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| Contact | Kaminamie-Town
Takasaki-City, Gunma 370-0801
JP
０９０１９０８７６２５
ferroelectricmaterials@jcom.home.ne.jp |
Goal | researcher |
| Skills | ferroelectric materials, thin film devices, sensors, capacitors |
| Qualifications | [I] FUTURE PLANS
~ Gas Sensor ~
The surface conductivity of semiconductor oxide depends on the composition of the surrounding gas atmosphere. It is a well-known phenomenon and has been used as gas sensing mechanism for many years. The surface conductivity of n-type semiconductor oxides is modified by adsorption/desorption of oxidizing gas species and related space charge effects: in oxidizing atmosphere, the oxide surface is covered by negatively charged oxygen adsorbates and the adjacent space charge region is electron-depleted, resulting in high resistance state; in reducing atmosphere, the oxygen adsorbates are removed by reaction with reducing gas species and electrons are re-injected into space charge regions, resulting in low resistance state.
Though the gas sensors based on semiconducting oxides have been known for at least 40 years, it is a major disadvantage that the detection limit has not reached 1-ppb-level. Practically, the measurement of ppb-level gaseous oxygen concentration is desirable in semiconductor industry. In semiconductor industry, inert gases are used for surface protection of materials or replacement of atmosphere. Since a slight amount of oxygen in inert gas results in surface degradation by oxidation, the in situ precise measurement of oxygen concentration in inert gas is required.
I have been made an effort to develop a gas sensor based on semiconductive ferroelectric or dielectric materials. Since the resistivity can change not only due to the increase/decrease of mobile electrons but also due to the change in the mobility of them.
If oxygen adsorption causes an increase in ionic displacement near the surface of semiconductive ferroelectric or dielectric thin films, the interaction between optical phonons and carrier conduction might be enhanced. The advantages of sensors using semiconductive ferroelectric or dielectric thin films are explained below. First, it is necessary to understand that the neutralization of electric charges upon oxygen absorption is mostly achieved by ionic polarization whether an oxide semiconductor, such as SnO2, which has conventionally been used for gas sensors, or a dielectric is used as a sensor material. Let us assume the density of oxygen absorbed to be 10E15 cm-2. However, assuming the number of electrons supplied from the sensing part of a sensor to oxygen absorbed (density of electrons multiplied by number of unit lattices equivalent to depletion layer width) to be 1.25 multiplied by 10E11 cm-2 (= 10E9 cm-2 multiplied by (50 nm/0.4 nm)), 1/8000 of the absorbed oxygen should be neutralized by electron supplied from a sensing material. Therefore, 7999/8000 of the absorbed oxygen should be neutralized by ionic polarization. Next, let us assume the case in which the electrical conduction of the sensor is controlled by polarons. Next, let us assume the case in which the electrical conduction of the sensor is controlled by polarons. In general, the conductance in materials is high under the control of large polarons, whereas it is low under the control of small polarons. Therefore, it is expected that the sensitivity of the sensor can be improved by changing the conductance using a dielectric in which both small and large polarons exist.
"Dielectric- or ferroelectric-material-based gas sensor and detection of transient phenomenon"
Miniaturized gas sensors or electronic noses to rapidly detect and differentiate trace amount of chemical agents are extremely attractive. The detection and monitoring of gases with solid state sensors has become a well established practice. Two major types of solid state gas sensor are already in widespread use: Resistive gas sensors based on semiconducting oxide films; Piezoelectric gas sensors based on quartz crystal microbalance (QCM) oscillators, or on surface acoustic wave (SAW) devices.
Resistive gas sensors based on semiconducting oxide films have been known for at least 40 years. The resistivity of film changes due to redox reactions of ambient gases, such as oxygen and reducing gases, with sensor materials resulting in the change in the number of electrons in sensor materials. Prominent example for semiconducting gas sensor is n-type SnO2, n-type WO3. Except for some H2S sensors such as nanophased WO3, it is a major disadvantage of all of semiconducting gas sensors that the detection limit is ppm-level.
Piezoelectric gas sensors based on QCM oscillators have been known for at least 30 years. A QCM is a standard term for microbalance mass sensors. Its operation principle is based on the resonant frequency shift caused by mass changes on the surface during the gas adsorption. In general, sensing films are coated on the surface. Though film-coated QCM sensors have enabled the detection of various gas molecules with sub-ppm-level to ppb-level detection limit, it is a major disadvantage of all of QCM-based gas sensors that the response is slow due to the slow diffusion of gas molecules into the sensing films.
I am planning to develop the oxygen sensor based on semiconductive ferroelectric or dielectric materials, and to develop a new detection method based on the detection of transient phenomena in the materials.
The electrical responses of the dielectric- or ferroelectric-material-based sensor are measured at room temperature using a pulsed current-voltage (I-V) measurement system. Before the measurement, two surface-cleaning processes are executed, namely, a heat treatment and a back-gating. The heat treatment is executed by the joule heating using the heating meander up to the enough temperature to annihilate the polarization in ferroelectric sensors, which is induced by oxygen adsorption. The heat treatment is followed by the back-gating. During the back-gating, the sensor is cooled down to room temperature, and electric field is applied to the surface electrodes while the heating meander is electrically grounded to suppress the polarization induced by oxygen adsorption. During the measurement, an in-plane bias field is applied between surface electrodes to detect the change of resistance in the sensor. When oxygen is present in the sample gas, the current decreases gradually since the adsorbing oxygen can accept the electron in the sensor, and the current reaches the minimum implying that the sensor surface is fully covered with a mono-molecular layer of oxygen. With increasing the oxygen concentration, the current reaches the minimum within a short time. The good sensitivity of 1-ppb-level and a fast response is a feature of the sensor and the newly developed detection scheme.
"Oxygen gas sensors based on SrTiO3 thin film"
In recent years, the interest in micromachined gas sensors has become more and more important because of their advantages such as small size, low-power consumption, and fast thermal response due to the thermally insulated structures of their hot areas. In the last few years, research activities have been concentrated toward the combination of thin-film technologies for the deposition of the sensitive materials and micromachining technologies for the fabrication of thermally insulated devices. These comprise supporting layers, heating elements, gas-sensitive layers, and electrodes. Recently, great attention has been devoted to gas sensors detecting gases at concentrations of ppb level, comparable with human noses.
The selection of gas-sensitive material is one of the most important aspects. Many different materials have been used such as SnO2 and WO3. Though SnO2 and WO3 are widely used as materials for gas sensors, their sensitivity does not reach the ppb-level. Aim of this plan is to find another material having enough sensitivity. I have chosen a perovskite-type oxide, which has good ferroelectric or dielectric property, as a candidate material for gas sensors.
The perturbation due to a gas molecular adsorption/desorption is quite weak. Therefore, the ferroelectric or dielectric property in the perovskite used as a gas-sensitive material should be sensitive to the weak perturbation. It is widely known that the dielectric properties in SrTiO3 thin films are extremely sensitive to external perturbations. For example, the anomalous local ferroelectricity can emerge in SrTiO3 thin films [1], though SrTiO3 single crystals are paraelectric with an ideal cubic structure at room temperature. Reference [1] discusses about the inhomogeneous ferroelectric phase induced by the inhomogeneous strain arising from the film/substrate interaction. The sensitive dielectric properties of SrTiO3 thin films can be utilized to detect weak external fields. Furthermore, nano-sized SrTiO3 has recently been recognized as an excellent material being capable of detecting O2 at room temperature [2, 3]. Thus, I have chosen SrTiO3 as a candidate material for gas sensors.
