Defect engineering toward the structures and dielectric behaviors of (Nb, Zn) co-doped SrTiO3 ceramics

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Original Article

Defect engineering toward the structures and dielectric behaviors of (Nb, Zn) co-doped SrTiO3 ceramics ⁎

Wengao Pana, Minghe Caoa, , Hua Haoa, Zhonghua Yaoa, Zhiyong Yua, Hanxing Liub a b

State Key Laboratory of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan, 430070, PR China International School of Materials Science and Engineering, Wuhan University of Technology, Wuhan, 430070, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Complex ions Defect engineering Dielectric properties SrTiO3ceramics

Defect engineering is applied to induce controllable changes in structures and dielectric behaviors in oxides. SrTiO3 ceramics modified by complex ions (Nb, Zn)x+ (x = 2, 3, 4 and 5) with variable valence states are prepared by standard solid state method. Phase structures, microstructures, defect structures and dielectric properties of the (Nb, Zn) co-doped SrTiO3 ceramics have been systematically investigated. In acceptor codoping, M2 ceramics (x = 2) exhibits extremely low dielectric loss (≤0.001). M3 ceramics (x = 3) shows low conductivity and excellent thermal stability. In equivalent co-doping, M4 ceramics (x = 4) possesses giant permittivity and low dielectric loss. In donor co-doping, M5 ceramics (x = 5) presents improved permittivity while deteriorated dielectric loss. Further investigations reveal that oxygen vacancy is beneficial to the localization of charge carriers and then the low tangent loss, while excess electrons not only contribute to the improved permittivity but also result in high dielectric loss. Appropriate concentrations of oxygen vacancy and electron are significantly important for the formation of multiple defect dipoles or defect clusters, and therefore the greatly modified dielectric properties. The findings may facilitate the ability to engineer the advanced electronic materials.

1. Introduction

sintering atmosphere [12], which have been widely attempted and proved to be effective by other researchers [13–16]. By regulating the no-stoichiometry, Zhou et al. obtained giant permittivity (∼5000) and low tangent loss (∼0.01) in Sr0.985Ce0.01Ti1 ± xO3 ceramics [17], which is also proved to be beneficial for improving the energy storage properties in Sr1 ± xTiO3 ceramics [18]. Substitution impurity ions for A and/or B sites is another widely used method to modulate the dielectric properties [19–22]. Modifying the sites with a single ion is usually too monotonous and insufficient to achieve excellent comprehensive dielectric properties. Ion substitution with multiple elements has attracted more and more attention and been progressively investigated. Recently, He et al. reported that niobium and aluminum co-doped SrTiO3 ceramics exhibit high permittivity (10,500) and low tangent loss (0.03) due to the related point defect mechanism [23]. In our previous work, SrTi1-x(Zn1/3Nb2/3)xO3(STZN) ceramics were prepared and improved dielectric properties were achieved [24]. Besides, In N2 sintering atmosphere, the STZN ceramics exhibited an obviously decreased dielectric loss while a maintained permittivity compared with the SrTi1-xNbxO3 ceramics [13], further implying the effectiveness and advantage of co-doping substitution. Co-doping substitution itself provides a more open and flexible

Advanced solid-state dielectrics are key enablers for potential applications in electronic circuit, microwave communication, renewable energy storage and high-power system [1–4]. Among various candidates, SrTiO3 (ST) based ceramics intrinsically possess many advantages, such as low dielectric loss, high breakdown strength and excellent temperature/frequency stabilities [5–9]. Whereas, it is still insufficient to meet the specific application requirements. Many efforts have been attempted to engineer or tailor the dielectric properties, mainly including process optimization, non-stoichiometry and ion substitution. Such as Qi et al. recently investigated the influences of sintering temperature on dielectric properties of Sr0.985Ce0.01TiO3 ceramics and found that giant permittivity and low tangent loss can only be achieved at a certain sintering temperature range, and the dielectric properties will be deteriorated due to destruction of defect clusters when the temperature further increased [10]. Chen et al. revealed dielectric relaxation peaks could be suppressed by annealing in a reducing atmosphere while reinforced or recreated in an oxidizing atmosphere [11]. Tkach et al. reported high tunability and colossal permittivity (∼105) in Sr1-1.5xYxTiO3 ceramics with the help of N2

⁎

Corresponding author. E-mail address: [email protected] (M. Cao).

https://doi.org/10.1016/j.jeurceramsoc.2019.09.027 Received 18 August 2019; Received in revised form 12 September 2019; Accepted 15 September 2019 0955-2219/ © 2019 Elsevier Ltd. All rights reserved.

