Site-occupancy of yttrium as a dopant in BaO-excess BaTiO3

Site-occupancy of yttrium as a dopant in BaO-excess BaTiO3

Materials Science and Engineering A335 (2002) 101– 108 www.elsevier.com/locate/msea Site-occupancy of yttrium as a dopant in BaO-excess BaTiO3 Ming-H...

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Materials Science and Engineering A335 (2002) 101– 108 www.elsevier.com/locate/msea

Site-occupancy of yttrium as a dopant in BaO-excess BaTiO3 Ming-Hong Lin a,*, Hong-Yang Lu b a

Department of Mechanical Engineering, National Kaohsiung Uni6ersity of Applied Sciences, Kaohsiung 80782, Taiwan, ROC b Institute of Materials Science and Engineering, National Sun Yat-Sen Uni6ersity, Kaohsiung 80424, Taiwan, ROC Received 31 July 2001; received in revised form 26 October 2001

Abstract The site-occupancy of yttrium in BaO-excess barium titanate (BaTiO3) has been studied following pressureless-sintering of Y2O3-doped compositions with the fast-firing technique in a conventional furnace. Grain growth inhibition to an average grain size (Gav) of :3 mm occurs at 0.30 mol% Y2O3 when the ceramic is sintered at temperatures below 1350 °C. Sintering at 1450 °C, however, requires 0.50 mol% Y2O3 for grain growth inhibition to become effective. The concentration difference of 0.30–0.50 mol% for grain growth inhibition is attributed to the temperature dependence of the Y2O3-doping level at which the transition from predominantly electron to vacancy compensation takes place. Y3 + may act as a donor or an acceptor when substituting, respectively for the Ba2 + - and Ti4 + -site. This conclusion is supported by conductivity measurements for fast-cooled samples at room temperature and by the grain growth anomaly observed during sintering. © 2002 Published by Elsevier Science B.V. Keywords: BaTiO3; Sintering; Donor; Acceptor; Grain growth

1. Introduction Barium titanate (BaTiO3) compositions doped with donors and acceptors have been investigated for defect reactions [1–3] when samples are either hot-pressed [4,5], or pressureless-sintered [6,7]. For doping with trivalent donor cations, commonly La3 + , Y3 + , Sb3 + , Nd3 + , and Sm3 + , substituting for the Ba2 + -site, the defect reactions have been proposed by Chan et al. [3]. The donor-doped BaTiO3 retains overall charge neutrality when the induced charge imbalance is compensated by a mixture of electrons and of cation vacancies [3,7 –9]. It usually requires donors, i.e. La3 + of only small amount [7,10] to make the insulator ceramic become semiconducting. Grain growth inhibition in donor-doped ceramics occurring at the higher donoroxide concentration of 0.20 – 0.30 mol% [5,7,10] coincides [7,10] with the loss of semiconductivity. The coincidence had been attributed [7] to the change of the charge compensating defect from electrons to cation vacancies as the donor concentration increased in the vicinity of the GGIT (grain growth inhibition threshold) [5]. The GGIT values for donor-oxides, * Corresponding author. E-mail address: [email protected] (M.-H. Lin).

D2O3 (e.g. La2O3) and D2O5 (e.g. Nb2O5) reported in the literature are in the range of 0.10 –0.30 mol%; they are accumulated in Table 1. Nevertheless, furnace cooling [1,7,10] has often been adopted in pressureless-sintering; semiconductivity is then affected by the inward diffusion of barium vacancies from the grain surfaces during cooling [1,9]. Besides, whether Y3 + in BaTiO3 behaves as the other trivalent cations is not conclusive. When the trivalent cation (T) substitutes for the B4 + -site T%B, it produces acceptor behaviour [16 –18]. The principal charge compensating defect for maintaining overall charge neutrality in the acceptor-doped BaTiO3 compositions is the oxygen vacancy [2]. Therefore, BaTiO3 ceramic can either be an insulator or a semiconductor at room temperature depending upon the type and concentration of additives [1,7,10], cooling rate [9], and balance between acceptor or donor impurities [10]. Cation substitutions for either the Ba2 + - or Ti4 + -site of donor-oxides La2O3 and Y2O3 has been simulated by computer and the respective solid-solution energy calculated by Lewis and Catlow [8]. Similar energy values of substituting for the Ba2 + - and Ti4 + -sites were indicated when incorporating Y3 + into BaTiO3. Its solubility although small was dependent upon the extent of the non-stoichiometry in the BaTiO3. However, recent