References:
[1] O. Tikhomirov, H. Jiang, and J. Levy, Phys. Rev. Lett., 89, 147601, 2002
[2] Y. Hu, O. K. Tan, J. S. Pan, H. Huang, and W. Cao, Sensors and Actuators B, 108, 244, 2005
[3] Y. Hu, O. K. Tan, W. Cao, and W. Zhu, IEEE Sensors Journal, 5, 825, 2005
"Oxygen gas sensors based on SrTiO3 thin film II."
Great attention has recently been devoted to gas sensors detecting gases at concentrations of ppb level, comparable with human noses: metal oxide semiconductor gas sensors, such as SnO2 and WO3, have been studied extensively because of their simple mechanisms in which the conductance is dependent on carrier concentrations, simple fabrication processes, and low cost. Moreover, thin films are the most suitable for the production of micro sensors: many researchers have studied and are still studying the fabrication methods of SnO2 or WO3 thin film sensors. Though SnO2 and WO3 are widely used as materials for gas sensors, a few problems remain: SnO2 has no enough sensitivity, has the cross-sensitivity to humidity, and shows the drift of the sensor performances; WO3 has the low cross-sensitivity to humidity, however, the sensitivity does not reach the ppb-level.
Aim of this work is to find another material having enough sensitivity.
SrTiO3 can be a promising candidate material for gas sensors: it is widely known material for O2 gas sensors operating at 973-1273 K; nano-sized SrTiO3 has recently been recognized as an excellent material being capable of detecting O2 at room temperature [1, 2]. Though SrTiO3 is a well-known material for O2 sensors, the detection mechanism in it has not been well understood as follows:
(a)In conventional SrTiO3 gas sensors operating at 973-1273 K, electrons, holes, and oxygen vacancies can be mobile. However, at room temperature, SrTiO3 has only n-type conductivity, and the mobility of holes or oxygen vacancies is rather small.
(b)In Ref. [1, 2], it is shown that the nano-sized SrTiO3 gas sensor has higher resistance in N2 than in N2/O2 = 80%/20% at room temperature, and concluded that the nano-sized SrTiO3 gas sensor has a p-type conductivity. It may be true since the insulation resistance is rather high, varying form 1.0-1.6%C3%971011 %CE%A9 in N2 to 2.0-4.0%C3%971010 %CE%A9 in N2/O2 = 80%/20%, and since the nano-sized SrTiO3 is fully electron-depleted and has Au contacts, and since the oxygen holes is observed by X-ray photoelectron spectroscopy.
(c)WO3, which can be viewed as a perovskite derivative and have a ferroelectric-like character, has n-type conductivity in the temperature range 290-370 K probably owing to oxygen vacancies and impurities, such as protons, accidentally injected into interstitial sites of the WO3 matrix during its preparing processes. The 100-500 nm thick WO3 thin films, in which electrons can be partially depleted near surfaces, have lower resistances in air than in vacuum [3]. However, in general, n-type semiconductors should have lower resistances in vacuum than in air since O2 accepts electron(s) from n-type semiconductors.
A scenario disabusing above-mentioned discords is likely the phonon-drag effect on the conduction of electrons due to the O2 adsorption:
(1)The electronegative O2 adsorbing on the film surface can rotate a component of electric dipole to the direction normal to the surface, leading to the generation of charged domain wall, since the domain wall tilting away from the permissible uncharged orientation causes the charged domain wall.
(2)The rotation of local electric dipole results in an elastic shear deformation.
(3)The deformation can propagate under an applied electric field to the direction parallel to the film surface leading to the generation of an acoustic wave.
(4)The acoustic wave may propagate accompanied by electrons.
It has been assumed that the electric conductivity can change in the region adjacent to the domain wall in LiNbO3 [4]. Furthermore, it is widely known that the dielectric properties in SrTiO3 thin films are extremely sensitive to external perturbations, e. g., the anomalous local ferroelectricity can emerge in SrTiO3 thin films [5], though SrTiO3 single crystals are paraelectric with an ideal cubic structure at room temperature. Reference [5] discusses about the inhomogeneous ferroelectric phase induced by the inhomogeneous strain arising from the film/substrate interaction. The sensitive dielectric properties of SrTiO3 thin films can be utilized to detect weak external fields.
Ferroelectric materials generally have domain fluctuations causing the change in conduction. For example, Lepikh has discussed about conduction temperature oscillations due to 90%C2%BA-domains fluctuation in Pb(Zr,Ti)O3 around 413 K (The Curie temperature of Pb(Zr,Ti)O3 ranges from 573 K to 673 K.) [6]. It should be noted that the thermal (pyroelectric) noise can be avoided by using the gas sensors at sufficiently lower temperatures than Curie temperature and/or the temperature in which the conduction temperature oscillations occur. When the conduction temperature oscillations occur during the gas sensing, a noise elimination process by a special electric circuitry may be needed. However, it is assumed that the self polarization due to the defects in perovskites and due to the O2 adsorption at the perovskite surfaces is sufficiently stronger than the perturbation due to the slight change in temperature, e. g., 1.1-1.4%C3%9710-4 Cm-2K-1 of pyroelectric coefficient in Pb(Zr,Ti)O3 [7] or 60%C3%9710-4 Cm-2K-1 of that in Pb(Sc1/2Ta1/2)O3 [8], as observed in the charge compensation effects in ultrathin films.
References:
[1] Y. Hu, O. K. Tan, J. S. Pan, H. Huang, and W. Cao, Sensors and Actuators B, 108, 244, 2005
[2] Y. Hu, O. K. Tan, W. Cao, and W. Zhu, IEEE Sensors Journal, 5, 825, 2005
[3] M. G. Hutchins, O. Abu-Alkhair, M. M. El-Nahass, and K. Abdel-Hady, J. Phys.: Condens. Matter, 18, 9987, 2006
[4] Y. Zhi, W. Qu, D. Liu, Z. Luan, Y. Zhou, and L. Liu, Appl. Phys. Lett., 89, 112912, 2006
[5] O. Tikhomirov, H. Jiang, and J. Levy, Phys. Rev. Lett., 89, 147601, 2002
[6] Y. I. Lepikh, Semiconductor Physics Quantum Electronics & Optoelectronics, 3, 308, 2000
[7] V. Ferrari, A. Ghisla, D. Marioli, and A. Taroni, IEEE Sensors Journal, 3, 212, 2003
[8] D. Damjanovic, P. Muralt, and N. Setter, IEEE Sensors Journal, 1, 191, 2001
~ UV Sensor ~
The space industry and science arena demand the development of a new generation of ultra-violet (UV) light sensors. The wavelength ranges of interest are 10-200 nm (6.2-124 eV) for solar physics studies, 200-280 nm (4.4-6.2 eV) for communications and missile warning, and 280-400 nm (3.1-4.4 eV) for atomospheric studies. Typical specifications are very low dark-current (< 10 pA), resistance to radiation doses, and high sensitivity.
