JP2018184336A - Composite ceramics and its production method - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a material which can exhibit excellent thermoelectric conversion characteristics especially when used as a material for a thermoelectric element.SOLUTION: This invention relates to composite ceramics, comprising: an oxide which has a Perovskite type crystal structure expressed by the following general formula (1): ABO; and a carbon material, wherein in the formula (1), A and B each express an element which can form a Perovskite type crystal structure, provided that at least one thereof is a metallic element; and δ expresses the amount of oxygen deficiencies, provided that it is the number of 0 or more but 1 or less.SELECTED DRAWING: Figure 2

Description

The present invention relates to a composite ceramic and a manufacturing method thereof. More specifically, the present invention relates to a composite ceramic that can be suitably used as a material for a thermoelectric conversion element and a method for producing the same.

Composite ceramics are composites of ceramics combined with fibers, particles, etc. to improve properties such as strength, and are used in various fields such as semiconductors, batteries, automobiles, information communication, industrial equipment, and medical care. .

By the way, in recent years, the needs for energy harvesting (energy harvesting), such as geothermal / hot spring thermal power generation using thermoelectric conversion elements, solar thermal power generation, waste heat power generation in factories and automobiles, etc., have increased due to the increase in ecology movement and energy saving. Therefore, attention has been paid to thermoelectric conversion elements having the ability to convert heat into electricity.

As materials for conventional thermoelectric conversion elements (thermoelectric materials), for example, materials such as Te-based and Se-based skutterudites and intermetallic compounds have been developed, but most of these materials are toxic and rare. Have the problem of instability (see Non-Patent Document 1, for example). In view of this situation, despite the fact that oxides generally have low carrier mobility and high thermal conductivity, the possibility of using oxides as thermoelectric materials has been studied. As typical p-type materials, there are NaCo ₂ O ₄ and Ca ₃ Co ₄ O _9, which are very promising particularly in a high temperature region because of the layered structure (see, for example, Non-Patent Documents 2 and 3). .

In addition, for example, SrTiO ₃ (STO) bulk ceramics having oxygen deficiency have been reported (for example, see Non-Patent Documents 4 to 7). However, it is usually prepared by annealing in a reducing gas (eg, H ₂ / Ar gas) at a high temperature after sintering, which is generally a sample because it is a process in which diffusion from the surface to the inside occurs. It is unavoidable that a gradient distribution of oxygen vacancies occurs.

K. Koumoto, R. Funahashi, E. Guilmeau, Y. Miyazaki, A. Weidenkaff, YF Wang, et al., Thermoelectric ceramics for energy harvesting, J. Am. Ceram. Soc. 96 (1) (2013) 1- twenty three I. Terasaki, Y. Sasago, K. Uchinokura, Large thermoelectric power in NaCo2O4 single crystals, Phys. Rev. B 56 (20) (1997) 12685-12687 AC Masset, C. Michel, A. Maignan, M. Hervieu, O. Toulemonde, F. Studer, et al., Misfit-layered cobaltite with an anisotropic giant magnetoresistance: Ca3Co4O9, Phys. Rev. B 62 (1) (2000 ) 166-175 J. Liu, CL Wang, WB Su, HC Wang, P. Zheng, JC Li, et al., Enhancement of thermoelectric efficiency in oxygen-deficient Sr1-xLaxTiO3-δ ceramics, Appl. Phys. Lett. 95 (16) ( 2009) 3 AV Kovalevsky, AA Yaremchenko, S. Populoh, A. Weidenkaff, JR Frade, Effect of a-site cation deficiency on the thermoelectric performance of donor-substituted strontium titanate, J. Phys. Chem. C 118 (9) (2014) 4596 -4606 L. Sagarna, S. Populoh, A. Shkabko, J. Eilertsen, AE Maegli, R. Hauert, et al., Influence of the oxygen content on the electronic transport properties of SrxEu1-xTiO3-d, J. Phys. Chem. C 118 (15) (2014) 7821-7831 AA Yaremchenko, S. Populoh, SG Patricio, J. Macias, P. Thiel, DP Fagg, et al., Boosting thermoelectric performance by controlled defect chemistry engineering in Ta-Substituted strontium titanate, Chem. Mat. 27 (14) (2015 ) 4995-5006

In order to obtain large electric power as a composite ceramic, particularly when used as a material for a thermoelectric conversion element, it has been desired to have excellent thermoelectric conversion characteristics.

This invention is made | formed in view of the said present condition, and it aims at providing the material which can exhibit the thermoelectric conversion characteristic which was excellent when used as a material of a thermoelectric conversion element especially.

The present inventor has studied various materials that can exhibit excellent thermoelectric conversion characteristics, particularly when used as a material of a thermoelectric conversion element, paying attention to a perovskite crystal structure having high temperature stability, and an oxidation having a perovskite crystal structure. Composite ceramics in which a carbon material was combined with the object was formed. The present inventors have the possibility that this composite ceramic material can exhibit a sufficiently excellent strength by introducing a carbon material, which can be applied to various applications, and is particularly used as a material for a thermoelectric conversion element. Found that the carrier density, electrical conductivity, and power factor are sufficiently large and that the thermal conductivity is sufficiently low, and that excellent thermoelectric conversion characteristics can be exhibited. I came up with the idea that I could do it. Furthermore, the present inventors have found that this composite ceramic is excellent in hardness and wear resistance and can be applied to various applications, and have reached the present invention.

That is, the present invention provides the following general formula (1):
ABO _{3 ± δ} (1)
(In the formula (1), A and B represent an element capable of forming a perovskite crystal structure, at least one is a metal element, and δ represents an oxygen deficiency and is a number of 0 or more and 1 or less. ) And a composite ceramic material of an oxide having a perovskite crystal structure and a carbon material.

The composite ceramic of the present invention is excellent in strength and wear resistance due to the composite of the carbon material, and can exhibit excellent thermoelectric conversion characteristics when used as a material of a thermoelectric conversion element.

