3TC5 ceramics surface was taken with a scanning electric microscope (JSM6330F, JEOL Co., Ltd.) operated at 15 keV. The ceramics specimen was polished by coarse grinding using emery papers and fine grinding using Cilloidal silica solution Lot#7199 (California) and diamond paste (grain size: 0.25 mm, Engis Corp.). The thermal etching at 1300 1C for 30 min was then performed with a heating rate of 10 1C/min and natural cooling. Finally, the conducting Pt–Pd atoms were sputtered on the specimen surface for SEM observations. Average grain size was estimated on the polished plane using Fullman’s method [40]. The BL3TC5 ceramic foil for TEM analysis was prepared by conventional dimpling and Ar ion beam milling techniques. SAI (selected area image) and SAD (selected area diffraction) observations on single grain were performed using a transmission electronic microscopy (CM300-TWIN, Philips) operated at 300 kV. The Miller indexes of crystal planes corresponding to diffraction spots were determined by the camera constant L  l ¼ 13.53 mm A˚ [41]. 2.3. Crystal structure observations and determination X-ray powder diffraction data were collected at RT between 201 p 2y p 1201 with a scan width of 0.021 using a X-ray diffractometer (Rint 2200, Rigaku) with Cu Ka radiation. Lattice parameters and unit cell volume were

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calculated by MS Modeling (Accelerys, Inc.) using the Reflex Package and Cu Ka1 radiation (l ¼ 1.540562 A˚). XRD raw data were stripped of Cu Ka2 (l ¼ 1.54439 A˚) contribution and smoothed with Reflex Indexing, background was removed by Bru¨ckner method [42] using Reflex Powder Indexing technique. The subsequent Gaussian fit was performed by ‘Sma4Win’ software to determine accurate peak positions for structural refinement calculations. In the calculations of ‘Pawley’ method of powder indexing, Pseudo-Voigt function for peak profiles was adopted. 2.4. Dielectric measurements The sputtered Au electrodes (5  5 mm2) were used for dielectric measurements of the polished ceramic disk (10.3 mm in diameter, 0.8 mm in thickness), as shown in Fig. 1b. Under ac electric fields with 1, 10, 102, 103, 104, and 105 Hz and with an ac electric field amplitude (EP) of 10 V/mm, temperature dependence of dielectric permittivity e0 (T) was measured at a heating rate of 2 1C/min from 100 to 200 1C using an AD-3521 FFT analyzer (A&D Co. Ltd.). The e0 (T) curves at 1 kHz were also measured from Ep ¼ 0.5 to 10 V/mm at a heating rate of 2 1C/min. Upon cooling, e0 (T) at 1 kHz was recorded from 200 to 100 1C under EP ¼ 10 V/mm and a cooling rate of 2 1C/min. 2.5. Ferroelectric measurements FHLs were measured for BL3TC5 using a conventional Sawyer–Tower circuit system with a 0.05 Hz sine wave and a field amplitude of 800 V/mm from 30 to 120 1C, over a field amplitude range of 200–1200 V/mm at 38 1C (i.e., Tm of e0 (T) on heating and at 1 kHz), respectively. 2.6. DSC measurements DSC measurements were performed with a DSC Q100 differential scanning calorimeter (Shimadzu co. Ltd.), with nitrogen as pure gas at a flow rate of 50 ml/min. Heat flow data between 20 and 160 1C at a heating/cooling rate of 5 1C/min in steps of 0.2 s were recorded by TA Instruments Universal Analysis 2000 software.

in terms of the 3rd and 4th line of Mn2+ standard sextet markers (g3 ¼ 2.032 and g4 ¼ 1.981). 2.8. Raman scattering investigations Raman spectra were measured at RT over a frequency range of 90–1300 cm1 on the ceramic powders using a Dispersive Raman spectrometer (Nicolet Almega XR, Thermo Electron Corporation, USA), with a green excitation light of 532 nm line of Nd:YAG laser focused on a spot of 0.3 mm in diameter. The laser power level was limited to 10% of normal output of 25 mW in order to avoid laser heating effect. The scattered light dispersed by the spectrometer was detected by an electron-cooled CCD detector operating at 50 1C, with a 2400 groove/mm grating. The spectral resolution is less than 3 cm1. 3. Results 3.1. Microstructure analyses A SEM micrograph of BL3TC5 is shown in Fig. 2. The average grain size is 0.9 mm. It is obvious that the codoping with 3% La and 5% Ce in BaTiO3 suppresses the grain growth, forming a homogenous and fine-grained ceramics, and shows an advantageous grain size refinement over Ba1xLaxTi1x/4O3 ceramics (23 mm) reported by Morrison et al. [9] and Ba(Ti1yCey)O3 ceramics (3 mm) by Chen et al. [37]. TEM observations on BL3TC5 foil are shown in Fig. 3. The SAD pattern in Fig. 3b is from the electron diffraction on the single grain (Fig. 3a), and Miller indexes corresponding to the diffraction spots were identified. The several spots were not in periodic orientation and are unrelated to possible defect order orientations, which is due to the fact that the selected area at the visual iris aperture limit was greater than this grain size so that a part of diffraction of neighboring grains was also collected. The SAI and SAD observations indicate that neither domain structure nor core-shell structured grains were formed in BL3TC5. It is known that BaTiO3 shows a tetragonal symmetry and ferroelectric domain structure at RT [5], and

2.7. EPR measurements EPR spectra were measured at RT using a JES-FR30T spectrometer (JEOL, Japan) at X band frequency (9.425 GHz). For EPR measurements, approximately 12 mg of powder was heat-treated at 500 1C for 2 h to remove any possibility of CO2 and H2O gas absorption. The instrumental parameters are microwave power ¼ 4 mW and magnetic field modulation ¼ 5 Gauss (G). The instrumental gain setting (gain ¼ 200) was chosen by reference to the EPR signal of single-phase 3% Eu-doped BaTiO3 sample [43]. The giromagnetic value was calculated

Fig. 2. SEM micrograph of thermally etched surface of BL3TC5 ceramics.