UV detection has traditionally been accomplished by photomultiplier tubes (PMTs), narrow-band gap semiconductors such as Si. PMTs exhibit high gain and low noise, and can be visible-blind. However, they are fragile and bulky. The narrow-band gap semiconductors, such as Si, can offer the advantages of compact solid-state devices. However, they show high dark-current, require the filters to block out visible and infrared light, and are prone to be damaged irreversibly.
Semiconductors with large bandgap such as GaN or Al1-xGaxN can be promising candidate materials for UV photodetectors, and have been actually used for UV-sensitive photodiodes. They show low dark-current, are visible-blind, and provide radiation hardness. GaN- or Al1-xGaxN-based UV photodetectors have the possibility to use any ratio for the content of gallium and aluminum. Therefore, it is possible to have a semiconductor with any bandgap in the range of 3.4-6.2 eV. Al0.15Ga0.85N-based UV photodetectors with a band gap of 3.5 eV have realized UV-A photodetectors [1], which can be used to detect near-UV (UV-A) (3.1-3.9 eV) radiation on the earth%E2%80%99s surface being hazardous to the human body. Al0.45Ga0.55N-based UV photodetectors with a bandgap of 4.2 eV have realized solar-blind UV imagers [2], which can be used for UV-B (3.9-4.4 eV) detection such as ozone layer monitoring.
Perovskite-type titanates, such as strontium titanate (SrTiO3), are possible UV-sensitive materials, and can provide easier fabrication than the GaN-based one does. Furthermore, their dielectric or ferroelectric characteristics can tune their Schottky barrier not only by the trapping-detrapping of carriers but also by the reorientation of electric dipoles. Thus, the dielectric or ferroelectric perovskite-type titanates are anticipated to provide high gain under the UV light illumination.
References:
[1] T. K. Ko, S. C. Shei, S. J. Chang, Y. K. Su, Y. Z. Chou, Y. C. Lin, C. S. Chang, W. S. Chen, C. K. Wang, J. K. Sheu, and W. C. Lai, IEEE Sensors Journal, 6, 964, 2006
[2] G. Mazzeo, J. -L. Reverchon, J. -Y. Duboz, and A. Dussaigne, IEEE Sensors Journal, 6, 957, 2006
"Effects of LaAlO3 or YAlO3 seed layers on electrical/electronic properties of SrTiO3 thin films grown onto the seed layers"
The purpose of this paper is to outline the effects of LaAlO3 or YAlO3 seed layers having antiferrodistortive AlO6 octahedra on electrical/electronic properties of SrTiO3 thin films grown onto the seed layers. The phase transition between an antiferrodistortive state to a ferroelectric one is expected to occur accompanied by the increase in conductivity and/or the increase in dielectric permittivity.
A physical system, that crosses the boundary between two phases, changes its properties, for example, melt or freeze. This macroscopic change is driven by microscopic fluctuations. In a quantum system, a quantum fluctuation drives a phase transition. The condensed matter having strong vibronic coupling (or rovibronic one, that stands for rotational-vibrational-electronic coupling), such as perovskite-type materials, can be promising candidates for the electronic applications using quantum fluctuations and using entanglement/disentanglement changing under external fields.
The perovskite-type materials such as stroutium titanate, barium titanate, and barium strontium titanate have attracted much attention owing to their superconductivity, quantum paraelectricity, ferroelectricity, and dielectric tunability, respectively. Particularly, SrTiO3 has gained additional scientific and practical importance owing to its eccentric phase transition, which may be driven by quantum paraelectricity, induced by the irradiation of ultraviolet light and by the application of DC electric field [1-3]. The crystal structures, lattice oscillations, and electric/electronic properties of SrTiO3 single crystals have been investigated over and over in the last five decades. Early studies contributed essentially to the understanding of physics in SrTiO3 single crystals as follows:
(1) Crystal structure: At room temperature, they have an ideal cubic perovskite structure with space group Pm3{bar}m. On cooling below 300 K, ferroelectric microordered regions appear at 115 K, which are induced by polar dopant impurities in the cubic phase of the compound [4, 5]. Around 110-105 K, the cubic-tetragonal antiferrodistortive phase transition occurs. This phase transition originates from the tilting of TiO6 octahedra in such a way that we have antiphase rotation of these octahedra in successive layers along one cubic axis [6]. This modification is described in the Glazer classification [7] by the a0a0c- tilt system. The final result is a tetragonal I4/mcm structure with two SrTiO3 molecular units in the primitive unit cell.
(2) Dielectric permittivity: Below 45 K, the permittivity increases steeply, and displacement versus electric field curve (D-E curve) exhibits hysteresis [8]. Below 10 K, the permittivity saturates as a result of quantum paraelectricity [9]. Around 60-130 K and around 5-20 K, the dielectric loss increases anomalously [10].
(3) Electric Conductivity: The electric conductivity increases with decreasing temperature. However, it decreases anomalously around 80K and around 47 K [11].
(4) The effect of ultraviolet light irradiation on dielectric permittivity: By the irradiation of ultraviolet light under a DC electric field, the dielectric permittivity increases gradually below 100 K, and steeply below 50 K [1, 2].
(5) The effect of ultraviolet light irradiation on electric conductivity: By the irradiation of ultraviolet light, the electric conductivity increases steeply below 105 K, and it saturates below 50 K [3].
(6) The effect of near-infrared light irradiation on electric conductivity and on second harmonic generation (SHG): By the irradiation of near-infrared Nd:YAG laser light (1.17 eV), the electric conductivity increases steeply below 35 K. Furthermore, gigantic SHG along [110] direction occurs by the irradiation of near-infrared laser light under a DC electric field [12].
From above results, it is speculated that the phase transition between the antiferrodistortive state to the ferroelectric one can be induced by the irradiation of ultraviolet light and by the application of DC electric field accompanied by the increase of electric conductivity and dielectric permittivity. Furthermore, it is speculated that an acoustic wave generates by the irradiation of near-infrared light and by the application of DC electric field accompanied by the increase of electric conductivity.
The crystal structures, lattice oscillations, and electric/electronic properties of SrTiO3 are strongly affected by twin structures causing canted domain structures [11], by intentionally or unintentionally doped chemical impurities causing ferroelectric microordered regions [4, 5, 13] or the coherent antiphase rotation of TiO6 octahedra as a result of intentional doping of Ca [5, 14, 15], by compressive or tensile strain causing ferroelectricity by removing the center of inversion symmetry of the crystals [16, 17], and by the application of electric field causing ferroelectricity and soft mode hardening [18]. Considering the application of SrTiO3 for electronic devices, thin films are useful. However, when we use thin films as electronic devices, we should avoid introducing uncontrollable defects such as twin structures and oxygen vacancies into thin films, even if they can partially introduce frustrated structures causing gigantic changes of electric/electronic properties. Furthermore, heterogeneity associated with intentional chemical doping can broaden the phase transition. The heterogeneity is therefore detrimental to obtain the sharp gigantic changes of electric/electronic properties. In contrast, a coherent compressive or a tensile strain seems to be useful.
