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Colossal permittivity and dielectric relaxation of (Li, In) Co-doped ZnO ceramics
Accepted Manuscript Colossal permittivity and dielectric relaxation of (Li, In) Co-doped ZnO ceramics Dong Huang, Zhifu Liu, Yongxiang Li, Yun Liu PII:
S0925-8388(16)34031-2
DOI:
10.1016/j.jallcom.2016.12.113
Reference:
JALCOM 40035
To appear in:
Journal of Alloys and Compounds
Received Date: 24 June 2016 Revised Date:
26 November 2016
Accepted Date: 9 December 2016
Please cite this article as: D. Huang, Z. Liu, Y. Li, Y. Liu, Colossal permittivity and dielectric relaxation of (Li, In) Co-doped ZnO ceramics, Journal of Alloys and Compounds (2017), doi: 10.1016/ j.jallcom.2016.12.113. 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 Colossal Permittivity and Dielectric Relaxation of (Li, In) Co-doped ZnO Ceramics Dong Huang a, b, c , Zhifu Liu a, *, Yongxiang Li a, b, *, Yun Liu d a
CAS Key Lab of Inorganic Functional Materials and Devices, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China School of Physical Science and Technology, ShanghaiTech University, Shanghai 201210, China c
d
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b
Department of Physics, The University of Hong Kong, Hong Kong, China
Research School of Chemistry, The Australian National University, ACT 2601, Australia
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Abstract:
In this study, a colossal permittivity up to 3800 and a low dielectric loss of 0.11 at 1 kHz
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have been obtained from the (Li, In) co-doped ZnO ceramic [Zn(1-2x)(Li, In)xO] when x was 0.5%. Electric modulus spectroscopy and impedance analysis were used to investigate the origin of its high permittivity. Two relaxation peaks and a dielectric anomaly were observed in the temperature range of 293-363 K. According to the Debye relaxation theory, the lowand high- temperature relaxation peaks with activation energies of 0.09 eV and 0.29 eV, have been attributed to the hopping of singly and doubly charged oxygen vacancies, which are
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created by lithium and indium ions doping and oxygen deficiency during sintering process. After thermal treatment in an oxidizing atmosphere, the peaks related to the singly and doubly charged oxygen vacancies disappear and the permittivity reduces to ~460 at
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room-temperature. From the X-ray photoelectron spectra (XPS), the concentrations of oxygen vacancies decrease after the O2-annealing process. The results reveal that the oxygen defects would be the main origin of the colossal permittivity of co-doped ZnO at room-temperature
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range.
Keywords: Colossal Permittivity, Dielectric Relaxation, Co-doped ZnO, Oxygen Vacancy.
*
Corresponding author emails: [email protected]; [email protected] 1
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Introduction Dielectric materials with colossal permittivity (CP) have drawn considerable attention
due to the emerging requirement of electronic device miniaturization and energy storage system, such as the multi-layer ceramic capacitors (MLCC) and ceramic power capacitors. To
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date, several kinds of materials with CP (relative permittivity larger than 103) have been discovered [1-5]. In BaTiO3-based perovskite materials, the CP phenomenon is attributed to two mechanisms, namely the ferroelectric effect and the internal barrier layer capacitance
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(IBLC) effect [1]. Although controversial, in materials like CaCu3Ti4O12 [2], (Li, Ti) co-doped NiO [3] and the copper oxide [4], IBLC effect seems to be the most plausible
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explanation. Recently, a new dielectric polarization mechanism, e.g. electron-pinned defect dipoles (EPDD), has been proposed to explain the CP phenomenon of In (acceptor) and Nb (donor) co-doped TiO2 [5]. Such CP strongly relies on the defect formation, local and average structure in each individual material [6]. CP phenomenon has also been observed in the doped
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NiO and transition-metal oxide CuO, which is explained by the mechanism different from that of co-doped TiO2. In the quest to discover new high performance CP dielectric materials, co-doping in simple oxides presents a promising approach with much scope and potential for
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providing a deep understanding of the CP mechanism.
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Recently, it has been reported that, in Mg-doped ZnO ceramics, the colossal permittivity can reach ~104 due to the IBLC mechanism, indicating that ZnO could be another promising material of colossal permittivity [7]. Wurtzite-type ZnO has a tetrahedral structure where Zn atoms are coordinated to four O atoms, and the d-electron of Zn atom hybridizes with the p-electron of oxygen atom. Much of the focus of ZnO has been on its semiconducting properties due to its n-type semiconducting nature originating from native defects. As for the non-divalent dopants, it is common to result in the instability of ionic charge distribution in the lattice. Furthermore, this charge unbalance would retard the dopant’s further incorporation 2
ACCEPTED MANUSCRIPT and create more native defects. So in terms of charge compensation, the donor and acceptor co-doping model is adopted to meet the charge balance that the dopants consist of one monovalent metal oxide and one trivalent metal oxide [8]. The n-type conductivity can be further enhanced by introducing donor dopants, e.g. Co, Mn, Al and In [9-12].
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First-principles calculation indicates that the carrier in In-doped ZnO is In′Zn [13, 14]. P-type ZnO could also be obtained with appropriate dopants to create holes in the materials [15], including substitution at the Zn-site with group-I elements (Li, Na, or K) [16-18] or O-site
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with group-V elements (N, P, As, or Sb) [19-22]. All these indicate that Li and In could be successfully incorporated into ZnO lattice. The co-doping of donor element, like In, and
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acceptor element, like Li, may lead to unique dielectric phenomenon. The relatively few studies on the exciting potential of doped ZnO ceramics and their dielectric behavior, especially regarding donor and acceptor co-doping, have stimulated the authors to study the dielectric properties of co-doped ZnO ceramics.
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In this work, (Li, In) co-doped ZnO ceramics have been synthesized. The influence of lithium and indium co-doping on the structural, morphological and dielectric properties of
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ZnO ceramics were investigated. CP was obtained from co-doped ZnO ceramics and the origin of CP phenomenon is discussed. Experimental details
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2.