It is a graph which shows the XRD pattern of pure STO, pure STO with oxygen deficiency (STO-H), and reduced graphene oxide (RGO) / STO complex. (A) pure STO, (b) pure STO-H, and (c) electron backscatter diffraction (EBSD) analysis of RGO / STO complex, and (d) pure STO, pure STO It is a figure which shows the particle size distribution of the oxide in -H and a RGO / STO composite. (A) TEM image of RGO / STO composite at low magnification, (b) TEM image of RGO / STO composite at high magnification, (c) RGO / STO composite of region surrounded by solid line in (b) HRTEM image of the body (showing RGO on the grain boundary) and HRTEM image of RGO / STO complex (showing RGO inside the particle) in the region surrounded by the broken line in (d) and (b) is there. It is a graph which shows (a) electrical conductivity, (b) Seebeck coefficient (Seebeck coefficient), and (c) power factor of pure STO, pure STO-H, and RGO / STO composite. FIG. 2 is a graph of electron energy loss spectroscopy (EELS) spectrum, (a) pure STO Ti-L edge, (b) pure STO OK edge, (c) RGO / STO composite Ti-L edge. And (d) shows the OK edge of the RGO / STO complex. It is a graph which shows (a) total thermal conductivity, (b) lattice thermal conductivity, and (c) ZT value of pure STO, pure STO-H, and RGO / STO composite. It is a graph which shows the zeta potential of modified STO and graphene oxide (GO) with respect to pH in an aqueous dispersion. (A) TEM image of the original STO powder at low magnification and (b) high magnification. It is a TEM image of GO / modified STO mixed (hybrid) powder at low magnification, (a) black arrow indicates the form of GO, and (b) black arrow indicates GO covering the surface of the modified STO particles. It is a graph which shows the Raman spectrum of each GO / STO mixed powder and RGO / STO composite. (A) Pure STO, (b) SEM image of RGO / STO composite. It is a graph which shows the electrical conductivity of pure STO. RGO / _Al 2 _{O 3} complex is a graph showing the electrical conductivity of the (carbon content 0.6 vol%). The RGO / STO composite is a low-angle annular dark field (LAADF) image in which (a) RGO is included in grain boundaries or (b) RGO is not included in grain boundaries. 6 is a graph plotting the average dynamic friction coefficient against the number of sliding times for the RGO / STO composites 2 to 5 produced in Examples 2 to 5 and the pure STO produced in Comparative Example 2. FIG. It is the result of having performed the Raman spectroscopic analysis about the RGO / STO composite 2 produced in Example 2. FIG. It is the result of having conducted the Raman spectroscopic analysis about the RGO / STO composite 3 produced in Example 3. FIG. It is the result of having performed the Raman spectroscopic analysis about the RGO / STO composite 4 produced in Example 4. FIG. It is the result of having performed the Raman spectroscopic analysis about the RGO / STO composite 5 produced in Example 5. FIG. It is the result of having performed the Raman spectroscopic analysis about the pure STO produced in the comparative example 2. FIG.

The present invention is described in detail below.
A combination of two or more preferred embodiments of the present invention described below is also a preferred embodiment of the present invention.

<Composite ceramics>
The composite ceramic of the present invention is a composite of an oxide having a perovskite crystal structure represented by the general formula (1) and a carbon material. By compounding a carbon material with the oxide, it is possible to generate pores in the composite to improve the thermoelectric conversion characteristics, and to increase the ceramic particle growth that normally occurs during the composite. The thermal conductivity can be kept low by keeping the particle size small.
The thermoelectric conversion characteristics of the composite ceramics are evaluated using a dimensionless figure of merit ZT value which is the product of the figure of merit Z (K ⁻¹ ) and the absolute temperature (K) in the thermoelectric conversion layer in the thermoelectric conversion element. The ZT value is represented by the following formula.
ZT = (S ² · σ / κ) · T
In the above formula, S represents the Seebeck coefficient (VK ⁻¹ ). If S is a positive value, it is a p-type element, and if it is a negative value, it is an n-type element. σ represents electrical conductivity (conductivity) (Ω ⁻¹ m ⁻¹ ). κ represents thermal conductivity (Wm ⁻¹ K ⁻¹ ). T represents the absolute temperature (K).
Theoretically, when waste heat is recycled by applying a thermoelectric conversion material having a large ZT value, the efficiency of fuel is greatly improved.
S ² · σ is also referred to as a power factor (PF). The larger the power factor (PF), the larger the ZT value and the better the thermoelectric conversion characteristics.
The ZT value and PF are measured and calculated by the method of the embodiment.

The composite ceramic of the present invention may be composited using one kind of carbon material and one oxide each having a perovskite crystal structure, or may be composited using two or more kinds. Also good. Moreover, other components other than the carbon material and the oxide having a perovskite crystal structure may be included.

The carbon material preferably has, for example, a graphene skeleton. By using a carbon material having a graphene skeleton, the composite ceramic of the present invention is more suitable as a material for a thermoelectric conversion element. This is presumably because the carbon material having a graphene skeleton has a large aspect ratio and can easily form a good dispersion state in a small amount.
The carbon material having the graphene skeleton has carbon (C) bonded by sp ² bond, and is not particularly limited as long as the carbon is arranged in a plane, but has carbon bonded to oxygen (O). It is preferable. In other words, the carbon material in the composite ceramic of the present invention is preferably graphite oxide. More preferably, it is graphene oxide in which oxygen is bonded to carbon of graphene.
In general, graphene refers to a sheet composed of a single layer in which carbon atoms bonded by sp ² bonds are arranged in a plane, and a graph laminated with a large number of graphene sheets is called graphite. The carbon material having the above and the graphene oxide in the present invention include not only a sheet consisting of only one layer but also those having a structure in which about 2 to 100 layers are laminated. The same applies to the graphene crystal described later.
Graphene oxide having such a laminated structure can be obtained by treating graphite with a known oxidizing agent, for example.
The carbon material according to the present invention may further have a functional group such as a carboxyl group, a hydroxyl group, a sulfur-containing group, and an alicyclic epoxy group.

The carbon material having the graphene skeleton preferably has a ratio of the number of carbon atoms to the number of oxygen atoms of 0.7 to 100. It is preferable for the ratio of the number of carbon atoms to the number of oxygen atoms to be in such a range because the carbon material exhibits good conductivity.
The ratio of the number of carbon atoms to the number of oxygen atoms of the carbon material having a graphene skeleton can be confirmed by the ratio of the total peak area of the O1s region and the total peak area of the C1s region obtained by XPS measurement.

The carbon material preferably has a specific surface area of 1 m ² / g or more, and more preferably 10 m ² / g or more. By using a carbon material having such a specific surface area, high dispersibility can be maintained when it is combined with an oxide. The upper limit of the specific surface area is not particularly limited, but may be, for example, 2500 m ² / g or less.
The specific surface area can be measured by a specific surface area measuring device by a nitrogen adsorption BET method.

The carbon material preferably has an average particle size of 1000 μm or less, and particularly preferably 500 μm or less. The average particle size is preferably 5 nm or more, and particularly preferably 1000 nm or more.
The average particle diameter can be measured by using a particle size distribution measuring device or observing using an electron microscope and an atomic force microscope.

Examples of the shape of the carbon material include fine powder, powder, granule, granule, scale, polyhedron, rod, and curved surface. In addition, the particles having an average particle size in the above-described range are, for example, a method of pulverizing particles with a ball mill or the like, dispersing the obtained coarse particles in a dispersant to obtain a desired particle size, and drying to solidify, In addition to the method of selecting the particle diameter by sieving the coarse particles, etc., it is possible to optimize the preparation conditions at the stage of producing the particles to obtain (nano) particles having a desired particle diameter.

The carbon material may be reduced at the time of compounding and reduced in the composite ceramic of the present invention. For example, when the carbon material is graphite oxide, in the composite ceramic of the present invention, the graphite oxide may be reduced to form reduced graphite oxide.