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XRD results [19] showed that the substitution of yttrium for titanium (with a higher solubility limit of 6.10 mol%), becoming an acceptor, was more favourable than that for barium (with a smaller solubility limit of : 0.75 mol%). The incorporation of the trivalent rare-earth cations into BaTiO3 as a function of the (Ba/Ti)-ratio [16] was studied by measuring the conductivity under different oxygen partial pressures. It was indicated that larger ions occupying predominantly the Ba2 + -site gave the donor behaviour while smaller ions preferring the Ti4 + site exhibited the acceptor behaviour. The intermediatesize cations can substitute for either the Ba2 + - or the Ti4 + -site depending on the (Ba/Ti)-ratio of the initial powder. The Y3 + -cation in the sixfold coordination having an ionic size of :0.090 nm [20] appears to be an intermediate one when incorporated into the BaTiO3 lattices [18]. It acted as a donor in TiO2-excess compositions and as an acceptor in BaO-excess [17]. Xue et al. [21] observed grain growth inhibition and the conductivity anomaly with Y2O3-doping in BaTiO3 when the ceramics were either pressureless-sintered, or hotpressed in a graphite die. They reported that Y3 + can be forced into either the barium or the titanium sublattices depending on the (Ba/Ti)-ratio in the initial powder composition. It was further suggested that Y3 + behaved as a donor in the hypo-stoichiometric compoTable 1 Grain growth inhibition threshold for donor cations in sintered BaTiO3 compositions Dopants (D2O3 or D2O5)

Powders A/B ratio

GGIT (mol%)

Reference

Nb’Ti

0.994

0.16

[11]

Ta’Ti Bi’Ba

La’Ba

Ce’Ba

Nd’Ba Gd’Ba

Y’Ba

*, Not specified.

0.994

0.16

[11]

0.994 * *

0.14 0.30 0.30

[11] [12] [13]

0.994 * * *

0.13 0.17 0.16 0.15

[11] [13] [12] [10]

0.994 * *

0.11 0.20 0.14

[11] [13] [12]

0.994 *

0.15 0.11

[11] [12]

0.994 * * *

0.18 0.09 0.10 0.15

[11] [12] [13] [14]

0.994 0.990 0.990

0.28 0.30 0.20

[11] [11] [15]

sitions of Ba/Ti B 1 and an acceptor in the hyper-stoichiometric of Ba/Ti \ 1. We have investigated the electrical conductivity and coarsening of the BaO-excess hyper-stoichiometric BaTiO3 by pressureless-sintering of Y2O3-doped compositions adopting the fast-firing technique in a conventional furnace. The conductivity and average grain size (Gav) of the sintered ceramics are consistent with the Y3 + site-occupancy exhibiting both the donor and acceptor behaviour.

2. Experimental procedure Commercial BaO-excess (Ba/Ti =1.013) non-stoichiometric BaTiO3 powder (Ticon® HPB) supplied by TAM Ceramics (Ferro, Niagara Falls, NY) was used. The initial powder was manufactured by solid-state reaction followed by jet milling to obtain an average particle size of 1.30 mm. It contains impurities of SiO2 (B 0.03 wt.%), Al2O3 (B 0.04 wt.%) and Fe2O3 (B0.05 wt.%) as specified by the manufacturer. The Y3 + -dopant was added by dissolving Y(NO3)3·6H2O (reagent grade, Merck, Darmstadt, Germany) in deionized water using a magnetic stirrer. The initial BaTiO3 powder was blended and mixed with suitable amounts of dopant. It was then mixed with 1 wt.% PVB (polyvinyl butyral) binder in absolute alcohol before milling in a plastic jar with nylon-coated steel balls for 2 h. The mixed slurry was oven-dried, deagglomerated by using an agate mortar and pestle, and passed through  74 mm sieve. An appropriate amount of powder was die-pressed to discs of 10 mm in diameter in a WC-inserted steel die by applying a uniaxial pressure of 100 MPa. Green discs were then sintered in air using a conventional tube furnace. Fast-firing was performed by pushing the sample into and withdrawal from the hot-zone in a short time (of  3–5 min) in order that the high temperature defect equilibrium can be retained for further study. Samples were accommodated in a Pt-lined high-purity Al2O3 boat attached to a metal rod. The sintered density was determined by applying the Archimedes technique where distilled water was used as the immersion medium. Sintering kinetic curves were established by plotting the final sintered density vs. sintering time dwell at temperature. The room temperature conductivity was measured by the two-point probe method using an electrometer (model 617, Keithley, Cleveland, OH) after In–Ga electrodes have been applied on polished sample surfaces. The identification of crystalline phases was made by X-ray diffractometry (XRD, Siemens D5000, Karlsruhe, Germany) using CuKa radiation operating at 30 kV and 2 mA with a Ni filter. Sintered samples were mechanically ground and polished with SiC grits successively before diamond lapping to 1 mm surface