Reagent-grade (≥ 99.99%) starting materials of Li2CO3, In2O3 and ZnO (all from Aladdin, China) were weighted according to the stoichiometric proportion of Zn(1-2x)(Li, In)xO (x = 0.25%, 0.5%, 0.75%, 1.00%, 1.50%, and 2.00%, respectively). Raw materials were mixed in ethanol by ball milling with Zirconia balls in a polypropylene bottle for 6 h at 400 rpm. The mixture was then dried and calcined at 1323 K for 2 h. Subsequently, the calcined sintered mixture was ball-milled again for 6 h. After drying, 6 wt% polyvinyl acetate (PVA) binder was added and the mixture granulated by sieving through a 60-mesh screen to produce 3
ACCEPTED MANUSCRIPT starting powders. The powders were uniaxially pressed into discs (13 mm in diameter and 1.5 mm in thickness) in a stainless steel die at a pressure of 100 MPa. Parts of the samples were pressed in a cold isostatic pressing (CIP) machine at a pressure of 200 MPa for 2 min. Sintering was then carried out at different temperatures (1498 K, 1523 K, 1548 K, 1573 K)
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for 12 h with a heating rate of 3 oC/min. The samples were allowed to cool down naturally to the ambient temperature.
The doped ZnO and un-doped ZnO powder were sintered at 1548 K for 12h and mixed
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with 20 wt% pure silicon powders. The crystal structures of samples were investigated via
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X-ray diffraction (XRD) with a Cu-Kα radiation (Bruker D8 Advance, Germany). Surfaces of samples were lapped and ground with SiC powders, then polished with 0.1 µm-Al2O3 powders to a mirror-like surface. The polished samples were thermally etched for 2 h at 1173 K. The microstructure was examined by a Supra 55 (Carl Zeiss, Germany) scanning electron microscope (SEM) equipped with an energy dispersive spectrometer (EDS, X-MaxN, Oxford,
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UK). Electrodes were made by printing silver paste on both sides of the disk-type samples and then fired at 973 K for 2 h in order to remove the polymeric component. The dielectric properties of the ceramics were measured by a WK-6500B Precision Impedance Analyzer
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(Wayne-Kerr Electronics Ltd., London, UK) in the frequency range of 20-106 Hz at a
3.
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temperature range from 150 K to 400 K, respectively. Results and discussion Fig. 1 shows the XRD pattern of Zn1-2x(Li, In)xO powders sintered at 1548 K. The pure silicon (Si) powder is used as reference for XRD measurement. The (Li, In)-doped ZnO diffraction peaks can be indexed to the wurtzite ZnO (JPCDS No. 36-1451) with a hexagonal structure of P63mc space group. No peak of secondary phases was observed in the XRD patterns of all the samples. The XRD peak located between 36° and 37°, attributed to the (101) plane, firstly shifts slightly to low angles with increasing doping amount. The peak shifts to 4
ACCEPTED MANUSCRIPT high angles when the doping amount exceeds 1%. The radii of Zn2+, Li+, and In3+ are 0.074 nm, 0.068 nm, and 0.094 nm, respectively. Since one Li+ and one In3+ may replace 2Zn2+ in terms of valence balance, the average radius is about 0.081 nm, which is larger than that of Zn2+. The substitution of Zn2+ by the dopants Li+ and In3+ would lead to the extending of the
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lattice parameter, which results in the shift of (101) peak to lower angle. When the doping amount x exceeds 1%, part of the In3+ dopant may exist in the form of In2O3 and form solid solution with ZnO [23]. Since the amount of In2O3 is very small, its diffraction peaks cannot
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be observed. However, its existence may encompass the ZnO lattice and leads to the shift of (101) peak to higher angle [24]. On the other hand, Li+ can still substitute Zn2+ when the
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doping concentration is higher than 1%. The substitution of Zn2+ by Li+ would lead to the decrease of the lattice parameter. As a result, the (101) diffraction peak shifts to higher angle [25].
Fig. 2 shows SEM micrographs of surfaces of Zn1-2x(Li, In)xO ceramics. The grain
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structure is relatively homogeneous throughout the sample, and grain boundaries are clearly seen. The ceramics are very dense, with grain sizes in the range of 8 to 14 µm. As the doping content of Li and In increases, the grain size decreases gradually, with more small grains
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appearing at the grain-boundary triple junctions. EDS surface scanning analysis shows that
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the In3+ is homogeneously distributed (Fig. 2h). However, Li+ cannot be discussed as it is undetectable in EDS.
Fig. 3a shows the dependence of permittivity and dielectric loss on the doping concentration. A high relative permittivity of 3800 and a comparatively low dielectric loss (tg δ) of about 0.12 can be obtained when the doping concentration x = 0.5%. The permittivity of co-doped ZnO increases when the sintering temperature increases from 1500 to 1600 K. Meanwhile, the dielectric loss increases sharply as well. Consequently, the sintering temperature of 1548 K is regarded as optimum because all the samples can be sintered at this 5
ACCEPTED MANUSCRIPT temperature whilst having relatively low dielectric loss. In order to understand the dielectric properties caused by the defects, the frequency dependences of permittivity and dielectric loss at various temperatures from 293 K to 363 K were measured as shown in Fig. 4a. The oblique line appearing in dielectric loss curves at
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frequency below 100 Hz, is attributed to the electrode polarization, which indicates the formation of Schottky barrier at the sample-electrode interface [26]. However, the peaks in the high-frequency region can’t be accurately clarified due to the existence of the background
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interference. Notably, when the temperature increases to 363 K, the peak is almost merged by the background.