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653

Fig. 3. TEM micrographs of BL3TC5 ceramic foil: (a) SAI; (b) SAD.

Ce doping in BaTiO3 easily forms a core-shell structured grain with an unreacted tetragonal ferroelectric BaTiO3 core, nonferroelectric modified shells, and gradient regions with a varying dopant concentration [44]. SEM and TEM results indicate that the co-doping with La and Ce leads to homogeneous distribution in BL3TC5 grains, and suppression of ferroelectricity in the BaTiO3 phase at RT.

Powder XRD patterns at RT for BT, BL3T, BTC5, and BL3TC5 are shown in Fig. 4. Structural evolution can be clearly observed from the main peak (1 0 1, 1 1 0) and the characteristic peaks (0 0 2, 2 0 0) in Fig. 5, and variations in lattice parameters (a, c) and unit cell volume (V0) with RE type doping in Fig. 6. The data of BL3T in Fig. 6 stem from Morrison et al. [10], which can be considered a sufficient incorporation of La into the BaTiO3 lattice. The cubic BaTiO3 data (a ¼ 4.031 A˚, V0 ¼ 65.50 A˚3) in Fig. 6 are from a JCPDS card [45]. BL3T maintains a tetragonal structure. During the process of powder and ceramic sintering at 1300 1C, a progressive increase in lattice parameter a and decrease in c can be observed, accompanying the approach of (0 0 2) and (2 0 0) peaks (Fig. 5), indicating that La doping has the tendency to make BaTiO3 shift from tetragonal to cubic symmetry. On the other hand, V0 of our BL3T sample (Fig. 5) is less than that obtained by Morrison et al. (Fig. 6), indicating that La is not easily sufficiently incorporated into the perovskite grains when sintering short-term at 1300 1C, and a certain amount of unincorporated La would be segregated along grain boundaries [46,47]. Although the exclusive occupation of Ba sites by La ions has been widely accepted [9,10,33–35] and the ionic radius of 12-CN La3+ (1.36 A˚) is less than that of Ba2+ (1.61 A˚) [36], the substitution of 3% La3+ for Ba2+ ions on interstitial space of BO6 octahedra skeleton does not lead to a marked change in V0 (Fig. 6).

Intensity (arb. unit)

3.2. Structural evolution and complex defect model

20

30

40

50

60

70

80

90

100

2θ (degree) Fig. 4. Powder XRD spectra. C and P in brackets represent ceramics and powder sintered, respectively.

It is well known that diffusion of RE cations on the B (Ti) site of the ABO3 perovskite lattice is slower than the A (Ba) site because the former produces larger lattice strain

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654

30.5

31

31.5

32

32.5

44

2θ (deg.)

45

46

2θ (deg.)

Fig. 5. Evolution of XRD peak with dopant type. The spectra were treated as: background removed, curve smoothed, and contribution from Cu Ka2 stripped by MS Modeling Reflex Package.

4.09 67.0

of Ce in the BaTiO3 lattice requires higher sintering temperatures (4 1400 1C) [29]. Incorporation of Y, Dy, Ho, and Er at the B site also shows a behavior similar to Ce [28,30–32]. BTC5 shows a slightly tetragonal structure approaching pseudo-cubic. The apparent increase in a, c and V0 (Fig. 6) accompanying large XRD peaks shifts to the low angle of BT (Fig. 5) is attributed to the expansion of crystal cells resulting from the occupation of Ti sites with larger Ce ions in the BaTiO3 host lattice. A pronounced broadening in full width at half maximum (FWHM) of symmetric (1 1 1) peak (0.311) in BTC5 can be observed as compared to other samples (p 0.221), indicating that Ce doping leads to a larger lattice strain, and further destroying the long-range order of the BaTiO3 lattice. The co-doping with La and Ce makes the diffraction peaks of BL3TC5 shift suddenly to higher diffraction angles relative to those of BTC5; XRD peaks gradually become narrow and symmetric, and peak intensity increases (Fig. 5). BL3TC5 has a cubic structure, which is not an ideal cubic perovskite structure and involves intrinsic disorder in the cubic phase [48,49], meaning that a cubic structure on the average is characterized by the broadening of the (2 0 0) peak in BL3TC5—the FWHM: 0.251—as contrasted with the same peak in BT (0.191) and BL3T (0.151), i.e., the symmetry of BL3TC5 is lower than cubic, but rather statistically cubic based on XRD peak shapes and calculations. Table 1 lists X-ray powder diffraction data of BL3TC5. The co-doping with La and Ce results in a sudden decline in V0 on comparing BL3TC5 with BTC5. The studies on defect chemistry in (Ba1xLax) (Ti1x/4yCey)O3 [39] indicate that Ce4+ ions isovalently

4.07

4.05 65.0 4.03

V0 (Å3)

a, c (Å)

66.0

64.0 4.01 63.0 3.99 BT

BL3T

BTC5

BL3TC5

Fig. 6. Variation in lattice parameters (a, c) and unit cell volume (V0) with dopant type. BT data are from cubic BaTiO3 [45] and our tetragonal BT ceramic sample, respectively. BL3T data conform to that obtained by Morrison et al. [10]. BT: BaTiO3; BL3T: (Ba10.03La0.03) Ti10.03/4O3; BTC5: Ba(Ti10.05Ce0.05)O3; and BL3TC5: (Ba10.03La0.03) (Ti10.03/40.05Ce0.05)O3.