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Fig. 1. Sintering kinetic curves for Y2O3-doped BaO-excess compositions at 1290 °C.

roughness for microstructural observations. Both the reflected light optical and scanning electron microscope (SEM, JEOL JEM6400, Tokyo, Japan) were used to observe the polished sections. Grain boundaries were delineated by thermal etching at 1200 °C, or chemical etching using 1% HF solution where appropriate. Average grain size was determined on SEM micrographs adopting the linear interception technique [22]. Thin foils for transmission electron microscopy (TEM) were prepared by the standard procedure of slicing the sintered disc to :200 mm thickness using a diamond-embedded saw, ultrasonic drilling to 3 mm discs, mechanical polishing to 1 mm of surface roughness and : 20 mm thickness, dimple-grinding and Ar+-ion beam thinning to electron transparency. Observations were performed in a JEOL AEM3010 operating at 300 kV.

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1290 °C for 4 h. The final sintered density obtained was decreased to Dzrel : 85% for samples with 1.0 mol% Y2O3. Grain growth inhibition is striking in samples doped with Y2O3 and sintered at 1350 °C for 50 h where the final sintered density of Dzrel \ 95% has been reached. Exaggerated grain growth represented by intragranularly located spherical pores and abnormally large grains (of several hundred micrometres in Fig. 2(a)) having concave boundaries is evident from the 0.15 mol% Y2O3-doped samples. Effective grain growth inhibition took place abruptly at the doping level of 0.30– 0.50 mol% Y2O3 (Fig. 2(b) and Fig. 4). No intragranularly located pores can be detected from samples doped with ] 0.30 mol% Y2O3 as shown in Fig. 2(b)–(d). The variation of the Gav with Y2O3-dopant level is described in more detail in the Section 3.2. Grain pullouts and (so) the fragility of samples [4,6] were not experienced and the crack-like grain boundaries [6] were absent from the polished sections of Y2O3-doped samples. The hygroscopic second-phase of Ba2TiO4 [4,6,23] has not been detected by TEM. However, intergranularly located second-phase particles have appeared in samples with 1.0 mol% Y2O3 addition. The larger ones are delineated by SEM-SEI (secondary electron image) as indicated by arrows in Fig. 2(d). One ( 1.2 mm) of these Y2Ti2O7 (JCPDS 27-0982) second-phase particles identified by TEM is shown in Fig. 3(a) with the corresponding SADP (selected-area diffraction pattern) shown in Fig. 3(b), although not detected by XRD due to its minor quantity. It has also been identified by both TEM [23] of the TiO2-excess samples and XRD [19] in the Y2O3-doped BaTiO3 compositions.