Please cite this article as: Wengao Pan, et al., Journal of the European Ceramic Society, https://doi.org/10.1016/j.jeurceramsoc.2019.09.027

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Fig. 1. (a) XRD pattern of samples. (b) The enlarged scale of (111) peak. (c) Refinement result of M4 sample.

9.857 GHz. Thermally stimulated depolarization current (TSDC) measurement was performed using pA electrometer (6517A, Keithley, USA). The DC polarization condition was 36∼108 V/mm at 190 °C for 10 min, and the heating rate for current measurement was 5 °C/min. To obtain possible information on local structural symmetry and the vibration and rotation state of chemical bonds, room temperature Raman spectra were recorded by confocal micro-Raman spectrometer (LabRam HR Evolution) with 633 nm internal Ar ion laser source. Low temperature (-150-100 °C) dielectric responses were measured by HP 4284A LCR meter. High temperature (25–350 °C) dielectric properties and complex impedance were evaluated using Agilent 4980A LCR meter.

approach to tailor the dielectric properties. However, to the best of our knowledge, almost all researches focus on substituting host lattice for the complex ions with equivalent chemical valence state, and then just changing the co-doping concentration, which have been aggressively reported in piezoelectric and TiO2 material systems [25–28]. Actually, rather than a fixed one, chemical valence state of complex ions can be flexibly adjusted by regulating the ratio between co-doping ions. With the variation of complex valence state, it is expectable that substitutiontypes (donor doping, equivalent doping and acceptor doping), relative concentration of charge carriers (such as oxygen vacancy and electron) and therefore the resulting dielectric behaviors will also be changed. In viewpoint of defect design, it is also feasible that substitution of complex ions with nonequivalent valence state would result in an optimized dielectric properties on the condition that the total charge balance is kept [28]. In this paper, a novel sight of co-doping substitution is presented. Chemical valence states of complex ions (Nb, Zn)x+ are elaborately designed by manipulating the ratio between Nb5+ and Zn2+ ions. Dense SrTiO3 ceramics modified by the complex ions (Nb, Zn)x+ were prepared, and systematic investigations on the phase structures, microstructures, defect structures and dielectric behaviors of the ceramics are reported. Which not only greatly broaden the scopes for materials design, but also provide a completely new strategy to manipulate the defect chemistry and the related functional characteristics.

3. Results and discussion Fig. 1(a) shows the XRD pattern of samples, which indicates (Nb, Zn)x+ ions substitution barely changes the main diffraction peaks compared to the standard SrTiO3 (PDF#35-0734). All samples are identified as pure SrTiO3 without any detectable impurity phase. Splits of the (111) peak corresponding to tetragonality are absent, as shown in Fig. 1(b), implying all samples maintain the cubic crystal structure. Besides, compared to the standard position of (111) peak, all samples shift toward lower 2θ angles, underlying the incorporation of (Nb, Zn)x+ ions to Ti sites. The specific structural parameters are calculated by XRD refinement. Fig. 1(c) presents the refinement result of M4 sample and all the detailed structural parameters are shown in Table 1 as well as in Fig. 3(a), which further confirm the lattice expansion due to the ion substitution. Fig. 2(a) exhibits the lattice constant and grain size of the sintered ceramics. As the complex valence state of (Nb, Zn)x+ ions increases from 2+ to 5+, the substitution type transforms from acceptor doping to equivalent doping and then donor doping. The lattice constant gradually increases, while the grain size slightly increases first and then decreases, achieving the maximum value at x = 3. The specific section morphology can be seen in Fig. 3, and some pores are obviously observed in the grain section in M2 and M3 ceramics. Actually, defect reaction will occur due to the ion substitution at high sintering temperature. Vacancy defects are believed to facilitate mass transports, thus promoting the grain growth and the formation of holes in grain region [12,30]. Evident from the XRD result and the SEM morphology (Fig. 3), it seems that defect reactions do not proceed as the nominal composition. In M2 ceramics, Zn2+ ions are excessive. If the excess