(such as 6-CN Ce4+: 0.87 A˚, Ti4+: 0.605 A˚ [36]) and hence needs higher energy incorporation than the latter. As a consequence, the separated BaCeO3 and BaTiO3 phases are observed when sintering at 1200 1C [37] and a solid solution

Table 1 Power XRD data of BL3TC5 ceramics (S.G. Pm3 m) (Cu Ka1 radiation, l ¼ 1.540562 A˚) dobs (A˚)

dcal (A˚)

I/I0

h

k

l

4.0242 2.8453 2.3230 2.0121 1.7995 1.6427 1.4225 1.3409 1.2725 1.2131 1.1615 1.1161 1.0755 1.0060 0.9761 0.9484 0.9228 0.8998

4.0241 2.8454 2.3232 2.0120 1.7996 1.6428 1.4227 1.3410 1.2725 1.2133 1.1616 1.1161 1.0755 1.0060 0.9760 0.9485 0.9229 0.8998

15 100 24 28 7 31 13 2 9 4 4 1 9 1 1 4 1 3

1 1 1 2 2 2 2 3 3 3 2 3 3 4 4 3 3 4

0 1 1 0 1 1 2 0 1 1 2 2 2 0 1 3 3 2

0 0 1 0 0 1 0 0 0 1 2 0 1 0 0 0 1 0

(System: Cubic) (a ¼ 4.024 A˚, V0 ¼ 65.16 A˚3).

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substitute for Ti sites (B sites), La3+ ions aliovalently substitute for Ba sites (A sites) and induce B-site Ti-vacancy defects. The charge balance compensation mechanism may be described as follows: BaO þ CeO2 ! BaBa þ CeTi þ 3Oo 2La2 O3 þ 3TiO2 !

4LaBa

655

homogeneous disorder of defects in the host BaTiO3 lattice. 3.3. Influence of temperature, frequency, and electric field on dielectric properties

(1) 0000

þ 3TiTi þ VTi þ 120o .

(2)

Discounting very low concentration background impurities (i.e., impurities in the starting materials), the defects are composed of Ce4+ ions at Ti sites (CeTi), La3+ ions at Ba 0000 sites (LaBa ), and Ti vacancies (VTi ) induced by La3+ doping. Defect notation proposed by Kro¨ger and Vink [50] was adopted. The above SEM, TEM and XRD results indicate clearly that BL3TC5 is a single-phase solid solution with an average cubic structure and a homogenous distribution of La and Ce dopants in the BaTiO3 host lattice. It is a reasonable explanation for the sudden decline of BTC5 V0 that Ti vacancies and La ions at Ba sites next to one or several CeO6 octahedra counteract BO6 octahedra expansion caused by the substitution of Ce for Ti site. Fig. 7 depicts the two defect complex models of La3+–Ce4+ (i.e., 0000 Ba7/8La1/8CeO3 cell) and VTi –Ce4+. It is apparent that the formation of the defect cluster may relieve the lattice strain associated with the individual lattice defects, and form a

Fig. 7. Unit cell of ideal cubic BaTiO3 (a) and the model of defect complexes of (Ba7/8La1/8)CeO3 (b) and CeO6–VTi complex (c) for BL3TC5.

Upon heating and cooling, the temperature dependences of dielectric permittivity e0 (T) and dielectric loss tan d(T) were measured under an ac electric field at 1 kHz with Ep ¼ 10 V/mm for BL3TC5 and BTC5 as shown in Fig. 8, where the BL3T data are plotted as derived from Ref. [9]. BL3TC5 shows stronger ‘DPT’ (g ¼ 1.81) when compared to BL3T (gE1) with FPT and BTC5 (g ¼ 1.47) with ‘DPT’, and a marked rising of e0 (e0 RT410 000). For relaxor-type dielectrics, the critical exponent g defined as 1=0  1=0 m ¼ ðT  T C Þg =C (where C is a constant, 1pgp2) is often adopted to quantify the diffuseness of ferroelectric phase transition with temperature [51]. The limiting values g ¼ 1 and 2 respond to normal Curie–Weiss law (FPT) and relaxor-type quadratic law (DPT). The Tm values of BL3TC5 on heating and cooling are 38 and 22 1C, respectively, showing larger dielectric thermal relaxation (DTm,HC ¼ 16 1C) effect than BTC5 (DTm,HC ¼ 5 1C). This feature of BL3TC5 implies that a very high dielectric permittivity may be retained over a broader range around RT. e0 (T) of BL3TC5 meets EIA specification ‘Y5V’ whether on heating or cooling. To our knowledge, highinsulating BL3TC5 is the highest ‘Y5V’ dielectrics among RE-doped BaTiO3 ceramics so far. tand(T) in Fig. 8b also shows dielectric thermal relaxation, and the tand peak (0.05) of BL3TC5 shifts rapidly to the higher-temperature side of Tm of BTC5. Upon heating, e0 (T) and tan d(T) in Fig. 9 were measured at 1 kHz over a weak ac electric field (i.e., so-called ‘‘zero field’’) range from 0.5 to 10 V/mm for BL3TC5, showing that the dielectric behavior of BL3TC5 is almost independent of applied electric field within this zero field range. The e0 (T) characteristics at different frequencies for BL3TC5 in Fig. 10a show that with increasing frequency, e0 m decreases and Tm shifts slightly to higher temperature from 35.0 1C at 1 Hz to 39.9 1C at 100 kHz, characteristic of ‘relaxor’ material. Although BL3TC5 has similar regularity in the frequency dispersion of e0 to high-k PMN [Pb(Mg1/3Nb2/3)O3] [52–54], the diffuseness of e0 with frequency (DT0 mf ¼ Tm,100 kHz–Tm,10 Hz ¼ 4 1C) for BL3TC5 is far lower than that of PMN (DT0 mf ¼ 20 1C) [52], indicating that BL3TC5 possesses relatively higher dielectric-frequency stability over a low-frequency range (p106 Hz). The tan d(T)’s in Fig. 10b show that when T 4100 1C, e0 rapidly increases with decreasing frequency; when To100 1C, e0 slightly increases with decreasing frequency, keeping almost constant similar to Fig. 9b. In the case of zero field, BL3TC5 shows a lower tan d (o 0.05) in the ‘Y5V’ temperature range compared to PMN (0.1 at 1 kHz) [52] and Zr-doped PMN (0.06–0.11 at 100 Hz) [54]. BL3TC5 is thus quite qualified for application in ‘Y5V’ceramic capacitor in the future.

