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Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics
Accepted Manuscript Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics Da-Yong Lu, Yan-Yan Peng, Xin-Yu Yu, Xiu-Yun Sun PII:
S0925-8388(16)31083-0
DOI:
10.1016/j.jallcom.2016.04.119
Reference:
JALCOM 37313
To appear in:
Journal of Alloys and Compounds
Received Date: 15 January 2016 Revised Date:
8 April 2016
Accepted Date: 12 April 2016
Please cite this article as: D.-Y. Lu, Y.-Y. Peng, X.-Y. Yu, X.-Y. Sun, Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.04.119. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Journal of Alloys and Compounds
Ms. Ref. No.: JALCOM-D-16-00766
Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics
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Da-Yong Lu a,*, Yan-Yan Peng a, Xin-Yu Yu a, b, Xiu-Yun Sun c,* Research Center for Materials Science and Engineering, Jilin Institute of Chemical Technology, Jilin 132022, China
b
College of Chemistry, Jilin University, Changchun 130021, China
c
College of Chemistry and Pharmaceutical Engineering, Jilin Institute of Chemical Technology, Jilin 132022, China
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a
ABSTRACT
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(Ba1−xLax)(Ti1−xTbx)O3 (0.03 ≤ x ≤ 0.20) (BLTT) and (Ba1−xLax)(Ti1−xTbx)O3-0.03nTb (x = 0.03, n = 0, 1, 3, 5, 7) (BL3TT-nT) ceramics were prepared using a mixed oxides method. The structure, microstructure, dielectric properties, valence state, and defect chemistry of BLTT and BL3TT-nT were investigated using X-ray diffraction (XRD), Raman spectroscopy, scanning electron microscopy (SEM), electron paramagnetic resonance (EPR), X-ray photoelectron spectroscopy (XPS), and dielectric temperature and frequency measurements. The solid solution limits
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of Tb in BLTT and BL3TT-nT were determined by XRD to be x = 0.15 and n = 3, respectively. The incorporation of the extra Tb ions in BLTT enhanced the relative density of the ceramic from 84 % to 96 %, accompanied by an improvement in dielectric permittivity. The defect chemistry of BLTT is discussed. Tb ions coexist in mixed-valence states, at the Ba-site as Tb3+, and at the Ti-site as Tb3+/Tb4+. The dielectric-peak temperature (Tm) decreased linearly
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at a rate of −19 °C/%(La-Tb) for BLTT. Due to its higher room-temperature permittivity (ε′ = 2450) and lower loss (tan δ = 0.029), BLTT with x = 0.07 can be considered as a promising dielectric for X7U applications. The existence
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of a small number of metastable Tb4+ ions at Ti-sites and the predominant presence of La3+-Tb3+ complexes at double sites are both responsible for the lower tan δ of BLTT.
Key-words: Ceramics; Mixed-valence; Defect complexes; Dielectric response; Electron paramagnetic resonance; X-ray photoelectron spectroscopy
* Corresponding author. Tel: +86 0432 62815308 E-mail Address: [email protected], [email protected] (D.-Y. Lu) XYSun@ jlict.edu.cn (X.-Y. Sun)
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1. Introduction
In the field of BaTiO3-based dielectric ceramic materials, La3+ ions are well-known to exclusively
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substitute at Ba-sites [1–3]. This site-occupation preference can create favorable conditions for the double-site substitution. Thus, the dielectric properties of BaTiO3 ceramics can be improved by means of co-doping with La and Mg [4], Sb [5], Sn [6], Zr [7], Ca [8], etc. La and Ce co-doped BaTiO3 ceramics
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with a high-permittivity (i.e., high-k) Y5V specification (−82% ≤ (ε'−ε'RT)/ε'RT ≤ +22% in a temperature range −30 to 85 °C) were discovered in 2006 [9]. Since then, double rare earth co-doping, such as doping
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with La/Tm [10] or La/Dy [11], has attracted attention from researchers because of the achievement of two typical high-k Y5V and temperature-stable X7R specifications (|(ε'−ε'RT)/ε'RT| ≤ 15% in a range −55 to 125 °C) [10,11]. The formation of La •Ba − Tm 'Ti or La •Ba − Dy'Ti complexes is predominant and is responsible for these dielectric properties of BaTiO3 ceramics co-doped with double rare earths.
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Another type of double-site substitution by double rare earths is called the self-compensation mode of R •Ba − R 'Ti (where R is a rare earth element) such as Eu •Ba − Eu 'Ti [12] and Dy•Ba − Dy'Ti [13].
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However, when Ba/Ti > 1 Eu can occupy Ba-sites predominantly in the mixed-valence states of Eu3+ (4f6)/Eu2+ (4f7). An interesting phenomenon is that Eu ions can exist stably as Eu3+ in Eu-doped BaTiO3
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with Ba/Ti = 1, where no valence change occurs [12]. In 2015, Tb ions were proven by EPR to be able to exist at the Ti-site, in part as Tb4+ (4f7); and in the case of Ba/Ti ≠ 1, Tb ions exhibit self-adjustable site occupations between the Ba-site as Tb3+ and the Ti-site as Tb4+ [14]. •
In this work, (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) ceramics in the self-compensation mode of La Ba − Tb Ti '
were prepared using a conventional mixed oxide method. The dielectric properties and point defects of BLTT were investigated. La and Tb were selected to be co-doped into BaTiO3, based on the following four considerations: (1) La donors play a key role in improving the dielectric permittivity of BaTiO3 ceramics [2] 2
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and Tb doping in BaTiO3 can result in low-temperature dielectric stability [14]. BaTiO3 ceramics co-doped with La and Tb were therefore expected to achieve high-k stability; (2) the exclusive occupations of larger
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La3+ ions at Ba-sites can expel smaller Tb ions into Ti-sites to form La •Ba − Tb 'Ti complexes, which is beneficial for the formation of single-phase ceramics; (3) Tb3+ (4f8) ions at Ti-sites have the ability to trap electrons in ceramics can thus lower dielectric loss by means of the transfer of Tb3+ to metastable Tb4+ (4f7);
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and (4) EPR of Tb4+ (4f7) [14] can be used to probe the valence stability of Tb3+ (4f8) at Ti-sites. The high compatibility of La and Tb in the BaTiO3 lattice [1–3,14] and the optimization of dielectric performance are
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common characteristics of these two elements. The self-adjustability of site occupations by Tb ions [14] was used to incorporate extra Tb ions into the BLTT lattice, to obtain highly dense ceramics. The dielectric properties of these ceramics were investigated.
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2. Experimental procedures
Ceramic raw materials were reagent-grade BaCO3, La2O3, TiO2, and Tb4O7 powders. The ceramics were prepared according to the nominal compositions (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03, 0.05, 0.07, 0.10, 0.15,
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0.20) (BLTT) and (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0, 1, 3, 5, 7) (BL3TT-nT) using a mixed oxides method with the same conditions as described elsewhere [8]. The sinter temperature (Ts) is based on
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the following considerations: (1) sufficient incorporation of La ions into the BaTiO3 lattice, which can contribute to improvement of the dielectric permittivity, needs a higher Ts and longer sintering time, for example, Ts = 1350 °C for 3 days for La-doped BaTiO3 ceramics [2]; (2) our experiments showed that very high Ts (=1400 °C) for La-doped BaTiO3 ceramics results in the formation of the liquid phase. The addition of Tb ions in BaTiO3 can solidify ceramics and inhibit appearance of the liquid phase [14]; (3) For BLTT with x = 0.07, our experiments showed that the ceramic density was improved 20 % when Ts was increased
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from 1300 to 1400 °C, accompanied by a change in color from white to light yellow. Thus, the final sintering conditions for La and Tb co-doped BaTiO3 ceramics were set as 1400 °C for 12 h in air.
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Structural characterization of ceramics was performed by powder X-ray diffraction (XRD) using a DX-2700 X-ray diffractometer (Dandong Haoyuan) employing Cu Kα1 radiation (λ = 1.540562 Å). Lattice parameters and unit cell volume were calculated by a MS Modeling software package (Accelrys, Inc.)
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using Rietveld refinement in the Reflex Package. Raman spectra were measured using a LabRAM XploRA Raman spectrometer (Horiba Jobin Yvon) with a 532 nm laser. Microstructural properties were determined
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using an EVOMA 10 scanning electron microscope (SEM) (Zeiss) operated at 15 keV. The average grain size (GS) was obtained from measurements of the number of grains per unit area on a polished surface NA using Fullman’s method [15], GS =
6 . Electron paramagnetic resonance (EPR) spectra were measured
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at room temperature using an A300-10/12 X-band spectrometer (Bruker) operating at 9.84 GHz. X-ray
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photoelectron spectra (XPS) measurements were performed at room temperature using an ESCALAB 250 X-ray photoelectron spectrometer (Thermo Electron Co.). The XPS data raw were processed by smoothing multiply times. The dielectric and conducting properties of the bulk ceramics were investigated at 1 kHz
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using a Concept 41 Dielectric/Impedance spectrometer (Novocontrol) with an applied voltage of 1 V. The accuracy in the measurements of dielectric permittivity, dielectric loss, resistivity, and temperature control
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are less than 5 %, 3 × 10−5, 3 × 10−5, and ±0.3 °C, respectively. Apart from the above systematic errors of the dielectric/impedance spectrometer, there are two factors to affecting the accuracy in measurement of dielectric permittivity: (1) Processing technique of electrodes: Ceramics were polished into disks (Φ10.2 mm, 0.8 mm in thickness) for dielectric measurements. Squared electrodes of 5 mm × 5 mm were applied on both surfaces of the ceramic disks. To minimize the effect of the air layer between the ceramic disk surface and contact
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electrode and to increase the conductivity, thin layers of Au atoms were sputtered on the surfaces of the polished ceramic disks, and then, the silver paste was spread uniformly on the thin Au electrode. These
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ceramic disks were heat-treated at 200 °C to form compact contact electrodes with good electrical conductivity. For these treatments, the accuracy in the measurement of dielectric permittivity of ceramics can be further improved, as shown in Fig. 1. It can be seen that the measured permittivity of ceramics with
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Ag electrodes formed by the silver paste, is greatly decreased because the existence of the air layer results in decreased capacitance. The sputtered Au atoms assist in squeezing out the air layer. On the other hand,
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compared to the circular electrodes (85 mm2) completely covering both surfaces of the ceramic disks (Φ10.2 mm), the use of smaller square electrodes (5 mm × 5 mm) has little influence on the measured permittivity, as shown in Fig. 2. The change in maximum permittivity (ε'm) is less than 4 %. (2) The edge effect of a parallel-plate capacitor: Scott and Curtis emphasize the necessity for accurate
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determination of the capacitance C of a parallel-plate capacitor comprised of two like rectangular plates with vacuum (air) as dielectric [16,17]. For a capacitor including dielectrics, the capacitance C and increment (∆) in
C can be expressed as follows:
d
+
2ε r ε 0 L πL ln( + 1) + 1 π d
2d πL ln( + 1) + 1 πL d
(1)
(2)
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∆=
ε r ε 0 L2
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C=
where L is the side length of square electrodes, εr the relative permittivity, ε0 the permittivity of free space, and d the separation between the plates. After the edge effect is considered, the measured permittivity should be reduced by 2.5 % compared to the actual permittivity. Thus, the uncertainty in the measurement of dielectric permittivity is less than 6.5 % after the effect of electrode shape and systematic errors of the dielectric/impedance spectrometer are included.