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The peaks of the defect relaxation can be clearly clarified in Modulus-Frequency spectra with the complex dielectric constant converted into its complex modulus (M' + iM'') [27-29]. The solid line in Fig. 4c indicates the fitted data, from which the activation energy was calculated and found to be 0.29 eV for the high temperature region (HTR), i.e. above 333 K,
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and 0.09 eV for the low temperature region (LTR), i.e. below 333 K. The peaks shift to higher temperature when the measuring frequency increases, indicating a thermally activated Debye-like relaxation. The relaxation parameters can be calculated according to the
Ea
k BT
)
(1)
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f = f 0 exp(−
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Arrhenius law as follows:
where f0 is pre-exponential factor, Ea is the activation energy of defect relaxation, kB is the Boltzmann constant, and T is absolute temperature. In the Arrhenius plot, the deviation shows and the presence of different activation energies for the different temperature regions are due to the change of relaxation mechanisms. For this phenomenon, a similar anomaly of dielectric relaxation was found in ZnO ceramics, in which the activation energy of which are 0.30 eV and 0.09 eV [30, 31]. One peak with the activation energy of 0.29 eV (Peak-2) has been reported in a few literatures, including in ZnO 6
ACCEPTED MANUSCRIPT film, which is caused by oxygen vacancy [30-32]. The other one (Peak-1, 0.09 eV) has been attributed to the interstitial oxygen [30, 31]. What’s more, these peaks would disappear after O2-annealing process in Fig. 4d. However, the interstitial oxygen is very high in formation energy and electrically inactive, according to the first principles calculation [33]. If Peak-1
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were associated with interstitial oxygen, the peak would still exist after the O2-annealing process. In O2-annealing process, excess oxygen atoms in the ZnO lattice can be accommodated in the form of oxygen interstitials. However, the result is totally inconsistent
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with the expectation as Fig. 4d shows. So it is concluded that this peak would be also related to the oxygen vacancy in the form of oxygen vacancy clusters. It has been reported that the
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energy of moving these clusters is less than the energy of moving individual oxygen vacancies [34-36].
Oxygen vacancies can trap charge carriers [34], such as the electrons. The electrons are localized in oxygen vacancies and can jump between the ionized oxygen vacancies, giving
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rise to a dipolar effect [36-38]. It is well-known that oxygen vacancies contribute to polarization in the forms of singly and doubly charged states in low- and high- temperature ranges. Two relaxation peaks have been reported to be respectively due to the hopping
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motions of singly and doubly charged oxygen vacancies in CaNb2O6 [39], La2Ti2O7 [26] and
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TiO2 [40]. The temperature where oxygen vacancies change from singly charged state to doubly charged state were reported to be about 800 K in SrTiO3 [41] and 700 K in NiTiO3 [42]. Vanheusden et al. have discovered that, in the pure ZnO, there is the transition between singly charged oxygen vacancies and doubly charged oxygen vacancies. It was demonstrated that the density of VO• increased as the temperature went down. In other words, the density of free carriers would increase as the temperature goes up, according to the equation: VO• → VO•• + e′ .
Vanheusden has discovered that at 294 K, the densities of VO• and free-carriers
has much differences [43], which indicates that the VO• → VO•• + e′ happens. So in this 7
ACCEPTED MANUSCRIPT discussion, the Peak-1 (0.09 eV) may be associated with the singly oxygen vacancy. With increasing temperature above 333 K, the weakly bonded electrons can be thermally activated becoming free ones, leading to the formation of doubly charged oxygen vacancy. The defect equation is shown below:
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333 K VO• ← →VO•• + e′
(2)
Regarding to the source of the singly and doubly charged oxygen vacancies, there are two main ways. Firstly, the oxygen vacancies could be created with the introduction of Li+ and
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In3+ ions into Zn-sites. Li+ ions would create the oxygen vacancies. However, Li+ doping into the Zn-site would enhance the Madelung energy by 13.56 eV, resulting the instability of ionic
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charge distribution in lattice [44]. For charge-compensation, the Li+ ions’ would occupy the interstitial sites. In this way, this would reduce the Li+ ions’ solubility. However, the introduction of In3+ into Zn2+-sites are considered as donors, which can reduce the Madelung energy by 9.73 eV [44]. The formation of the defect InZn• would be near the LiZn′ sites,
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which enhance the stability of the ionic charge distribution. So the In3+ ions doping can enhance the incorporation of acceptor Li+ ions because of the strong acceptor-donor attractive interaction, which indicates more oxygen vacancies would be created. Secondly, during
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sintering process at high temperature with low oxygen partial pressure, the oxygen vacancies
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could be generated due to the loss of oxygen from the crystal lattice. The defect reaction equations are shown as follows: ′ + VO•• + 3OO Li2O + 2ZnO → 2 LiZn
(3)
In2 O3 + 2 ZnO → 2 In Zn• + VZn′′ + 5OO
(4)
OO ↔
1 O2 + VO•• + 2e′ 2
(5)
In addition, after the O2-annealing process, the permittivity and the dielectric loss decrease dramatically to 460 and 0.04 respectively at 1 kHz in the inset of Fig. 4a, which 8
ACCEPTED MANUSCRIPT indicates the dielectric property is mainly associated with the oxygen vacancies. To verify that the O2-annealing process strongly affects the oxygen defects in the (Li, In) co-doped ZnO system, X-ray photoelectron spectra (XPS) of the as-prepared and O2-annealed sample are shown in Fig. 5. The XPS spectra for O 1s peaks depict a wide and asymmetric peak that is
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further deconvoluted into three peaks centered at different binding energies indicating the presence of different O species. Generally speaking, the O1s peak has been observed in the binding energy (BE) region of 529-535 eV. The peak around 529-530 eV has been attributed
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to lattice oxygen and the peak around 530.7-531.6 eV has been attributed to oxygen in the region of non-stoichiometric oxides. Furthermore, for the moisture, its binding energy lies
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between 531.5 eV (OH) and 533 eV (H2O) is due to the chemisorbed oxygen [45]. The scan of O1s spectra in Fig. 5a and 5b exhibit a peak of OІ (530.08 eV) which is attributed to the oxidized metal ions in the samples, viz O-In and O-Zn, in the lattice [46], a peak of OII (530.95 eV) is caused by the O deficiency and a peak of OIII (532.01 eV) is due to the
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surface adsorbed oxygen species such as OH or O2 [45, 47]. It is noted that the intensity ratio of the OII/OI decreases from 0.62 to 0.55 after O2-annealing process, which indicates that the concentrations of oxygen vacancies decrease after O2-annealing process [48].