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0.10

1.6

0.08

tanδ

ε '(x 104)

1.2

0.8

0.06 0.04

0.4

0.02

0.0 -100

-50

0 50 100 Temperature (°C)

150

0.00 -100

200

-50

0 50 100 Temperature (°C)

150

200

Fig. 8. Temperature dependences of (a) dielectric permittivity (e0 ) and (b) dielectric loss (tand) measured at 1 kHz and EP ¼ 10 V/mm for BL3TC5 and BTC5 ceramics (on heating and cooling). The BL3T data are derived from Ref. [9] for comparison.

0.10

1.6

0.08

tanδδ

ε '(x 104)

1.2

0.8

0.06 0.04

0.4

0.02

0.0 -100

-50

0 50 100 Temperature (°C)

150

0.00 -100

200

-50

0 50 100 Temperature (°C)

150

200

Fig. 9. Temperature dependences of (a) dielectric constant (e0 ) and (b) dielectric loss (tand) measured at 1 kHz and different electric fields (0.5–10 V/mm) for BL3TC5 ceramics (on heating).

0.10

1.6

0.08

tanδδ

ε '(x 104)

1.2

0.8

0.06 0.04

0.4

0.02

0.0 -50

0

50 100 Temperature (°C)

150

0.00 -100

-50

0 50 100 Temperature (°C)

150

200

Fig. 10. Temperature dependences of (a) dielectric constant (e0 ) and (b) dielectric loss (tand) measured at different frequencies from 1 Hz to 100 kHz and 10 V/mm for BL3TC5 ceramics (on heating).

3.4. Ferroelectric characteristics Fig. 11 shows the RT FHLs measured under a strong ac electric field of EP ¼ 600 V/mm for BT, BTC5, and BL3TC5 ceramics. The Ce doping in BaTiO3 suppresses strong ferroelectricity of BT (Pr ¼ 9 mC/cm2), but the higher remanent polarization in BTC5 (Pr ¼ 3 mC/cm2) is attributed to a slight average tetragonal structure at RT. The co-doping leads to a nonlinear saturation in FHL

(PrE0) for BL3TC5, which is associated with a disordered average cubic structure and Tm close to RT. Fig. 12 shows variation in FHL measured at EP ¼ 800 V/mm with temperature for BL3TC5. With increasing temperature, Pr decreases, and FHL shows a nonlinear saturation without the electric polarization (PrE0) at 30 1C near Tm (38 1C at 1 kHz), then a very weak FHL occurs at 60 1C, finally changing into linear saturation when TX90 1C. The chance in Pr(T) of BL3TC5 is different

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657

40 FHL EP = 600 V/mm

Pr (μ C/cm2)

20

BT BL3TC5

BTC5 0

-20

-40 -0.8

-0.4

0 E (kV/mm)

0.4

0.8

Fig. 11. Ferroelectric hysteresis loops (FHL) at room temperature for BT, BTC5, and BL3TC5 ceramics. Measuring ac electric field: EP ¼ 600 V/mm, f ¼ 0.05 Hz.

Fig. 13. Variation of electric hysteresis loops and Pr at 38 1C with applied electric field for BL3TC5.

with a strong EP ¼ 2000 V/mm, and Pr increases with EP (Fig. 13) and stayed weaker. 3.5. DSC phase transition investigations

Fig. 12. Variation of (a) electric hysteresis loop and (b) remanent polarization (Pr) at 800 V/mm with temperature for BL3TC5.

from that of BT. The latter shows a sudden disappearance above TC [1]. Thus, the behavior of BL3TC5 in the vicinity of RT can be considered as characteristic of a relaxor material. After measurement in strong electric field, the FHLs measured at 38 1C and different EPS for BL3TC5 are shown in Fig. 13. FHL did not return to initial nonlinear saturation at Tm (Fig. 12) after the sample was polarized

It is well known that the crystal phase transition is accompanied by an endothermal peak on heating and an exothermal peak on cooling. Hence, the phase transition of materials can be monitored by means of the change of DSC heat flow. Fig. 14 shows the DSC traces for orthorhombic– tetragonal (o–t) and tetragonal–cubic (t–c) phase transitions on heating and cooling. The same DSC thermal relaxation effect as e0 (T) was observed for BT and BL3T. Here, only t–c phase transition is considered. The DSC results of BL3T indicate that the prolongation of sintering time and the increase of sintering temperature can shift TC to a low temperature that is more pronounced than sintering for undoped BaTiO3. The DSC peak specially shifts to low temperature more rapidly when clamped (i.e., cold-pressed pellet sintered) (119.8 1C) than when powder sintered (125.6 1C in Fig. 14a), while peak intensity decreases markedly. It is reported that when pellets with same compositions as our BL3T were sintered at 1350 1C for 3 days, e0 (T) shows an FPT behavior with a very sharp dielectric peak at TmE60 1C [10]. This means that the t–c temperature of the BaTiO3 host cells can be largely shifted to low temperature and stabilized by Ti vacancies to compensate for lattice defects induced by a sufficient incorporation of La3+ at Ba sites [10]. However, after calcinations at 1300 1C for our BL3T, a marginal decrease