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3. Results
Powder XRD patterns of (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) ceramics (BLTT) are shown in Fig. 3.
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BLTT with x ≤ 0.15 exhibited a single-phase perovskite structure (Fig. 3(a)), indicating that La and Tb ions were completely incorporated into the BaTiO3 lattice. A small amount of a secondary phase of Tb4O7 was observed for BLTT with x = 0.20 whereas the main (1 0 0) peak of Tb4O7 did not appear at x = 0.15. (Fig. 3
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inset). The solid solution limit of Tb in BLTT was therefore determined to be greater than x = 0.15 by XRD. The x = 0.03 sample had a tetragonal structure, characterized by two separate (0 0 2) and (2 0 0) peaks,
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whereas for when x ≥ 0.05, BLTT converted to a cubic structure, as marked by a symmetric (2 0 0) peak (Fig. 3(b)).
Powder XRD patterns of (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0 – 7) ceramics (BL3TT-nT) are shown in Fig. 4. BL3TT-nT with n = 0 is the same sample as the above BLTT with x = 0.03 (BL3TT).
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There was no indication of any secondary phase separating out until n = 3. The solid solution limit of the extra Tb ions in BL3TT was determined to be n = 3, i.e., 3 % La and 12 % Tb ions can be incorporated
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simultaneously into the BaTiO3 lattice. Tb4O7 appeared as a secondary phase in the main perovskite phase for n ≥ 5. BL3TT-nT was tetragonal for n = 0 and 1 but converted to a cubic structure for n = 3 (Fig. 4(b)).
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The variations in lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for BLTT and BL3TT-nT are shown in Fig. 5. Both V0s of BLTT and BL3TT-nT are greater than the V0 of the tetragonal BaTiO3 (JCPDS Cards No. 5-626), and increase as x or n increases. In particular, V0 of BLTT with x = 0.15 is greater than that of the cubic BaTiO3 (JCPDS Cards No. 31-174). These results indicate the partial Ti-site occupation of Tb ions [14]. The Raman spectra of BLTT and BL3TT-nT at room temperature are shown in Fig. 6. The Raman spectrum of the cubic BaTiO3 generally exhibits three typical bands peaking at 265, 520, and 720 cm-1 6
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because of the intrinsic disorder in the cubic phase [18–20]. These bands also appeared in BLTT and BL3TT-nT. An additional band appeared and its peak position red-shifted from 818 to 805 (or 808) cm-1
number of Tb3+ ions incorporated into the BaTiO3 lattice [14,21].
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with increasing x or n. This band originates from the Raman charge effect [21], indicating an increase in the
The SEM images of BLTT and BL3TT-nT are shown in Fig. 7. The variations in average grain size (GS)
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as a function of x for BLTT and as a function of n for BL3TT-nT are shown in Fig. 8. BLTT exhibited a fine-grained microstructure and a uniform grain size distribution. The GS of BLTT increased from 0.8 to
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1.8 µm with increasing x. The relative density (ρr) here is defined as the ratio of the volumetric mass density to the theoretical density. The variations in ρr as a function of x for BLTT and as a function of n for BL3TT-nT are shown in Fig. 9. BL3TT-nT exhibited a lower porosity and a larger GS with respect to BLTT (Figs. 7–9). This suggests that Tb ions are not good at inhibiting grain-growth because GS = 3.6 µm for
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BL3TT-nT with n = 3 (Fig. 8). The ρr of BL3TT-nT was 96 % when n ≥ 1, whereas BLTT was less than 84 % when x ≤ 0.10 (Fig. 9). This result revealed that the incorporation of the extra Tb ions in BL3TT can greatly enhance the ceramic density.
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The temperature dependences of dielectric permittivity (ε'), loss (tan δ), and resistivity (ρ) for BLTT and BL3TT-nT are shown in Fig. 10. The dielectric peak of BLTT gradually broadened with increasing x and
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shifted toward lower temperatures. A weaker dielectric peak in the vicinity of 30 °C for the x = 0.03 BLTT sample (Fig. 10(a)) corresponded to the orthorhombic-tetragonal (t-c) phase transition (T2) of BaTiO3 [22]. The dielectric-peak temperature (Tm), representing the temperature at which maximum permittivity (ε'm) occurred, decreased linearly at a rate of −19 °C/%(La-Tb) (Fig. 10 inset). BLTT with x = 0.07 and BL3TT-nT with n = 3 satisfied the X7U dielectric specification, i.e., −56 % ≤ |ε'−ε'RT| ≤ +22 % in the range of −55 to 125 °C (Figs. 10(a) and (d)). BLTT and BL3TT-nT with n ≤ 1 exhibited a lower tan δ (≤ 0.034)
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below 75 °C (Figs. 10(b) and (e)) and their resistivity (ρ) was 107 − 108 Ω cm (Figs. 10(c) and (f)). A higher room-temperature value of tan δ = 0.058 and a very low value of ρ = 5 × 106 Ω cm for BL3TT-nT
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with n = 3 relates to a higher concentration of the extra Tb ions. The addition of the extra Tb ions in BL3TT resulted in an increase in ε'.
The frequency (f) dependences of ε'RT and tan δ were measured for two X7U dielectrics, BLTT with x =
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0.07 and BL3TT-nT with n = 3, as shown in Fig. 11, where the inset depicts the ρ-f curve. The ε' of the two X7U dielectrics remained almost constant in the f range of 10 to105 Hz. BLTT with x = 0.07 exhibited a
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higher frequency stability compared with BL3TT-nT with n = 3 (Fig. 11(a)). Both tan δ and ρ were sensitive to f and decreased as f was increased(Fig. 11(b) and inset).
The room-temperature EPR spectra of BLTT and BL3TT-nT are shown in Fig. 12. A broad signal with g ~6.5 can be attributed exclusively to Tb4+ (4f7) Kramers ions at Ti-sites [14]. A weaker Tb4+-related signal
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could be observed for BLTT and BL3TT-nT. To qualitatively understand the extent of occupation of Ti-ites by Tb4+ , we measured an EPR spectrum of the precious sample of Ba(Ti1−xTbx)O3 with x = 0.05 (BTT5) [14] at room temperature, and compared this spectrum with the EPR spectrum of (Ba1−xLax)(Ti1−xTbx)O3
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(BLTT) with x = 0.05, as seen in Fig. 13. Tb3+ Kramers ions are EPR-silent in theory. Tb ions in BTT5 were found to exist predominantly at Ti-sites as Tb4+, and only a small number of Tb3+ ions substituted at
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Ba-sites [14], assuming that the intensity of the g
~6.5 signal is proportional to the concentration of
Ti-site Tb4+ ions. If all of the Ti-site Tb ions were Tb3+, the Tb4+-related signal would disappear. It can be seen from Fig. 13 that the Tb4+-related signal with g ~6.5 was very weak for the x = 0.05 BLTT sample,as ×
compared to BTT5. This comparison reveals two results: (1) the number of Ti-site Tb4+ ions ( Tb Ti ) in BLTT is smaller; and (2) Tb ions in BLTT exist predominantly as Ti-site Tb3+ ( Tb 'Ti ) because to a certain extent the ratio of La/Tb = 1 can limit the transfer of Tb ions from Ti-sites to Ba-sites. The comparison in
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Fig. 13 further clarifies that La •Ba − Tb 'Ti defect complexes in the BLTT lattice cannot be completely •
×
formed and a small number of La Ba − Tb Ti complexes may coexist. For this reason and owing to the
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amphoteric behavior of Tb ions, a small number of Tb ions inevitably entered Ba-sites as Tb3+ ( Tb •Ba ) to preserve the lattice electroneutrality [14]. The presence of the g = 2.004 signal associated with Ti vacancies [18,23,24] (Fig. 12) in all of the samples further verifies the occupation of Ba-sites by Tb3+ because of the
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relation 4Tb •Ba → VTi'''' .
The signal at g = 1.974, which appeared only in BLTT with x = 0.03, was attributed to ionized
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'' ) [13,23,25]. This indicates that intrinsic Ba as well as Ti vacancies may coexist in Ba-vacancy defects ( VBa
the main perovskite lattice.
Mn impurity usually exists in the starting materials, as detected in Eu-doped BaTiO3 ceramics [12]. The appearance of Mn2+ sextet signal in BLTT with x = 0.03 in Fig. 12 implies the reduction of Mn4+ or Mn3+ in
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ceramics to Mn2+. This phenomenon was also observed for (Ba1−xLax)(Ti1−xDyx)O3 with x = 0.03, after ''
which Mn2+ signal disappeared with increasing x [11]. An interesting phenomenon is that VBa -related signal is accompanied by the appearance of the Mn2+ sextet signal, and these signals disappear
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simultaneously when x > 0.03 or n > 1. This fact suggests that as x is increased, Tb ions may transfer from Ti-ites to Ba-ites in BLTT to fill in Ba vacancies.
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The mixed-valence states of Tb4+/Tb3+ in BLTT and BL3TT-nT were further confirmed by XPS. The Tb 4d XPS spectra of BLTT with x = 0.10 and 0.15 and BL3TT-nTb with n = 1 and 3 were measured and smoothed, as shown in Fig. 14. The Tb 4d core-level lines in compounds generally show multiplet structures, which are similar to the main and satellite lines in pure metals [26,27]. The peaks with several higher core-level binding energies originate from Tb4+ 4d core-level lines [26], and in particular, a peak at ~155 eV can be considered as an indicator of Tb4+ [26–28]. The peaks at ~153, ~150, and ~148 eV can be
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used to determine whether a compound has Tb3+ ions or not [26–30]. On the basis of these investigations, three peaks at 155.9, 153.2 and 150.7 eV in Fig. 14 were attributed to contributions from multiplet
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structures of Tb4+ in BLTT, while the three peaks at 153.2, 150.7 and 147.8 eV in Fig. 14 were attributed to multiplet structures of Tb3+. The two intermediate peaks at 153.2 and 150.7 eV cannot be unambiguously distinguished because each peak is a superposition of Tb4+ and Tb3+ lines. The XPS configuration of BL3TT-nT is similar to that of BLTT. The extra Tb ions in BL3TT-nT mainly mainly form •
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•
×
self-compensated Tb Ba − Tb Ti complexes except for a small number of Tb Ba − Tb Ti complexes [31], '
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as seen in Fig. 15 by the Tb4+ EPR comparison between BL3TT-nT (n = 3) and (Ba1−xTbx)(Ti1−xTbx)O3 (x = 0.05) with a nominal self-compensation mode.
