Ceramics International 42 (2016) 17792–17797

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Microstructure and electrical properties of BaTiO3-(Bi0.5M0.5)TiO3

(M¼Li, Na, K, Rb) ceramics with positive temperature coefficient ofresistivity

Mengmeng Yang a,b, Zhijian Peng a,n, Chengbiao Wang a, Xiuli Fu b,nn

a School of Engineering and Technology, China University of Geosciences, Beijing 100083, PR Chinab School of Science, Beijing University of Posts and Telecommunications, Beijing 100876, PR China

a r t i c l e i n f o

Article history:Received 17 June 2016Received in revised form28 July 2016Accepted 17 August 2016Available online 18 August 2016

Keywords:PTCR ceramicsLead freeElectrical propertiesCurie temperature

x.doi.org/10.1016/j.ceramint.2016.08.10742/& 2016 Elsevier Ltd and Techna Group S.r

esponding author.responding author.ail addresses: pengzhijian@cugb.edu.cn (Z. Penbupt.edu.cn (X. Fu).

a b s t r a c t

Lead-free 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics with positive temperaturecoefficient of resistivity (PTCR) were prepared by conventional solid sintering reaction using high-purityreagents of metal oxides and carbonates. The effect of (Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) doping on themicrostructure and electrical properties of the samples were investigated. X-ray diffraction analysis in-dicated that all the samples were of a single perovskite structure with tetragonal phase, and the cal-culated c/a value increased first and then decreased. The Raman shift at 305 cm�1 became strong firstand then declined when M in the dopant (Bi0.5M0.5)TiO3 was changed from Li to Rb. The Curie tem-perature of the samples displayed a tendency to increase first and then decrease, in which BaTiO3-(Bi0.5K0.5)TiO3 ceramics presented the highest value. Moreover, after the incorporation of alkali ions, thePTCR value (defined by the resistivity jump with the ratio of the maximum to minimum) and roomtemperature resistivity of the samples decreased first and then increased, in which BaTiO3-(Bi0.5Na0.5)TiO3 ceramics possessed the lowest room temperature resistivity.

& 2016 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

1. Introduction

It is well-known that BaTiO3 (BT) based semiconducting cera-mics exhibit positive temperature coefficient of resistivity (PTCR)characteristics [1–3]. The resistivity of the samples sharply risesnear the Curie temperature (Tc), which can be applied in severalkinds of electrical devices, such as temperature sensor, self-con-trolled heater and device for over-current protection [4]. Althoughpure BT ceramics might be an insulator because of their high en-ergy gap (about 3.1 eV) if they are prepared by sintering in air, theycan be converted into a semiconductor through doping with var-ious transitional metals ions, such as trivalent (Y3þ , La3þ) andpentavalent (Nb5þ , Ta5þ) donor impurities, or sintering in redu-cing atmosphere, causing the loss of oxygen in the ceramics [5,6].In addition, the mechanism of PTCR effect of semiconductingceramics based on BT can be explained by a double Schottky bar-rier formed at the grain boundary [2,3]. So, potential barriersshould be formed at the grain boundaries of BT ceramics, which

.l. All rights reserved.

g),

can be realized by developing a surface layer of acceptor state onthe BT grains, such as the absorbed oxygen, acceptor ions, and/orcation vacancy [7–10].