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of TC suggests that La is not sufficiently incorporated into the BaTiO3 lattice. No DSC peaks in BTC5 and BL3TC5 were observed, which is associated with the following order/disorder model for relaxor materials. In fact, for the incorporation of any RE into BaTiO3 ceramics, the higher dopant concentration (above 4%) usually leads to the disappearance of DSC peak. 3.6. EPR investigations EPR was employed to investigate the defects and impurities in BL3TC5. The main impurities in the starting materials, which we had to consider, are Sr (o0.5%) in the reagent-grade BaCO3 and Fe (o0.005%). We prepared two BaTiO3 ceramics with an excess of TiO2 (0.5%) to compensate for Sr in BaCO3 and without excess TiO2. There is almost no difference in EPR response between both BaTiO3 ceramics. Because the EPR signal associated with Ba-vacancy defects was also observed for the BaTiO3 sample without excess TiO2, indicating that the presence of the impurities Sr does not induce an A-site excess. In considering that an excess of TiO2 probably induces Bavacancy defects, an excess of TiO2 was not used for all compositions including BaTiO3 ceramics. Although the impurities Mn and Cr are usually present in BT ceramics, we did not detect clear EPR signals associated with Mn2+ (3d5, 6S5/2) sextet and Cr3+ (3d3, 4F3/2) at RT [55] and even at 150 1C in the cubic phase [43]. Thus, these microimpurities can be ignored. Fig. 15 shows the RT EPR spectra. A symmetric signal at g ¼ 1.974 with a line width of 18 G for BT and BL3T ceramics and an asymmetric signal at g ¼ 2.004 with a line width of 19 G for BL3T, BTC5, and BL3TC5 were observed. Oxygen vacancies (Vo) in the perovskite lattice can be ignored based on these facts: (1) sintering in air provides

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Fig. 14. DSC traces for the phase transition of BT, BL3T, BTC5, and BL3TC5 (a) on heating, (b) on cooling. C and P (in brackets) represent ceramics and powder sintered, respectively.

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sufficient oxygen; (2) the high insulating behavior of the ceramic samples (r4108 O cm)—the presence of oxygen vacancies under electric field easily causes conducting behavior owing to its movability; (3) no resonances associated with Fe3+–Vo complex at g ¼ 11.6, 5.54, and 2.47 [55–57] were clearly detected when Fe was incorporated into B site as Fe3+ (3d5, 6S5/2) [55–58], although Fe3+ is also responsible for g ¼ 2.00 [57,59,60]. Correspondingly, the two signals in Fig. 15 are unrelated to Fe3+ impurities. Fe impurities (o0.005%) in our samples are infinitely small and may be ignored when compared to higher concentrations of La and Ce dopants. The Ce doping in BTC5 and the co-doping in BL3TC5 0 lead to the disappearance of the signal at g ¼ 1.974 (VBa ) in

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dynamic symmetry in a much smaller region (correlation length below 2–3 nm) within a much shorter time (o 1 ns) [65]. Raman spectra obtained at RT for BT, BL3T, BTC5, and BL3TC5 are shown in Fig. 16. The first-order Raman spectra of tetragonal BT (Fig. 16a), similar to those of polycrystalline BaTiO3 samples reported by other authors [66–69], show two asymmetric, broad and intense bands associated with A1 (TO2) and A1 (TO3) optical modes, a sharp band [i.e., the ‘‘silent’’ mode A1+E(TO+LO) from cubic F2u] and a weak band [A1 (LO3)+E (LO3)], peaking at 261, 514, 305, and 715 cm1, respectively. The phonon modes corresponding to Raman bands are assigned as proposed by Dixit et al. [70] in terms of some reports on BaTiO3 single crystals [71–80]. The observed anti-resonance effect at 180 cm1 as an interference feature, is attributed by Scalabrin et al. to a coupling between the sharp A1 (TO1) and broad A1 (TO2) modes [78]. A weak swelling at 470 cm1 is attributed to overlapping E (TO3)+E (LO2) optical modes. Only a very weak band