The Ti 2p XPS spectra of BLTT and BL3TT-nTb are shown in Fig. 16. Ti ions in BLTT and BL3TT-nTb were confirmed to maintain a stable oxidation state of 4+ based on the following four reasons: (1) the two
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prominent peaks were identified as Ti 2p3/2 at ~458.1 eV and Ti 2p1/2 at ~463.9 eV, respectively, and coincide with the reported values of Ti 2p XPS for BaTiO3 and SrTiO3 [32] with the Ti4+-ion profile; (2) the difference between Ti 2p3/2 and Ti 2p1/2 orbit-spin splitting binding energies in BLTT and BL3TT-nTb (5.8
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eV) is the nearly same as that of BaTiO3 (5.6 eV) [33]; (3) no peak related to Ti3+ 2p is evident [34]; and (4) the ∆FWHM (full width at half maximum) value (1.7 - 1.8 eV) of Ti 2p3/2 main peak is less than that of
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BaTiO3 (< 2.0 eV) [35], which also suggests that Ti3+ or mixed-valence Ti4+/Ti3+ do not exist in BLTT and BL3TT-nTb.
4. Discussion
4.1. Defect chemistry
For the single-phase (Ba1−xLax)(Ti1−xTbx)O3 (0.03 ≤ x ≤0.15) ceramics (BLTT), the XRD results reveal
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the predominant occupation of Tb ions at Ti-sites (Fig. 5). The above EPR and XPS results indicate that the •
•
''
point defects in BLTT with x = 0.03 are Ba-site La3+ ( La Ba ), Ba-site Tb3+ ( TbBa ), Ba vacancies ( VBa ),
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Ti-site Tb3+ ( Tb'Ti ), Ti-site Tb4+ ( Tb×Ti ), and Ti vacancies ( VTi'''' ). Mn impurities in ceramics were neglected here because of very low concentrations. At a lower doping level of x = 0.03, both Ba and Ti vacancies may coexist in the main perovskite lattice (Fig. 12).
•
''
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To a certain extent Tb ions can transfer from Ti-sites to Ba-sites during sintering and upon cooling. When x ≥ 0.05, Ba vacancies are filled up by Tb Ba and the VBa -related EPR signal disappears (Fig. 12). Tb
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ions in BLTT exist predominantly as Tb 'Ti (Fig. 13), forming La •Ba − Tb 'Ti defect complexes because of the exclusive site occupation of La3+ at Ba-sites.
2La2O3 + Tb 4O7 → 4La •Ba + 4Tb'Ti + 12OO + 1/2O2 ↑
(3)
A small number of Tb ions may exist at Ba-sites in the form of Tb3+ ( Tb •Ba ) because of the amphoteric ×
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behavior of Tb ions [14,31]. On the other hand, some Ti-site Tb ions exist as Tb4+ ( TbTi ), implying that half of the Tb4+ ions in the starting material Tb4O7 cannot be reduced completely into Tb3+ ions. However, the number of Tb×Ti is very small as shown by the EPR comparison (Fig. 13) and XPS observations (Fig. •
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14). The appearance of a small number of Tb Ba and Tb×Ti may expressed by
2Tb4O7 + BaO → 4Tb•Ba + Ba Ba + 4Tb×Ti + VTi'''' + 15OO •
(4)
''
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The incorporation of Tb Ba into Ba-sites through VBa results in an excess of Ba-site cations and •
•
further in the steady existence of VTi'''' (Fig. 12), caused by the extra La Ba and Tb Ba ,in all of the •
samples except for La Ba − Tb Ti defect complexes. '
4La•Ba → VTi''''
(5)
4Tb•Ba → VTi''''
(6)
•
Thus, La Ba − Tb Ti defect complexes cannot be completely formed in BLTT and a small number of '
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Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes may coexist in the BaTiO3 lattice. The defect chemistry of the single-phase (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0, 1, 3) ceramics
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(BL3TT-nT) is similar to that of BLTT. The extra Tb ions in BL3TT-nT exist as Ba-site Tb3+ ( Tb•Ba ), Ti-site Tb3+ ( Tb'Ti ), and Ti-site Tb4+ ( Tb×Ti ) [31]. This behavior is similar to the predominant •
•
self-compensation mode of Tb ions in BaTiO3 [31]. A large number of La Ba − Tb Ti and Tb Ba − Tb Ti '
'
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defect complexes and a small number of La •Ba − Tb×Ti and Tb •Ba − Tb×Ti defect complexes may coexist in BL3TT-nT.
4.2.1. Linear shift of Tm with x for BLTT
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4.2. Dielectric properties
•
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The above analyses of the defect chemistry of BLTT indicate that La Ba − Tb Ti defect complexes are '
predominant, and the number of Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes is very small. The •
•
shift of Tm with x is therefore mainly determined by La Ba − Tb Ti defect complexes. Tb Ba − Tb Ti and '
'
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La •Ba − Tb×Ti complexes have a minor effect on a shift in Tm. The Tm decreases linearly at a rate of
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−19 °C/%(La-Tb) for BLTT (Fig. 10(b) inset). This is mainly attributed to a linear increase in the quantity •
of La Ba − Tb Ti defect complexes with an increase in x. '
For BL3TT-nT, the Tm decreases nonlinearly with n (Fig. 10(e) inset). Similar to the analyses on BLTT, •
•
the shift of Tm with n is mainly determined by the predominant La Ba − Tb Ti and Tb Ba − Tb Ti defect '
'
complexes. The concentration of La ions is fixed at 3 %, whereas the number of Tb ions increases with n. •
•
Thus, the nonlinear shift in Tm is attributed to the different dielectric responses of La Ba and Tb Ba .
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4.2.2. Lower dielectric loss of BLTT
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All of the BLTT samples show lower tan δ below 75 °C (Fig. 10(e)). Because the Mn2+ EPR signal appears only in BLTT with x = 0.03, it is concluded that the reduction of Mn4+ or Mn3+ to Mn2+ impurities •
in ceramics is not the main reason for lowering of tan δ. The predominant La Ba − Tb Ti defect complexes '
and the small number of Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes in BLTT can suppress the •
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•
donor effect of La Ba and Tb Ba and the acceptor effect of Tb 'Ti . On the other hand, the lower tan δ is
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probably due to the existence of a small number of metastable Tb×Ti ions. To a certain extent, each Ti-site Tb4+ ion in the ceramic has some ability to trap an electron and can thus be reduced to Tb 'Ti , which •
concatenates with Tb Ba to form the self-compensation mode of Tb•Ti − Tb'Ti .
Tb×Ti + e′ → Tb'Ti
(7)
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The above two reasons can contribute to the reduction in dielectric loss of BLTT.
4.2.3. Comparison between two X7U dielectrics and Effect of the extra Tb ions on dielectric properties of
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BLTT
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Both BLTT with x = 0.07 (BLTT7) and BL3TT-nT with n = 3 (BL3TT-3T) satisfy the X7U dielectric specifications. Compared with BLTT7, BL3TT-3T exhibits a higher ε'RT (= 5990) and tan δ (= 0.058) (Figs. 10(b) and (e)) at room temperature. The SEM results show that the relative density (ρr) of BL3TT-3T is •
96 %, whereas that of BLTT7 is 84 % (Figs. 7 and 8). Because La Ba plays a significant role in improving ε' among rare-earth donors [2], the improvement in ε'RT caused by the addition of the extra Tb ions in BLTT •
cannot be attributed to the Tb Ba − Tb Ti defect complexes, but is most likely related with the rapid '
13
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improvement in ceramic density of BL3TT. At a higher Tb doping level of n = 3, the higher tan δ of •
•
BL3TT-3T implies that the donor effect of Tb Ba is greater than that La Ba . Due to its lower loss (tan δ =
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0.029), the X7U BLTT7 is a promising dielectric for X7U applications
5. Conclusions
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(Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) and (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics were prepared using a mixed oxides method. BLTT with 0.03 ≤ x ≤0.15 and
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BL3TT-nT with 0 ≤ n ≤ 3 exhibit a single-phase perovskite structure. The unit cell volume increases with increasing x or n. The XRD and Raman results provide evidence of the predominant occupation of Tb ions at Ti-sites. The incorporation of the extra Tb ions in BLTT with x = 0.03 can greatly enhance the ceramic density of BLTT, showing an amazing relative density of 96 %. A study of the defect chemistry indicates •
•
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that the point defects in BLTT with x ≥ 0.05 are Ba-site La3+ ( La Ba ), Ba-site Tb3+ ( Tb Ba ), Ti-site Tb3+ •
( Tb'Ti ), Ti-site Tb4+ ( Tb×Ti ), and Ti vacancies ( VTi'''' ). A large number of La Ba − Tb Ti defect complexes •
×
•
'
and a small number of La Ba − Tb Ti and Tb Ba − Tb Ti defect complexes may coexist in BLTT. The Tm '
•
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decreases linearly at a rate of −19 °C/%(La-Tb) for BLTT. The predominant formation of La Ba − Tb Ti '
defect complexes is responsible for the linearly shifting Tm rate. The existence of a small number of •
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metastable Tb×Ti ions and the predominant presence of La Ba − Tb Ti defect complexes are both '
responsible for the lower dielectric loss of BLTT. BLTT with x = 0.07 satisfies X7U dielectric specifications. BLTT with x = 0.07 because of its higher room-temperature permittivity (ε′ = 2450) and lower loss (tan δ = 0.029), can be considered as a promising dielectric for X7U applications.
Acknowledgements
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This work was supported by the projects of the National Natural Science Foundations of China (Grant No. 21271084) and of Jilin Province (Grant No. 20160101290JC).
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References
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Z. Phys. B 96 (1995) 451–454.
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45 (1982) 14–39.
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Figure captions
Fig. 1. Temperature dependence of the dielectric permittivity (ε') for a BaTiO3-based sample A with (a) sputtered Au/Ag electrodes and with (b) Ag electrodes, measured at 1 kHz using an RCL meter (Fluke PM6306).
Fig. 2. Temperature dependence of ε' for a BaTiO3-based sample B with (a) square electrodes of 5 mm × 5
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mm and with (b) circular electrodes of 85 mm2, measured at 1 kHz using an RCL meter (Fluke PM6306). Fig. 3. XRD patterns of (a) (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BLTT.
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Fig. 4. XRD patterns of (a) (Ba0.97La0.03)(Ti0.97Tb0.03)O3 – 0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BL3TT-nT.