The composite ceramic of the present invention preferably has a carbon content of 5% by volume or less. By reducing the carbon content in this way, it is possible to sufficiently suppress the formation of a network of carbon materials, to maintain the Seebeck coefficient at a sufficiently high value, and to improve the thermoelectric conversion characteristics of the composite ceramic of the present invention. Can be. The carbon content is more preferably 3% by volume or less, and still more preferably 1% by volume or less.
Further, the composite ceramic of the present invention is capable of uniformly distributing the carbon material in the composite while demonstrating good electrical conductivity, and an element made using the composite ceramic has good strength. In view of the above, the carbon content is preferably 0.1% by volume or more, more preferably 0.2% by volume or more, and still more preferably 0.3% by volume or more.

Also from the viewpoint of hardness and wear resistance, the composite ceramic of the present invention preferably has a carbon content of 3% by volume or less. By reducing the carbon content in this way, it is possible to improve the hardness and wear resistance of the composite ceramic of the present invention without excessively suppressing the sintering of the ceramics. From the viewpoint of hardness and wear resistance, the carbon content is more preferably 1% by volume or less, and still more preferably 0.5% by volume or less.
In addition, if the composite ceramic of the present invention has too little carbon content, the ceramics may be sufficiently sintered and the hardness and wear resistance may be lowered. It is preferably at least 01% by volume, more preferably at least 0.03% by volume, and even more preferably at least 0.05% by volume.
The said carbon content is measured by the method of an Example.

The oxide according to the present invention has the following general formula (1):
ABO _{3 ± δ} (1)
(In the formula (1), A and B represent an element capable of forming a perovskite crystal structure, at least one is a metal element, and δ represents an oxygen deficiency and is a number of 0 or more and 1 or less. And a perovskite crystal structure represented by:

In the above formula (1), A and B represent elements that can form a perovskite crystal structure, and are not particularly limited as long as one is a metal element, but each preferably represents a metal element.
A preferably represents an alkaline earth metal element or a rare earth metal element. Examples of the alkaline earth metal element include Ca, Sr, Ba and the like.
Examples of rare earth metal elements include Sc, Y, and lanthanoids.

In the above formula (1), B preferably represents Ti, V, Cr, Mn, Fe, Co, Ni, Zr, Nb, Mo, Hf, Ta, or W, and Ti, V, Mn, Fe More preferably, it represents Co, Ni, Zr, Nb, or Mo.

Among these, the oxide is represented by SrTiO _{3 ± δ} (δ represents the amount of oxygen deficiency and is a number of 0 or more and 1 or less) because of its large effective carrier mass, high temperature stability, and tolerance to doping. It is particularly preferable to have a perovskite crystal structure. This oxide, as an n-type thermoelectric material, is excellent in effective carrier mass, high-temperature stability, and doping tolerance, but exhibits the effect of the present invention more remarkably by being compounded in the composite ceramic of the present invention. it can. As a typical perovskite with ABO ₃ composition, undoped STO is an intrinsic insulator. Electrical conductivity is imparted by introducing a different dopant into the A site or the B site or by reduction treatment for generating oxygen vacancies. In the composite ceramic of the present invention, for example, oxygen vacancies are generated during composite formation, so that the carrier density is particularly high and high electrical conductivity can be exhibited.

The average particle size of the oxide in the composite ceramic of the present invention is preferably 10 μm or less. When the average particle size of the composite ceramic of the present invention is suppressed in this way, the thermal conductivity can be made sufficiently lower, and it becomes more suitable as a thermoelectric conversion material. The average particle diameter is more preferably 5 μm or less, still more preferably 2 μm or less, and particularly preferably 1 μm or less.
The lower limit of the average particle diameter is not particularly limited, but is usually 0.1 μm or more.
The said average particle diameter can be measured by EBSD by the method of an Example.

In the composite ceramic of the present invention, the content of the oxide having the perovskite crystal structure is preferably 95 to 99.9% by volume. By setting it as such a content rate, composite ceramics will become the thing excellent in the thermoelectric conversion characteristic. The ratio of the oxide having the perovskite crystal structure is more preferably 97 to 99.8% by volume, and still more preferably 99 to 99.7% by volume.

From the viewpoint of hardness and wear resistance, the composite ceramic of the present invention preferably contains 97 to 99.99% by volume of the oxide having the perovskite crystal structure. By setting it as such a content rate, composite ceramics will be excellent in hardness and abrasion resistance. From the above viewpoint, the ratio of the oxide having the perovskite crystal structure is more preferably 99 to 99.97% by volume, and still more preferably 99.5 to 99.95% by volume.

In the composite ceramic of the present invention, the ratio I (2D) / I (G) of the peak intensity I (2D) of the 2D band and the peak intensity I (G) of the G band measured by Raman spectroscopy is 0.1 or more. Is preferred. Such a composite ceramic is one in which the thin layer carbon material is more uniformly dispersed in the composite ceramic, and therefore, the thermoelectric conversion characteristics of the present invention can be exhibited more remarkably. The ratio I (2D) / I (G) is more preferably 0.2 or more, and further preferably 0.3 or more.
The upper limit of the ratio I (2D) / I (G) is not particularly limited, but is usually 2 or less.
The ratio I (2D) / I (G) is calculated by measuring a Raman spectrum by the method of the example.

The composite ceramic of the present invention preferably has a carrier density at 300K of 5 × 10 ¹⁹ cm ⁻³ or more. The carrier density is more preferably 1.0 × 10 ²⁰ cm ⁻³ or more.
The upper limit of the carrier density is not particularly limited, but is usually 1.0 × 10 ²² cm ⁻³ or less.
The carrier density is measured by the method of the example.

The composite ceramic of the present invention preferably has an electric conductivity at 473 K of 2000 S / m or more (2000 A / V · m or more), more preferably 3000 S / m or more.
The upper limit of the electrical conductivity is not particularly limited, but is usually 100,000 S / m or less.
The electrical conductivity is measured by the method of the example.

The composite ceramic of the present invention preferably has a total thermal conductivity κ _total, which is the sum of the lattice thermal conductivity κ _l and the electronic thermal conductivity κ _{e at} 373 K, of 10 W / m · K or less, and 8 W / It is more preferably m · K or less, and further preferably 6 W / m · K or less.
The lower limit of the _total thermal conductivity κ _total is not particularly limited, but is usually 0.1 W / m · K or more.
The total thermal conductivity κ _total is measured by the method of the example.

In the composite ceramic of the present invention, the power factor of thermoelectric conversion at 473 K is preferably 4 × 10 ⁻⁴ W / m · K ² or more, and preferably 5 × 10 ⁻⁴ W / m · K ² or more. Is more preferable.
The upper limit of the power factor is not particularly limited, but is usually 1 × 10 ⁻² W / m · K ² or less.
The power factor is measured and calculated by the method of the embodiment.