In recent years SrTiO3 thin films have shown new aspects as follows:
(7) Strong two photon absorption at room temperature: The thin films deposited on quartz substrates have larger 2 photon absorption coefficient under the irradiation of near-infrared (1.17 eV) laser light, even at 300 K, than that of SrTiO3 single crystals [19]. Reference [19] discusses grain boundaries as an origin of the phenomenon.
(8) Soft mode hardening and Raman-active TO2 and TO4 mode at room temperature: The soft mode in the film deposited on SrRuO3-electrode/LaAlO3-substrate is hardened. Reference [20] discusses the strain introduced by the substrate as an origin of the phenomenon. Furthermore, TO2 (c.a. 170 cm-1) and TO4 (c.a. 550 cm-1) appear at 300 K in thin films, though TO2 and TO4 in single crystals are very slight even at 5 K [20, 21]. Reference [21] discusses the defects, such as oxygen vacancies and grain boundaries, as origins of the phenomenon.
(9) The coupling between the infrared active TO1 (soft mode) and TO2 mode at room temperature: In SrTiO3 single crystals, infrared active soft mode (TO1) and TO2 mode are coupled with increasing temperature, however the coupling is not perfect even at 1200 K [22]. In contrast, hardened-TO1 and TO2 are already coupled (imperfectly) even at 300 K in SrTiO3 thin films deposited on sapphire substrates [23].
(10) Room-temperature ferroelectricity in strained SrTiO3 thin films and enhanced ferroelectricity in strained BaTiO3 thin films: Room-temperature ferroelectricity is observed in SrTiO3 thin films in which tensile strain is introduced by DyScO3 substrates [24]. The induced ferroelectricity in SrTiO3 thin films deposited on sapphire (0001) substrates had been observed below 125 K by Petzelt et al. [26]. Enhanced ferroelectricity is observed in BaTiO3 thin films in which compressive strain is introduced by GdScO3 or DyScO3 substrates [25].
Above results encourage me to utilize non-linear gigantic electric/electronic properties changes of SrTiO3 at room temperature by introducing coherent strain into them.
LaAlO3 and YAlO3 seem to be promising candidates as seed layers on which SrTiO3 thin films are made as follows:
(a) LaAlO3 and YAlO3 can introduce compressive in-plane strain into SrTiO3 layers deposited on them. On the analogy of a photo-induced resistivity changes in (Bi,Ca)MnO3 thin films [27], the in-plane conductivity changes more drastically in a compressively strained thin film than in a tensile strained one.
(b) LaAlO3 and YAlO3 have antiferrodistortive tilted AlO6 octahedra at 300 K. These modifications are described in the Glazer classification [7] by the a-a-a- tilt system for LaAlO3, and by the a-b+a- tilt system for YAlO3. They are therefore expected to introduce the antiferrodistortion into SrTiO3 layers deposited on them.
(c) By conducting infrared variable angle spectroscopic ellipsometry on LaAlO3 (100) and YAlO3 (100) single crystals, one can see the bending modes at 428 cm-1 for LaAlO3 and at 451 cm-1 for YAlO3 (B1u(7), B1u (8), B2u(5), B2u (6), B3u(7), and B3u (8) [28]), and the stretching modes at 660 cm-1 for LaAlO3 and at 660 cm-1 for YAlO3 (B1u(10), B2u(7), B2u(8), and B3u(10) [28]). A new mode appears at 490 cm-1 (B1u (9) and B3u(9) [28]) in YAlO3 owing to its lowered lattice symmetry. The bending modes and stretching modes in LaAlO3 and YAlO3 single crystals are harder than those of powdered cubic SrTiO3 (bending at 395 cm-1, stretching at 610 cm-1), powdered tetragonal BaTiO3 (bending at 400 cm-1, stretching at 545 cm-1), powdered orthorhombic CaTiO3 (bending around 360 cm-1 [broad], stretching at 540 cm-1 with a shoulder at 700 cm-1), and powdered ilmenite MgTiO3 (bending at 350 cm-1 and at 475 cm-1, stretching at 600 cm-1 [broad]) [29]. This is the general trend in AlO6 octahedra. In BaTiO3 single crystals and SrTiO3 single crystals the bending mode does not appear [29]. However, the bending modes appear in LaAlO3 and YAlO3 single crystals. It is therefore assumed that the grain-boundaries-induced frustrated structures in powdered SrTiO3, in powdered BaTiO3, in powdered CaTiO3, and in powdered MgTiO3, are already present in LaAlO3 or YAlO3 single crystals. The bending mode tends to appear around morphotropic phase boundary in (K,Na,Li)NbO3 as a prologue for the structural transformation from orthorhombic to tetragonal symmetry [30].
YAlO3 seems to be more promising as a seed layer owing to following reasons:
(d) On the analogy of a photoelectron spectroscopic study [31] and a theoretical investigation [32] on LaTiO3 and YTiO3, it is expected that SrTiO3 on YAlO3 has lower electric conductivity at the electronic ground state since the degenerated t2g orbitals may be lifted to dxy and doubly degenerated dyz and dzx owing to its lattice distortion. Furthermore, SrTiO3 on YAlO3 may have orbital ordering alternating dyz and dzx (This may occur only at cryogenic temperatures). The lift of t2g orbitals and their orbital ordering can lower the electric conductivity at the electronic ground state.
(e) It is expected that the interaction between second nearest-neighbor TiO6 octahedra in SrTiO3 on YAlO3 is stronger than that in SrTiO3 on LaAlO3, and that the dipoles in SrTiO3 on YAlO3 is more frustrated than that in SrTiO3 on LaAlO3. In more frustrated structures, the electronic configuration or orbital ordering may melt easily by lattice oscillation excitations at the electronic excited state, and may have high electric conductivity at the electronic excited state.
(f) By conducting atomic force microscopic studies on LaAlO3 (100) and YAlO3 (100) single crystals, one can see that LaAlO3 has c.a. 0.38-nm-high steps and c.a. 200-nm-wide terraces. In contrast, neither the obvious step nor the terrace can be observed at the surface of YAlO3. However, the aggregates aligning rhombohedrally or hexagonally can be observed on the YAlO3 (100) surface, since Y2(CO3)3 [or Y2O3 and Y2(CO3)3] tends to aggregate near the steps; this implies that the rhombohedrally or hexagonally strained structure is constructed in orthorhombic YAlO3. If the transition between a rhombohedral phase and a tetragonal one occurs via an orthorhombic one, YAlO3 seems to have a quasi-transition-state symmetry. The transition between an orthorhombic phase and a trigonal one has been observed in (Nd,Sm)AlO3 solid solution [33-35]. However, the heterogeneity associated with intentional chemical doping can broaden phase transition. The introducing of coherent strain is therefore more preferable.