3.2. Grain growth inhibition 3. Results

3.1. Sintering beha6iour and microstructure Y2O3-doping showed an adverse effect on the sintered density of BaO-excess 1.013 BaTiO3 powder. Adding Y2O3 has hindered the densification. The effect is characterized by a significant decrease of the final sintered density by Dzrel :18% from zrel :97 to 79% on increasing the doping level to 0.80 mol% Y2O3 when samples were sintered at 1215 °C for 100 h. The hindrance of densification is clearly demonstrated by the sintering curves at 1290 °C (as shown in Fig. 1) for the Y2O3-doped BaO-excess compositions. Two groups of curves obtained for high (0.65– 1.0 mol%) and low (0.15– 0.50 mol%) Y2O3-content, respectively can be distinguished from Fig. 1. Higher doping levels (of \ 0.65 mol%) have further impeded densification (as demonstrated by the lower group whose data were denoted by open symbols), the final sintered densities are lower by at least Dzrel :10% after sintering at

Fig. 4 gives the variation of the Gav with Y2O3-doping level when samples were sintered at temperatures between 1215 and 1450 °C for 50 h. Grain growth inhibition starts from the threshold of 0.30 mol% Y2O3 for sintering at temperatures below 1450 °C. There exists a concentration range of 0.30–0.50 mol% Y2O3 where the Gav is effectively suppressed to 3–5 mm at 0.30 mol% initially before grain growth recurs to Gav : 100 mm at 0.50 mol%. This range of dopant concentration is a window of GGIT [5] as indicated in Fig. 4. However, the GGIT window did not emerge until the sintering temperature had been raised to ] 1270 °C (as shown in Fig. 4). Increasing Gav only becomes appreciable for Y2O3-doping levels higher than 0.50 mol% at such sintering temperatures. For samples sintered at the lower temperatures of 1215 and 1250 °C, the Gav has levelled off and remains at 3 mm from 0.30 to 1.0 mol% Y2O3 (as shown in Fig. 4). The latter concentration is the highest doping level investigated in this study.

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The window of GGIT can be clearly identified from samples sintered at 1450 °C. The abrupt reduction of Gav from  100 to  3 – 5 mm has occurred twice at 0.50 and 1.0 mol% Y2O3, respectively (as shown in Fig. 4). The reduction of Gav at the two distinctive doping levels is very significant in terms of grain growth inhibition [5]. The exaggerated grain growth often found in the undoped TiO2-excess BaTiO3 compositions [5,6] has completely been arrested and that leads consequently to a homogeneous microstructure. The refined microstructure is shown typically in Fig. 2(b) and (d). The resumption of grain growth manifested by the increase of Gav to 6 mm in samples doped with Y2O3 ]0.65 mol% (as indicated by arrow in Fig. 4) is recognized from sintering at 1350 °C (or ]1270 °C), only that

the variation is not as appreciable as that shown for samples sintered at 1450 °C. For sintering at lower temperatures of 1215 and 1250 °C, grain growth was hindered to an Gav of  3 mm by adding Y2O3 higher than 0.30 mol%. The sintered microstructure remains suppressed at such an Gav, and no window of GGIT has emerged in Fig. 4. This is similar to the TiO2-excess BaTiO3 samples doped with other donor-oxides [7,10] of e.g. La2O3 when sintered at 1350 °C.

3.3. Semiconducti6ity The samples fast-fired at 1350 °C for 50 h changed colour from bluish in the lightly doped (of 0.15 mol% Y2O3) to a darker hue as the doping level increased to

Fig. 2. Microstructures of Y2O3-doped BaO-excess samples sintered at 1350 °C for 50 h: (a) 0.15 mol%; (b) 0.50 mol%; (c) 0.65 mol%; and (d) 1.0 mol% (SEM-SEI).

M.-H. Lin, H.-Y. Lu / Materials Science and Engineering A335 (2002) 101–108

Fig. 3. Second-phase grains in 1.0 mol% Y2O3-doped BaO-excess composition from sintering at 1350 °C/50 h: (a) BF image; and (b) SADP (Z: [110]) (TEM).

Fig. 4. Grain size dependence upon Y2O3-doping level in the BaO-excess samples sintered at 1215 –1450 °C.