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Fig. 5. Variations in the lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for (a) BLTT and (b) BL3TT-nT ceramic powders. The V0s of the tetragonal BaTiO3 (t-BaTiO3, JCPDS Cards No. 5-626) and the cubic BaTiO3 (c-BaTiO3, JCPDS Cards No. 31-174) are drawn as a dashed line for comparison. Fig. 6. Room-temperature Raman spectra of (a) BLTT and (b) BL3TT-nT with a single-phase structure. Fig. 7. SEM images of thermally etched surface of BLTT with (a) x = 0.03, (b) x = 0.05, (c) x = 0.07, (d) x = 0.10; and of BL3TT-nT with (e) n = 1, (f) n = 3. Fig. 8. Variations in average grain size (GS) as a function of x for BLTT and as a function of n for
18
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BL3TT-nT. Fig. 9. Variations in relative density (ρr) as a function of x for BLTT and as a function of n for BL3TT-nT. Fig. 10. Temperature dependences of (a) and (d) dielectric permittivity (ε'), (b) and (e) dielectric loss (tan
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δ), and (c) and (f) resistivity (ρ) for BLTT and BL3TT-nT, measured at 1 kHz.
Fig. 11. Frequency dependences of (a) ε' and (b) tan δ for the x = 0.07 BLTT and the n = 3 BL3TT-nT ceramics, measured at 25 °C from f = 10 to105 Hz. The inset depicts the ρ-f curve.
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Fig. 12. Room-temperature EPR spectra of (a) BLTT and (b) BL3TT-nT.
Fig. 13. Room-temperature EPR comparison between (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) and Ba(Ti1−xTbx)O3
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[14] with x = 0.05.
Fig. 14. Tb 4d XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3. Fig. 15. Room-temperature EPR comparison between BL3TT-nT with n = 3 and (Ba1−xTbx)(Ti1−xTbx)O3 with x = 0.05 [28].
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Fig. 16. Ti 2p XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
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Au + Ag electrodes
1.0
Ag electrodes
0.5
0.0
-50
0
50
100 o
T ( C)
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ε' (× 104)
BaTiO3-based ceramic A
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1.5
150
200
AC C
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Fig. 1. Temperature dependence of the dielectric permittivity (ε') for a BaTiO3-based sample A with (a) sputtered Au/Ag electrodes and with (b) Ag electrodes, measured at 1 kHz using an RCL meter (Fluke PM6306).
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2
1.0
A = 25 mm
A = 85 mm2
0.5
0.0
-50
0
50 100 o T ( C)
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4
ε' (× 10 )
BaTiO3-based ceramic B
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1.5
150
200
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Fig. 2. Temperature dependence of ε' for a BaTiO3-based sample B with (a) square electrodes of 5 mm × 5 mm and with (b) circular electrodes of 85 mm2, measured at 1 kHz using an RCL meter (Fluke PM6306).
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(211)
(200)
(200)
32
x = 0.15 (200)
30
(220) (300) (310) (311) (222)
x = 0.20
*
x = 0.10 (200)
(200)
x = 0.07
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x = 0.10
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x = 0.15
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(110)
(200)
x = 0.20
2θ (o)
(210)
(111)
(100)
(100) from Tb 4O7
x = 0.15
28
(110)
Tb4O7 x = 0.20
Intensity (arb. units)
(b)
(Ba1-x Lax)(Ti1-x Tbx)O3 (100)
(110)
(a)
x = 0.07
x = 0.05
20
40
60
2θ (o)
80
(002) (200)
(113)/(311) (222)
(212) (103)/(310)
x = 0.03 (202)
(211)
(201)
(200)/(002)
(111)
(100)
(101)/(110)
x = 0.05
x = 0.03
44
45
46
2θ (o)
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Fig. 3. XRD patterns of (a) (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BLTT.
(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3
(110)
(200)
(211) *
*
*
*
*
(311)
*
n=5
(220) (300) (310) (311) (222)
(200) (210)
(220)
(110)
n=7
(200)
(100)
(100)
n=5
n=3
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*
n=3
(002) (200)
Intensity (arb. units)
T b4O7
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n=7 *
*
(b)
- 0.03nTb
(200)
(a)
(200)
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n=1
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(002) (200)
n=1
n=0
n=0
0 20
40
60
2θ (o)
80
44
45
46
2θ (o)
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Fig. 4. XRD patterns of (a) (Ba0.97La0.03)(Ti0.97Tb0.03)O3 – 0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BL3TT-nT.
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67.2
4.08
(a)
4.04
o
64.8
4.02 o
3
t-BaT iO3 ( V0 = 64.41 A )
64.0 c
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4.00
a 63.2 0.02
0.04
0.06
0.08
0.10
0.12
0.14
x (Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb o
4.12
3
c-BaT iO3 ( V0 = 65.50 A )
65.6
V 0 (A3)
3.98 0.16
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66.4
(b)
V0
4.08
o
64.8
o
3
t-BaT iO3 ( V0 = 64.41 A )
64.0
o
a, c (A)
V 0 (A3)
o
c-BaT iO3 (V0 = 65.50 A3)
65.6
a, c (A)
4.06
V0
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(Ba1-xLa x)(Ti 1-xTb x)O3
66.4
o
4.04
c
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a
63.2
0
2
4
6
4.00 8
n
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Fig. 5. Variations in the lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for (a) BLTT and (b) BL3TT-nT ceramic powders. The V0s of the tetragonal BaTiO3 (t-BaTiO3, JCPDS Cards No. 5-626) and the cubic BaTiO3 (c-BaTiO3, JCPDS Cards No. 31-174) are drawn as a dashed line for comparison.
808 811
720
520
n=3
818
n=1
200
400
600
(a)
1000
1200
805
720
520
805
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(Ba1-xLax)(Ti 1-xTb x)O3 265
Intensity (a. u.)
800
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n=0
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(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb
Intensity (a. u.)
(b)
265
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818
805
x = 0.15
200
400
600
x = 0.10 x = 0.07 x = 0.05
800
1000
1200
Raman shift(cm -1)
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Fig. 6. Room-temperature Raman spectra of (a) BLTT and (b) BL3TT-nT with a single-phase structure.
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(b)
(c)
(d)
(f)
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(e)
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(a)
Fig. 7. SEM images of thermally etched surface of BLTT with (a) x = 0.03, (b) x = 0.05, (c) x = 0.07, (d) x = 0.10; and of BL3TT-nT with (e) n = 1, (f) n = 3.
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n 0
1
2
3
4
BL3TT-nT 2
1
0 0.05
0.1
x
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BLTT
0.15
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GS ( µ m)
3
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Fig. 8. Variations in average grain size (GS) as a function of x for BLTT and as a function of n for BL3TT-nT.
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n 0
1
2
3
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100
ρr (%)
BL3TT-nT 90
80
70 0.05
0.10
x
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BLTT
0.15
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Fig. 9. Variations in relative density (ρr) as a function of x for BLTT and as a function of n for BL3TT-nT.
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8000
8000
n =1
4000
2000
2000
0 0.30
0 0.30 120
(e) -19 C/%(La-Tb)
0.20 0
tanδ
-60 0.02
0.10
0.04
0.06
0.08
80
o
60
120
40
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tanδ
0.20
n =3
o
o
T m ( C)
(b)
0.10
0
0.10
x
0.00 2
0
1
2
3
n
0.00 2
(f)
8
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(c) ρ (×108 Ωcm)
n =0
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4000
(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb
6000
ρ (×10 Ωcm)
ε′
6000
= 0.03 = 0.05 = 0.07 = 0.10
ε′
x x x x
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(d)
(Ba1-xLa x)(Ti 1-xTb x)O3
T m ( C)
(a)
1
0 -50
0
50
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T (oC)
100
150
200
1
0 -50
0
50
100
150
200
T (oC)
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Fig. 10. Temperature dependences of (a) and (d) dielectric permittivity (ε'), (b) and (e) dielectric loss (tan δ), and (c) and (f) resistivity (ρ) for BLTT and BL3TT-nT, measured at 1 kHz.
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10000
(a)
o
T = 25 C
BLTT with x = 0.07
8000
BL3TT with n = 3
ε′
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6000 4000 2000
0.10
103 8 6 4 2 0 101
102
104
103
105
104
105
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tan δ
0.15
8
(b)
ρ (×10 Ωcm)
102
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0 101 0.20
f (Hz)
0.05
0.00 101
102
103
104
105
f (Hz)
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Fig. 11. Frequency dependences of (a) ε' and (b) tan δ for the x = 0.07 BLTT and the n = 3 BL3TT-nT ceramics, measured at 25 °C from f = 10 to105 Hz. The inset depicts the ρ-f curve.
(b)
4+
TbT i
g = 2.004 (VT i)
Intensity
n=3
n=1
Mn
(a)
2000 TbT4+i
3000
g = 2.004 (VT i)
x = 0.15
Intensity
4000
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1000
2+
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g = 1.974 (VBa)
n=0
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x = 0.07 x = 0.05
x = 0.03
(VBa) g = 1.974
1000
2000
3000
4000
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H (G)
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Fig. 12. Room-temperature EPR spectra of (a) BLTT and (b) BL3TT-nT.
TbTi4+
Receiver Gain: 1.00 ×10
4
Intensity (a. u)
g = 6.58
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Ba(Ti 1-xTb x)O3 with x = 0.05 [14]
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(Ba 1-xLax)(Ti 1-xTb x)O3 with x = 0.05
VT i g = 2.004
1000
2000
3000
4000
5000
6000
H (G)
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Fig. 13. Room-temperature EPR comparison between (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) and Ba(Ti1−xTbx)O3 [14] with x = 0.05.
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Fig. 14. Tb 4d XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
4+
Intensity (a. u)
TbTi
Receiver Gain: 1.00 ×10
4
(Ba 1-xTbx)(Ti 1-xTbx)O3 with x = 0.05 [28]
(Ba 0.97La 0.03)(Ti 0.97Tb0.03)O3-0.03nTb
VT i g = 2.004
1000
2000
3000
4000
5000
6000
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with n = 3
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Fig. 15. Room-temperature EPR comparison between BL3TT-nT with n = 3 and (Ba1−xTbx)(Ti1−xTbx)O3 with x = 0.05 [28].
5.8 eV
n=3
Ti 2p3/2 = 1.8eV
Ti 2p1/2 464.0
FWHM
n=1
5.8 eV
463.9
FWHM = 1.7 eV
x = 0.15
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Ti 2p1/2
FWHM = 1.7eV
463.9
Intensity (a. u.)
(a) BLTT
465
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460 458.1
455
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= 1.8eV
464.0
FWHM
458.2
Intensity (a. u.)
(b) BL3TT- nT
458.2
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x = 0.10
455
460
465
Binding energy (eV)
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Fig. 16. Ti 2p XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
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Highlights
► Dielectric properties and defect chemistry of (Ba1−xLax)(Ti1−xTbx)O3 were studied.
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► Solid solution limit of Tb in BLTT is determined by XRD to be greater than x = 0.15. ► Tb ions coexist in the mixed valence states of Ba-site Tb3+ and Ti-site Tb3+/Tb4+.
► The dielectric-peak temperature (Tm) decreased linearly at a rate of −19 °C/%(La-Tb).
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► The x = 0.07 sample is a promising dielectric for X7U applications.