2. Experimental details Complex metal ions (Nb, Zn)x+ were separately tuned as Zn2+, (Zn2/3Nb1/3)3+, (Zn1/3Nb2/3)4+ and Nb5+ ions for doping SrTiO3 ceramics. SrTi1-0.015x/4(Nb, Zn)0.015O3 (x = 2, 3, 4 and 5) ceramics were prepared by traditional solid state method with reagent-grade commercial powders of SrCO3, TiO2, Nb2O5 and ZnO. Samples modified by the corresponding valence states are hereafter abbreviated as M2, M3, M4 and M5, respectively. Mixtures were stoichiometrically weighed and ball-milled for 24 h in ethanol medium, then dried and calcined to form the main crystalline phase. The calcined powders were again milled and redried. Afterwards, pellets with 12 mm in diameter and 1 mm in thickness were uniaxially pressed (200 MPa) using 5 wt% PVA as a binder and then sintered at 600 °C for 2 h to remove the organic impurity. Next, green pellets were sintered at 1500 °C for 2 h in air atmosphere. Finally, dense ceramics were obtained and polished for further measurement, and Ag electrodes were coated before characterizing the electrical properties. Phase structure was detected by X-ray diffraction (XRD) (X'Pert PRO, PANalytical, Holland). Structural parameters were calculated by XRD refinement using GSAS (General Structure Analysis System) software with graphic interface [29]. Section morphology was observed by scanning electron microscope (SEM) (JEOL JSM-6700 F). X-ray photoelectron spectroscopy (XPS) (ESCALAB 250Xi, Thermo Fisher Scientific) was involved to analyze element valence state in samples. Room temperature electron paramagnetic resonance (EPR) spectra were collected by an X-band spectrometer (E500, Bruker, German) operating at

Table 1 Unit-cell parameters and refinement results.

Space group lattice constant (Å) Volume (Å3) Rwp

2

M2

M3

M4

M5

Pm-3m 3.9064 59.6129 4.45%

Pm-3m 3.9067 59.6244 5.91%

Pm-3m 3.9067 59.6230 6.39%

Pm-3m 3.9069 59.7789 6.27%

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Fig. 2. Grain size and lattice constant (a), permittivity and dielectric loss (b) of the sintered ceramics.

Zn2+ ions enter the lattice to form interstitial ions for charge compensation, the cell volume should be obviously expanded, resulting in great changes in lattice constant. Besides, in this case, oxygen vacancies are no longer needed for charge compensation, thus growth of grain size and emergence of holes in grain section should be inhibited. While both XRD and SEM results are opposite to the hypotheses. In M5 ceramics, TiO2 is insufficient, in other words, SrO is excessive. On the one hand, if the Ti vacancies are formed to maintain electrical neutrality, the lattice constant would decrease and the grain growth should be accelerated, which also go against the experimental results. On the other hand, it has been reported that grain size is restricted by excess SrO while promoted by excess TiO2 in SrTiO3 ceramics [31]. In our case, grain size of M5 ceramics is evidently the smallest, indicating that, instead of Ti insufficient, the excess SrO is most probably present and restrains the grain growth. Generally, it is energetically favorable for excess ZnO (M2 ceramics) and excess SrO (M5 ceramics) to exist as secondary phase although they cannot be detected due to the extremely low amount. Appropriate contents of secondary phase are reported to serve as low loss unit dispersed into the ceramic matrix and greatly benefit the decreased dielectric loss [32]. It can be seen from the Fig. 2(b) that M2 ceramics indeed exhibits a very low dielectric loss (∼0.1%), and the dielectric loss of M5 ceramics really decreases compared with a stoichiometric composition SrTi0.985Nb0.015O3 (εr = 4610, tan δ = 0.11) in our

preliminary studies. From the perspective of defect chemistry, it is widely accepted that oxygen vacancies are expected in acceptor substitution and electrons are provided in donor substitution [9,33]. In our case, oxygen vacancies are created because of the incorporation of Zn2+ ions into the host Ti4+ sites, as described in Eq. (1), and electrons are introduced by the Nb5+ donor doping, as shown in Eq. (2). SrTiO3