(g) By conducting Raman spectroscopic studies under the irradiation of visible laser light (for example, 532 nm = 2.33 eV. The band gap of LaAlO3 is 5.6 eV, and that of YAlO3 is 3.53 eV.) on LaAlO3 (100) and YAlO3 (100) single crystals, LaAlO3 shows Eg(1) at 33 cm-1 [36], A1g(1) at 125 cm-1 [36], Eg(2) at 152 cm-1 [37], and Eg(4) at 487 cm-1 [36], and YAlO3 shows in-phase stretching Ag(1) at 150 cm-1, in-phase stretching B2g(1) at 157 cm-1, in-phase rotation Ag(2) at 197 cm-1, in-phase bending B2g(2) at 219 cm-1, out-of-phase bending Ag(3) at 278 cm-1, out-of-phase bending B2g(3) at 283 cm-1, out-of-phase rotation Ag(4) at 345 cm-1, out-of-phase bending B3g(3) at 403 cm-1, mixed vibrations of Y and O: Ag(5) at 412 cm-1, in-phase rotation B1g(4) at 540 cm-1, mixed vibrations of Y and O: B2g(6) at 552 cm-1, mixed vibrations of Y and O: Ag(7) at 553 cm-1, and out-of-phase rotation B3g(4) at 555 cm-1 [28]. One can see the atomic displacements in orthorhombic LaMnO3 and YMnO3 ceramics in ref. [38]. The presence of in-phase and out-of-phase modes in YAlO3 implies that YAlO3 has a frustrated structure. In SrTiO3 single crystals, one can see only the second order peaks from c.a. 200 cm-1 to 500 cm-1 and from 600 cm-1 to 800 cm-1 at 300 K since it has an ideal cubic perovskite structure at 300 K. At 8 K, SrTiO3 single crystals, which have tetragonal I4/mcm structure with two SrTiO3 molecular units in the primitive unit cell, show Eg(TO1) at 16 cm-1, A1g(TO1) at 46 cm-1, Eg at 146 cm-1, B2g at 232 cm-1, B1g at 450 cm-1. As for Raman spectra in tetragonal BaTiO3 single crystals, which have tetragonal P4mm structure, E(TO1) at 38 cm-1, the combination of E(TO1) and A1(TO2) at 154 cm-1, A1(TO1) around 176-178 cm-1, E(TO2) around 178-180 cm-1, E(LO2) around 178-180 cm-1, the combination of A1(LO2) and E(LO2) at 182 cm-1, A1(LO2) at 189 cm-1, the combination of E(TO2) and A1(TO1) at 217 cm-1, A1(TO1) around 274-276 cm-1, B1 around 303-304 cm-1, the combination of E(TO4) and E(LO4) at 306 cm-1, E(TO4) around 306-308 cm-1, E(LO4) around 307-308 cm-1, E(LO1) at 466 cm-1, the combination of A1(LO1) and E(LO1) at 469 cm-1, A1(LO1) at 471 cm-1, E(TO3) around 488-498 cm-1, the combination of E(TO3) and A1(TO3) at 495 cm-1, A1(TO3) around 515-518 cm-1, E(LO3) around 708-722 cm-1, the combination of A1(LO3) and E(LO3) at 716 cm-1, and A1(LO3) 720-725 cm-1 have been observed [39, 40]. If YAlO3 seed layers can introduce antiferrodistortion into SrTiO3 thin films, the Raman spectra of SrTiO3 thin films deposited on YAlO3 are like that of I4/mcm structure not like that of P4mm structure.
Brookite TiO2 and anatase TiO2 are also interesting as alternatives to SrTiO3. Anatase TiO2 thin films show the luminescence of self-trapped excitons [41]. The origin of self-trapping of excitons is the lattice relaxation with the electron-phonon coupling energy of 0.9 eV. In rutile thin films as well as in crystals, such emission of self-trapped excitons is not observed. However, it is difficult to obtain highly crystallized brookite or anatase TiO2 thin films. This will be a future work.
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"Comment on the giant photodielectricity in SrTiO3"
Recently, the giant photodielectricity in SrTiO3 at cryogenic temperatures has been discovered.
[1] Photo-induced inhomogeneities may attribute to the giant photodielectricity:
(1) Photocarriers are localized as small polarons with a huge dipole moment owing to an electron-phonon coupling, resulting in the creation of photoinduced polar domains accompanying spatial lattice distortion. Furthermore, the small polaron has a finite lifetime, leading to the redistribution of polar domains through the generation and annihilation of them.
(2) Generally, the electron polaron doping enhances an antiferrodistortive instability accompanying the rotation of TiO6 octahedra, while the hole polaron doping suppresses it. Both the electron and the hole doping suppress the ferroelectric instability. Such a frustration can attribute to the redistribution of polar domains.
(3) Electron correlation increases with the photocarrier generation. This can weaken electron-phonon coupling regionally owing to the screening of Coulomb interaction between small polarons, while can strengthen it owing to Jahn-Teller distortion. Such a frustration can attribute to the redistribution of polar domains. The photoinduced inhomogeneity seems to be an intrinsic effect.
[2] Under the application of electric field, the motion of photocarriers captured by domain boundaries associated with the motion of polar domains may induce the macroscopic ordering motion of large number of atoms and electrons. This macroscopic ordering motion may attribute to the giant photodielectricity.
Generally, the materials with antiferrodistortive MO6 (M = transition metals) lattice structures can have in-phase and out-of-phase MO6 rotation modes, while the materials having the ground-states with cubic structures cannot have those structures. If a slight antiferrodistortion can be introduced into SrTiO3 at room temperature, the giant photodielectricity may be obtained at room temperature.
"Photo-induced phase transition due to a domain boundary generation in perovskites"
Perovskites can be very promising candidate materials for sensor applications, since an exciter such as a molecule or a light quantum can generate a domain boundary in them under a DC electric field causing a phase transition. In the present paper, I concentrate on the detection of light.
Near a meta-stable morphotropic phase boundary, the phase transition tends to occur easily. In fine-grained BaTiO3, an orthorhombic phase and a tetragonal one coexist at room temperature because of internal strain [1]. However, the phase transition in BaTiO3 is too slow - It means that an excited phase is long-lived in BaTiO3. - to utilize it for the detection of matter moving fast. In contrast, strained SrTiO3, in which an antiferrodistortive phase and a ferroelectric one coexist, is expected to realize the detection of matter moving fast.
Epitaxially grown SrTiO3 thin films deposited on LaAlO3 single crystal substrates had been studied by Ryen et al. [2 - 4]; a columnar structure were found to consist of individual subgrains connected by antiphase boundaries or small-angle tilt boundaries, which are associated with interfacial dislocations at the film/substrate interface. Relatively clean boundaries can be promising candidates for the detection of a weak exciter: grain boundaries in YBa2Cu3O7 polycrystalline films, artificial grain boundaries in YBa2Cu3O7 step-edge junctions on MgO substrates, and YBa2Cu3O7 grain-boundary junctions on NdGaO3 or LaAlO3 bicrystal substrates had been investigated as millimeter-wave/submillimeter-wave radiation detectors [5]. To realize the detection of one molecule or one light quantum, it is preferable that the surface of sensing layer is coherently antiferrodistortive without a partial dislocation, such as a threading dislocation and a grain boundary. Threading dislocations or grain boundaries can emerge from misfit dislocations near the film/substrate interface. If the film thickness does not exceed the critical value, misfit dislocations do not occur and the misfit is accommodated by elastic strain. The SrTiO3 film thickness should therefore be less than or equal to 20 nm. If it is possible, the precise control of Sr/Ti ratio should be conducted since the irregular Sr/Ti ratio results in the occurrence of misfit dislocations, as described in refs. [6] and [7].