S0925-8388(16)31083-0
DOI:
10.1016/j.jallcom.2016.04.119
Reference:
JALCOM 37313
To appear in:
Journal of Alloys and Compounds
Received Date: 15 January 2016 Revised Date:
8 April 2016
Accepted Date: 12 April 2016
Please cite this article as: D.-Y. Lu, Y.-Y. Peng, X.-Y. Yu, X.-Y. Sun, Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics, Journal of Alloys and Compounds (2016), doi: 10.1016/j.jallcom.2016.04.119. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Journal of Alloys and Compounds
Ms. Ref. No.: JALCOM-D-16-00766
Dielectric properties and defect chemistry of La and Tb co-doped BaTiO3 ceramics
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Da-Yong Lu a,*, Yan-Yan Peng a, Xin-Yu Yu a, b, Xiu-Yun Sun c,* Research Center for Materials Science and Engineering, Jilin Institute of Chemical Technology, Jilin 132022, China
b
College of Chemistry, Jilin University, Changchun 130021, China
c
College of Chemistry and Pharmaceutical Engineering, Jilin Institute of Chemical Technology, Jilin 132022, China
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a
ABSTRACT
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(Ba1−xLax)(Ti1−xTbx)O3 (0.03 ≤ x ≤ 0.20) (BLTT) and (Ba1−xLax)(Ti1−xTbx)O3-0.03nTb (x = 0.03, n = 0, 1, 3, 5, 7) (BL3TT-nT) ceramics were prepared using a mixed oxides method. The structure, microstructure, dielectric properties, valence state, and defect chemistry of BLTT and BL3TT-nT were investigated using X-ray diffraction (XRD), Raman spectroscopy, scanning electron microscopy (SEM), electron paramagnetic resonance (EPR), X-ray photoelectron spectroscopy (XPS), and dielectric temperature and frequency measurements. The solid solution limits
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of Tb in BLTT and BL3TT-nT were determined by XRD to be x = 0.15 and n = 3, respectively. The incorporation of the extra Tb ions in BLTT enhanced the relative density of the ceramic from 84 % to 96 %, accompanied by an improvement in dielectric permittivity. The defect chemistry of BLTT is discussed. Tb ions coexist in mixed-valence states, at the Ba-site as Tb3+, and at the Ti-site as Tb3+/Tb4+. The dielectric-peak temperature (Tm) decreased linearly
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at a rate of −19 °C/%(La-Tb) for BLTT. Due to its higher room-temperature permittivity (ε′ = 2450) and lower loss (tan δ = 0.029), BLTT with x = 0.07 can be considered as a promising dielectric for X7U applications. The existence
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of a small number of metastable Tb4+ ions at Ti-sites and the predominant presence of La3+-Tb3+ complexes at double sites are both responsible for the lower tan δ of BLTT.
Key-words: Ceramics; Mixed-valence; Defect complexes; Dielectric response; Electron paramagnetic resonance; X-ray photoelectron spectroscopy
* Corresponding author. Tel: +86 0432 62815308 E-mail Address: [email protected], [email protected] (D.-Y. Lu) XYSun@ jlict.edu.cn (X.-Y. Sun)
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1. Introduction
In the field of BaTiO3-based dielectric ceramic materials, La3+ ions are well-known to exclusively
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substitute at Ba-sites [1–3]. This site-occupation preference can create favorable conditions for the double-site substitution. Thus, the dielectric properties of BaTiO3 ceramics can be improved by means of co-doping with La and Mg [4], Sb [5], Sn [6], Zr [7], Ca [8], etc. La and Ce co-doped BaTiO3 ceramics
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with a high-permittivity (i.e., high-k) Y5V specification (−82% ≤ (ε'−ε'RT)/ε'RT ≤ +22% in a temperature range −30 to 85 °C) were discovered in 2006 [9]. Since then, double rare earth co-doping, such as doping
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with La/Tm [10] or La/Dy [11], has attracted attention from researchers because of the achievement of two typical high-k Y5V and temperature-stable X7R specifications (|(ε'−ε'RT)/ε'RT| ≤ 15% in a range −55 to 125 °C) [10,11]. The formation of La •Ba − Tm 'Ti or La •Ba − Dy'Ti complexes is predominant and is responsible for these dielectric properties of BaTiO3 ceramics co-doped with double rare earths.
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Another type of double-site substitution by double rare earths is called the self-compensation mode of R •Ba − R 'Ti (where R is a rare earth element) such as Eu •Ba − Eu 'Ti [12] and Dy•Ba − Dy'Ti [13].
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However, when Ba/Ti > 1 Eu can occupy Ba-sites predominantly in the mixed-valence states of Eu3+ (4f6)/Eu2+ (4f7). An interesting phenomenon is that Eu ions can exist stably as Eu3+ in Eu-doped BaTiO3
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with Ba/Ti = 1, where no valence change occurs [12]. In 2015, Tb ions were proven by EPR to be able to exist at the Ti-site, in part as Tb4+ (4f7); and in the case of Ba/Ti ≠ 1, Tb ions exhibit self-adjustable site occupations between the Ba-site as Tb3+ and the Ti-site as Tb4+ [14]. •
In this work, (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) ceramics in the self-compensation mode of La Ba − Tb Ti '
were prepared using a conventional mixed oxide method. The dielectric properties and point defects of BLTT were investigated. La and Tb were selected to be co-doped into BaTiO3, based on the following four considerations: (1) La donors play a key role in improving the dielectric permittivity of BaTiO3 ceramics [2] 2
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and Tb doping in BaTiO3 can result in low-temperature dielectric stability [14]. BaTiO3 ceramics co-doped with La and Tb were therefore expected to achieve high-k stability; (2) the exclusive occupations of larger
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La3+ ions at Ba-sites can expel smaller Tb ions into Ti-sites to form La •Ba − Tb 'Ti complexes, which is beneficial for the formation of single-phase ceramics; (3) Tb3+ (4f8) ions at Ti-sites have the ability to trap electrons in ceramics can thus lower dielectric loss by means of the transfer of Tb3+ to metastable Tb4+ (4f7);
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and (4) EPR of Tb4+ (4f7) [14] can be used to probe the valence stability of Tb3+ (4f8) at Ti-sites. The high compatibility of La and Tb in the BaTiO3 lattice [1–3,14] and the optimization of dielectric performance are
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common characteristics of these two elements. The self-adjustability of site occupations by Tb ions [14] was used to incorporate extra Tb ions into the BLTT lattice, to obtain highly dense ceramics. The dielectric properties of these ceramics were investigated.
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2. Experimental procedures
Ceramic raw materials were reagent-grade BaCO3, La2O3, TiO2, and Tb4O7 powders. The ceramics were prepared according to the nominal compositions (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03, 0.05, 0.07, 0.10, 0.15,
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0.20) (BLTT) and (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0, 1, 3, 5, 7) (BL3TT-nT) using a mixed oxides method with the same conditions as described elsewhere [8]. The sinter temperature (Ts) is based on
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the following considerations: (1) sufficient incorporation of La ions into the BaTiO3 lattice, which can contribute to improvement of the dielectric permittivity, needs a higher Ts and longer sintering time, for example, Ts = 1350 °C for 3 days for La-doped BaTiO3 ceramics [2]; (2) our experiments showed that very high Ts (=1400 °C) for La-doped BaTiO3 ceramics results in the formation of the liquid phase. The addition of Tb ions in BaTiO3 can solidify ceramics and inhibit appearance of the liquid phase [14]; (3) For BLTT with x = 0.07, our experiments showed that the ceramic density was improved 20 % when Ts was increased
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from 1300 to 1400 °C, accompanied by a change in color from white to light yellow. Thus, the final sintering conditions for La and Tb co-doped BaTiO3 ceramics were set as 1400 °C for 12 h in air.
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Structural characterization of ceramics was performed by powder X-ray diffraction (XRD) using a DX-2700 X-ray diffractometer (Dandong Haoyuan) employing Cu Kα1 radiation (λ = 1.540562 Å). Lattice parameters and unit cell volume were calculated by a MS Modeling software package (Accelrys, Inc.)
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using Rietveld refinement in the Reflex Package. Raman spectra were measured using a LabRAM XploRA Raman spectrometer (Horiba Jobin Yvon) with a 532 nm laser. Microstructural properties were determined
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using an EVOMA 10 scanning electron microscope (SEM) (Zeiss) operated at 15 keV. The average grain size (GS) was obtained from measurements of the number of grains per unit area on a polished surface NA using Fullman’s method [15], GS =
6 . Electron paramagnetic resonance (EPR) spectra were measured
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at room temperature using an A300-10/12 X-band spectrometer (Bruker) operating at 9.84 GHz. X-ray
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photoelectron spectra (XPS) measurements were performed at room temperature using an ESCALAB 250 X-ray photoelectron spectrometer (Thermo Electron Co.). The XPS data raw were processed by smoothing multiply times. The dielectric and conducting properties of the bulk ceramics were investigated at 1 kHz
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using a Concept 41 Dielectric/Impedance spectrometer (Novocontrol) with an applied voltage of 1 V. The accuracy in the measurements of dielectric permittivity, dielectric loss, resistivity, and temperature control
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are less than 5 %, 3 × 10−5, 3 × 10−5, and ±0.3 °C, respectively. Apart from the above systematic errors of the dielectric/impedance spectrometer, there are two factors to affecting the accuracy in measurement of dielectric permittivity: (1) Processing technique of electrodes: Ceramics were polished into disks (Φ10.2 mm, 0.8 mm in thickness) for dielectric measurements. Squared electrodes of 5 mm × 5 mm were applied on both surfaces of the ceramic disks. To minimize the effect of the air layer between the ceramic disk surface and contact
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electrode and to increase the conductivity, thin layers of Au atoms were sputtered on the surfaces of the polished ceramic disks, and then, the silver paste was spread uniformly on the thin Au electrode. These
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ceramic disks were heat-treated at 200 °C to form compact contact electrodes with good electrical conductivity. For these treatments, the accuracy in the measurement of dielectric permittivity of ceramics can be further improved, as shown in Fig. 1. It can be seen that the measured permittivity of ceramics with
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Ag electrodes formed by the silver paste, is greatly decreased because the existence of the air layer results in decreased capacitance. The sputtered Au atoms assist in squeezing out the air layer. On the other hand,
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compared to the circular electrodes (85 mm2) completely covering both surfaces of the ceramic disks (Φ10.2 mm), the use of smaller square electrodes (5 mm × 5 mm) has little influence on the measured permittivity, as shown in Fig. 2. The change in maximum permittivity (ε'm) is less than 4 %. (2) The edge effect of a parallel-plate capacitor: Scott and Curtis emphasize the necessity for accurate
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determination of the capacitance C of a parallel-plate capacitor comprised of two like rectangular plates with vacuum (air) as dielectric [16,17]. For a capacitor including dielectrics, the capacitance C and increment (∆) in
C can be expressed as follows:
d
+
2ε r ε 0 L πL ln( + 1) + 1 π d
2d πL ln( + 1) + 1 πL d
(1)
(2)
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∆=
ε r ε 0 L2
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C=
where L is the side length of square electrodes, εr the relative permittivity, ε0 the permittivity of free space, and d the separation between the plates. After the edge effect is considered, the measured permittivity should be reduced by 2.5 % compared to the actual permittivity. Thus, the uncertainty in the measurement of dielectric permittivity is less than 6.5 % after the effect of electrode shape and systematic errors of the dielectric/impedance spectrometer are included.