Moreover, Tc is a critical parameter for the application of PTCRceramics, but the Tc value of the already known BT ceramics is onlyroughly 120 °C. In history, the Tc value of BT-based PTCR ceramicswas manipulated by substituting Pb2þ for Ba2þ in case that theywere used at a temperature above 120 °C [11]. However, becauseof the high toxicity of lead oxide during the preparation of Pb-contained products, which is why they have been forbidden inapplication by laws, considerable efforts have been devoted todeveloping lead-free BT-based PTCR ceramics [12]. Over the pastdecades, the use of Bi0.5Li0.5TiO3 (BLT), Bi0.5Na0.5TiO3 (BNT), andBi0.5K0.5TiO3 (BKT) in BT-based ceramics have made great progressfor developing lead-free PTCR ceramics with high Tc. In literature,Huo et al. [13] first reported that BaTiO3-Bi0.5Na0.5TiO3 (BT-BNT)ceramics were promising lead-free PTCR materials with Tc above130 °C. Later, Shimada [4] and Yasushi et al. [14] both successfullyprepared BT-BNT ceramics, confirming their excellent PTCR effect.In addition, Leng [15] and Yuan et al. [16] indicated that BaTiO3-Bi0.5K0.5TiO3 (BT-BKT) system also possessed good PTCR effect withTc higher than 130 °C. And Pu et al. [17] reported that the Tc valueof BaTiO3-Bi0.5Li0.5TiO3 (BT-BLT) can be enhanced by 30 °C over BT

M. Yang et al. / Ceramics International 42 (2016) 17792–17797 17793

ceramics with doping of BLT. However, there is still no reportabout the incorporation effect of Bi0.5Rb0.5TiO3 (BRT) into thesemiconducting BT ceramics. What is more, because Li, Na, K andRb belong to the same main-group (IA) in the Mendeleev periodictable of elements, it is reasonable to inquire what is the differenceif (Bi0.5M0.5)TiO3 (BMT, M¼Li, Na, K, Rb) with different alkali ionsare doped into BT ceramics, although the alkali ions may have si-milar chemical and physical properties but different sizes of ionicdiameter. However, there is no literature on the comparison be-tween BMT doped BT ceramics.

Therefore, in this work, a series of BT based ceramics dopedwith BMT of different alkali ion were prepared, and the influenceof BMT doping on the phase composition, microstructure andelectrical properties of the obtained BT based ceramics wasinvestigated.

2. Experimental procedure

2.1. Sample preparation

The nominal composition of the samples was 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb). The samples were preparedby conventional solid-state reaction. And the used starting mate-rials were commercially available BaCO3, TiO2, Bi2O3, Li2CO3,Na2CO3, K2CO3 and Rb2CO3 powders with a purity higher than99.9 wt%. In typical process, the powders were mixed by ball-milling with zirconia as grinding media in absolute alcohol for24 h. After milling, the obtained slurries were dried off at 80 °C inan oven open to air. Then, the resultant powder chunks weregrinded and sieved. After that, the obtained powders were cal-cined for 4 h at 900 °C in a Muffle furnace with a heating rate of5 °C /min. After granulation with 5 wt% polyvinyl alcohol, theprepared composite powders were pressed into green disks with asize of 10�10�2 mm3 under a pressure of 60 MPa. For the binderremoval, the green disks were heated in a Muffle furnace at a rateof 2 °C/min, and held at 500 °C for 2 h. Then the samples weresintered in an alumina tube furnace (SJG-16, China). During thesintering process, the alumina tube was first evacuated and flu-shed repeatedly with high-purity N2 gas so as to eliminate theremnant gases in it. Then, the furnace was heated up from roomtemperature to 1340 °C and held at this temperature for 4 h in a N2

flow of 300 sccm. After sintering, the furnace was cooled down to1000 °C for 2 h in air. After that, the furnace was cooled downnaturally to room temperature simply by shutting down theelectricity of the heating system of the furnace. Finally, an alumi-num paste was coated on both surfaces of the samples and treatedat 570 °C to obtain good Ohmic contacts between the electrodesand sample.

2.2. Materials characterization

Room-temperature phase identification of the sintered sampleswas carried out by X-ray diffraction (XRD, Model: D8-Advance,Germany) with CuKα radiation (λ¼1.54178 Å) through a con-tinuous scanning mode at a speed of 4°/min. The lattice parameterwas calculated based on the XRD results. Raman scattering spec-troscopy was adopted to detect the local structure in the samplesusing an instrument of Horiba HR800 spectrometer. The micro-structure was examined by using scanning electron microscopy(SEM, Model: LEO-1530, Germany). Then the grain morphology,size, and distribution could be evaluated. The resistivity of thespecimens as a function of temperature from room temperature toabout 300 °C was measured using a Keithley 2410 Digital SourceMeter while an electric stove was used to heat the samples. Theresistivity (ρ, Ω cm) can be calculated by the following equation:

πρ= ( )R d

h4 1

2

where R (Ω) is the measured resistance, d (cm) is the electrodediameter, and h (cm) is the thickness of samples. Dielectric per-mittivity (εr) was measured by an impedance analyzer (Agilent4294A) at a bias voltage of 500 mV and 1 MHz frequency withinthe temperature range of 30–300 °C.