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BT and the appearance of a new signal at g ¼ 2.004 (VTi ) (Fig. 15). In accordance with the Dunbar et al. [61] and Kolodiazhnyi and Petric [62] observations and results that the excess positive charge of donor ions is compensated by cation vacancies, we indexed the signals of g ¼ 1.974 and 2.004 as Ba vacancies and Ti vacancies, respectively 0 0000 (assumedly singly ionized VBa and VTi ). Thus, our EPR results are in accord with the Lewis and Catlow [63] and Buscaglia et al. [58] deductions from atomic simulation calculations of defect energies on RE dopant incorporation in BaTiO3, i.e., barium-vacancy compensation is less favored than titanium-vacancy compensation. Ce may enter Ti sites as Ce4+ (4f 0, EPR silent) [29,37,38] or Ba sites as Ce3+ (4f 1, 2F5/2) [11–13,29], depending on the starting composition. Although BTC5 was prepared according to Ba(Ti10.05Ce0.05)O3, it might be inevitable that a small amount of Ce enters Ba sites as Ce3+, leading to A-rich behavior and further B-site Ti vacancies to maintain charge equilibrium [12,61]. A pronounced broadening of EPR response at 2500 G and the appearance of a g ¼ 2.004 signal in BTC5 [Fig. 15d] may be considered as an indication of Ce entering Ba sites in part as Ce3+ (1.34 A˚ [36]). The co-doping in BL3TC5 leads to disappearance of broadening at 2500 G in BTC5, which implies that La3+ has priority over Ce just as Ce3+ in entering Ba sites for solid solution formation of BL3TC5, further forming an A-rich mode and preventing Ce from occupying Ba sites as Ce3+ and forcing it into Ti sites as Ce4+. There is no signal for BL3T powder sintered at 1100 1C, indicating that La is not quite incorporated into the BaTiO3 lattice at this temperature. The selection of sintering temperature (1300 1C) may induce the appearance of Ti vacancies (Fig. 15b), but the g ¼ 2.004 signal intensity is lower than that in BTC5 or BL3TC5. A stronger signal with g ¼ 2.004 was observed in pellet sintered BL3T (Fig. 15(c)). However, the appearance of a very weak g ¼ 1.974 signal suggests that a small amount of Ba vacancies coexist with Ti vacancies. Our experiments demonstrated that the g ¼ 1.974 signal will disappear when La is increased to 0.04. Kornienko et al. observed a strong EPR line with g ¼ 1.963 in La-doped BaTiO3 at RT and assigned it to Ti3+–La3+ complex [64]. This signal was not observed in our BL3T samples, which suggests that the reduction of Ti4+ to Ti3+ did not occur. The above XRD and EPR studies on BL3T indicate that the selection of sintering temperature and contact with mixed grains may increase incorporation of La into the BaTiO3 lattice and furthermore induce the Ti vacancies. Ti vacancies may contribute to diffusion of Ce4+ ions into perovskite and 0000 the formation of VTi –Ce4+ defect complexes in BL3TC5.

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3.7. Raman scattering investigations The close relationship between ferroelectricity and lattice dynamics makes Raman spectroscopy (RS) a valuable technique for the study of ferroelectric materials. Raman scattering spectra can give some information on a local and

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at 985 cm1 in Fig. 17 is characteristic of the second-order Raman Spectra for BaTiO3 [81,82]. When BaTiO3 passes through the t–c phase transition point (TC), all the active Raman modes in the tetragonal phase (P4 mm) will be inactive in the perfect cubic phase (Pm3 m) (see Fig. 7a) owing to forbidding of Raman selected rules [83]. However, the two broad A1 (TO) bands at 260 and 515 cm1 persist into the cubic phase above TC, which is attributed to intrinsic disorder in the cubic phase [48,49,76,78,84], and they become broad and weak with increasing temperature [82,84]. The two bands at 305 and 715 cm1 disappear above TC, which is considered the signature of t–c phase transition. Raman spectrum of powder sintered BL3T (Fig. 16b) is similar to that of BT. By comparison with BT, the main features in pellet sintered BL3T (Fig. 16c) are: (1) all peaks broaden; (2) the anti-resonance effect at 180 cm1 weakens, and a new band at 183 cm1 [A1 (LO1)] appears; (3) the A1 (TO2) band red-shifts to 253 cm1; (4) the band intensity at 305 cm1 decreases, indicating that the substitution of 3% La at Ba sites does not thoroughly destroy long-range order of the ferroelectric BaTiO3; (5) a clear band at 836 cm1 associated with La appears. Kchikech and Maglione first reported a band at 840 cm1 in La-doped BaTiO3, with an intensity increasing linearly with La content [85]. Hence, this band may be considered as an indication of the amount of La incorporated into the BaTiO3 lattice. The evolution of the band at 836 cm1 reveals that clamped sintering (Fig. 16c) will incorporate more La into the BaTiO3 lattice than powder sintering (Fig. 16b), and contributes greatly to the decrease of tetragonality, in good agreement with XRD, DSC, and