YAlO3 seed layer may stabilize the ground state in the SrTiO3 thin film against unwanted perturbations, like ninety-degree boundaries in mixed a/c-axis-oriented Yba2Cu3O7 [5] do, because the tilt angle of AlO6 octahedron in YalO3 is larger than that in LaAlO3.
As for a long-range interaction causing the phase transition, I am not sure that its origin is whether the Coulomb interaction or the interaction between a low frequency optic phonon and an acoustic one so-called the phonon waterfall effect [8]. Stock et al. found that the coupling between a transverse acoustic phonon and a transverse optic phonon, which is observed in SrTiO3 single crystals, is weak in Pb(Mg1/3Nb1/3)O3 single crystals, and that the coupling between a transverse acoustic phonon and a diffuse quasielastic component is strong in Pb(Mg1/3Nb1/3)O3 single crystals [9]. The investigation of the origin of the phase transition will be a future work.
References:
[1] G. Arlt, D. Hennings, and G. de With, J. Appl. Phys., 58, 1619, 1985
[2] L. Ryen, E. Olsson, L. D. Madsen, C. N. Johnson, X. Wang, S. N. Jacobsen, U. Helmersson, S. Rudner, and L. -D. Wernlund, Microelectronic Eng., 29, 309, 1995
[3] L. Ryen, E. Olsson, L. D. Madsen, X. Wang, C. N. L. Edvardsson, S. N. Jacobsen, U. Helmersson, S. Rudner, and L. -D. Wernlund, J. Appl. Phys., 83, 4884, 1998
[4] L. Ryen, E. Olsson, L. D. Madsen, C. N. L Johnson, X. Wang, S. N. Jacobsen, U. Helmersson, L. -D. Wernlund, and S. Rudner, Mat. Res. Soc. Symp. Proc., 401, 369, 1996
[5] H. Hilgenkamp and J. Mannhart, Rev. Mod. Phys., 74, 2002, 485
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[8] S. Kamba, M. Kempa, V. Bovtun, J. Petzelt, K. Brinkman, and N. Setter, J. Phys.: Condens. Matter, 17, 3965, 2005
[9] C. Stock, H. Luo, D. Viehland, J. F. Li, I. P. Swainson, R. J. Birgeneau, and G. Shirane, J. Phys. Soc. Jpn., 74, 3002, 2005
"Dynamic modulations of SrTiO3 lattice structures induced by weak external fields"
The generation of a domain boundary, such as a ferroelectric domain in antiferrodistortive surroundings or the rotation of a local electric dipole, is expected to cause gigantic changes of electrical/electronic properties in a perovskite.
It is assumed that an antiferrodistortive instability and a ferroelectric one can coexist in SrTiO3, and that increased hydrostatic compressive pressure enhances the antiferrodistortive instability while suppressing the ferroelectric one [1]. At room temperature, SrTiO3 single crystal has an ideal cubic perovskite structure with space group Pm3{bar}m. On cooling below 300 K, the cubic-tetragonal antiferrodistortive phase transition occurs around 110-105 K. This phase transition originates from the antiphase rotation of TiO6 octahedra along one cubic axis [2]. This modification is described in the Glazer classification by the a0a0c- tilt system [3]. The final result is a tetragonal I4/mcm structure with two SrTiO3 molecular units in the primitive unit cell. The antiferrodistortive phase transition may occur in SrTiO3 thin films deposited on LaAlO3 or on YAlO3 at room temperature. The antiferrodistortive transition near the surface of SrTiO3 single crystals occurs around 150 K, i.e., 40-45 K higher than that of bulk [4]. Furthermore, LaAlO3 single crystals and YAlO3 single crystals are known to have antiferrodistortive tilted AlO6 octahedra at room temperature.
The local charge compensation at the surface may induce a ferroelectric domain in antiferrodistortive SrTiO3 thin films deposited on LaAlO3 or on YAlO3. PbTiO3 thin films having the tetragonal symmetry can form monodomain structures in which the polarization has a component perpendicular to the film plane because the surface charge can be passivated by polar molecules such as OH adsorbates and because the bottom charge can be passivated by conductive electrodes, while forming 180 degrees stripe domains without the charge passivation [5].
LaAlO3 single crystal substrates and YAlO3 single crystal substrates are very expensive. Therefore, I intend to introduce a structural frustration by another method such as the surface charge compensation or the irradiation of light on SrTiO3 thin films having tensile strain. PbTiO3 thin films having tensile strain can form regularly-spaced polar domains in which the polarization is rotated away from the substrate normal, characterizing a low-symmetry phase not observed in the bulk PbTiO3 [6]. Similar low-symmetry phases are believed to be responsible for the large piezoelectric responses observed in Pb(Zr,Ti)O3 and related systems. Room-temperature ferroelectricity is observed in SrTiO3 thin films, in which tensile strain is introduced by DyScO3 substrates [7]. When multidomain structures are introduced into SrTiO3 thin films, weakly clamped polarization components in minor domains can rotate under a weak external field.
References:
[1] W. Zhong and D. Vanderbilt, Phys. Rev. Lett., 74, 2587, 1995
[2] G. Shirane and Y. Yamada, Phys. Rev., 177, 858, 1969
[3] A. M. Glazer, Acta Crystallogr., B 28, 3384, 1972
[4] Z. Salman, R. F. Kiefl, K. H. Chow, M. D. Hossain, T. A. Keeler, S. R. Kreitzman, C. D. P. Levy, P. I. Miller, T. J. Parolin, M. R. Pearson, H. Saadaoui, J. D. Shultz, M. Smadella, D. Wang, and W. A. MacFarlane, Phys. Rev. Lett., 96, 147601, 2006
[5] D. D. Fong, A. M. Kolpak, J. A. Eastman, S. K. Streiffer, P. H. Fuoss, G. B. Stephenson, D. M. Kim, K. J. Choi, C. B. Eom, I. Grinberg, and A. M. Rappe, Phys. Rev. Lett., 96, 127601, 2006
[6] G. Catalan, A. Janssens, G. Rispens, S. Csiszar, O. Seeck, G. Rijnders, D. H. A. Blank, and B. Noheda, Phys. Rev. Lett., 96, 127602, 2006
[7] J. H. Haeni, P. Irvin, W. Chang, R. Uecker, P. Reiche, Y. L. Li, S. Choudhury, W. Tian, M. E. Hawley, B. Craigo, A. K. Tagantsev, X. Q. Pan, S. K. Streiffer, L. Q. Chen, S. W. Kirchoefer, J. Levy, and D. G. Schlom, Nature, 430, 758, 2004
[II] WORK EXPERIENCES, CONCENTRATIONS & ACHIEVEMENTS:
I have worked at 4 electronics companies.
I have researched on dielectric materials and materials used for batteries. The all of materials researched are transition metal oxides.
(4) 2003-present: Taiyo Yuden Co. Ltd.