3. Results and discussion

3.1. Composition, structure and morphology

Fig. 1 illustrates the XRD patterns of the as-prepared BT-BMTceramics with different alkali ions. All the diffraction peaks of thesamples match well with the published data of the host BaTiO3

(JCPDS card no. 81-2202), implying that all the BT-BMT ceramicsremain in a single phase with a perovskite ABO3-type structurewith tetragonal symmetry throughout the composition range. Inaddition, through an elaborate observation on 2θ angle around31.5° of the BT-BMT ceramics, it can be found that the (101) peakshifted to higher 2θ angle after BNT and BKT were doped while itwould shift to lower 2θ angle after BLT and BRT were doped. Thisphenomenon could be explained as follows. In BT-BMT ceramics,the doped Bi3þ , Naþ , Kþ and Rbþ cations will replace the site ofBa2þ ions according to Pauling Rule [18]. And because the ionicradii of Bi3þ (0.096 nm), Naþ (0.102 nm), and Kþ (0.133 nm) weresmaller than that of Ba2þ (0.134 nm), but Rbþ (0.152 nm) wasbigger than Ba2þ (0.134 nm), so the (101) peak in BT-BNT and BT-BKT ceramics would shift to higher 2θ angle, and that in BT-BRTshift to lower 2θ angle. For BT-BLT, Pu et al. [17] ever pointed outthat Liþ ions might not enter the crystal lattice but segregate atgrain boundary. Here we supposed that Liþ ions might occupy theTi4þ sites, because the radius of Liþ ion ( +rLi ¼0.076 nm; co-ordination number, CN¼6) is close to that of Ti4þ

( +rTi4 ¼0.076 nm; CN¼6), but much smaller than that of Ba2þ

( +rBa2 ¼0.134 nm; CN¼12) [19]. Thus, it is reasonable that Liþ ionswould preferentially substitute for the sites of Ti4þ ions in thelattice of BT ceramics, which would result in (101) peak shifting tolower 2θ angle.

In addition, the diffraction peak at 2θ angle of around 45° wasasymmetric and split into (002) and (200) peaks, indicating thatall the samples showed a tetragonal perovskite ABO3 structure[16]. The splitting of the (200)/(002) lines increased first and thendecreased, when the samples were doped with Liþ , Naþ , Kþ andRbþ . The increase means an enhanced tetragonality (c/a, whichcan usually denote the tetragonality directly) [20]. As presented inTable 1, the c/a values calculated from the XRD data for the sam-ples are 1.0110, 1.0117, 1.0119 and 1.0116, which also increased firstand then decreased, when the M in the doped BMT was changedfrom Li to Rb with increasing ionic radius. This result is consistentwith that of the splitting of (200)/(002) lines.

Raman spectroscopy is more sensitive to structural change thanXRD, which can distinguish more easily the crystalline phasespossessing different symmetries but similar lattice constants. TheRaman spectra for the as-prepared BT-BMT ceramics are presentedin Fig. 2. The sharp peak at 305 cm�1 and broad band at 722 cm�1

are specifically attributed to the tetragonal phase of BaTiO3 [21].The intensity of the sharp peak at 305 cm�1 increased first andthen decreased when they were doped with Liþ , Naþ , Kþ andRbþ , reaching a maximum for BT-BKT, confirming the change oftetragonality of the samples as mentioned above [18].