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3. Results
Powder XRD patterns of (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) ceramics (BLTT) are shown in Fig. 3.
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BLTT with x ≤ 0.15 exhibited a single-phase perovskite structure (Fig. 3(a)), indicating that La and Tb ions were completely incorporated into the BaTiO3 lattice. A small amount of a secondary phase of Tb4O7 was observed for BLTT with x = 0.20 whereas the main (1 0 0) peak of Tb4O7 did not appear at x = 0.15. (Fig. 3
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inset). The solid solution limit of Tb in BLTT was therefore determined to be greater than x = 0.15 by XRD. The x = 0.03 sample had a tetragonal structure, characterized by two separate (0 0 2) and (2 0 0) peaks,
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whereas for when x ≥ 0.05, BLTT converted to a cubic structure, as marked by a symmetric (2 0 0) peak (Fig. 3(b)).
Powder XRD patterns of (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0 – 7) ceramics (BL3TT-nT) are shown in Fig. 4. BL3TT-nT with n = 0 is the same sample as the above BLTT with x = 0.03 (BL3TT).
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There was no indication of any secondary phase separating out until n = 3. The solid solution limit of the extra Tb ions in BL3TT was determined to be n = 3, i.e., 3 % La and 12 % Tb ions can be incorporated
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simultaneously into the BaTiO3 lattice. Tb4O7 appeared as a secondary phase in the main perovskite phase for n ≥ 5. BL3TT-nT was tetragonal for n = 0 and 1 but converted to a cubic structure for n = 3 (Fig. 4(b)).
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The variations in lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for BLTT and BL3TT-nT are shown in Fig. 5. Both V0s of BLTT and BL3TT-nT are greater than the V0 of the tetragonal BaTiO3 (JCPDS Cards No. 5-626), and increase as x or n increases. In particular, V0 of BLTT with x = 0.15 is greater than that of the cubic BaTiO3 (JCPDS Cards No. 31-174). These results indicate the partial Ti-site occupation of Tb ions [14]. The Raman spectra of BLTT and BL3TT-nT at room temperature are shown in Fig. 6. The Raman spectrum of the cubic BaTiO3 generally exhibits three typical bands peaking at 265, 520, and 720 cm-1 6
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because of the intrinsic disorder in the cubic phase [18–20]. These bands also appeared in BLTT and BL3TT-nT. An additional band appeared and its peak position red-shifted from 818 to 805 (or 808) cm-1
number of Tb3+ ions incorporated into the BaTiO3 lattice [14,21].
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with increasing x or n. This band originates from the Raman charge effect [21], indicating an increase in the
The SEM images of BLTT and BL3TT-nT are shown in Fig. 7. The variations in average grain size (GS)
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as a function of x for BLTT and as a function of n for BL3TT-nT are shown in Fig. 8. BLTT exhibited a fine-grained microstructure and a uniform grain size distribution. The GS of BLTT increased from 0.8 to
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1.8 µm with increasing x. The relative density (ρr) here is defined as the ratio of the volumetric mass density to the theoretical density. The variations in ρr as a function of x for BLTT and as a function of n for BL3TT-nT are shown in Fig. 9. BL3TT-nT exhibited a lower porosity and a larger GS with respect to BLTT (Figs. 7–9). This suggests that Tb ions are not good at inhibiting grain-growth because GS = 3.6 µm for
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BL3TT-nT with n = 3 (Fig. 8). The ρr of BL3TT-nT was 96 % when n ≥ 1, whereas BLTT was less than 84 % when x ≤ 0.10 (Fig. 9). This result revealed that the incorporation of the extra Tb ions in BL3TT can greatly enhance the ceramic density.
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The temperature dependences of dielectric permittivity (ε'), loss (tan δ), and resistivity (ρ) for BLTT and BL3TT-nT are shown in Fig. 10. The dielectric peak of BLTT gradually broadened with increasing x and
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shifted toward lower temperatures. A weaker dielectric peak in the vicinity of 30 °C for the x = 0.03 BLTT sample (Fig. 10(a)) corresponded to the orthorhombic-tetragonal (t-c) phase transition (T2) of BaTiO3 [22]. The dielectric-peak temperature (Tm), representing the temperature at which maximum permittivity (ε'm) occurred, decreased linearly at a rate of −19 °C/%(La-Tb) (Fig. 10 inset). BLTT with x = 0.07 and BL3TT-nT with n = 3 satisfied the X7U dielectric specification, i.e., −56 % ≤ |ε'−ε'RT| ≤ +22 % in the range of −55 to 125 °C (Figs. 10(a) and (d)). BLTT and BL3TT-nT with n ≤ 1 exhibited a lower tan δ (≤ 0.034)
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below 75 °C (Figs. 10(b) and (e)) and their resistivity (ρ) was 107 − 108 Ω cm (Figs. 10(c) and (f)). A higher room-temperature value of tan δ = 0.058 and a very low value of ρ = 5 × 106 Ω cm for BL3TT-nT
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with n = 3 relates to a higher concentration of the extra Tb ions. The addition of the extra Tb ions in BL3TT resulted in an increase in ε'.
The frequency (f) dependences of ε'RT and tan δ were measured for two X7U dielectrics, BLTT with x =
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0.07 and BL3TT-nT with n = 3, as shown in Fig. 11, where the inset depicts the ρ-f curve. The ε' of the two X7U dielectrics remained almost constant in the f range of 10 to105 Hz. BLTT with x = 0.07 exhibited a
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higher frequency stability compared with BL3TT-nT with n = 3 (Fig. 11(a)). Both tan δ and ρ were sensitive to f and decreased as f was increased(Fig. 11(b) and inset).
The room-temperature EPR spectra of BLTT and BL3TT-nT are shown in Fig. 12. A broad signal with g ~6.5 can be attributed exclusively to Tb4+ (4f7) Kramers ions at Ti-sites [14]. A weaker Tb4+-related signal
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could be observed for BLTT and BL3TT-nT. To qualitatively understand the extent of occupation of Ti-ites by Tb4+ , we measured an EPR spectrum of the precious sample of Ba(Ti1−xTbx)O3 with x = 0.05 (BTT5) [14] at room temperature, and compared this spectrum with the EPR spectrum of (Ba1−xLax)(Ti1−xTbx)O3
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(BLTT) with x = 0.05, as seen in Fig. 13. Tb3+ Kramers ions are EPR-silent in theory. Tb ions in BTT5 were found to exist predominantly at Ti-sites as Tb4+, and only a small number of Tb3+ ions substituted at
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Ba-sites [14], assuming that the intensity of the g
~6.5 signal is proportional to the concentration of
Ti-site Tb4+ ions. If all of the Ti-site Tb ions were Tb3+, the Tb4+-related signal would disappear. It can be seen from Fig. 13 that the Tb4+-related signal with g ~6.5 was very weak for the x = 0.05 BLTT sample,as ×
compared to BTT5. This comparison reveals two results: (1) the number of Ti-site Tb4+ ions ( Tb Ti ) in BLTT is smaller; and (2) Tb ions in BLTT exist predominantly as Ti-site Tb3+ ( Tb 'Ti ) because to a certain extent the ratio of La/Tb = 1 can limit the transfer of Tb ions from Ti-sites to Ba-sites. The comparison in
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Fig. 13 further clarifies that La •Ba − Tb 'Ti defect complexes in the BLTT lattice cannot be completely •
×
formed and a small number of La Ba − Tb Ti complexes may coexist. For this reason and owing to the
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amphoteric behavior of Tb ions, a small number of Tb ions inevitably entered Ba-sites as Tb3+ ( Tb •Ba ) to preserve the lattice electroneutrality [14]. The presence of the g = 2.004 signal associated with Ti vacancies [18,23,24] (Fig. 12) in all of the samples further verifies the occupation of Ba-sites by Tb3+ because of the
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relation 4Tb •Ba → VTi'''' .
The signal at g = 1.974, which appeared only in BLTT with x = 0.03, was attributed to ionized
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'' ) [13,23,25]. This indicates that intrinsic Ba as well as Ti vacancies may coexist in Ba-vacancy defects ( VBa
the main perovskite lattice.
Mn impurity usually exists in the starting materials, as detected in Eu-doped BaTiO3 ceramics [12]. The appearance of Mn2+ sextet signal in BLTT with x = 0.03 in Fig. 12 implies the reduction of Mn4+ or Mn3+ in
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ceramics to Mn2+. This phenomenon was also observed for (Ba1−xLax)(Ti1−xDyx)O3 with x = 0.03, after ''
which Mn2+ signal disappeared with increasing x [11]. An interesting phenomenon is that VBa -related signal is accompanied by the appearance of the Mn2+ sextet signal, and these signals disappear
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simultaneously when x > 0.03 or n > 1. This fact suggests that as x is increased, Tb ions may transfer from Ti-ites to Ba-ites in BLTT to fill in Ba vacancies.
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The mixed-valence states of Tb4+/Tb3+ in BLTT and BL3TT-nT were further confirmed by XPS. The Tb 4d XPS spectra of BLTT with x = 0.10 and 0.15 and BL3TT-nTb with n = 1 and 3 were measured and smoothed, as shown in Fig. 14. The Tb 4d core-level lines in compounds generally show multiplet structures, which are similar to the main and satellite lines in pure metals [26,27]. The peaks with several higher core-level binding energies originate from Tb4+ 4d core-level lines [26], and in particular, a peak at ~155 eV can be considered as an indicator of Tb4+ [26–28]. The peaks at ~153, ~150, and ~148 eV can be
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used to determine whether a compound has Tb3+ ions or not [26–30]. On the basis of these investigations, three peaks at 155.9, 153.2 and 150.7 eV in Fig. 14 were attributed to contributions from multiplet
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structures of Tb4+ in BLTT, while the three peaks at 153.2, 150.7 and 147.8 eV in Fig. 14 were attributed to multiplet structures of Tb3+. The two intermediate peaks at 153.2 and 150.7 eV cannot be unambiguously distinguished because each peak is a superposition of Tb4+ and Tb3+ lines. The XPS configuration of BL3TT-nT is similar to that of BLTT. The extra Tb ions in BL3TT-nT mainly mainly form •
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•
×
self-compensated Tb Ba − Tb Ti complexes except for a small number of Tb Ba − Tb Ti complexes [31], '
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as seen in Fig. 15 by the Tb4+ EPR comparison between BL3TT-nT (n = 3) and (Ba1−xTbx)(Ti1−xTbx)O3 (x = 0.05) with a nominal self-compensation mode.