The composite ceramics of the present invention may be able to exhibit excellent strength due to the introduction of carbon, and may be applied to various uses. Among them, thermoelectric conversion materials; pump parts; cutting tools; seals Rings; bearings; pulverized parts such as balls and blades of ball mills; medical instruments; surgical instruments; jigs; electrical equipment; parts for devices; preferably used as building material products, and more preferably used as thermoelectric conversion materials. In the case of a thermoelectric conversion material, for example, it can be suitably used for one or both of an n-type thermoelectric conversion material and a p-type thermoelectric conversion material, but among them, it is more preferably used as an n-type thermoelectric conversion material. Normally, depending on the type of oxide used for the composite, it is determined whether it is an n-type thermoelectric conversion material or a p-type thermoelectric conversion material.
The thermoelectric conversion material is a material generally used for a thermoelectric conversion layer or the like in a thermoelectric conversion element that converts heat into electric power, and may be the composite ceramic itself of the present invention, other resins, Various additives and the like may be further included as necessary.

(Method for producing composite ceramics of the present invention)
The composite ceramic of the present invention can be produced by mixing the carbon material and the oxide having the perovskite crystal structure by using various mixing methods. From the viewpoint of suppressing the heating and preventing the burning of carbon sufficiently without the need for heating for a long period of time, the manufacturing method of the present invention has a carbon material and a perovskite crystal structure. It is preferable to include a step of sintering an oxide, and it is more preferable to include a step of sintering by a pulse current pressure sintering method. In other words, the present invention is a method for producing a composite ceramic, which comprises the following general formula (1):
ABO _{3 ± δ} (1)
(In the formula (1), A and B represent an element capable of forming a perovskite crystal structure, at least one is a metal element, and δ represents an oxygen deficiency and is a number of 0 or more and 1 or less. It is also a method for producing a composite ceramic, comprising a step of sintering an oxide having a perovskite type crystal structure represented by) and a carbon material by a pulse current pressure sintering method.

The temperature in the sintering step is preferably 500 ° C. or higher. More preferably, it is 700 degreeC or more, More preferably, it is 1000 degreeC or more. In addition, the temperature is preferably 2000 ° C. or lower, and more preferably 1500 ° C. or lower. The time for performing the sintering step is preferably 1 minute to 5 hours. More preferably, it is 2 minutes-3 hours, More preferably, it is 3 minutes-1 hour.
Sintering may be performed in air or in an inert gas atmosphere such as nitrogen, helium, or argon. Further, the pressure condition is not particularly limited, and it can be performed under pressure, normal pressure, and reduced pressure.

The production method of the present invention preferably further includes a step of uniformly mixing the carbon material and the oxide having a perovskite crystal structure by heteroaggregation before the sintering step. The mixing step can be performed, for example, by continuously or intermittently injecting a water dispersion of a negatively charged carbon material into water in which a positively charged oxide is dispersed. This step is particularly effective when the oxide can form a positive surface charge (zeta potential) sufficient in water to attract the negatively charged carbon material.
As the raw material oxide carry a positive surface charge in water, before the mixing step, oxide (3-aminopropyl) NH ²⁺ on oxide surfaces by reacting an amino group-containing silane compounds such as trimethoxysilane It is preferable to graft a group having a group having a positive surface charge such as.

<Thermoelectric conversion element>
A thermoelectric conversion element is generally an element that converts heat into electric power, which is produced by combining two types of thermoelectric conversion materials, p-type and n-type.
As described above, when a large electric power is obtained using a thermoelectric conversion material, a thermoelectric conversion element in which both an n-type thermoelectric conversion material and a p-type thermoelectric conversion material are sandwiched between an insulating material on a high temperature side and an insulating material on a low temperature side. Will be configured. The thermoelectric conversion material of the present invention includes the composite ceramic of the present invention described above, and can be used for one or both of an n-type thermoelectric conversion material and a p-type thermoelectric conversion material.
Such a thermoelectric conversion element including the thermoelectric conversion material of the present invention is also one aspect of the present invention.

The thermoelectric conversion element is usually configured to include an electrode in addition to the thermoelectric conversion material. The material of the electrode is not particularly limited, and examples thereof include transparent electrodes such as indium tin oxide (ITO) and indium zinc oxide (IZO), metals, and conductive carbon materials.

The thermoelectric conversion element of the present invention can be used in various forms. Examples of applications of the thermoelectric conversion element of the present invention include geothermal / hot spring thermal generators, solar thermal generators, waste heat generators such as factories and automobiles, body temperature, etc. Examples thereof include generators such as generators, various electric products using the generator as at least one power source, electric motors, and artificial satellites. In addition, you may comprise the thermoelectric conversion apparatus which mounts many thermoelectric conversion elements and can obtain desired electric power.

The present invention will be described in more detail with reference to the following examples. However, the present invention is not limited to these examples. Unless otherwise specified, “part” means “part by weight” and “%” means “mass%”.

Various characteristic values were measured by the following methods.
(Transmission electron microscope (TEM) observation, electron energy loss spectroscopy (EELS), and low-angle annular dark field (LAADF))
This was performed using an atomic resolution analytical electron microscope (JEM-ARM200F) manufactured by JEOL.
(Electron backscatter diffraction (EBSD))
A field emission scanning electron microscope (JSM-7100F) manufactured by JEOL Ltd. was used.

(Field emission scanning electron microscope (SEM) observation)
This was performed using a scanning electron microscope (JSM-6500F) manufactured by JEOL.
(Micro Raman spectrum)
Measurement was performed using a SOLAR TII Nanofider (Tokyo Instruments Inc.) having a 100 × objective lens with an incident laser beam having a wavelength of 532 nm.

(X-ray diffraction (XRD) pattern)
The crystal structure, space group, and lattice constant were calculated using a powder X-ray diffractometer Smartlab 9 kw diffractometer manufactured by Rigaku Corporation having a Cu Kα radiation source (λ = 0.15418 nm).
(Electrical conductivity, Seebeck coefficient and power factor (PF))
It measured and calculated by the static direct current method based on the inclination of a temperature difference curve using the thermoelectric conversion characteristic evaluation apparatus (ZEM-3, ULVAC-RIKO Co., Ltd. product) under a low-pressure helium atmosphere.

(Carrier density, carrier mobility)
It measured using the hall | hole measurement system (The Toyo Technica Co., Ltd. make, ResiTest8400 series) in air | atmosphere at room temperature.
(Thermal diffusivity)
Measurement was performed by a laser flash method using a laser flash analyzer (LFA457, Netzsch).

(specific heat)
It measured using the differential scanning calorimeter (DSC404F3, Netzsch).
(Carbon content in the composite)
Measurement was performed using a carbon / sulfur analyzer (844 series, manufactured by LECO).
(Zeta potential)
Measurement was performed with a red laser (wavelength: 633 nm) by a laser Doppler microelectrophoresis technique using Zetasizer Nano ZS (manufactured by Malvern). The pH value of the sample was adjusted with an MPT-automatic titrator (Malvern) using an aqueous HCl solution (0.01M) and an aqueous NaOH solution (0.01M). Prior to measurement, the sample was sonicated for 0.5 hour to uniformly disperse.