ZnO → ZnTi′ ′ + Vo•• + Oo× SrTiO3

• Nb2 O5 → 2NbTi + 2e′ + 5Oo×

(1) (2)

More oxygen vacancies should be contained in M2 and M3 ceramics. Mass transports are greatly facilitated by the oxygen vacancies, forming the pores in grain regions and promoting the growth of grain size [34]. Contrarily, less oxygen vacancies are involved in M4 and M5 ceramics, mass transports are suppressed and grain growth is restricted. Permittivity exhibits a similar tendency with the grain size, while the optimum is obtained at x = 4, as shown in Fig. 2(b). Particularly, compared with pure SrTiO3 ceramics (marked as ST in figures, hereafter), the permittivity of M2 ceramics slightly increases while the dielectric loss greatly decreases, implying acceptor substitution is effective in suppressing dielectric loss while incapable in dramatically improving permittivity. The mismatched peak positions between permittivity and grain size indicate that the greatly modified dielectric properties at room

Fig. 3. SEM fracture morphology of the sintered ceramics. 3

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Fig. 4. (a) Raman spectra of the samples measured at room temperature, and inset is local enlarged view. (b) XPS spectrum of Ti 2p peak of M4 ceramics, and insert shows the contrast of Ti 2p3/2 peak for all samples.

Fig. 5. (a) Room temperature EPR spectrum of M4 ceramics. (b) TSDC spectra of M4 ceramics at different polarization electric field.

Therefore, it can be inferred that more Ti3+ ions are involved in M4 ceramics. M4 ceramics exhibits greatly modified dielectric properties and is selected as typical sample for more sophisticated analysis. EPR is a powerful tool to identify the ions with unpaired electronic structure, which is adopted to provide direct evidences of point defect, such as oxygen vacancy and Ti3+ ion in this work. Room temperature EPR spectrum of M4 ceramics based on g-factor is recorded in Fig. 5(a). The g-factors of 1.932 and 1.976 are the signal of typical paramagnetic centre with an electron spin S = 1/2 and an axial crystal field symmetry [37]. In SrTiO3, Ti3+ is a paramagnetic ion with S = 1/2, certainly locating in a sixfold coordination with axial or at least pseudoaxial symmetry. Therefore, the two peaks corresponding to g = 1.932 and 1.976 are assigned to the anisotropic Ti3+ ions with 3d1 electrons trapped on the lattice [37]. The signals at g = 1.982 and 2.003 can be attributed to the electrons trapped on oxygen vacancies [38]. The peak at g = 2.003 is significantly broadened, which means the electronic relaxation time becomes longer. TSDC is an effective approach to establish the relationship between microscopic defect structures and macroscopic electrical properties. To better understand the polarization form of these point defects, TSDC spectra of M4 ceramics measured at various external polarization electric field are collected, as shown in Fig. 5(b). With increase of polarization electric field (as shown by the arrow direction), TSDC peak positions are almost fixed and the peak intensities increase gradually, revealing that the point defects tend to associate with each other and form defect dipoles [39]. It is also energetically favorable for the point defects to aggregate into clusters due to the reduction in distortion energy [40,41]. Besides, the broad TSDC band can be well fitted by several similar peaks, which most probably means that multiple defect dipoles or defect clusters are involved in M4 ceramics and co-contribute to the improved dielectric properties. Temperature stabilities of permittivity and dielectric loss of