The Ti 2p XPS spectra of BLTT and BL3TT-nTb are shown in Fig. 16. Ti ions in BLTT and BL3TT-nTb were confirmed to maintain a stable oxidation state of 4+ based on the following four reasons: (1) the two
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prominent peaks were identified as Ti 2p3/2 at ~458.1 eV and Ti 2p1/2 at ~463.9 eV, respectively, and coincide with the reported values of Ti 2p XPS for BaTiO3 and SrTiO3 [32] with the Ti4+-ion profile; (2) the difference between Ti 2p3/2 and Ti 2p1/2 orbit-spin splitting binding energies in BLTT and BL3TT-nTb (5.8
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eV) is the nearly same as that of BaTiO3 (5.6 eV) [33]; (3) no peak related to Ti3+ 2p is evident [34]; and (4) the ∆FWHM (full width at half maximum) value (1.7 - 1.8 eV) of Ti 2p3/2 main peak is less than that of
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BaTiO3 (< 2.0 eV) [35], which also suggests that Ti3+ or mixed-valence Ti4+/Ti3+ do not exist in BLTT and BL3TT-nTb.
4. Discussion
4.1. Defect chemistry
For the single-phase (Ba1−xLax)(Ti1−xTbx)O3 (0.03 ≤ x ≤0.15) ceramics (BLTT), the XRD results reveal
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the predominant occupation of Tb ions at Ti-sites (Fig. 5). The above EPR and XPS results indicate that the •
•
''
point defects in BLTT with x = 0.03 are Ba-site La3+ ( La Ba ), Ba-site Tb3+ ( TbBa ), Ba vacancies ( VBa ),
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Ti-site Tb3+ ( Tb'Ti ), Ti-site Tb4+ ( Tb×Ti ), and Ti vacancies ( VTi'''' ). Mn impurities in ceramics were neglected here because of very low concentrations. At a lower doping level of x = 0.03, both Ba and Ti vacancies may coexist in the main perovskite lattice (Fig. 12).
•
''
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To a certain extent Tb ions can transfer from Ti-sites to Ba-sites during sintering and upon cooling. When x ≥ 0.05, Ba vacancies are filled up by Tb Ba and the VBa -related EPR signal disappears (Fig. 12). Tb
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ions in BLTT exist predominantly as Tb 'Ti (Fig. 13), forming La •Ba − Tb 'Ti defect complexes because of the exclusive site occupation of La3+ at Ba-sites.
2La2O3 + Tb 4O7 → 4La •Ba + 4Tb'Ti + 12OO + 1/2O2 ↑
(3)
A small number of Tb ions may exist at Ba-sites in the form of Tb3+ ( Tb •Ba ) because of the amphoteric ×
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behavior of Tb ions [14,31]. On the other hand, some Ti-site Tb ions exist as Tb4+ ( TbTi ), implying that half of the Tb4+ ions in the starting material Tb4O7 cannot be reduced completely into Tb3+ ions. However, the number of Tb×Ti is very small as shown by the EPR comparison (Fig. 13) and XPS observations (Fig. •
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14). The appearance of a small number of Tb Ba and Tb×Ti may expressed by
2Tb4O7 + BaO → 4Tb•Ba + Ba Ba + 4Tb×Ti + VTi'''' + 15OO •
(4)
''
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The incorporation of Tb Ba into Ba-sites through VBa results in an excess of Ba-site cations and •
•
further in the steady existence of VTi'''' (Fig. 12), caused by the extra La Ba and Tb Ba ,in all of the •
samples except for La Ba − Tb Ti defect complexes. '
4La•Ba → VTi''''
(5)
4Tb•Ba → VTi''''
(6)
•
Thus, La Ba − Tb Ti defect complexes cannot be completely formed in BLTT and a small number of '
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Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes may coexist in the BaTiO3 lattice. The defect chemistry of the single-phase (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0, 1, 3) ceramics
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(BL3TT-nT) is similar to that of BLTT. The extra Tb ions in BL3TT-nT exist as Ba-site Tb3+ ( Tb•Ba ), Ti-site Tb3+ ( Tb'Ti ), and Ti-site Tb4+ ( Tb×Ti ) [31]. This behavior is similar to the predominant •
•
self-compensation mode of Tb ions in BaTiO3 [31]. A large number of La Ba − Tb Ti and Tb Ba − Tb Ti '
'
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defect complexes and a small number of La •Ba − Tb×Ti and Tb •Ba − Tb×Ti defect complexes may coexist in BL3TT-nT.
4.2.1. Linear shift of Tm with x for BLTT
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4.2. Dielectric properties
•
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The above analyses of the defect chemistry of BLTT indicate that La Ba − Tb Ti defect complexes are '
predominant, and the number of Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes is very small. The •
•
shift of Tm with x is therefore mainly determined by La Ba − Tb Ti defect complexes. Tb Ba − Tb Ti and '
'
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La •Ba − Tb×Ti complexes have a minor effect on a shift in Tm. The Tm decreases linearly at a rate of
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−19 °C/%(La-Tb) for BLTT (Fig. 10(b) inset). This is mainly attributed to a linear increase in the quantity •
of La Ba − Tb Ti defect complexes with an increase in x. '
For BL3TT-nT, the Tm decreases nonlinearly with n (Fig. 10(e) inset). Similar to the analyses on BLTT, •
•
the shift of Tm with n is mainly determined by the predominant La Ba − Tb Ti and Tb Ba − Tb Ti defect '
'
complexes. The concentration of La ions is fixed at 3 %, whereas the number of Tb ions increases with n. •
•
Thus, the nonlinear shift in Tm is attributed to the different dielectric responses of La Ba and Tb Ba .
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4.2.2. Lower dielectric loss of BLTT
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All of the BLTT samples show lower tan δ below 75 °C (Fig. 10(e)). Because the Mn2+ EPR signal appears only in BLTT with x = 0.03, it is concluded that the reduction of Mn4+ or Mn3+ to Mn2+ impurities •
in ceramics is not the main reason for lowering of tan δ. The predominant La Ba − Tb Ti defect complexes '
and the small number of Tb •Ba − Tb 'Ti and La •Ba − Tb×Ti defect complexes in BLTT can suppress the •
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•
donor effect of La Ba and Tb Ba and the acceptor effect of Tb 'Ti . On the other hand, the lower tan δ is
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probably due to the existence of a small number of metastable Tb×Ti ions. To a certain extent, each Ti-site Tb4+ ion in the ceramic has some ability to trap an electron and can thus be reduced to Tb 'Ti , which •
concatenates with Tb Ba to form the self-compensation mode of Tb•Ti − Tb'Ti .
Tb×Ti + e′ → Tb'Ti
(7)
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The above two reasons can contribute to the reduction in dielectric loss of BLTT.
4.2.3. Comparison between two X7U dielectrics and Effect of the extra Tb ions on dielectric properties of
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BLTT
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Both BLTT with x = 0.07 (BLTT7) and BL3TT-nT with n = 3 (BL3TT-3T) satisfy the X7U dielectric specifications. Compared with BLTT7, BL3TT-3T exhibits a higher ε'RT (= 5990) and tan δ (= 0.058) (Figs. 10(b) and (e)) at room temperature. The SEM results show that the relative density (ρr) of BL3TT-3T is •
96 %, whereas that of BLTT7 is 84 % (Figs. 7 and 8). Because La Ba plays a significant role in improving ε' among rare-earth donors [2], the improvement in ε'RT caused by the addition of the extra Tb ions in BLTT •
cannot be attributed to the Tb Ba − Tb Ti defect complexes, but is most likely related with the rapid '
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improvement in ceramic density of BL3TT. At a higher Tb doping level of n = 3, the higher tan δ of •
•
BL3TT-3T implies that the donor effect of Tb Ba is greater than that La Ba . Due to its lower loss (tan δ =
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0.029), the X7U BLTT7 is a promising dielectric for X7U applications
5. Conclusions
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(Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) and (Ba0.97La0.03)(Ti0.97Tb0.03)O3–0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics were prepared using a mixed oxides method. BLTT with 0.03 ≤ x ≤0.15 and
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BL3TT-nT with 0 ≤ n ≤ 3 exhibit a single-phase perovskite structure. The unit cell volume increases with increasing x or n. The XRD and Raman results provide evidence of the predominant occupation of Tb ions at Ti-sites. The incorporation of the extra Tb ions in BLTT with x = 0.03 can greatly enhance the ceramic density of BLTT, showing an amazing relative density of 96 %. A study of the defect chemistry indicates •
•
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that the point defects in BLTT with x ≥ 0.05 are Ba-site La3+ ( La Ba ), Ba-site Tb3+ ( Tb Ba ), Ti-site Tb3+ •
( Tb'Ti ), Ti-site Tb4+ ( Tb×Ti ), and Ti vacancies ( VTi'''' ). A large number of La Ba − Tb Ti defect complexes •
×
•
'
and a small number of La Ba − Tb Ti and Tb Ba − Tb Ti defect complexes may coexist in BLTT. The Tm '
•
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decreases linearly at a rate of −19 °C/%(La-Tb) for BLTT. The predominant formation of La Ba − Tb Ti '
defect complexes is responsible for the linearly shifting Tm rate. The existence of a small number of •
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metastable Tb×Ti ions and the predominant presence of La Ba − Tb Ti defect complexes are both '
responsible for the lower dielectric loss of BLTT. BLTT with x = 0.07 satisfies X7U dielectric specifications. BLTT with x = 0.07 because of its higher room-temperature permittivity (ε′ = 2450) and lower loss (tan δ = 0.029), can be considered as a promising dielectric for X7U applications.
Acknowledgements
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This work was supported by the projects of the National Natural Science Foundations of China (Grant No. 21271084) and of Jilin Province (Grant No. 20160101290JC).
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References
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Figure captions
Fig. 1. Temperature dependence of the dielectric permittivity (ε') for a BaTiO3-based sample A with (a) sputtered Au/Ag electrodes and with (b) Ag electrodes, measured at 1 kHz using an RCL meter (Fluke PM6306).
Fig. 2. Temperature dependence of ε' for a BaTiO3-based sample B with (a) square electrodes of 5 mm × 5
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mm and with (b) circular electrodes of 85 mm2, measured at 1 kHz using an RCL meter (Fluke PM6306). Fig. 3. XRD patterns of (a) (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BLTT.
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Fig. 4. XRD patterns of (a) (Ba0.97La0.03)(Ti0.97Tb0.03)O3 – 0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BL3TT-nT.
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Fig. 5. Variations in the lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for (a) BLTT and (b) BL3TT-nT ceramic powders. The V0s of the tetragonal BaTiO3 (t-BaTiO3, JCPDS Cards No. 5-626) and the cubic BaTiO3 (c-BaTiO3, JCPDS Cards No. 31-174) are drawn as a dashed line for comparison. Fig. 6. Room-temperature Raman spectra of (a) BLTT and (b) BL3TT-nT with a single-phase structure. Fig. 7. SEM images of thermally etched surface of BLTT with (a) x = 0.03, (b) x = 0.05, (c) x = 0.07, (d) x = 0.10; and of BL3TT-nT with (e) n = 1, (f) n = 3. Fig. 8. Variations in average grain size (GS) as a function of x for BLTT and as a function of n for
18
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BL3TT-nT. Fig. 9. Variations in relative density (ρr) as a function of x for BLTT and as a function of n for BL3TT-nT. Fig. 10. Temperature dependences of (a) and (d) dielectric permittivity (ε'), (b) and (e) dielectric loss (tan
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δ), and (c) and (f) resistivity (ρ) for BLTT and BL3TT-nT, measured at 1 kHz.