EPR results. Especially, the band at 832–840 cm1 exists in Ba1x/2(Ti1xNbx)O3 (BTN) [86] and Ba(Ti1yCay)O3 [87] ceramics with aliovalent substituents, while it does not appear in (Ba1xSrx)TiO3 [68] and Ba(Ti1xZrx)O3 (BTZ) [70,88–90] ceramics with isovalent substituents. Therefore, this mode may be attributed to an internal deformation of the BO6 octahedron caused by the charge difference of different types of ions at equivalent sites in BaTiO3. Correspondingly, as in the case of La3+ at Ba sites, the weak peak at 823 cm1 in BTC5 may be attributed to Ce entering Ba sites in part as Ce3+ (Fig. 16d), while it has no connection with Ce4+ at Ti sites. This result is in accordance with the deduction from EPR. The most marked change in BTC5 is that the 261 cm1 and ‘‘silent’’ 305 cm1 modes of BT are merged into a sharp band peaking at 293 cm1 (Fig. 16d), evidently different from BTZ5 (5% Zr-doped BaTiO3) in clear ‘‘silent’’ mode at RT [70,88–90]. The XRD result calculated by peak positions indicates that BTC5 has a tetragonal structure, while lattice parameters of BTC5 are far greater than that of cubic BaTiO3 (Fig. 6). The disappearance of the ‘silent’ mode in BTC5 might be connected to the more pronounced expansion of TiO6 octahedra in BTC5 (Ce4+: 0.87 A˚, Ti4+: 0.605 A˚, Zr4+: 0.72 A˚ [36]). The ordered polar regions in BTC5 are indicated by the nondecreasing, broad A1 (LO3) band at 711 cm1, and the appearance of a new anti-resonance effect at 169 cm1 [A1 (TO1)+E (TO2)+E (LO1)] analogous to orthorhombic or rhombohedral perovskite phases [71,91]. In addition, a thinningdown of Raman bands in BTC5 with respect to BT was observed (Fig. 16 and the inset). A further broadening of all Raman bands was observed for BL3TC5 (Fig. 16e and the inset), indicating increasing average cubicity. The persistence of the broad A1 (LO3) mode at 717 cm1 reveals that the short-range polar ordering and lattice strain induced predominantly by RE dopants are present in BL3TC5, because the severalper-mille concentration of the background impurities is not high enough to induce such a strong A1 (LO3) band if BL3TC5 had an ideal cubic structure in light of the temperature-dependent Raman response of lightly doped BaTiO3 [92]. The characteristic band associated with La in BL3T persists in BL3TC5, peaking at 830 cm1. The most marked feature in BL3TC5 is both the disappearance of the ‘‘silent’’ mode at 305 cm1 and the asymmetric broadening shape of the A1 (TO2) band extending from 300 to 100 cm1. The former implies breaking of the long-range ordered ferroelectric phase. The red-shifted A1 (TO2) mode at 251 cm1, which is merged in BTC5, reappears as a main peak in BL3TC5, and the two peaks at 182 and 293 cm1 in BTC5 become two shoulders of the A1 (TO2) band in BL3TC5, revealing the existence of local polar ordering and B-site Ce4+ defects, respectively. A stronger band peaking at 111 cm1 appears clearly in BL3TC5. This band also appear in BTC5 (108 cm1) and BL3T (107 cm1) as a shallow protuberance. Farhi et al. and Dobal et al. attributed this band, which was observed

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at 117 cm1 in BTN [86] and at 119 cm1 in BTZ [88,89], to a mass effect related with Nb and Zr ions motion against oxygen octahedra [86,88]. Similarly, the mass effect is tenable for BTC5 as well. However, such an interpretation 0000 becomes unclear because it also exists in our BL3T with VTi 00 and BLT with VBa reported elsewhere [85]. It is apparent that the increase in the peak intensity at 111 cm1 for BL3TC5 is related to the co-doping with La and Ce, rather than cation vacancies. Because the clear anti-resonance effect at 126 cm1 is only formed in BTC5 and BL3TC5 with ‘DPT’, this interference effect at 126 cm1 caused by the coupling between two A1 (TO) modes with different types of BO6 octahedra may be considered as an indication of ‘DPT’. In a word, Raman scattering investigations reveal that BL3TC5 is a solid solution consisting of longrange lattice disorder and local polar order. 4. Discussion The e0 (T) of BL3TC5 shows a strong DPT. The frequency dispersion of e0 (T) on the low-temperature side of Tm indicates that BL3TC5 is a ‘‘relaxor’’-type material. We ever proposed a microscopic-scale compositional fluctuation model based on the interaction between LaBa 0000 or VTi and CeTi defects for interpreting the high-k ‘DPT’ phenomena and the rapid shift of Tm with La content in BLTC [39]. The recent studies have shown that composition fluctuations are not required to produce a relaxor state, but rather a nano-domains (or polar micro-regions) model (i.e., order/disorder model) that is responsible for relaxor behavior [93]. Samara summarized the nature of relaxor materials wonderfully [93], i.e., random lattice disorder produced by chemical substitution in ABO3 perovskites can lead to the formation of dipolar impurities and defects that have profound influence on the static and dynamic properties of these materials. Because of the high polarizability of the ABO3 host lattice associated with the soft ferroelectric mode, dipolar entities polarize regions around them forming polar nano-domains whose size is determined by the temperature-dependent correlation length (rc) of the host. When these dipolar entities possess more than one equivalent orientation, they may undergo dielectric relaxation in an applied ac field. In the very dilute limit (o0.1 at%), each polar domain behaves as a noninteracting dipolar entity. At higher concentrations of disorder, the polar domains can interact leading to more complex relaxational behavior with distribution of relaxation times. For BL3TC5, XRD, SEM, and TEM results indicate the formation of defect clusters and a homogeneous disorder of defects in the host BaTiO3 lattice; EPR studies demonstrate existence of Ti vacancies; RS and FHL investigations reveal the existence of local polar ordering. La3+ ion at the Ba site inevitably induces off-center Ti ion locations at eight positions slightly shifted on the cube diagonal. Ce4+ ion in CeO6 octahedron may also locate at an off-center position along the cube diagonal or (0 0 1)