Fig. 3 displays typical SEM images on the surfaces of the pre-pared BT-BMT ceramics. The average grain sizes of BT-BNT, BT-BKTand BT-BRT ceramics are 6.52, 9.01 and 7.34 mm, which are all

Fig. 1. XRD patterns of the as-prepared 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics: (a) BT-BLT, (b) BT-BNT, (c) BT-BKT, and (d) BT-BRT.

Table 1Lattice parameters of the prepared 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na,K, Rb) ceramics.

Samples a c c/a

BT 3.993 4.036 1.0108BT-BLT 3.9959 4.0399 1.0110BT-BNT 3.9921 4.0390 1.0117BT-BKT 3.9930 4.0408 1.0119BT-BRT 3.9997 4.0460 1.0116

Fig. 2. Raman spectra of the as-prepared 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3

(M¼Li, Na, K, Rb) ceramics.

M. Yang et al. / Ceramics International 42 (2016) 17792–1779717794

smaller than that of BT-BLT ceramics (16.19 mm). For the BT-BNT,BT-BKT and BT-BRT ceramics, because the sizes of the doped Naþ ,Kþ and Rbþ ions are larger than that of Ti4þ ion, but close to thatof Ba2þ ion, so during sintering, they would replace the sites ofBa2þ ions partially [22]. It is known that low valence cation tendedto segregate onto grain boundaries of BT ceramics in high tem-perature equilibrium [23,24]. Therefore, the redundant Naþ , Kþ

and Rbþ ions would segregate at the grain boundaries in theprepared BT-BMT ceramics, forming a thin M-rich (M¼Na, K andRb) shell on the grains, inhibiting the grain growth, which is whythey have smaller grain sizes than BT-BLT ceramics. Moreover,Naþ , Kþ and Rbþ ions were easily volatilized at high tempature,and the evaporation amount of Naþ (as Na2O), Kþ (as K2O) andRbþ (as Rb2O) is related to the melting point of Na2O, K2O and

Rb2O, which is 1132, 770 and 400 °C, respectively. So, there aremore Kþ and Rbþ to volatilize than Naþ during sintering, and lessredundant Kþ or Rbþ to form a thin K- or Rb-rich shell. As a result,the average grain sizes of BT-BKT and BT-BRT ceramics are rela-tively lager than that of BT-BNT ceramics. In addition, the ionicradii of Kþ (0.133 nm) and Ba2þ (0.134 nm) are almost equal, sothere are more Kþ ions to replace the sites of Ba2þ according toPauling Rule [18]. That is, the K-rich layer in BT-BKT ceramics isthinner than Rb-rich one in BT-BRT. So, the BT-BKT ceramics havelager average grain size than BT-BRT. For BT-BLT ceramics, Liþ ionswould substitute for Ti4þ ions, and such process can be proceedednearly completely due to their almost equal ionic radii and samecoordination number in the ceramics [19]. In other word, therewere less Liþ ions to segregate at the grain boundaries to form Li-rich shell on the BT-BLT grains. So, the grain size of the BT-BLTceramics was lager than those of the BT-BNT, BT-BKT and BT-BRTceramics.

3.2. Electrical properties

Fig. 4 illustrates the temperature dependence of electrical re-sistivity for the as-prepared BT-BMT ceramic samples. It can beseen that all the samples presented typical PTCR effect with amarked resistivity jump (ρmax/ρmin4103) near the Tc. Moreover,the room-temperature resistivity (ρ25) of the BT-BMT ceramics hasa tendecy to rise when Naþ , Kþ and Rbþ were doped, but they areall lower than that of BT-BLT ceramics. For BT-BNT, BT-BKT and BT-BRT ceramics, the rise in ρ25 could be explained from the view ofthe existence of cation vacancies in the samples. When the sam-ples were sintered in reducing atmosphere, a large amount ofoxgen vacanices would be formed, and the evaporation of Ba2þ

site, Naþ (as Na2O), Kþ (as K2O) and Rbþ (as Rb2O) was induced athigh tempature, which can be expressed by the following equa-tions:

→ + + ′ ( )∙∙O V O e0.5 2 2O O 2

( )+ = ″ + + ↑ ( )⋅ × ⋅⋅Bi O V V Bi O2 3 2 3 3Ba O Ba O 2 3

( )′ + = ″ + + ↑ ( )× ⋅⋅Na O V V Na O0.5 0.5 0.5 4Ba O Ba O 2

( )′ + = ″ + + ↑ ( )× ⋅⋅K O V V K O0.5 0.5 0.5 5Ba O Ba O 2

( )′ + = ″ + + ↑ ( )× ⋅⋅Rb O V V Rb O0.5 0.5 0.5 6Ba O Ba O 2

Fig. 3. Typical SEM images on the surfaces of the prepared 0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics: (a) BT-BLT, (b) BT-BNT, (c) BT-BKT, and (d) BT-BRT.

Fig. 4. Temperature dependence of electrical resistivity for the as-prepared0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics.

Fig. 5. Temperature dependence of permittivity for the as-prepared0.912Ba0.97TiO3-0.088(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics.

M. Yang et al. / Ceramics International 42 (2016) 17792–17797 17795

As a result, a large amount of ″VBa would be formed, whichwould inhibit the reduction of Ti4þ to Ti3þ by neutralizing thecharge associated with ⋅⋅VO and reduce the effective bulk carrierconcentration [17], finally resulting in a higher room temperatureresistivity. In BT-BNT, BT-BKT and BT-BRT ceramics, the BT-BNTceramics had the lowest ρ25, due to that the melting point of Na2Ois the highest among Na2O, K2O and Rb2O. Resultantly, there arethe least Naþ ions to volatilize. In other word, BT-BNT ceramicshave the least amount of compensated oxgen vacanices. For BT-BLTceramics, Liþ ions would preferentially substitute for Ti4þ in BTlattices, forming acceptor defect LiTi

‴. And this defect wouldcompensate the residual charges formed by the donor Bi3þ ions oroxygen vacancies. The resultant complex defect (3 ⋅BiBa LiTi

‴) wouldreduce the semiconductivity of the grain boundary layers, thusincreasing their resistivity at room temperature [19].

Fig. 5 shows the temperature dependence of the permittivity(εr) at 1 MHz for the as-prepared BT-BMT ceramics. The BT-BLTceramics present a normal ferroelectric characteristic, and their εr-T curves display a sharp shape near the Curie point. However, forBT-BNT, BT-BKT and BT-BRT ceramics, the peaks of the εr-T curvesnear Tc became broad, much like those of relaxor ferroelectrics.This phenomenon was correlated with the substitution of Bi3þ

ions by Mþ (M¼Na, K, Rb), which distorted the unit cell andchanged the dipolar moment, thus inducing lattice strain [25]. Thecations co-occupy the same lattice of unit cell, and thus the che-mical composition and crystal structure are inhomogeneous innanometer scale, finally resulting in the relaxation feature [26]. Itis well known that the peak position of the εr-T curves corres-ponds to the Tc, which does not depend on the measurementfrequency [27]. From Fig. 5, it can be seen that the peaks of the as-prepared BT-BMT ceramics around 130–150 °C were higher than

M. Yang et al. / Ceramics International 42 (2016) 17792–1779717796

that of pure BT (120 °C). In other word, the Tc progressively movedtoward higher temperature after BT ceramics were doped withBMT, in which BT-BKT ceramics present the highest Tc. The in-crease of Tc had a close relationship with c/a value[20]. In a per-ovskite ABO3 structure, A-site ion is at the center of oxygen do-decahedron, and the oxygen ions in the oxygen dodecahedron alsoform oxygen octahedron, in which B-site ion is at the center. Thatis to say, oxygen octahedron was influenced directly and indirectlyby A-site ion [28]. So, although few of A-site ions were replaced,there would be a vast number of B–O bonds (here Ti–O bonds) tobe influenced, inducing the change of the tetragonality of thesamples. In this work, the melting point of Bi2O3 (820 °C) is verylow. So the Bi-O bond is very weak. As a result, the A–O bonds inoxygen dodecahedron would be weakened when Bi3þ replacedthe site of Ba2þ (A-site) in BaTiO3. The interactions between Ti4þ