[STO surface modification (1)]
Raw material strontium titanate (STO) powder (manufactured by Sigma-Aldrich, purity 99%, average size 100 nm) was added to toluene (Wako Pure Chemical Industries, Ltd., purity 99.99%), and 4 ml of (3-aminopropyl) ) Trimethoxysilane (APTS, Sigma-Aldrich, purity 97%) was added dropwise to the resulting STO solution. Next, the mixed solution was heated to 423 K and kept in an Ar atmosphere for 6 hours in a container equipped with a reflux device. The product was washed with toluene, vacuum filtered, and subsequently dried at 333 K to obtain STO (modified STO) powder whose surface was modified by grafting a group having an amino group onto the particle surface. The modified STO powder was positively charged in water, a maximum zeta potential of 45 mV was obtained at pH 3.3, and a stable colloid was formed (see FIG. 7).

(Example 1)
[Preparation of RGO / STO Complex 1]
Graphene oxide (GO) can be prepared by modified Hummers method (S. Park, JH An, RD Piner, I. Jung, DX Yang, A. Velamakanni, et al., Aqueous suspension and characterization of chemically modified graphene sheets, Chem. Mat. 20 (21) (2008) 6592-6594). STO powder (2.4 g) was placed in 500 ml of deionized water and the pH was adjusted to 3.3. A STO colloidal solution was formed by sonication for 2 hours. The GO colloid solution was added dropwise while stirring the STO colloid solution. The mixture was filtered and dried at 333K.

The GO colloid solution and the STO colloid solution were uniformly mixed by heteroaggregation. By thoroughly mixing the STO colloid solution and the GO colloid solution, each GO sheet was completely captured by the STO nanoparticles, and a GO / STO mixed powder in which GO was uniformly distributed was obtained (see FIG. 9). It can be seen that recombination and aggregation of GO nanosheets could be effectively suppressed by the heteroaggregation process. FIG. 8 is a TEM image of the original STO powder at (a) low magnification and (b) high magnification.

While applying GO / STO mixed (hybrid) powder at a pressure of 60 MPa in a vacuum atmosphere, a pulsed current pressure sintering apparatus (discharge plasma sintering apparatus) using a graphite die (Sinter 511-S, Sumitomo Coal Mining Co., Ltd.) ), And the temperature was controlled from room temperature to 1473 K at a heating rate of 100 K / min, and the mixture was fired and densified by holding at 1473 K for 5 minutes. During sintering, GO was reduced to RGO, and finally a disc (pellet) of RGO / STO composite 1 having a diameter of 10 mm and a thickness of 3 mm was obtained. An RGO / Al ₂ O ₃ composite (carbon content 0.6% by volume) was prepared in the same manner as RGO / STO composite 1 (carbon content 0.64% by volume). For comparison, pure STO ceramics (carbon content 0% by volume) were also sintered under the same conditions. Oxygen deficient STO was prepared by heat treating pure STO ceramics at 1573 K for 10 hours in H ₂ / Ar flowing gas.
The composite prepared by the method of Example 1 was of high quality, with the graphene in the matrix being uniformly dispersed and free of aggregation.

Three different samples of pure STO, oxygen deficient STO (also denoted as STO-H), and RGO / STO complex 1 were first analyzed by X-ray diffraction (XRD) to confirm the resulting crystal phase. did.
Below, RGO content (volume%), apparent density (g / cm ³ ), relative density (%), crystal structure, space group, lattice of pure STO, STO-H, and RGO / STO composite 1 Indicates a constant. The STO ceramic with oxygen deficiency is prepared by heat-treating pure STO ceramic in H ₂ / Ar for comparison.

As shown in Table 1 above, from the results of XRD measurements, not only all samples have the same cubic structure with a Pm3m space group, but also very similar lattice constants close to the values reported by the standard card. It was also shown that these materials do not have a significant change from the structure of the stoichiometric composition STO. Further, the GO / STO mixed powder becomes a dense RGO / STO composite 1 by being sintered by the pulse current pressure sintering method. In addition, since all the relative densities are 98% or more, it can be said that it is dense.

Also, as shown in FIG. 1, the XRD patterns of all the samples are in good agreement with the cubic perovskite STO (PDF®- # 35-0734) shown at the bottom, and during the sintering or reduction process, It has been shown that no metastasis has occurred. The absence of RGO-derived peaks in the complex is thought to be due to the low carbon content, but the presence of RGO converted from GO was identified by Raman spectra (FIG. 10). The ratio I (2D) / I (G) of the peak intensity I (2D) of the 2D band and the peak intensity I (G) of the G band as measured by Raman spectroscopy was 0.49.

One of the effects of RGO on the microstructure of ceramic composites is the suppression of particle growth, which is important for reducing the thermal conductivity of thermoelectric materials. As shown by EBSD, the average particle size of pure STO without RGO (also referred to simply as average size) is 13.31 μm, and the particles are mainly distributed in the range of average particle size of 10-18 μm ( FIG. 2 (a) and FIG. 2 (d)). In STO-H, the long-term annealing treatment after sintering spreads the particle distribution, and the average particle size is larger than that of pure STO, reaching 14 μm (FIGS. 2B and 2D). ). On the other hand, in the case of the RGO / STO composite 1, the average particle diameter of the oxide in the composite is 0.58 μm, which is 20 times or more smaller than that of pure STO or STO-H. Furthermore, although the average particle diameter of oxides in the composite is mainly concentrated at 0.6 μm, it includes particles smaller than 0.2 μm and larger than 1 μm. It is in a mixed state in the body (FIGS. 2 (c) and 2 (d)). This is effective in improving the ZT value of the complex according to the theory proposed by Zhao et al. (LD Zhao, BP Zhang, WS Liu, JF Li, Effect of mixed grain sizes on thermo-electric performance of Bi ₂ Te ₃ compound, J. Appl. Phys. 105 (2) (2009) 6).

Since the detailed structural information of RGO in the composite was hardly obtained by EBSD or scanning electron microscope (FIG. 11), the fine structure of RGO / STO composite 1 was measured by TEM and HRTEM as shown in FIG. Observed. The sufficiently dense and fine particle structure in the composite was reconfirmed by TEM observation (FIG. 3 (a)). In a higher-magnification bright-field image (Figure 3 (b)), it can be seen that RGO can be clearly distinguished from the surrounding STO matrix and that RGO is uniformly distributed in the composite and there is no significant aggregation. It was. Furthermore, there are two types of RGOs in the composite, those present at the grain boundaries and those contained within the particles (FIG. 3 (d)). The former RGO can effectively suppress the movement of grain boundaries, and is a main factor for suppressing grain growth. However, the lateral size of RGO is much larger than that of STO particles (Figure 8), and it is very difficult for the grain boundaries to completely encompass the entire RGO platelet, so that RGO is completely encapsulated in the particles. The latter type done is very rare and unique. The RGO contained inside the particles is a very small size that looks like debris (small pieces) (FIG. 3 (c)), and occurred between STO and RGO having a perovskite-type crystal structure during the sintering process. Probably produced as a result of the reaction. Another consequence of this reaction is that STO loses oxygen and becomes oxygen deficient STO ceramic, as demonstrated in the next section.