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It is well known that the low- and high- frequency semicircles in the Nyquist plot
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represent the different relaxations incited by the intergrain and intragrain effect, respectively [49]. In Fig. 6a, it can be seen that only one semicircle exists, and it is remarkably depressed by the bias voltage. This fact demonstrates that the dielectric properties are mainly due to the grain-boundary effect. After thermal treatment in an oxygen atmosphere, the relaxation peaks of both VO•• and V O• disappear. Besides, the Nyquist semicircle of O2-annealed sample cannot be affected by the bias voltage as shown in Fig. 6b. The permittivity of the ZnO ceramics decreases to ~460, and the dielectric loss decreases to 0.04 after thermal treatment in oxygen. The colossal permittivity of the (Li, In) co-doped ZnO ceramics has a connection 9
ACCEPTED MANUSCRIPT with these two defects VO•• and V O• , which contribute to the formation of grain boundary relaxation at room-temperature. The movement of defect charges would generate a micro-capacitance according to the grain boundary-layer mechanism, which is responsible for
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the colossal permittivity of the specimen.
Conclusions
Li+ and In3+ co-doped ZnO ceramics were prepared and their structural and dielectric
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properties were investigated. Li+ and In3+ have been successfully incorporated into the ZnO
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lattice. The dopant is distributed homogeneously in the ceramics and no obvious agglomeration has been found. CP is obtained from the co-doped ZnO ceramics and a permittivity of up to 3800 is obtained from Zn0.99(Li, In)0.005O ceramics. The dielectric modulus and impedance spectroscopy analysis indicate that the dielectric relaxation in the co-doped ZnO is associated with singly and doubly charged oxygen vacancies. The oxygen
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vacancies are mainly created with the introduction of lithium and indium ions into Zn-site and the sintering process. Furthermore, after thermal treatment in an oxidizing atmosphere for 2 h, the relaxation peaks of singly and doubly charged oxygen vacancies defects disappear
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and the permittivity reduces to ~460 at room-temperature. From the X-ray photoelectron
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spectra, the concentrations of oxygen vacancies decrease after O2-annealing process. It can be concluded that oxygen vacancies are the main origin of the CP phenomenon in Zn0.99(Li, In)0.005O ceramics at room-temperature range.
Acknowledgements The authors acknowledge the financial support from the Ministry of Science and Technology of China through the program of International S&T Cooperation (S2015ZR1108), 973-project (2015CB654604, 2015CB654605), Natural Science Foundation of China 10
ACCEPTED MANUSCRIPT (51572279) and the CAS/SAFEA International Partnership Program for Creative Research Teams. The authors would like to thank Dr. Adrian Trinchi of CSIRO, Materials Science & Engineering, Australia for his valuable discussion and kind help. References
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[34] X. Wang, X. Lu, C. Zhang, X. Wu, W. Cai, S. Peng, H. Bo, Y. Kan, F. Huang, J. Zhu, Oxygen-vacancy-related high-temperature dielectric relaxation in SrTiO3 ceramics, Journal of Applied Physics, 107 (2010) 114101. [35] C.C. Wang, C.M. Lei, G.J. Wang, X.H. Sun, T. Li, S.G. Huang, H. Wang, Y.D. Li, Oxygen-vacancy-related dielectric relaxations in SrTiO3 at high temperatures, Journal of Applied Physics, 113 (2013) 094103. [36] W. Li, K. Chen, Y. Yao, J. Zhu, Y. Wang, Correlation among oxygen vacancies in bismuth titanate ferroelectric ceramics, Applied Physics Letters, 85 (2004) 4717. [37] X. Sun, C. Wang, G. Wang, C. Lei, T. Li, L. Liu, A. Feteira, Low-temperature dielectric relaxations associated with mixed-valent structure in Na0.5Bi0.5Cu3Ti4O12, Journal of the American Ceramic Society, 96 (2013) 1497-1503. [38] W. Bai, G. Chen, J.Y. Zhu, J. Yang, T. Lin, X.J. Meng, X.D. Tang, C.G. Duan, J.H. Chu, Dielectric responses and scaling behaviors in Aurivillius Bi6Ti3Fe2O18 multiferroic thin films, Applied Physics Letters, 100 (2012) 082902. [39] H. Li, C. Wang, H. Wang, D. Zhang, N. Zhang, Low-frequency dielectric properties of CaNb2O6 ceramics over a wide temperature range, Ceramics International, 41 (2015) 14773-14779. 12
ACCEPTED MANUSCRIPT [40] C. Wang, N. Zhang, Q. Li, Y. Yu, J. Zhang, Y. Li, H. Wang, N. Alford, Dielectric relaxations in rutile TiO2, Journal of the American Ceramic Society, 98 (2015) 148-153. [41] C. Lee, J. Destry, J.L. Brebner, Optical absorption and transport in semiconducting SrTiO3, Physical Review B, 11 (1975) 2299-2310. [42] T. Li, C.C. Wang, C.M. Lei, X.H. Sun, G.J. Wang, L.N. Liu, Conductivity relaxation in NiTiO3 at high temperatures, Current Applied Physics, 13 (2013) 1728-1731. [43] K. Vanheusden, C.H. Seager, W.L. Warren, D.R. Tallant, J.A. Voigt, Correlation between photoluminescence and oxygen vacancies in ZnO phosphors, Applied Physics Letters, 68 (1996) 403.
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[44] T. Yamamoto, H. Katayama-Yoshida, Physics and control of valence states in ZnO by codoping method, Physica B: Condensed Matter, 302–303 (2001) 155-162. [45] J.F. Moulder, J. Chastain, R.C. King, Handbook of X-ray photoelectron spectroscopy: a reference book of standard spectra for identification and interpretation of XPS data, Physical Electronics Eden Prairie, MN1995.
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[46] M. Ahmad, H. Sun, J. Zhu, Enhanced photoluminescence and field-emission behavior of vertically well aligned arrays of In-doped ZnO Nanowires, ACS applied materials & interfaces, 3 (2011) 1299-1305. [47] K.G. Saw, N.M. Aznan, F.K. Yam, S.S. Ng, S.Y. Pung, New insights on the Burstein-Moss shift and band gap narrowing in indium-doped zinc oxide thin films, PloS one, 10 (2015) e0141180.
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[48] S.K. Gupta, R. Jindal, A. Garg, Microscopic investigations into the effect of surface treatment of cathode and electron transport layer on the performance of inverted organic solar cells, ACS Applied Materials & Interfaces, 7 (2015) 16418-16427.