temperature mainly result from the defect polarization instead of the interface polarization. With increase of chemical valence state of complex ions, concentration of oxygen vacancy decreases while that of electron increases, and the dielectric loss also increases gradually. High content of oxygen vacancy is beneficial to localization of charge carriers and the low dielectric loss (such as M2 ceramics). Whereas high content of electron not only contributes to polarization, but also results in a high conductivity and therefore a high dielectric loss (such as M5 ceramics). In suitable concentrations of oxygen vacancy and electron, electrons are full localized by the oxygen vacancies and the related defect complexes, which are responsible for the giant permittivity (∼9100) and low tangent loss (∼0.03) in M4 ceramics. To reveal the specific defect structures contributed to the optimized dielectric properties, many efforts have been tried. Raman is employed to detect the local structural symmetry, and room temperature Raman spectra ranged from 200 to 1000 cm−1 are collected in Fig. 4(a). All samples exihibt a similar trend, underlying the same crystal sturcture among them [35]. No first-order Raman modes can be expected in cubic structural SrTiO3 at room temperature. Therefore, the two broad scattering bands located at the regions of 220-450 cm−1 and 590-760 cm−1 are ascribled to the classical second-order Raman modes, as marked in the figure. Besides, appearance of the forbidden first-order Raman modes TO4 (542 cm−1) and LO4 (798 cm−1) overlapped with second order bands, indicates the relaxation of symmetry principles due to the presence of impurity ions and defect in the lattice [36]. Fig. 4(b) is the XPS spectrum of Ti 2p peak of M4 ceramics, in which trace signal of Ti3+ ions is detected, suggesting some Ti4+ ions have been reduced. The insert shows contrast of Ti 2p3/2 peak position for all samples, and the peak of M4 ceramics shift obviously towards lower binding energy compared with other ceramics. As we all know, binding energy is positively correlated with the ionic valence state and electronegativity, and the binding energy of Ti4+ is higher than that of Ti3+ ions.

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Fig. 6. Temperature dependent permittivity and dielectric loss of different ceramics. Inserts in (b), (d) and (f) are the local enlarged scales.

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Fig. 7. (a) Complex impedance spectroscopy of samples measured at 500 °C. Inserts are the equivalent circuit and local magnification view. (b) Frequency dependent conductivity of samples measured at 500 °C.

observed around −40 °C and the activation energy is confirmed as 0.42 eV, which can be ascribed to the coupling effect of the conduction electrons with the motion of the off-center B site ions [40]. Individual hopping motion of the B site ions makes the host lattice being polarized and increases the underlying dipole moment, which are reported to bring about the dielectric relaxation at higher frequency (∼1010 Hz) [42]. The dielectric relaxation at high temperature (> 300 °C) most probably results from interface polarization, such as internal barrierlayer capacitor effects between grain and grain boundary, and/or Schottky barrier depletion layer effect on the sample-electrode interface, as reported in rutile titanium dioxide [43,44]. Actually, it is also insufficient to completely exclude Maxwell-Wagner effect at the present limited frequency range, while the effect usually can be neglected when the temperature is not very high [43]. To better understand the electrical properties of samples, complex impedance and conductivity are characterized and plotted. Fig. 7(a) is the complex impedance spectroscopy measured at 500 °C. All spectra can be well simulated by two series connected parallel circuits, representing the grain and grain boundary, separately. Each parallel circuit includes a resistance, a capacitance and a CPE (constant phase element), which is adopted to correct the simulation results, as shown in the insert. Insulativity of samples can be roughly evaluated by comparing the semicircle radius of them. It can be inferred that insulativity of M3 ceramics is the best, while those of other samples are actually close. Another insert is the local enlarged scale, from which nonzero abscissa intercept of ST, M2 and M3 ceramics are observed, indicating the better insulativity at high frequency segment. In the insert, the incomplete small semicircles corresponding to the grains in ST and M2 ceramics may result from the limited measurement frequency range. Besides, it is worth noting that M2 ceramics have higher grain resistance than pure ST ceramics, which can be the reason why the M2 ceramics exhibit an extremely low dielectric loss. Fig. 7(b) presents the frequency dependent conductivity at 500 °C. All conductivities generally increase with the increase of frequency, indicating the observed conductivities are related to the hopping behavior of localized carriers [45]. Moreover, conductivity of M3 ceramics is obviously the lowest while that of M4 ceramics is the slightly higher compared with other three samples, which is consist with the results of complex impedance. Defect dipoles or defect clusters benefit the dielectric properties. However, at strong external incentives, such as high temperature and/ or high applied electric field, these defect dipoles or clusters seem to be metastable and will become carriers [40], resulting in high conductivity. Therefore, it is inferable that more defect dipoles or clusters are included in M4 ceramics and responsible for the strengthened conductivity at high temperature.