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direction due to the formation of (La3+–CeTi, VTi –CeTi) defect complexes and the evidence of Raman scattering. It is well known that well-ordered perovskite materials show an FPT behavior. With increasing disorder in B-site cations, materials begin to exhibit the classical ‘DPT’ of a 0000 ferroelectric relaxor [94]. La3+ at Ba sites, VTi , off-center Ce4+ at Ti sites, and deformed oxygen octahedra may form randomly dipolar fields, and further induce dipoles in a region, forming nano-domains. Supposing that 5% Ce4+ ions are uniformly distributed in the BaTiO3 lattice, and 0000 La3+ and VTi form complexes with Ce4+, it can be predicted that the possible nano-domain size and its correlation length in BL3TC5 will be less than 2 nm. Such a small size cannot be detected by SEM or TEM. The presence of a higher concentration of defects and randomly local dipolar fields in BL3TC5 suppresses the formation of long-range polar ordering in the highly polarizable BaTiO3 lattice. The maximum permittivity no longer corresponds substantially to the t–c phase transition of ‘‘relaxor’’-type BL3TC5, evidenced by peak (2 0 0) broadening, average cubic structure below Tm, disappearance of the ‘‘silent’’ mode at 305 cm1 and persistence of the A1 (LO3) Raman mode at 717 cm1, and existence of FHL at temperatures above Tm, but rather to the slowing down of dipolar dynamics [93,95]. Correspondingly, the breadth of the permittivity peak is a manifestation of the response of dipolar nano-domains, rather than so-called ‘‘DPT’’. Dielectric permittivity is usually measured in a zero field (Fig. 9). The dynamics of the nano-domains may be very slow since they tend to align their polarization with the local fields, leading to the frequency dispersion of e0 (T) in the low-frequency range owing to a delay in the correction between nano-domains. Under both zero field and a strong dc biasing field, the frequency dispersion of BL3TC5 will gradually disappear with increasing biasing field, and e’m decreases linearly (the related data are not given), which suggests that the application of strong biasing field will align and grow nano-domains towards macrodomains, breaking up randomly oriented and slowed down nanodomains, i.e., forming a stronger ferroelectric state. Thus, lattice disorder is an essential ingredient for the occurrence of relaxor behavior in BL3TC5. XRD, DSC, EPR, and RS results for solid solution formation of BL3TC5 indicate that La is not easily incorporated into the BaTiO3 lattice when sintering at 1300 1C for 3 h, but rather segregates in part along grain boundaries [46,47]. A sufficient substitution of La at Ba sites occurs after sintering at 1350 1C for 3 days, leading to Ti vacancies which are expected as charge compensating defects, a linear shifting rate (25 1C/at% La) of Tm with FPT below x ¼ 0.06 [10]. Ce cannot be incorporated sufficiently into the Ti sites of the BaTiO3 lattice below 1400 1C [29], and our EPR and RS results show that Ce enters Ba sites in part as Ce3+ (1480 1C for 24 h). A sufficient substitution of Ce at Ti site occurs at 1540 1C for 6 h, leading to a higher shifting rate of Tm with ‘DPT’

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(7 1C/at% Ce) [37,38] than our BTC (3 1C/at% Ce) [39]. XRD monitoring for BLTC at high dopant concentration (xp0:125; yp0:2) reveals that La2O3 and CeO2 may accelerate solid solution into the BaTiO3 lattice even at a low temperature of 1100 1C [39]. Thus, it can be deduced that the formation of defect cluster in BL3TC5 (1480 1C for 24 h) may relieve the lattice strain and decrease incorporated energy associated with the individual lattice defects, leading to a homogeneous solid solution.

and Dr. Zen-ei Tanaka of Yamagata Center of Indus. Tech., for their support and help. This work is partly supported by Fundamental Research Project (No. 20060515) of Department of Science and Technology of Jilin Province, China. Thanks are also due to Mr. Hisao Fukumori of Fukumori International Patent Co. Ltd. Japan for financial support.

References 5. Conclusions The co-doping with 3% La at Ba sites and 5% Ce at Ti sites in BaTiO3 (BL3TC5) may lead to a fine-grained ceramics (0.9 mm), strong DPT of BaTiO3 around Tm near RT, and pronounced raising of dielectric permittivity in a low-frequency range (o100 kHz), meeting EIA specification of ‘Y5V’. The frequency dispersion of e0 (T) indicates that BL3TC5 is a typical ‘‘relaxor’’-type material. The dielectric permittivity is almost independent of electric field over a zero field range of 0.5–10 V/mm. BL3TC5 exhibits some advantages such as larger dielectric relaxation (DTC,HC ¼ 16 1C at 1 kHz), higher diffuseness with temperature (g ¼ 1.81), smaller dielectric loss (o0.05), and lower diffuseness with frequency (DT0 mf ¼ Tm,100 kHz–Tm,10 Hz ¼ 4 1C). The forma0000 tion of the defect complexes La3+–Ce4+ and VTi –Ce4+ is responsible for decreasing incorporated energy and relieving lattice strain associated with the individual lattice defects, leading to a disordered average cubic structure and high-k ‘relaxor’ behavior. The Ti-vacancy compensation mode in BL3TC5 is evidenced by the ESR signal at g ¼ 2.004. RS and FHL investigations reveal the existence of local polar ordering in BL3TC5 with randomly disordered defect cluster. Lattice disorder is an essential ingredient for the occurrence of relaxor behavior in BL3TC5. The high-k relaxor nature of BL3TC5 is characterized by average cubic structure with long-range lattice disordering and local polar ordering, a slow change of the e0 (T) and Pr(T) curves around Tm, no phase transition observed by DSC, and a broad, redshifted A1 (TO2) Raman phonon mode at 251 cm1 accompanied by the disappearance of the ‘‘silent’’ mode at 305 cm1 and a clear anti-resonance effect at 126 cm1. The paraelectricity and high-k behavior around RT in zero fields suggest that BL3TC5 is a promising candidate for high-k ‘Y5V’ capacitor or DRAM application in the future. Acknowledgments The authors would like to thank Professors Dr. Mikio Sugano, Dr. Susumu Ikeda, Dr. Tateaki Ogata, Dr. Kiyohito Koyama, Dr. Noriyuki Kuramono, Dr. Takashi Kuriyama; associate professors Dr. Masahito Sano, Dr. Hidero Unuma, Dr. Tomonori Koda, Dr. Norio Matsuda; Lecturer Mr. Akiyoshi Onuki; assistants Dr. Katsutoshi Okazaki, Dr. Akihiro Nishioka, Pos-Dr. Nihoko Nishio; technical assistance by Dr. Tadaaki Satake, Mr. Hideshige Suzuki, Mr. Tutomu Murayama of Yamagata University,

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