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[49] C. Wang, L. Liu, Q. Li, S. Huang, J. Zhang, J. Zheng, C. Cheng, Revisiting the low-temperature dielectric properties of ZnO, Journal of Applied Physics, 116 (2014) 124101.
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Captions of Figures
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Fig. 1. The XRD patterns of Zn1-2x(Li, In)xO powders: x=0%, 0.25%, 0.5%, 0.75%, 1%, 1.5%, 2%, with Si as a reference.
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Fig. 2. The SEM micrographs of the as-sintered Zn1-2x(Li, In)xO ceramics after annealing process at 1273 K for 30 min with (a) x=0.25%, (b) 0.5%, (c) 0.75%, (d) 1%, (e) 1.5%, and (f) 2%; (g) the SEM micrograph of the as-sintered Zn0.99(Li, In)0.005O ceramic with the annealing process at 1173 K for 2 h; (h) the lateral scanning of the Indium distribution in ZnO0.99(Li, In)0.005O ceramic without the annealing process.
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Fig. 3. (a) the dielectric properties of Zn1-2x(Li, In)xO with x= 0.25%, 0.5%, 0.75%, 1%, 1.5%, 2% samples sintered at 1548 K; (b) the dielectric properties of Zn0.99(Li, In)0.005O sample sintered at the temperature ranging from 1498 K to 1573 K.
Fig. 4. The dielectric properties of Zn0.99(Li, In)0.005O ceramics at a measuring temperature range from 293 K to 363 K: (a) and the inset are the permittivity- and dielectric lossfrequency spectrum; (b) the M"-frequency spectrum; (c) the Arrhenius plots of peak 3 and peak 4; (d) the M"-frequency spectrum of the O2-annealed sample (973 K/2 h).
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Fig. 5. X-ray photoelectron spectrum of O1s (a) in the as-prepared and (b) O2-annealed sample
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Fig. 6. (a) the Nyquist plots recorded at room temperature under different dc bias, the inset is the equivalent circuit used to represent the electric properties of grain-boundary interface effect in ZnO ceramics; (b) the Nyquist plots recorded at room temperature under different dc bias after O2-annealing process at 973 K for 2 h.
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Highlights:
The (Li, In) co-dopants have been successfully incorporated into ZnO
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lattice. Colossal permittivity (CP) of 3800 is obtained for (Li, In) co-doped ZnO ceramics.
••
•
and V O .
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of VO
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The CP originates from the grain-boundary effect with defect dipoles
S0925-8388(16)34031-2
DOI:
10.1016/j.jallcom.2016.12.113
Reference:
JALCOM 40035
To appear in:
Journal of Alloys and Compounds
Received Date: 24 June 2016 Revised Date:
26 November 2016
Accepted Date: 9 December 2016
Please cite this article as: D. Huang, Z. Liu, Y. Li, Y. Liu, Colossal permittivity and dielectric relaxation of (Li, In) Co-doped ZnO ceramics, Journal of Alloys and Compounds (2017), doi: 10.1016/ j.jallcom.2016.12.113. 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 Colossal Permittivity and Dielectric Relaxation of (Li, In) Co-doped ZnO Ceramics Dong Huang a, b, c , Zhifu Liu a, *, Yongxiang Li a, b, *, Yun Liu d a
CAS Key Lab of Inorganic Functional Materials and Devices, Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China School of Physical Science and Technology, ShanghaiTech University, Shanghai 201210, China c
d
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b
Department of Physics, The University of Hong Kong, Hong Kong, China
Research School of Chemistry, The Australian National University, ACT 2601, Australia
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Abstract:
In this study, a colossal permittivity up to 3800 and a low dielectric loss of 0.11 at 1 kHz
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have been obtained from the (Li, In) co-doped ZnO ceramic [Zn(1-2x)(Li, In)xO] when x was 0.5%. Electric modulus spectroscopy and impedance analysis were used to investigate the origin of its high permittivity. Two relaxation peaks and a dielectric anomaly were observed in the temperature range of 293-363 K. According to the Debye relaxation theory, the lowand high- temperature relaxation peaks with activation energies of 0.09 eV and 0.29 eV, have been attributed to the hopping of singly and doubly charged oxygen vacancies, which are
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created by lithium and indium ions doping and oxygen deficiency during sintering process. After thermal treatment in an oxidizing atmosphere, the peaks related to the singly and doubly charged oxygen vacancies disappear and the permittivity reduces to ~460 at
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room-temperature. From the X-ray photoelectron spectra (XPS), the concentrations of oxygen vacancies decrease after the O2-annealing process. The results reveal that the oxygen defects would be the main origin of the colossal permittivity of co-doped ZnO at room-temperature
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range.
Keywords: Colossal Permittivity, Dielectric Relaxation, Co-doped ZnO, Oxygen Vacancy.
*
Corresponding author emails: [email protected]; [email protected] 1
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Introduction Dielectric materials with colossal permittivity (CP) have drawn considerable attention
due to the emerging requirement of electronic device miniaturization and energy storage system, such as the multi-layer ceramic capacitors (MLCC) and ceramic power capacitors. To
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date, several kinds of materials with CP (relative permittivity larger than 103) have been discovered [1-5]. In BaTiO3-based perovskite materials, the CP phenomenon is attributed to two mechanisms, namely the ferroelectric effect and the internal barrier layer capacitance
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(IBLC) effect [1]. Although controversial, in materials like CaCu3Ti4O12 [2], (Li, Ti) co-doped NiO [3] and the copper oxide [4], IBLC effect seems to be the most plausible
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explanation. Recently, a new dielectric polarization mechanism, e.g. electron-pinned defect dipoles (EPDD), has been proposed to explain the CP phenomenon of In (acceptor) and Nb (donor) co-doped TiO2 [5]. Such CP strongly relies on the defect formation, local and average structure in each individual material [6]. CP phenomenon has also been observed in the doped
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NiO and transition-metal oxide CuO, which is explained by the mechanism different from that of co-doped TiO2. In the quest to discover new high performance CP dielectric materials, co-doping in simple oxides presents a promising approach with much scope and potential for
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providing a deep understanding of the CP mechanism.