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0.65 mol% Y2O3. For doping level exceeds 0.65 mol% Y2O3, samples become insulators of brown to buff in appearance. The changing of sample colour is consistent with the conductivity measurements. Fig. 5 showing the dependence of conductivity upon Y2O3-doping level reveals that samples sintered below 1450 °C exhibit an abrupt reduction at 0.65 mol% Y2O3 in the conductivity by eight decades ( 108). Better room temperature conductivity of 10 − 2 –10 − 4 (V mm) − 1 occurs with the doping levels of B 0.50 mol% Y2O3. For samples sintered at 1450 °C, a decreasing trend of conductivity is still observed but no discontinuity as in samples sintered at temperatures of B1450 °C has been observed from increasing the Y2O3-doping. It appears that air-cooling from fast-firing in a conventional furnace has retained the high temperature defect equilibrium similar to those of Y2O3-doped TiO2-excess compositions [23]. Furthermore, the abrupt reduction of the semiconductivity over the range of 0.50–0.65 mol% Y2O3-doping (as shown in Fig. 5) is clearly demonstrated. This also suggests that the principal charge compensation mechanism at the higher temperature of 1450 °C is different from that in the lower temperature range of B 1450 °C. For Y2O3-doping of less than 0.50 mol%, however, both samples of large and small grains (Fig. 2(a) and (b)) exhibit similar semiconductivity (as given in Fig. 5). For those containing 0.15 mol% Y2O3, the conductivity is almost the same (within one order of magnitude) from samples sintered at either high temperature of 1450 °C (of Gav : 100 mm) or low of 1250 °C (of Gav : 13 mm). Nevertheless, it is decreased by tenfold at 0.50 mol% Y2O3 with increasing dopant content only for samples sintered at 1215 °C.

Fig. 5. Conductivity vs. Y2O3-doping level for BaO-excess samples sintered at 1215 – 1450 °C.

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4. Discussion

4.1. Grain growth inhibition mechanism Grain growth inhibition similar to the La2O3-donordoped TiO2-excess BaTiO3 compositions [7,10] has occurred by  0.30 mol% Y2O3-doping (referred to Fig. 4). This is also consistent with the GGIT values from other dopants given in Table 1, but a window of 0.30 –0.50 mol% where grain growth remains inhibited was observed. The Gav remains small at 3 mm for doping levels beyond the GGIT window. However, the refined microstructure is limited to samples sintered below 1270 °C, the recurrence of coarsening in samples sintered at higher temperatures (\1270 °C) appears to be contradictory to other donor-doped compositions [7,10]. This may be explained by adopting the solid-solution defect model [24]. The solid-solubility of BaO in BaTiO3 is below 100 ppm [25]. Consider that the valence-III donor cation Y3 + substitutes for the Ba2 + -site and generates extrinsic defects in the stoichiometric BaTiO3, three defect reaction equations [3] can be described below. X 2Y2O3 +3TiO2 “ 4 Y’Ba +V§%Ti +12OX O +3TiTi

(1a)

X 2Y2O3 +6TiO2 “ 4Y’Ba +2V¦Ba +18OX O +6TiTi

(1b)

2Y2O3 +4TiO2 “ 4Y’Ba +12O +4Ti +O2(g) +4e% X O

X Ti

(1c) TiO2 on the right-hand side of the equations required for mass balance does not represent the hyper-stoichiometric TiO2-excess composition. The cation vacancies of V¦Ba, V§%Ti, and electrons (e%) are created for charge compensation in the host lattice when the donor ion substitutes for the Ba2 + -site cation and forms Y’Ba. The principal charge compensating defect of V¦Ba has been suggested to be kinetically more favourable [26] although V§%Ti was energetically more likely to be the dominant species according to its formation energy calculated by Lewis et al. [27]. Since samples sintered at all temperatures, except 1450 °C, have become insulator of |: 10 − 11 (V mm) − 1 for Y2O3 \0.65 mol% (referred to Fig. 5), a transition range of the Y2O3-doping level may exist in which the principal charge compensation species changes from the electron (Eq. (1c)) to the (cation) vacancy (Eqs. (1a) and (1b)). This is indicated by the sharp reduction of conductivity, as given in Fig. 5, over the dopant concentration of 0.50– 0.65 mol% Y2O3. The gap between the sintering curves for samples doped, respectively with 0.50 and 0.65 mol% Y2O3 (referred to Fig. 1) may also suggest a change of the principal charge compensating defect from one (i.e. cation vacancy) [6,7] which aids to the densification to one (electron (Eq. (1c)) or oxygen vacancy (Eq. (2)) given below) which retards it. For samples sintered at higher temperatures of \1270 °C,