1 patent
6 papers published
[1] T. HARA, "Electronic structures near surfaces of perovskite type oxides", Materials Chemistry and Physics, 91, 243, 2005
Abstract
This work is intended to draw attention to the origin the electronic structures near surfaces of perovskite type oxides. Deep states were observed by ultraviolet photoelectron spectroscopic measurements. The film thickness dependent electronic structures near surfaces of (Ba0.5Sr0.5)TiO3 thin films were observed. As for the 117-308 nm-thick-(Ba0.5Sr0.5)TiO3 films, deep states were lying at 0.20eV, at 0.55eV, and at 0.85eV below the quasi-Fermi level, respectively. However, as for the 40nm-thick-(Ba0.5Sr0.5)TiO3 film, the states were overlapped. The A-site doping affected electronic structures near surfaces of SrTiO3 single crystals. No evolution of deep states in non-doped SrTiO3 single crystal was observed.%E3%80%80However, the evolution of deep states in La-doped SrTiO3 single crystal was observed.
http://www.sciencedirect.com/science?_ob=ArticleURL&_udi=B6TX4-4F4H9TY-9&_coverDate=06%2F15%2F2005&_alid=424994618&_rdoc=1&_fmt=&_orig=search&_qd=1&_cdi=5580&_sort=d&view=c&_acct=C000050221&_version=1&_urlVersion=0&_userid=10&md5=b43f7c9496a25b9d9910afd60158754f
[2] T. HARA, "Defect-related leakage behaviors and degradation mechanisms of (Ba,Sr)TiO3 films", Integrated Ferroelectrics, 70, 79, 2005
Abstract
The defect-related relaxation behaviors and the leakage of fresh and/or dc-electrically degraded specimens of Au/(Ba0.5Sr0.5)TiO3/Pt capacitors on the TiO2 coated sapphire substrates were investigated. The relaxation behaviors of (Ba0.5Sr0.5)TiO3 films are assumed to be the electron-detrapping in the depletion layer. The local electric field enhancement due to oxygen vacancies near the Pt/(Ba0.5Sr0.5)TiO3 interface, which is estimated from the Poisson equation, was assumed to be sufficiently high as a cause of tunneling conduction, and assumed to be a cause of dc-electrical degradation.
[3] T. HARA, "Electron-detrapping from localized states in the band gap of (Ba,Sr)TiO3", Solid State Communications, 132, 109, 2004
Abstract
We tried to relate the relaxation currents of Pt/62nm-thick-(Ba0.5Sr0.5)TiO3/Pt capacitors to the results of ultraviolet photoemission spectroscopic measurements for the 62nm-thick-(Ba0.5Sr0.5)TiO3/Pt specimens. The slowest relaxation (159-313s at applied voltages of 1.5-3V, and at a measuring temperature of 40%CB%9AC) and the relatively faster relaxations (4.92-5.43s and 0.25-0.46s) were assigned as the electron-detrapping from the localized state at 0.80eV below the quasi-Fermi level, from the localized state at 0.55eV below the quasi-Fermi level, and from the localized state at 0.30eV below the quasi-Fermi level, respectively. The decreasing of relaxation time in accordance with the increasing of bias voltage is probably due to the decreasing of depletion width. The decreasing of depletion width is probably due to the detrapping of electrons from deep localized states in accordance with the downward bending of quasi-Fermi level in the depletion layer, and due to the decreasing of relative dielectric constant in the depletion layer in accordance with the increasing of bias voltage.
http://www.sciencedirect.com/science?_ob=ArticleURL&_udi=B6TVW-4CXHGFN-1&_coverDate=10%2F01%2F2004&_alid=424994822&_rdoc=1&_fmt=&_orig=search&_qd=1&_cdi=5545&_sort=d&view=c&_acct=C000050221&_version=1&_urlVersion=0&_userid=10&md5=92717fcf01894cfde1a0dbacadd4b042
[4] T. HARA, "Electrical characteristics of (Ba,Sr)TiO3 films accounted by partially depleted model", Microelectronic Engineering, 75, 316, 2004
Abstract
We investigated the leakage current versus voltage (I-V) characteristics, the capacitance versus thickness of (Ba0.5Sr0.5)TiO3 film (C-t) characteristics, and the relaxation currents of sputtered (Ba0.5Sr0.5)TiO3 films with the thickness of 40-166nm. The I-V characteristics can be explained by the partially depleted model especially when the thickness of (Ba0.5Sr0.5)TiO3 film exceeds 62nm. The C-t characteristics indicate that the relative dielectric constant in the internal layer (out of the depletion layer) does not change by applied voltages. This can be explained by assuming that the electric field is concentrated at the partially depleted layer, and that the relative dielectric constant in the depletion layer decreases in accordance with the increasing of applied voltage. The relaxation currents may be explained by assuming that the relative dielectric constant in the depletion layer decreases in accordance with the increasing of applied voltage.
http://www.sciencedirect.com/science?_ob=ArticleURL&_udi=B6V0W-4CVR3VP-4&_coverDate=09%2F30%2F2004&_alid=424994935&_rdoc=1&_fmt=&_orig=search&_qd=1&_cdi=5657&_sort=d&view=c&_acct=C000050221&_version=1&_urlVersion=0&_userid=10&md5=321ce675fb781caf6d3bc48be92e88c8
[5] T. HARA, "Intentionally inserted oxygen depleted (Ba0.5Sr0.5)TiO3 layers as a model of DC-electrical degradation", IEEE Trans. on Device and Materials Reliability, 4, 670, 2004
Abstract
The relative dielectric constant versus voltage (%CE%B5r%E3%80%80-V) characteristics and the current density versus electric field (J-E) characteristics of (Ba0.5Sr0.5)TiO3 films, which have intentionally inserted oxygen depleted layers near the bottom electrodes, were investigated as a model of dc-electrical degradation phenomena. Our investigation demonstrated that the intentionally inserted oxygen depleted layer is the cause of the tunneling conduction.
http://ieeexplore.ieee.org/Xplore/login.jsp?url=/iel5/7298/30213/01388439.pdf?arnumber=1388439
[6] T. HARA, "Leakage behavior of DC electrically degraded (Ba,Sr)TiO3 thin films", IEEE Trans. on Device and Materials Reliability, 4, 268, 2004
Abstract
The phenomena of dc electrical degradation of (Ba0.5Sr0.5)TiO3 thin films was studied. From our experimental and analytical results of current versus voltage (I-V) characteristics, it was shown that the degraded devices exhibited analogous leakage behaviors with the devices which have thin intercalated (Ba0.5Sr0.5)TiO3 layers with intentionally introduced oxygen vacancies between cathodes and thick (Ba0.5Sr0.5)TiO3 layers without intentionally introduced oxygen vacancies. This could be explained by assuming that oxygen vacancies accumulate at the interfaces between the cathodes and the (Ba0.5Sr0.5)TiO3 films after fatigue.
http://ieeexplore.ieee.org/Xplore/login.jsp?url=/iel5/7298/29219/01318632.pdf?arnumber=1318632
(3)1998-2003: Kyocera Corp.,
(Ba,Sr)TiO3 Thin Film Capacitor deposited using sputtering technique:
A tetragonal phase in BaTiO3-rich clusters having compressive strain and an orthorhombic phase in SrTiO3-rich clusters having tensile strain can be present. Two phases can coexist in (Ba,Sr)TiO3.
(Ba,Sr)TiO3 thin films deposited on platinized sapphire substrates can have higher permittivity along the surface-normal direction than the films deposited on platinized Si substrates, since the in-plane tensile strain in films grown on sapphire substrates is smaller than that in films grown on Si substrates. It is assumed that the increase in in-plane tensile strain causes the increase in the population of orthorhombic phase.