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Recently, it has been reported that, in Mg-doped ZnO ceramics, the colossal permittivity can reach ~104 due to the IBLC mechanism, indicating that ZnO could be another promising material of colossal permittivity [7]. Wurtzite-type ZnO has a tetrahedral structure where Zn atoms are coordinated to four O atoms, and the d-electron of Zn atom hybridizes with the p-electron of oxygen atom. Much of the focus of ZnO has been on its semiconducting properties due to its n-type semiconducting nature originating from native defects. As for the non-divalent dopants, it is common to result in the instability of ionic charge distribution in the lattice. Furthermore, this charge unbalance would retard the dopant’s further incorporation 2
ACCEPTED MANUSCRIPT and create more native defects. So in terms of charge compensation, the donor and acceptor co-doping model is adopted to meet the charge balance that the dopants consist of one monovalent metal oxide and one trivalent metal oxide [8]. The n-type conductivity can be further enhanced by introducing donor dopants, e.g. Co, Mn, Al and In [9-12].
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First-principles calculation indicates that the carrier in In-doped ZnO is In′Zn [13, 14]. P-type ZnO could also be obtained with appropriate dopants to create holes in the materials [15], including substitution at the Zn-site with group-I elements (Li, Na, or K) [16-18] or O-site
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with group-V elements (N, P, As, or Sb) [19-22]. All these indicate that Li and In could be successfully incorporated into ZnO lattice. The co-doping of donor element, like In, and
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acceptor element, like Li, may lead to unique dielectric phenomenon. The relatively few studies on the exciting potential of doped ZnO ceramics and their dielectric behavior, especially regarding donor and acceptor co-doping, have stimulated the authors to study the dielectric properties of co-doped ZnO ceramics.
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In this work, (Li, In) co-doped ZnO ceramics have been synthesized. The influence of lithium and indium co-doping on the structural, morphological and dielectric properties of
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ZnO ceramics were investigated. CP was obtained from co-doped ZnO ceramics and the origin of CP phenomenon is discussed. Experimental details
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2.
Reagent-grade (≥ 99.99%) starting materials of Li2CO3, In2O3 and ZnO (all from Aladdin, China) were weighted according to the stoichiometric proportion of Zn(1-2x)(Li, In)xO (x = 0.25%, 0.5%, 0.75%, 1.00%, 1.50%, and 2.00%, respectively). Raw materials were mixed in ethanol by ball milling with Zirconia balls in a polypropylene bottle for 6 h at 400 rpm. The mixture was then dried and calcined at 1323 K for 2 h. Subsequently, the calcined sintered mixture was ball-milled again for 6 h. After drying, 6 wt% polyvinyl acetate (PVA) binder was added and the mixture granulated by sieving through a 60-mesh screen to produce 3
ACCEPTED MANUSCRIPT starting powders. The powders were uniaxially pressed into discs (13 mm in diameter and 1.5 mm in thickness) in a stainless steel die at a pressure of 100 MPa. Parts of the samples were pressed in a cold isostatic pressing (CIP) machine at a pressure of 200 MPa for 2 min. Sintering was then carried out at different temperatures (1498 K, 1523 K, 1548 K, 1573 K)
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for 12 h with a heating rate of 3 oC/min. The samples were allowed to cool down naturally to the ambient temperature.
The doped ZnO and un-doped ZnO powder were sintered at 1548 K for 12h and mixed
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with 20 wt% pure silicon powders. The crystal structures of samples were investigated via
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X-ray diffraction (XRD) with a Cu-Kα radiation (Bruker D8 Advance, Germany). Surfaces of samples were lapped and ground with SiC powders, then polished with 0.1 µm-Al2O3 powders to a mirror-like surface. The polished samples were thermally etched for 2 h at 1173 K. The microstructure was examined by a Supra 55 (Carl Zeiss, Germany) scanning electron microscope (SEM) equipped with an energy dispersive spectrometer (EDS, X-MaxN, Oxford,
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UK). Electrodes were made by printing silver paste on both sides of the disk-type samples and then fired at 973 K for 2 h in order to remove the polymeric component. The dielectric properties of the ceramics were measured by a WK-6500B Precision Impedance Analyzer
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(Wayne-Kerr Electronics Ltd., London, UK) in the frequency range of 20-106 Hz at a
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temperature range from 150 K to 400 K, respectively. Results and discussion Fig. 1 shows the XRD pattern of Zn1-2x(Li, In)xO powders sintered at 1548 K. The pure silicon (Si) powder is used as reference for XRD measurement. The (Li, In)-doped ZnO diffraction peaks can be indexed to the wurtzite ZnO (JPCDS No. 36-1451) with a hexagonal structure of P63mc space group. No peak of secondary phases was observed in the XRD patterns of all the samples. The XRD peak located between 36° and 37°, attributed to the (101) plane, firstly shifts slightly to low angles with increasing doping amount. The peak shifts to 4
ACCEPTED MANUSCRIPT high angles when the doping amount exceeds 1%. The radii of Zn2+, Li+, and In3+ are 0.074 nm, 0.068 nm, and 0.094 nm, respectively. Since one Li+ and one In3+ may replace 2Zn2+ in terms of valence balance, the average radius is about 0.081 nm, which is larger than that of Zn2+. The substitution of Zn2+ by the dopants Li+ and In3+ would lead to the extending of the
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lattice parameter, which results in the shift of (101) peak to lower angle. When the doping amount x exceeds 1%, part of the In3+ dopant may exist in the form of In2O3 and form solid solution with ZnO [23]. Since the amount of In2O3 is very small, its diffraction peaks cannot
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be observed. However, its existence may encompass the ZnO lattice and leads to the shift of (101) peak to higher angle [24]. On the other hand, Li+ can still substitute Zn2+ when the
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doping concentration is higher than 1%. The substitution of Zn2+ by Li+ would lead to the decrease of the lattice parameter. As a result, the (101) diffraction peak shifts to higher angle [25].
Fig. 2 shows SEM micrographs of surfaces of Zn1-2x(Li, In)xO ceramics. The grain
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structure is relatively homogeneous throughout the sample, and grain boundaries are clearly seen. The ceramics are very dense, with grain sizes in the range of 8 to 14 µm. As the doping content of Li and In increases, the grain size decreases gradually, with more small grains
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appearing at the grain-boundary triple junctions. EDS surface scanning analysis shows that
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the In3+ is homogeneously distributed (Fig. 2h). However, Li+ cannot be discussed as it is undetectable in EDS.