[Electrical characteristics of RGO / STO composite 1]
The stoichiometric composition STO is an insulator due to its large band gap (3.2 eV). However, at a high temperature exceeding 1273 K, neutral oxygen (denoted as O ₀ ^X ) on the regular oxygen site of STO tends to be desorbed thermodynamically, leaving oxygen vacancies with two electrons.
O ₀ ^X ← → V ₀ ”+ 2e + 1 / 2O ₂ (1)

In a reducing environment such as a vacuum, the desorption of oxygen is further promoted, and a slightly conductive STO ceramic (pure STO) is obtained after pulsed current pressure sintering. When measurements are repeated to reach steady performance (FIG. 12), pure STO after sintering has a very low electrical conductivity of 17.9 S / m at room temperature as shown in FIG. 4 (a). Thus, the electrical conductivity increases slightly with increasing temperature and finally reaches 176 S / m at 760K. When this method is adopted, the electrical conductivity is greatly improved, and the electrical conductivity of STO-H is 640 S / m at room temperature and 2025 S / m (maximum value) at 570 K, and oxygen vacancies compared to pure STO. Can be said to have increased considerably. Even more surprisingly, the electrical conductivity of the as-sintered RGO / STO composite 1 without any additional reduction process is even higher than STO-H and develops a conductivity of 1633 S / m at room temperature, At 470K, the maximum value is 4669 S / m.

Theoretically, the electrical conductivity σ is expressed as in the following formula (2). Therefore, the improvement of the electrical conductivity is caused by an increase in carrier mobility (μ), an increase in carrier density (n), or these Due to both increases.
σ = μne (2)
The electrical conductivity of the RGO / STO composite 1 has been improved by using a Hall measurement system to determine which constituent factors have the most influence, using pure STO, STO-H, and RGO. For / STO complex 1, the carrier density and mobility at room temperature (300 K) were measured. The results are shown in Table 2 below. M ₀ represents the mass of free electrons.

In contrast to pure STO (carrier density 3.64 × 10 ¹⁸ cm ⁻³ and mobility 0.29 cm ² V ⁻¹ s ⁻¹ ), in STO-H, the carrier density ( Both 4.54 × 10 ¹⁹ cm ⁻³ ) and mobility (0.49 cm ² V ⁻¹ s ⁻¹ ) increased. Furthermore, the carrier density (1.59 × 10 ²⁰ cm ⁻³ ) of the RGO / STO composite 1 is about 3.5 times that of STO-H, and the mobility (0.64 cm ² V ⁻¹ s ⁻¹ ). Is slightly larger than STO-H. Therefore, the improvement in the electrical conductivity of the RGO / STO composite 1 is caused by the increase in the carrier density over the mobility. These carrier density and mobility increases are clearly associated with RGO because the only difference between pure STO and RGO / STO complex 1 is the addition of RGO.

(Comparative Example 1)
[Preparation of RGO / Al ₂ O ₃ composite]
Using the same method as in Example 1, an RGO / Al ₂ O ₃ composite having the same carbon content (0.6 vol%) was produced. Unlike STO, Al ₂ O ₃ shows very stable insulation even under the condition that RGO is contained as a conductive component in the RGO / Al ₂ O ₃ composite, and RGO / Al ₂ at room temperature. The O ₃ composite had an electrical conductivity of only 84.2 S / m (FIG. 13). Table 3 below shows the carrier density and carrier mobility of the RGO / Al ₂ O ₃ composite at room temperature (300 K).

As shown in Table 3 above, although the RGO / Al ₂ O ₃ composite has the same carbon content (0.6 vol%) as the RGO / STO composite 1, its carrier density is 3.92. × 10 ¹⁸ cm ⁻³ , which was 2 digits lower than the carrier density of the RGO / STO composite 1. Therefore, it is considered that the high carrier density of the RGO / STO composite 1 is not due to RGO itself but oxygen vacancies generated in the composite having a perovskite crystal structure.

To prove our guess, electron energy loss spectroscopy (EELS) was used to examine the electronic structure changes of RGO / STO complex 1 and pure STO. As shown in FIG. 5, the fine structure of the OK edge in pure STO is exactly the same from the inside of the grain to the grain boundary. In contrast, the OK edge of the grain boundary in the composite became high strength at the high energy loss part. This is consistent with observations for oxygen deficient STO. The Ti-L edge presents stronger evidence by a clear distinction between the Ti ³⁺ and Ti ⁴⁺ formal valences. All oxygen vacancies can formally donate two electrons to the closest Ti d-band, and the number of oxygen is three times that of titanium, so the sensitivity of Ti is six times higher than that of O. In the RGO / STO composite 1, the EELS spectrum of the Ti-L edge at the grain boundary where RGO is present is clearly different from the inside of the particle without RGO (the region where there is no carbon signal), and a high concentration of Ti ³⁺ ions is present in the vicinity of RGO. It was confirmed to exist. Furthermore, since the low angle annular dark field (LAADF) signal can detect strain fields from oxygen vacancies by distortion of adjacent cation sites and subsequent dechanneling of the electron beam from the cation train, the LAADF image is RGO. Promoted oxygen vacancies (FIG. 14). Therefore, it can be concluded that the amount of oxygen deficiency in the complex mainly concentrated in the vicinity of RGO is much higher than that of pure STO. Considering the RGO pieces contained in the STO particles, the reason why the carrier density of the RGO / STO composite 1 is very high is that the RGO in the composite becomes a carbon source and is slightly in combination with oxygen atoms on the surface of the STO particles. It is considered that the reaction promoted the formation of oxygen vacancies in the STO complex. In this process, some RGO loses the size necessary to prevent grain boundary migration and is surrounded by STO particles.

<Thermal conductivity and ZT value>
The Seebeck coefficient with respect to the temperature of pure STO, STO-H and RGO / STO complex 1 is shown in FIG. Despite being two orders of magnitude higher in electrical conductivity than pure STO, the Seebeck coefficient of RGO / STO composite 1 is only reduced by 16.6%. In the case of a degenerate semiconductor, two important parameters related to the Seebeck coefficient are a carrier density n and an effective mass m ^* at a predetermined temperature, and are represented by the following formula (3).
S = (8π ² k _B ² / 3eh ² ) m ^* T (π / 3n) ^2/3 (3)
In the above formula (3), k _B represents Boltzmann constant, h represents Planck's constant, e represents carrier charge, and T represents absolute temperature. Thus, if the effective mass does not change, an increase in carrier density usually results in a decrease in Seebeck coefficient. However, the effective mass of STO increased due to the introduction of oxygen vacancies (Table 2 above), and the value of the Seebeck coefficient could be generally maintained. As a result, the power factor of the RGO / STO composite 1 is an order of magnitude higher than that of pure STO and about twice as high as that of STO-H at 520K as shown in FIG.