Fig. 11. Frequency dependences of (a) ε' and (b) tan δ for the x = 0.07 BLTT and the n = 3 BL3TT-nT ceramics, measured at 25 °C from f = 10 to105 Hz. The inset depicts the ρ-f curve.
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Fig. 12. Room-temperature EPR spectra of (a) BLTT and (b) BL3TT-nT.
Fig. 13. Room-temperature EPR comparison between (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) and Ba(Ti1−xTbx)O3
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[14] with x = 0.05.
Fig. 14. Tb 4d XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3. Fig. 15. Room-temperature EPR comparison between BL3TT-nT with n = 3 and (Ba1−xTbx)(Ti1−xTbx)O3 with x = 0.05 [28].
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Fig. 16. Ti 2p XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
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Au + Ag electrodes
1.0
Ag electrodes
0.5
0.0
-50
0
50
100 o
T ( C)
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ε' (× 104)
BaTiO3-based ceramic A
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1.5
150
200
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Fig. 1. Temperature dependence of the dielectric permittivity (ε') for a BaTiO3-based sample A with (a) sputtered Au/Ag electrodes and with (b) Ag electrodes, measured at 1 kHz using an RCL meter (Fluke PM6306).
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2
1.0
A = 25 mm
A = 85 mm2
0.5
0.0
-50
0
50 100 o T ( C)
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4
ε' (× 10 )
BaTiO3-based ceramic B
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1.5
150
200
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Fig. 2. Temperature dependence of ε' for a BaTiO3-based sample B with (a) square electrodes of 5 mm × 5 mm and with (b) circular electrodes of 85 mm2, measured at 1 kHz using an RCL meter (Fluke PM6306).
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(211)
(200)
(200)
32
x = 0.15 (200)
30
(220) (300) (310) (311) (222)
x = 0.20
*
x = 0.10 (200)
(200)
x = 0.07
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x = 0.10
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x = 0.15
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(110)
(200)
x = 0.20
2θ (o)
(210)
(111)
(100)
(100) from Tb 4O7
x = 0.15
28
(110)
Tb4O7 x = 0.20
Intensity (arb. units)
(b)
(Ba1-x Lax)(Ti1-x Tbx)O3 (100)
(110)
(a)
x = 0.07
x = 0.05
20
40
60
2θ (o)
80
(002) (200)
(113)/(311) (222)
(212) (103)/(310)
x = 0.03 (202)
(211)
(201)
(200)/(002)
(111)
(100)
(101)/(110)
x = 0.05
x = 0.03
44
45
46
2θ (o)
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Fig. 3. XRD patterns of (a) (Ba1−xLax)(Ti1−xTbx)O3 (x = 0.03 – 0.20) (BLTT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BLTT.
(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3
(110)
(200)
(211) *
*
*
*
*
(311)
*
n=5
(220) (300) (310) (311) (222)
(200) (210)
(220)
(110)
n=7
(200)
(100)
(100)
n=5
n=3
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*
n=3
(002) (200)
Intensity (arb. units)
T b4O7
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n=7 *
*
(b)
- 0.03nTb
(200)
(a)
(200)
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n=1
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(002) (200)
n=1
n=0
n=0
0 20
40
60
2θ (o)
80
44
45
46
2θ (o)
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Fig. 4. XRD patterns of (a) (Ba0.97La0.03)(Ti0.97Tb0.03)O3 – 0.03nTb (n = 0 – 7) (BL3TT-nT) ceramics. (b) Gaussian-fitting to the XRD peaks in the vicinity of 45° for BL3TT-nT.
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67.2
4.08
(a)
4.04
o
64.8
4.02 o
3
t-BaT iO3 ( V0 = 64.41 A )
64.0 c
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4.00
a 63.2 0.02
0.04
0.06
0.08
0.10
0.12
0.14
x (Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb o
4.12
3
c-BaT iO3 ( V0 = 65.50 A )
65.6
V 0 (A3)
3.98 0.16
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(b)
V0
4.08
o
64.8
o
3
t-BaT iO3 ( V0 = 64.41 A )
64.0
o
a, c (A)
V 0 (A3)
o
c-BaT iO3 (V0 = 65.50 A3)
65.6
a, c (A)
4.06
V0
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(Ba1-xLa x)(Ti 1-xTb x)O3
66.4
o
4.04
c
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a
63.2
0
2
4
6
4.00 8
n
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Fig. 5. Variations in the lattice parameters (a, c) and unit-cell volume (V0), as a function of x, for (a) BLTT and (b) BL3TT-nT ceramic powders. The V0s of the tetragonal BaTiO3 (t-BaTiO3, JCPDS Cards No. 5-626) and the cubic BaTiO3 (c-BaTiO3, JCPDS Cards No. 31-174) are drawn as a dashed line for comparison.
808 811
720
520
n=3
818
n=1
200
400
600
(a)
1000
1200
805
720
520
805
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(Ba1-xLax)(Ti 1-xTb x)O3 265
Intensity (a. u.)
800
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n=0
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(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb
Intensity (a. u.)
(b)
265
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818
805
x = 0.15
200
400
600
x = 0.10 x = 0.07 x = 0.05
800
1000
1200
Raman shift(cm -1)
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Fig. 6. Room-temperature Raman spectra of (a) BLTT and (b) BL3TT-nT with a single-phase structure.
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(b)
(c)
(d)
(f)
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(a)
Fig. 7. SEM images of thermally etched surface of BLTT with (a) x = 0.03, (b) x = 0.05, (c) x = 0.07, (d) x = 0.10; and of BL3TT-nT with (e) n = 1, (f) n = 3.
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n 0
1
2
3
4
BL3TT-nT 2
1
0 0.05
0.1
x
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BLTT
0.15
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GS ( µ m)
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Fig. 8. Variations in average grain size (GS) as a function of x for BLTT and as a function of n for BL3TT-nT.
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2
3
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100
ρr (%)
BL3TT-nT 90
80
70 0.05
0.10
x
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BLTT
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Fig. 9. Variations in relative density (ρr) as a function of x for BLTT and as a function of n for BL3TT-nT.
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8000
8000
n =1
4000
2000
2000
0 0.30
0 0.30 120
(e) -19 C/%(La-Tb)
0.20 0
tanδ
-60 0.02
0.10
0.04
0.06
0.08
80
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60
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40
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n =3
o
o
T m ( C)
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0
0.10
x
0.00 2
0
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2
3
n
0.00 2
(f)
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(c) ρ (×108 Ωcm)
n =0
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4000
(Ba0.97La 0.03)(Ti 0.97Tb 0.03)O3 - 0.03nTb
6000
ρ (×10 Ωcm)
ε′
6000
= 0.03 = 0.05 = 0.07 = 0.10
ε′
x x x x
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(d)
(Ba1-xLa x)(Ti 1-xTb x)O3
T m ( C)
(a)
1
0 -50
0
50
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T (oC)
100
150
200
1
0 -50
0
50
100
150
200
T (oC)
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Fig. 10. Temperature dependences of (a) and (d) dielectric permittivity (ε'), (b) and (e) dielectric loss (tan δ), and (c) and (f) resistivity (ρ) for BLTT and BL3TT-nT, measured at 1 kHz.
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10000
(a)
o
T = 25 C
BLTT with x = 0.07
8000
BL3TT with n = 3
ε′
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103 8 6 4 2 0 101
102
104
103
105
104
105
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8
(b)
ρ (×10 Ωcm)
102
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0 101 0.20
f (Hz)
0.05
0.00 101
102
103
104
105
f (Hz)
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Fig. 11. Frequency dependences of (a) ε' and (b) tan δ for the x = 0.07 BLTT and the n = 3 BL3TT-nT ceramics, measured at 25 °C from f = 10 to105 Hz. The inset depicts the ρ-f curve.
(b)
4+
TbT i
g = 2.004 (VT i)
Intensity
n=3
n=1
Mn
(a)
2000 TbT4+i
3000
g = 2.004 (VT i)
x = 0.15
Intensity
4000
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(VBa) g = 1.974
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2000
3000
4000
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Fig. 12. Room-temperature EPR spectra of (a) BLTT and (b) BL3TT-nT.
TbTi4+
Receiver Gain: 1.00 ×10
4
Intensity (a. u)
g = 6.58
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Ba(Ti 1-xTb x)O3 with x = 0.05 [14]
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VT i g = 2.004
1000
2000
3000
4000
5000
6000
H (G)
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Fig. 13. Room-temperature EPR comparison between (Ba1−xLax)(Ti1−xTbx)O3 (BLTT) and Ba(Ti1−xTbx)O3 [14] with x = 0.05.
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Fig. 14. Tb 4d XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
4+
Intensity (a. u)
TbTi
Receiver Gain: 1.00 ×10
4
(Ba 1-xTbx)(Ti 1-xTbx)O3 with x = 0.05 [28]
(Ba 0.97La 0.03)(Ti 0.97Tb0.03)O3-0.03nTb
VT i g = 2.004
1000
2000
3000
4000
5000
6000
M AN U
H (G)
SC
with n = 3
RI PT
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AC C
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TE D
Fig. 15. Room-temperature EPR comparison between BL3TT-nT with n = 3 and (Ba1−xTbx)(Ti1−xTbx)O3 with x = 0.05 [28].
5.8 eV
n=3
Ti 2p3/2 = 1.8eV
Ti 2p1/2 464.0
FWHM
n=1
5.8 eV
463.9
FWHM = 1.7 eV
x = 0.15
M AN U Ti 2p3/2 458.1
Ti 2p1/2
FWHM = 1.7eV
463.9
Intensity (a. u.)
(a) BLTT
465
SC
460 458.1
455
RI PT
= 1.8eV
464.0
FWHM
458.2
Intensity (a. u.)
(b) BL3TT- nT
458.2
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x = 0.10
455
460
465
Binding energy (eV)
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Fig. 16. Ti 2p XPS spectra of (a) BLTT with x = 0.10 and 0.15, and (b) BL3TT-nTb with n = 1 and 3.
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Highlights
► Dielectric properties and defect chemistry of (Ba1−xLax)(Ti1−xTbx)O3 were studied.
RI PT
► Solid solution limit of Tb in BLTT is determined by XRD to be greater than x = 0.15. ► Tb ions coexist in the mixed valence states of Ba-site Tb3+ and Ti-site Tb3+/Tb4+.
► The dielectric-peak temperature (Tm) decreased linearly at a rate of −19 °C/%(La-Tb).
AC C
EP
TE D
M AN U
SC
► The x = 0.07 sample is a promising dielectric for X7U applications.