Fig. 3a shows the dependence of permittivity and dielectric loss on the doping concentration. A high relative permittivity of 3800 and a comparatively low dielectric loss (tg δ) of about 0.12 can be obtained when the doping concentration x = 0.5%. The permittivity of co-doped ZnO increases when the sintering temperature increases from 1500 to 1600 K. Meanwhile, the dielectric loss increases sharply as well. Consequently, the sintering temperature of 1548 K is regarded as optimum because all the samples can be sintered at this 5
ACCEPTED MANUSCRIPT temperature whilst having relatively low dielectric loss. In order to understand the dielectric properties caused by the defects, the frequency dependences of permittivity and dielectric loss at various temperatures from 293 K to 363 K were measured as shown in Fig. 4a. The oblique line appearing in dielectric loss curves at
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frequency below 100 Hz, is attributed to the electrode polarization, which indicates the formation of Schottky barrier at the sample-electrode interface [26]. However, the peaks in the high-frequency region can’t be accurately clarified due to the existence of the background
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interference. Notably, when the temperature increases to 363 K, the peak is almost merged by the background.
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The peaks of the defect relaxation can be clearly clarified in Modulus-Frequency spectra with the complex dielectric constant converted into its complex modulus (M' + iM'') [27-29]. The solid line in Fig. 4c indicates the fitted data, from which the activation energy was calculated and found to be 0.29 eV for the high temperature region (HTR), i.e. above 333 K,
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and 0.09 eV for the low temperature region (LTR), i.e. below 333 K. The peaks shift to higher temperature when the measuring frequency increases, indicating a thermally activated Debye-like relaxation. The relaxation parameters can be calculated according to the
Ea
k BT
)
(1)
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f = f 0 exp(−
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Arrhenius law as follows:
where f0 is pre-exponential factor, Ea is the activation energy of defect relaxation, kB is the Boltzmann constant, and T is absolute temperature. In the Arrhenius plot, the deviation shows and the presence of different activation energies for the different temperature regions are due to the change of relaxation mechanisms. For this phenomenon, a similar anomaly of dielectric relaxation was found in ZnO ceramics, in which the activation energy of which are 0.30 eV and 0.09 eV [30, 31]. One peak with the activation energy of 0.29 eV (Peak-2) has been reported in a few literatures, including in ZnO 6
ACCEPTED MANUSCRIPT film, which is caused by oxygen vacancy [30-32]. The other one (Peak-1, 0.09 eV) has been attributed to the interstitial oxygen [30, 31]. What’s more, these peaks would disappear after O2-annealing process in Fig. 4d. However, the interstitial oxygen is very high in formation energy and electrically inactive, according to the first principles calculation [33]. If Peak-1
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were associated with interstitial oxygen, the peak would still exist after the O2-annealing process. In O2-annealing process, excess oxygen atoms in the ZnO lattice can be accommodated in the form of oxygen interstitials. However, the result is totally inconsistent
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with the expectation as Fig. 4d shows. So it is concluded that this peak would be also related to the oxygen vacancy in the form of oxygen vacancy clusters. It has been reported that the
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energy of moving these clusters is less than the energy of moving individual oxygen vacancies [34-36].
Oxygen vacancies can trap charge carriers [34], such as the electrons. The electrons are localized in oxygen vacancies and can jump between the ionized oxygen vacancies, giving
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rise to a dipolar effect [36-38]. It is well-known that oxygen vacancies contribute to polarization in the forms of singly and doubly charged states in low- and high- temperature ranges. Two relaxation peaks have been reported to be respectively due to the hopping
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motions of singly and doubly charged oxygen vacancies in CaNb2O6 [39], La2Ti2O7 [26] and
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TiO2 [40]. The temperature where oxygen vacancies change from singly charged state to doubly charged state were reported to be about 800 K in SrTiO3 [41] and 700 K in NiTiO3 [42]. Vanheusden et al. have discovered that, in the pure ZnO, there is the transition between singly charged oxygen vacancies and doubly charged oxygen vacancies. It was demonstrated that the density of VO• increased as the temperature went down. In other words, the density of free carriers would increase as the temperature goes up, according to the equation: VO• → VO•• + e′ .
Vanheusden has discovered that at 294 K, the densities of VO• and free-carriers
has much differences [43], which indicates that the VO• → VO•• + e′ happens. So in this 7
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333 K VO• ← →VO•• + e′
(2)
Regarding to the source of the singly and doubly charged oxygen vacancies, there are two main ways. Firstly, the oxygen vacancies could be created with the introduction of Li+ and
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In3+ ions into Zn-sites. Li+ ions would create the oxygen vacancies. However, Li+ doping into the Zn-site would enhance the Madelung energy by 13.56 eV, resulting the instability of ionic
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charge distribution in lattice [44]. For charge-compensation, the Li+ ions’ would occupy the interstitial sites. In this way, this would reduce the Li+ ions’ solubility. However, the introduction of In3+ into Zn2+-sites are considered as donors, which can reduce the Madelung energy by 9.73 eV [44]. The formation of the defect InZn• would be near the LiZn′ sites,
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which enhance the stability of the ionic charge distribution. So the In3+ ions doping can enhance the incorporation of acceptor Li+ ions because of the strong acceptor-donor attractive interaction, which indicates more oxygen vacancies would be created. Secondly, during
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sintering process at high temperature with low oxygen partial pressure, the oxygen vacancies
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could be generated due to the loss of oxygen from the crystal lattice. The defect reaction equations are shown as follows: ′ + VO•• + 3OO Li2O + 2ZnO → 2 LiZn
(3)
In2 O3 + 2 ZnO → 2 In Zn• + VZn′′ + 5OO
(4)
OO ↔
1 O2 + VO•• + 2e′ 2
(5)
In addition, after the O2-annealing process, the permittivity and the dielectric loss decrease dramatically to 460 and 0.04 respectively at 1 kHz in the inset of Fig. 4a, which 8
ACCEPTED MANUSCRIPT indicates the dielectric property is mainly associated with the oxygen vacancies. To verify that the O2-annealing process strongly affects the oxygen defects in the (Li, In) co-doped ZnO system, X-ray photoelectron spectra (XPS) of the as-prepared and O2-annealed sample are shown in Fig. 5. The XPS spectra for O 1s peaks depict a wide and asymmetric peak that is
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further deconvoluted into three peaks centered at different binding energies indicating the presence of different O species. Generally speaking, the O1s peak has been observed in the binding energy (BE) region of 529-535 eV. The peak around 529-530 eV has been attributed
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to lattice oxygen and the peak around 530.7-531.6 eV has been attributed to oxygen in the region of non-stoichiometric oxides. Furthermore, for the moisture, its binding energy lies
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between 531.5 eV (OH) and 533 eV (H2O) is due to the chemisorbed oxygen [45]. The scan of O1s spectra in Fig. 5a and 5b exhibit a peak of OІ (530.08 eV) which is attributed to the oxidized metal ions in the samples, viz O-In and O-Zn, in the lattice [46], a peak of OII (530.95 eV) is caused by the O deficiency and a peak of OIII (532.01 eV) is due to the
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surface adsorbed oxygen species such as OH or O2 [45, 47]. It is noted that the intensity ratio of the OII/OI decreases from 0.62 to 0.55 after O2-annealing process, which indicates that the concentrations of oxygen vacancies decrease after O2-annealing process [48].