different ceramics are given in Fig. 6. In first sight, temperature dependent dielectric properties of M2 ceramics exhibit a similar various tendency with those of pure SrTiO3 ceramics. However, some distinct differences can be observed after careful comparison. Overall, the permittivity of M2 ceramics increases while the dielectric loss (from -150 to 200 °C) decreases compared with those of pure SrTiO3 ceramics. Besides, the relaxation peaks in M2 ceramics are appropriately pushed to higher temperature direction. As a result, an extremely low dielectric loss (< 0.001) is achieved in M2 ceramics at a wide temperature range (Fig. 6(d)). Here, it should be explanation that, part of small negative dielectric loss data at higher frequencies in Fig. 6(b) and (d) are omitted, which may due to the fact that these dielectric loss are too low (very close to zero) to be precisely measured. In M3 ceramics, the relaxation peaks continue moving towards high temperature direction and further broaden the thermal stable region. Moreover, a set of relaxation peaks (marked as Peak A) is observed around temperature range of -150-50 °C, as shown in the insert of Fig. 6(f). The similar relaxation peaks can also be found in the permittivity and dielectric loss spectra of M4 and M5 ceramics. To further understand the relaxation mechanism, the activation energies are calculated by the Arrhenius formula: fr =f0 exp(-Ea/kT)

(3)

Where fr is the dielectric relaxation frequency, f0 is the pre-exponential factor, Ea is the activation energy for relaxation, k is the Boltzmann constant, and T is the measurement temperature (K). In M3 ceramics, the activation energy relaxation peak of Peak A is confirmed as 0.24 eV. The similar phenomenon has also been observed in Sr1-1.5xBixTiO3 ceramics and attributed to first-ionization of oxygen vacancies [11]. In M4 ceramics, there are three obvious dielectric relaxation peaks. Due to the similar appearance temperature range and the close activation energy (0.21 eV) compared with M3 ceramics, the first one can be assigned to Peak A and believed to originate from the same mechanism. The second one (Peak B) located around 80 °C with relaxation activation energy Ea=0.64 eV may result from second ionization of oxygen vacancies, according to the literature [40]. Creations of oxygen vacancy are usually accompanied by the electron charge density redistribution. Here, two kinds of electron transport processes can be expected. One is that electrons transfer from ionized oxygen vacancies to Ti4+ ions, as assigned to Peak A and Peak B. Anther process is the hopping movement of electrons between Ti3+ and Ti4+ ions among different B sites, i.e. reorientation of Ti3+ ions with respect to oxygen vacancies [42]. Relaxation activation energy of the third one (Peak C) is 0.75 eV. Ang et al. have reported the motions of oxygen vacancy related defect complex or defect cluster are characterized by activation energy of 0.74-0.86 eV [11], thus Peak C could be associated with this relaxation mechanism. These local polar structures interact with each other via the highly polarizable host crystal lattice. In M5 ceramics, except the appearance of Peak A, Peak B and Peak C, another relaxation peak (Peak D) is

4. Conclusions For the first time, defect engineering and the resulting effects on the 6

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phase structures, microstructures, defect structures and dielectric properties of SrTiO3 ceramics modified by complex ions (Nb, Zn)x+ (x = 2, 3, 4 and 5) have been systematically investigated. Results show that, chemical valence states of the complex ions have significant influences on the structures and dielectric behaviors. With increase of the valence states, the lattice constant increases while the grain size slightly increases and then greatly decreases. The permittivity achieves the maximum value in M4 ceramics whereas the dielectric loss continues increasing. Acceptor co-doping is beneficial to low dielectric loss (≤0.001), high insulativity and excellent temperature stability. Equivalent co-doping leads to giant permittivity and low dielectric loss. Donor co-doping results in improved permittivity as well as increased tangent loss. Further investigations indicate that, more point defects (such as Ti3+ and oxygen vacancy etc.) are involved and associated with each other to form various defect dipoles or defect clusters, cocontributing to the greatly modified dielectric properties in M4 ceramics. The findings offer a completely novel approach to engineer new electronic functionality.

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This work was supported by the NSFC-Guangdong Joint Funds of the Natural Science Foundation of China (No. U1601209), National Key Basic Research Program of China (973 Program) (No. 2015CB654601), Technical Innovation Special Program of Hubei Province (2017AHB055), State Key Laboratory of Advanced Technology Materials Synthesis and Processing (Wuhan University of Technology) (2018-KF-11), National Natural Science Foundation of China (51872213) and the Fundamental Research Funds for the Central Universities (2019-YB-010).

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