The films deposited under high pressures and low RF powers have better crystallinitiy than that in the films deposited under low pressures and high RF powers. However, the in-plane tensile strain can be increased under relatively high pressures and low RF powers. The DC-electrical degradation appears to be accelerated in films with large in-plane tensile strain. This may be due to the faster diffusion of oxygen vacancies in films with large in-plane tensile strain than that in films with small in-plane tensile strain [Thermally induced tensile stress may be applicable to fuel cells. It is widely known that the thermally induced tensile stress is the origin of an enhanced oxygen diffusion rate in ultrathin gate oxides. Furthermore, an enhanced electrochemical forming in RRAMs may be attributed to the thermally induced tensile stress.].
In contrast, no degradation in permittivity was observed even in films with large in-plane tensile strain. This is probably due to the polarization orientation distribution in (Ba,Sr)TiO3, in which 18 variants may be permitted by symmetry, and due to its resistive character against the phase separation. It is assumed that the degradation in permittivity observed in tetragonal dielectrics is due to the pinning of 180 degrees domain wall by charged oxygen vacancies at the dielectric-electrode interface.
Pb(Mg1/3,Nb2/3)O3-PbTiO3 Thin Film Capacitor deposited using sol-gel technique:
Minor substitutions of Ti4+ for Mg2+/Nb5+ in Pb(Mg1/3,Nb2/3)O3 allow the coexistence of a rhombohedral phase, a tetragonal one introduced, and an orthorhombic/monoclinic one introduced. This modification of Pb(Mg1/3,Nb2/3)O3 with Ti4+ results in high relative dielectric constant of 4000 even in thin films.
The dc-electrical degradation of insulation resistance was observed. Contrary to (Ba,Sr)TiO3, the degradation in permittivity was observed in Pb(Mg1/3,Nb2/3)O3-PbTiO3. This is probably due to the separation between a tetragonal phase and a rhombohedral one in Pb(Mg1/3,Nb2/3)O3-PbTiO3, namely, due to the pinning of 180 degrees domain wall by charged oxygen vacancies at the dielectric-electrode interface.
56 patents
(2)1997-98: Sony Energytech Co. Ltd.,
(1)1993-97: Izumo Murata Mfg. Co. Ltd.,
(Sr,Pb,Ca)TiO3 Ceramic Capacitor:
Minor substitutions of Pb2+ for Sr2+ in SrTiO3 allow the tetragonal deformation. And minor substitutions of Ca2+ for Sr2+ in SrTiO3 allow the TiO6 tilting. This modification of SrTiO3 with Pb2+ and Ca2+ results in temperature-independent high relative dielectric constant of 2000.
The polarization orientation distribution in (Sr,Pb,Ca)TiO3, where 18 variants are permitted by symmetry, are not as important as it would be, e.g., for tetragonal BaTiO3, where only six variants are permitted by symmetry. This effect can improve dielectric responses at high frequencies.
Under dc electrical loading for a few seconds, the slight decrease in permittivity was observed in (Sr,Pb,Ca)TiO3. This is probably due to the longitudinal ferroelectric polarization. Heat treatments above 400 K rejuvenate the permittivity in (Sr,Pb,Ca)TiO3.
The permittivity in (Sr,Pb,Ca)TiO3 increases with increasing quasi-transversally compressive pressure. However, the temperature dependence in permittivity increases with increasing the pressure. This is probably due to the increase in the population of the tetragonal phase.
Under dc electrical loading at high humidity, the degradation of insulation resistance was observed. %E3%80%80The degradation is due to the accumulation of oxygen vacancies at the cathode-dielectric interface accompanying the H2 generation.
SrTiO3-SrZrO3 Ceramic Capacitor:
Not Sr(Ti,Zr)O3 but SrTiO3&SrZrO3 mixture used in high voltage and high reliability capacitor applications.
I think that Sr(Ti,Zr)O3 has a band gap value close to that of SrTiO3.
BaTiO3-Ba(Ti1-x,Nb0.66x,Co0.33x)O3 Ceramic Capacitor with a core-shell structure:
Minor substitutions of Nb5+/Co2+ for Ti4+ in BaTiO3 at the surfaces of BaTiO3-grains allow the coexistence of a tetragonal phase and an orthorhombic one in BaTiO3-cores. This is due to the decrease in grain size. This modification of BaTiO3 with Nb5+ and Co2+ results in temperature-independent high relative dielectric constant of 4000.
SiO2 dopant play the role of sintering aid to enhance densification of BaTiO3-Ba(Ti1-x,Nb0.66x,Co0.33x)O3 without change in the stoichiometry. Transmission electron microscopic study can observe the SiO2 precipitating at triple junctions.
Mn2+ dopant incorporated into Ti4+ site of BaTiO3 plays an acceptor role, which compensate for the number of conduction electron donated by oxygen vacancies forming Ba(Ti1-x, Mnx)O3-x contributing to the improvement of insulation resistance and tan%CE%B4 especially when the ceramic is sintered in a reducing condition. Not only Mn2+ but Mn3+/4+ can be present when the ceramic is sintered in an oxdizing condition.
Nb5+ dopant incorporated into Ti4+ site of BaTiO3 plays a donor role providing conduction electrons when low-dose doping is conducted. However, high-dose doping of Nb5+ can generate Ba vacancies contributing to the improvement of insulation resistance especially when the ceramic is sintered in an oxidizing condition. The generation of Ba vacancies results in obtaining the high (soft) piezoelectricity. It should be noted that The Ba vacancies cannot induce high piezoelectricity in a tetragonal BaTiO3. Therefore, a tetragonal and an orthorhombic phase should coexist to obtain the Ba-vacancies-induced high piezoelectricity.
(Ba, Ca)(Ti, Zr, Sn)O3 Ceramic Capacitor:
Minor substitutions of Ca2+ for Ba2+ in BaTiO3 allow the coexistence of tetragonal and orthorhombic phases. This modification of BaTiO3 with Ca2+ decreases the temperature-dependence in the relative dielectric constant. Minor substitutions of Zr4+ (or Sn4+) for Ti4+ in BaTiO3 allow the coexistence of tetragonal, orthorhombic, and rhombohedral phases, known as the pinching effect. This modification of BaTiO3 with Zr4+ (or Sn4+) results in the enhancement of relative dielectric constant at room temperature. The polarization orientation distribution in (Ba, Ca)(Ti, Zr, Sn)O3, where 26 variants are permitted by symmetry, are not as important as it would be, e.g., for tetragonal BaTiO3, where only six variants are permitted by symmetry. This effect can improve dielectric responses at high frequencies.
BaTiO3-BaZrO3 Ceramic Capacitor:
Not Ba(Ti,Zr)O3 but BaTiO3&BaZrO3 mixture used in high voltage and igh reliability capacitor applications.
Under ac electrical loading, the degradation of insulation resistance was observed.%E3%80%80The degradation is due to the accumulation of oxygen vacancies at both electrode-dielectric interfaces.
1 patent
[III] ACADEMIC QUALIFICATIONS:
1991(BS) & 1993(MS), School of Science, Department of Chemistry, Hiroshima University |
Updated: 2008-04-16 |
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