and its neighboring O2� in oxygen octahedron were so strong thatTi4þ would leave its site, inducing the increase of tetragonality,and Ti4þ could not resume its seat unless the tetragonal ferro-electric was wrecked at higher temperature [13]. Resultantly, theTc of the BT-BMT ceramics was clearly enhanced with the increaseof tetragonality, and it had the same change tendency with tet-ragonality, which increased first and then decreased in the presentwork, when M in the doped BMT was changed from Li to Rb.

The ratio of maximum to minimum resistivity (ρmax/ρmin) forthe as-prepared BT-BMT ceramics decreased first and then in-creased. It is known that the height of potential barriers on grainboundaries and the density of the grain surface acceptor state ofBT-based ceramics are the two critical factors in charge of the PTCRcharacteristic [29]. In this work, the ceramics were sintered in a N2

atmosphere, in which the residual O2 concentration was very low.Therefore, a large amount of oxygen vacancies would be formed inthe ceramics after sintering. During the re-oxidation process,oxygen vacancies would be filled up, and oxygen atoms would firstdiffused at grain boundaries followed by accessing into the grains[29]. In literature, Sinclair et al. [30,31] suggested that the ab-sorbed oxygen were the predominant acceptor state at the grainboundary, and the resistivity of grain boundary took the dominantposition in the overall resistivity of BT-based PTCR ceramics whenthe samples were sintered in reducing and re-oxidizing atmo-spheres. For BT-BNT, BT-BKT and BT-BRT ceramics, the averagegrain sizes increased first and then decreased. And it is known thatif the samples had smaller grain sizes, they would have more grainboundaries. More grain boundaries among the small grains wouldcause oxygen vacancy to be compensated by oxygen easily, whichwould lead to higher height of potential barriers, and finally re-sulting in the rising of ρmax/ρmin rapidly [32]. Because BT-BNTceramics had the smallest grain size, so its ρmax/ρmin value wouldbe the maximum. For BT-BLT ceramics, Liþ ions would substitutefor Ti4þ , forming the acceptor defects of LiTi‴ [19]. These defectswould increase the state density of the acceptors at grain bound-aries, and play a leading role in improving the PTCR effect byjoining with the existing acceptors such as the adsorbed oxygen.As the adsorbed oxygen formed and the acceptor segregated atgrain boundaries, both the surface charge density and height ofpotential barriers increased, which resulted in the most excellentPTCR effect of BT-BLT among all the BT-BMT ceramics.

4. Conclusions

BaTiO3-(Bi0.5M0.5)TiO3 (M¼Li, Na, K, Rb) ceramics were pre-pared by conventional solid sintering reaction using high-purityreagents of metal oxides and carbonates. The sintering of thesamples was carried out in a reducing atmosphere of N2 at 1340 °Cfollowed by oxidation at 1000 °C in air.

The tetragonality denoted by c/a values calculated from XRD

data increased first and then decreased, and the Raman shift at305 cm�1 became strong first and then declined. In all the pre-pared BT-BMT ceramics, the BT-BLT ceramics had the largest andBT-BNT possessed the smallest grain sizes, and the average grainsizes of BT-BNT, BT-BKT and BT-BRT ceramics increased first andthen decreased.

All the samples presented typical PTCR characteristics. The ra-tio of maximum to minimum resistivity and room-temperatureresistivity decreased first and then increased, in which BT-BNTceramics had the lowest room temperature resistivity and BT-BLTceramics had the best PTCR performance. The Tc of BT-BMT cera-mics was all enhanced after the doping of BMT, and had a tendecyto increase first and then decrease when M was changed from Li toRb.

Acknowledgements

This work was supported by the National Natural ScienceFoundation of China (grant nos. 61274015, 11274052 and51172030), and Excellent Adviser Foundation in China Universityof Geosciences from the Fundamental Research Funds for theCentral Universities.

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