however, Gav starts to increase again for doping level beyond 0.65 mol% Y2O3 (shown in Fig. 4) is not in accord with the argument that coarsening should remain suppressed after the threshold. Unlike the refined microstructure of highly La2O3doped samples [5,7,10] (of [La2O3]\ 0.20 mol%; square bracket for concentration [7]) where the Gav remains small at Gav : 3–5 mm, grain growth has recurred and Gav increased to  100 mm in samples doped with \ 0.50 mol% Y2O3 and sintered at 1450 °C (referred to Fig. 4). At higher Y2O3-doping levels, then, the correlation between the increased grain size (shown in Fig. 4) and the decreased conductivity (shown in Fig. 5) does not become immediately clear. Since both the grain growth inhibition at lower dopant levels (of 50.50 mol% Y2O3) and the rate-determining species for densification should not alter with the increasing dopant concentration, the recurrence of grain growth at Y2O3 \ 0.65 mol% and inhibition at Y2O3 \ 0.80 mol% to Gav : 3 mm has invalidated the mechanism of (dz/ dG)-ratio control based on the solid-solution defect model. A grain growth inhibition mechanism, different from the enhanced (dz/dG)-ratio [7,10,24], that accounts for both the decreasing barium vacancy (V¦Ba) concentration [6] and decreased conductivity is thought to have existed for the Y2O3-doped samples. It is probably that two distinctive mechanisms of grain growth inhibition may account for the observations of the GGIT window with grain growth inhibition at Y2O3 \ 0.30 mol%, recurrence at Y2O3 \ 0.65 mol% and inhibition again at Y2O3 \ 0.80 mol% to Gav :3 mm; one mechanism for samples doped with 50.50 mol% Y2O3 and the other for higher doping levels of 1.0 mol% Y2O3 (or \ 0.80 mol%). For doping level exceeds 0.50 mol% Y2O3, Y3 + becomes an acceptor with the principal charge compensating defect of oxy’’ gen vacancy (VO) at \ 1270 °C. The transition from a predominantly donor behaviour to an acceptor behaviour takes place at :0.50 mol% Y2O3. Y3 + substituting for both the Ba2 + - and Ti4 + -site has indeed been reported [16–18,21]. The reduction of cation vacancies (Eqs. (1a) and (1b)) and the increase of oxygen vacancies (Eq. (2)) concurrently with the increasing Y2O3-doping level beyond 0.50 mol% result in grain growth restoration, as revealed by Fig. 4. This is because the principal charge compensating defect of cation vacancy (V¦Ba of Eq. (1b)) from the donor behaviour responsible for improving the (dz/dG)-ratio has been significantly reduced in concentration when ’’ replaced by VO (Eq. (2)) of the acceptor behaviour. For Y2O3-doping greater than 0.80 mol% grain growth inhibition would then occur by a mechanism completely different from the solid-solution defect model. It may be the second-phase pinning as supported by the presence of Y2Ti2O7 (shown Fig. 2(d) and Eq. (3)) in sintered samples.

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4.2. Mixed charge compensation with Y2O3 doping The incorporation of Y3 + into BaTiO3 being more complex than other Ba2 + -site donor-dopants of La3 + [7,10], Nd3 + [17], and Gd3 + [17,28] has been recognized, because of the substantial discrepancy of its ionic size with both Ba2 + and Ti4 + [8,17], and the variable site-occupancy [18,14]. Electron (e%), cation vacancies of V¦Ba and V§%Ti are three possible charge compensating defects for donor-doping, described, respectively by Eqs. (1a), (1b) and (1c). The defect reaction for Y3 + as an acceptor and charge compensated principally by ’’ oxygen vacancy (VO) is described by, ’’

X 2BaO+Y2O3 “2Y%Ti +VO +2BaX Ba +5OO

(2)