The thermal diffusivity (D), specific heat (Cp), and density (d) were measured to determine the total thermal conductivity, which is shown in FIG. 6 (a). According to the Wiedemann-Franz law, the electronic thermal conductivity (κe) and the lattice thermal conductivity (κl) are distinguished.
κe = σLT (4)
In the above formula (4), L represents the Lorentz number (2.44 × 10 ⁻⁸ V ² / K ² ). It can be seen that in pure STO, STO-H and RGO / STO composite 1, the contribution of electronic thermal conductivity is almost negligible, while lattice thermal conductivity plays an important role (Table 4 below). Table 4 below shows the total thermal conductivity κtotal, electronic thermal conductivity κe, and lattice thermal conductivity κl of pure STO, STO-H, and RGO / STO composite 1 at 760K.

The lattice thermal conductivity κl of pure STO was 7.7 W / mK at 370 K, decreased with increasing temperature, and reached 5.62 W / mK at 760 K. In the case of STO-H, the lattice thermal conductivity κl is higher than that of pure STO, and the effect of increasing the particle size by long-time heat treatment in hydrogen is apparent. In contrast, the RGO / STO composite 1 has a lattice thermal conductivity κl of only 5.7 W / mK at 370K and 3.59 W / mK at 760K. This low lattice thermal conductivity κl is believed to be due to a particle size 20 times smaller than that of pure STO, which is highly suppressed by the RGO packing phase.

The temperature dependence of the dimensionless figure of merit (ZT) for pure STO, STO-H and RGO / STO composite 1 is shown in FIG. 6 (c). The ZT value of RGO / STO complex 1 is higher than that of pure STO and STO-H, 0.03 at room temperature, 0.09 at 760K, and more than twice that of STO-H. It was. This improvement in ZT is considered to be mainly attributable to the operational effects of the three examples. That is, i) Incorporation of RGO into an oxide having a perovskite crystal structure, such as STO, can greatly improve the electrical conductivity of the composite by generating oxygen vacancies. ii) The increased conductivity can sufficiently maintain the Seebeck coefficient by simultaneously increasing the effective mass of electrons. iii) The thermal conductivity of the RGO / STO composite 1 is reduced due to the highly constrained particle growth brought about by RGO.

In the present invention, a novel perovskite ceramic with oxygen deficiency formed is shown. Instead of heat treatment in a reducing atmosphere after sintering, graphene oxide (GO) dispersed uniformly in STO powder is added to the STO during the discharge plasma sintering (pulse current pressure sintering) process. It is thought that oxygen deficiency is more likely to occur. GO reduced at high temperature (RGO) acted as a carbon source to further reduce STO at a sintering temperature of 1473 K, increasing oxygen vacancies and residual graphene fragments contained in the final STO compact. The obtained RGO / oxygen-deficient STO complex has a greatly improved power factor and reduced thermal conductivity compared to oxygen-deficient STO prepared by reduction with hydrogen for a long time, and is capable of being transferred to the A or B site. Even without doping, the ZT value improved to 0.09 at 760K.
The present invention also provides a method for quickly, greenly and effectively producing oxygen-deficient perovskite ceramics for improving thermoelectric conversion characteristics.

<Preparation Example 1 of graphene oxide (GO) colloidal solution>
Concentrated sulfuric acid (special grade, manufactured by Wako Pure Chemical Industries, Ltd.) 28.75 parts and natural graphite (Z-25, scaly graphite, manufactured by Ito Graphite Industries Co., Ltd.) 1.00 parts were added to a corrosion resistant reactor to obtain a mixed solution. While stirring the mixed solution, potassium permanganate (special grade reagent, manufactured by Wako Pure Chemical Industries, Ltd.) was introduced 20 times into the mixed solution at 15 minute intervals. The single charge of potassium permanganate was 0.15 parts, and the total charge was 3.00 parts. After completion of the addition of potassium permanganate, the mixture was heated to 35 ° C. and aged for 2 hours while maintaining the liquid temperature to obtain a reaction slurry (slurry oxide-containing slurry). Subsequently, 20.00 parts of the reaction slurry was added to another container containing 75.58 parts of ion-exchanged water with stirring, and further 30% hydrogen peroxide solution (special grade reagent, manufactured by Wako Pure Chemical Industries, Ltd.). After adding 1.08 parts, stirring was continued for 30 minutes and then stirring was stopped. After the stirring was stopped, the mixture was allowed to stand overnight and separated from the precipitate layer to the supernatant layer, and then the supernatant was extracted. Thereafter, ion exchange water having the same volume as the supernatant extracted to wash the precipitate was added to the container, stirred for 30 minutes, stopped stirring, allowed to stand for 5 hours or more, and then the supernatant was extracted again. The above-mentioned washing with ion-exchanged water and extraction of the supernatant were repeated until the pH of the supernatant was 2 or higher. An appropriate amount of ion-exchanged water was added to the resulting precipitate layer, and then graphite oxide was dispersed by a homogenizer treatment to obtain a GO colloid solution A (graphene oxide colloid solution A). The GO concentration in the obtained GO colloid solution was 1.0 part.

<Preparation Example 2 of GO Colloid Solution>
GO colloid solution B (graphene oxide colloid solution B) in the same manner as described in Preparation Example 1 except that natural graphite as a raw material was changed to natural graphite (Z-5F, scale-like graphite, manufactured by Ito Graphite Industries). Got. The GO concentration in the obtained GO colloid solution was 1.1 parts.

[STO surface modification (2)]
Strontium titanate (STO) powder (manufactured by Sigma-Aldrich, purity 99.0%, average particle size 100 nm) and 3-aminopropyltrimethoxysilane (Sigma-Aldrich, purity 97.0%) as toluene (Wako) First grade, manufactured by Wako Pure Chemical Industries, Ltd.). Next, the mixture was heated to a temperature of 380 K and maintained for 6 hours under a N ₂ atmosphere while refluxing. The product was washed with isopropyl alcohol (reagent special grade, manufactured by Wako Pure Chemical Industries, Ltd.), pure water, ethanol (reagent special grade, manufactured by Wako Pure Chemical Industries, Ltd.) in this order, and a tabletop centrifuge (manufactured by Kokusan Co., Ltd., H-36α). ) And centrifuged at 5000 rpm for 5 minutes and collected. After collection, vacuum drying was performed at 333 K to obtain a surface-modified STO (modified STO).