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It is well known that the low- and high- frequency semicircles in the Nyquist plot
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represent the different relaxations incited by the intergrain and intragrain effect, respectively [49]. In Fig. 6a, it can be seen that only one semicircle exists, and it is remarkably depressed by the bias voltage. This fact demonstrates that the dielectric properties are mainly due to the grain-boundary effect. After thermal treatment in an oxygen atmosphere, the relaxation peaks of both VO•• and V O• disappear. Besides, the Nyquist semicircle of O2-annealed sample cannot be affected by the bias voltage as shown in Fig. 6b. The permittivity of the ZnO ceramics decreases to ~460, and the dielectric loss decreases to 0.04 after thermal treatment in oxygen. The colossal permittivity of the (Li, In) co-doped ZnO ceramics has a connection 9
ACCEPTED MANUSCRIPT with these two defects VO•• and V O• , which contribute to the formation of grain boundary relaxation at room-temperature. The movement of defect charges would generate a micro-capacitance according to the grain boundary-layer mechanism, which is responsible for
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the colossal permittivity of the specimen.
Conclusions
Li+ and In3+ co-doped ZnO ceramics were prepared and their structural and dielectric
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properties were investigated. Li+ and In3+ have been successfully incorporated into the ZnO
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lattice. The dopant is distributed homogeneously in the ceramics and no obvious agglomeration has been found. CP is obtained from the co-doped ZnO ceramics and a permittivity of up to 3800 is obtained from Zn0.99(Li, In)0.005O ceramics. The dielectric modulus and impedance spectroscopy analysis indicate that the dielectric relaxation in the co-doped ZnO is associated with singly and doubly charged oxygen vacancies. The oxygen
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vacancies are mainly created with the introduction of lithium and indium ions into Zn-site and the sintering process. Furthermore, after thermal treatment in an oxidizing atmosphere for 2 h, the relaxation peaks of singly and doubly charged oxygen vacancies defects disappear
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and the permittivity reduces to ~460 at room-temperature. From the X-ray photoelectron
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spectra, the concentrations of oxygen vacancies decrease after O2-annealing process. It can be concluded that oxygen vacancies are the main origin of the CP phenomenon in Zn0.99(Li, In)0.005O ceramics at room-temperature range.
Acknowledgements The authors acknowledge the financial support from the Ministry of Science and Technology of China through the program of International S&T Cooperation (S2015ZR1108), 973-project (2015CB654604, 2015CB654605), Natural Science Foundation of China 10
ACCEPTED MANUSCRIPT (51572279) and the CAS/SAFEA International Partnership Program for Creative Research Teams. The authors would like to thank Dr. Adrian Trinchi of CSIRO, Materials Science & Engineering, Australia for his valuable discussion and kind help. References
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Captions of Figures
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Fig. 1. The XRD patterns of Zn1-2x(Li, In)xO powders: x=0%, 0.25%, 0.5%, 0.75%, 1%, 1.5%, 2%, with Si as a reference.
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Fig. 2. The SEM micrographs of the as-sintered Zn1-2x(Li, In)xO ceramics after annealing process at 1273 K for 30 min with (a) x=0.25%, (b) 0.5%, (c) 0.75%, (d) 1%, (e) 1.5%, and (f) 2%; (g) the SEM micrograph of the as-sintered Zn0.99(Li, In)0.005O ceramic with the annealing process at 1173 K for 2 h; (h) the lateral scanning of the Indium distribution in ZnO0.99(Li, In)0.005O ceramic without the annealing process.
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Fig. 3. (a) the dielectric properties of Zn1-2x(Li, In)xO with x= 0.25%, 0.5%, 0.75%, 1%, 1.5%, 2% samples sintered at 1548 K; (b) the dielectric properties of Zn0.99(Li, In)0.005O sample sintered at the temperature ranging from 1498 K to 1573 K.
Fig. 4. The dielectric properties of Zn0.99(Li, In)0.005O ceramics at a measuring temperature range from 293 K to 363 K: (a) and the inset are the permittivity- and dielectric lossfrequency spectrum; (b) the M"-frequency spectrum; (c) the Arrhenius plots of peak 3 and peak 4; (d) the M"-frequency spectrum of the O2-annealed sample (973 K/2 h).
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Fig. 5. X-ray photoelectron spectrum of O1s (a) in the as-prepared and (b) O2-annealed sample
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Fig. 6. (a) the Nyquist plots recorded at room temperature under different dc bias, the inset is the equivalent circuit used to represent the electric properties of grain-boundary interface effect in ZnO ceramics; (b) the Nyquist plots recorded at room temperature under different dc bias after O2-annealing process at 973 K for 2 h.
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Highlights:
The (Li, In) co-dopants have been successfully incorporated into ZnO
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lattice. Colossal permittivity (CP) of 3800 is obtained for (Li, In) co-doped ZnO ceramics.
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and V O .
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of VO
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The CP originates from the grain-boundary effect with defect dipoles