’’ O

Oxygen vacancy (V ), as argued precedingly, may become the principal charge compensating defect when Y2O3-dopant exceeds 0.65 mol%. Y3 + substituting for Ba2 + becomes a donor and 4+ becomes an acceptor [8,18,21] is self-compensatTi ing. Self-compensation in the donor-and-acceptor codoped BaTiO3 has been calculated theoretically [8] as well as demonstrated experimentally [7,10]. The reduced conductivity of the Y2O3-doped samples is therefore due to self-compensation. The ceramic still contains the charge compensating electron (Eq. (1c)), albeit of a much reduced concentration as evidenced by the decreasing conductivity in Fig. 5. The increasing oxygen vacancy (Eq. (2)) is responsible for the restoration of coarsening and (so) the increased Gav (shown in Fig. 4). Therefore, Y2O3 should be treated as a less effective (donor-) dopant in terms of promoting densification and semiconductivity in doped-BaTiO3 samples. Samples added with 0.30 mol% Y2O3 and sintered at 1450 °C have an Gav of :100 mm and the maximum conductivity has been reached before grain growth inhibition occurred. On the other hand, those with similar Gav :100 mm from 0.600.80 mol% Y2O3 (shown in Fig. 4) exhibit not only lower conductivities but also a decreasing trend towards higher Y2O3-doping levels (shown in Fig. 5). The decreasing but still higher conductivities before grain growth suppression occurred for the second time may be accounted for by a partition of Y3 + being a donor or an acceptor at 1450 °C, favouring electron compensation. The conductivity declines to | =10 − 10 (V mm) − 1 when the acceptor behaviour prevails at 1.0 mol% Y2O3. Y3 + becomes self-compensating and the defect reaction equation below shows a mixture of the donor- and acceptor-behaviour: 3Y2O3 +4TiO2 + 2BaO

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structure has become heterogeneous that the conductivity may then be determined by the relative proportion of Y3 + site-occupancy as the Ba2 + -site donor to that as the Ti4 + -site acceptor for Y2O3 \ 0.65 mol%. The recurrence of grain growth inhibition at \0.80 mol% Y2O3 doping, however, cannot be ascribed to the acceptor behaviour adopting the solid-solution defect model. Y3 + acting as an acceptor at \0.80 mol% Y2O3 in the BaO-excess composition could have exsolved TiO2 to form second-phase Y2Ti2O7 (shown in Fig. 3). These particles exert pinning on grain boundaries by which grain growth is inhibited equally effectively as that imparted by the enhanced (dz/dG)ratio [24]. This argument necessitates Y3 + site-occupancy as an acceptor in a relatively larger proportion (Eq. (2)). Ba2TiO4 in the undoped, BaO-excess samples [6] has not been found by TEM indicates that the initially BaO-excess composition, upon Y2O3-doping, has been modified towards stoichiometry. A dopant concentration of 0.65 mol% Y2O3 would have resulted in the cation stoichiometry (i.e. Ba/(Y + Ti) =1.0) should Y3 + as an acceptor dissolves completely in the initially BaO-excess 1.013 BaTiO3. Grain growth inhibited again at 1.0 mol% Y2O3 may be attributed to the second-phase Y2Ti2O7, particularly when the solid-solubility of Y3 + of 1.5 at.% (0.75 mol% Y2O3) [14] in BaTiO3 has been exceeded. The reinstatement of such an inhibition effect is more pronounced for \0.80 mol% Y2O3-doped samples at higher temperatures of 1450 °C with higher solid-solubility. 5. Conclusions Grain growth inhibition to an Gav of :3 mm in the Y2O3-doped BaO-excess BaTiO3 compositions sintered at 51350 °C occurs at the threshold of 0.30 mol% Y2O3. The transition from semiconductor to insulator occurs since Y3 + acts both as donor and acceptor substituting, respectively for the Ba2 + - and Ti4 + -site in BaTiO3. Acknowledgements Funding of this research by the National Science Council of Taiwan through NSC-82-0405-E-110-029, 83-0405-E-110-007 and 89-2215-E-151-001, is acknowledged. Thanks are due to Dr Mike S.H. Chu of TAM Ceramics for supplying the BaTiO3 powder, and an anonymous referee for helpful comments.

X X “4Y’Ba +4e% + 2Y%Ti +VO +2BaX Ba +4TiTi +17OO ’’

+ O2(g)

(3)

The sintered BaTiO3 ceramic consists of grains with Y3 + acting as a donor or an acceptor. Since the defect

References [1] J. Daniels, K.H. Hardtl, D. Hennings, R. Wernicke, Philips Res. Rep. 31 (Part 1 – 5) (1976) 487.

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