(Example 2)
<Preparation of RGO / STO complex 2>
Modified STO (2.4 g) was added to pure water (500 ml), and 35% hydrochloric acid (for precision analysis, manufactured by Kanto Chemical Co., Inc.) was added dropwise and sonicated several times to adjust the pH to 3.3. A colloidal solution was obtained. The GO colloid solution A obtained in Preparation Example 1 was added dropwise while stirring the obtained STO colloid solution. The GO in the GO colloid solution was added dropwise to 100 parts of STO and GO in a total of 1.0 parts. The resulting mixture was collected by centrifugation at 5000 rpm for 5 minutes using a tabletop centrifuge. After the collection, vacuum drying was performed at 333 K to obtain a GO / STO mixed powder.

While applying a pressure of 60 Mpa in a vacuum atmosphere to the GO / STO mixed powder, using a graphite die, a pulse energization pressure sintering apparatus (discharge plasma sintering apparatus) (manufactured by Sinterland, LABBOX ^TM- 125C) The temperature was controlled from room temperature to 1623 K at a temperature increase rate of 100 K / min, and the mixture was fired and densified by holding at 1623 K for 5 minutes. During sintering, GO was reduced to RGO, and finally, a disc (pellet) of RGO / STO composite 2 having a diameter of 10 mm and a thickness of 2 mm was obtained.

Example 3
<Preparation of RGO / STO complex 3>
Except for changing the amount of the colloidal GO colloid solution so that the GO in the GO colloid solution becomes 2.0 parts with respect to the total of 100 parts of STO and GO, the diameter is 10 mm and the thickness is 2 mm by the method described in Example 2. A disc of RGO / STO composite 3 was prepared.

(Example 4)
<Preparation of RGO / STO complex 4>
The diameter of the colloidal solution to be dropped is changed so that the GO in the colloidal GO solution is 3.0 parts with respect to the total of 100 parts of STO and GO. A disc of RGO / STO composite 4 was prepared.

(Example 5)
<Preparation of RGO / STO complex 5>
The GO colloid solution to be dropped is changed to the GO colloid solution B prepared in Preparation Example 2, and the amount of the GO colloid solution to be dropped is set to 2.0 with respect to a total of 100 parts of STO and GO. A disc of RGO / STO composite 5 having a diameter of 10 mm and a thickness of 2 mm was prepared by the method described in Example 2 except that the part was changed to be part.

(Comparative Example 2)
STO powder (manufactured by Sigma-Aldrich, purity 99.0%, average particle size 100 nm) is applied with a pulsed current pressure sintering device (discharge plasma sintering device) using a graphite die while applying a pressure of 60 Mpa in a vacuum atmosphere. The temperature is controlled from room temperature to 1623K with a temperature rising rate of 100K / min., And it is fired and densified by holding at 1623K for 5 minutes. Finally, a pure STO disk having a diameter of 10 mm and a thickness of 2 mm is obtained. It was.

<Measurement of carbon content in RGO / STO composite>
The carbon amount in the RGO / STO composites 2 to 5 produced in Examples 2 to 5 was measured using the carbon / sulfur analyzer described in the examples (844 series, manufactured by LECO). The measurement was performed three times, and the average value was defined as the carbon content. The results are shown in Table 5.

<Hardness measurement>
Using a semi-Vickers hardness tester (HSV-20, manufactured by Shimadzu Corp.), the Vickers hardness of the RGO / STO composites 2-5 prepared in Examples 2-5 and the pure water STO prepared in Comparative Example 2 was measured. The test load and load holding time when measuring the hardness of the RGO / STO composite were 9.8 N and 5 s, and the test load and load holding time when measuring the hardness of pure STO were 9.8 N and 30 s. The results are shown in Table 6.

The RGO / STO composites 2 to 5 were confirmed to have improved hardness as compared with pure STO.

<Friction coefficient and wear resistance measurement>
The friction coefficient and wear resistance of the RGO / STO composites 2 to 5 produced in Examples 2 to 5 and the pure STO produced in Comparative Example 2 were measured using a static and dynamic friction measuring machine (TL201Ts, manufactured by Trinity Labs). evaluated. The results are shown in Table 7 and FIG.
Conditions: sliding distance 10 mm, sliding speed 5 mm / s, load 400 g, number of sliding times 30 reciprocations, contactor SUJ φ6 mm ball

It can be seen that the RGO / STO composites 2 to 5 have a smaller coefficient of friction than pure STO, and the increase in the coefficient of friction is small even after repeated sliding. From this, it was confirmed that the RGO / STO composites 2 to 5 have improved wear resistance.

<Raman spectroscopy>
Presence of the carbon (RGO) component of the RGO / STO composites 2-5 prepared in Examples 2-5 and the pure STO prepared in Comparative Example 2 using micro Raman (NRS-3100, manufactured by JASCO Corporation) confirmed. The results are shown in FIGS. Presence of carbon was determined by whether the G band (peak around 1600 cm ^-1) and D-band (peak near 1350 cm ^-1), G 'band (peak near 2700 cm ^-1) is present.
Conditions: Use of 532 nm laser, objective lens 20 times, CCD capture time 1 second, total 64 times (resolution = 4 cm ⁻¹ )

In RGO / STO composites 2-5, the presence of each peak derived from carbon (RGO) can be confirmed, indicating that RGO is present in the ceramic. No carbon component could be confirmed with pure STO.

Claims (9)

The following general formula (1):
ABO _{3 ± δ} (1)
(In the formula (1), A and B represent an element capable of forming a perovskite crystal structure, at least one is a metal element, and δ represents an oxygen deficiency and is a number of 0 or more and 1 or less. ) And a composite ceramic of a carbon material and an oxide having a perovskite crystal structure. In said Formula (1), A represents an alkaline-earth metal element or a rare earth metal element, The composite ceramics of Claim 1 characterized by the above-mentioned. In said Formula (1), B represents Ti, V, Cr, Mn, Fe, Co, Ni, Zr, Nb, Mo, Hf, Ta, or W, The claim 1 or 2 characterized by the above-mentioned. The composite ceramic described. The composite ceramic according to claim 1, wherein the composite ceramic has a carbon content of 5% by volume or less. The composite ceramic is characterized in that the ratio I (2D) / I (G) of the peak intensity I (2D) of the 2D band and the peak intensity I (G) of the G band measured by Raman spectroscopy is 0.1 or more. The composite ceramic according to any one of claims 1 to 4. The composite ceramic according to any one of claims 1 to 5, wherein an average particle diameter of an oxide in the composite ceramic is 10 µm or less. The composite ceramic according to claim 1, wherein the carbon material is graphite oxide. The composite ceramic according to claim 1, wherein the composite ceramic is used as a thermoelectric conversion material. A method for producing a composite ceramic, comprising:
The production method comprises the following general formula (1):
ABO _{3 ± δ} (1)
(In the formula (1), A and B represent an element capable of forming a perovskite crystal structure, at least one is a metal element, and δ represents an oxygen deficiency and is a number of 0 or more and 1 or less. A method for producing a composite ceramic, comprising a step of sintering an oxide having a perovskite type crystal structure represented by the above and a carbon material by a pulse current pressure sintering method.

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