WO2021215403A1 - Lithium ion-conductive glass ceramic - Google Patents

Abstract

The present invention provides a lithium ion-conductive glass ceramic which is able to be easily produced industrially, while having low grain boundary resistivity, high lithium ion conductivity and high reduction resistance. This lithium ion-conductive glass ceramic is able to be used as a solid electrolyte for an all-solid-state secondary battery and the like. The above are achieved by means of a lithium ion-conductive glass ceramic which contains, in mol% based on oxides, from 10% to 20% of an Li2O component, from 30% to 40% of a P2O5 component, from 40% to 50% of a ZrO2 component, from 0% to 4% of a Y2O3 component, from 0% to 3% of an Al2O3 component and from 0% to 2% of a GeO2 component, while comprising a crystal phase that has a rhombohedral NASICON structure, wherein the lattice constant of the a-axis of this crystal phase as identified by means of X-ray diffractometry and Rietveld analysis is 8.872 Å or more.

Description

Lithium ion conductive glass ceramics

The present invention relates to lithium ion conductive glass ceramics.

In recent years, lithium-ion secondary batteries having a high energy density and capable of charging and discharging have been widely used in applications such as power supplies for electric vehicles and power supplies for mobile phone terminals.
Most of the lithium ion secondary batteries currently on the market use a liquid electrolyte (electrolyte solution) in order to have a high energy density. As the electrolytic solution, a solution obtained by dissolving a lithium salt in an aprotic organic solvent such as a carbonic acid ester or a cyclic ester is used.

However, in a lithium ion secondary battery that uses a liquid electrolyte (electrolyte solution), there is a risk that the electrolyte solution will leak out. Further, an organic solvent generally used for an electrolytic solution is a volatile flammable substance, and there is a problem that it is not preferable in terms of safety.

Therefore, it has been proposed to use a solid electrolyte instead of a liquid electrolyte (electrolyte) such as an organic solvent as the electrolyte of the lithium ion secondary battery. Further, the development of an all-solid-state secondary battery in which a solid electrolyte is used as the electrolyte and all other components are also solid is underway.
Typical characteristics required for the solid electrolyte of the all-solid-state secondary battery include lithium ion conductivity, reduction resistance, stability during firing, and stability over time.

For example, Patent Document 1, as a solid electrolyte is interposed between the negative electrode containing a positive electrode and the active material containing an active _{material, Li 1 + x + z M} x (Ge 1-y Ti y) 2-x Si z P _{Contains a crystal phase of 3-z} O ₁₂ (however, one or more selected from 0 ≦ x ≦ 0.8, 0 ≦ y ≦ 1.0, 0 ≦ z ≦ 0.6, M = Al, and Ga). Solid electrolytes, including glass ceramics, are disclosed. Further, Non-Patent Document 1 discloses _{NASICON type Li 1.15} Y _0.15 Zr _1.85 (PO ₄ ) ₃ having high lithium ion conductivity at room temperature, which is used as a solid electrolyte.

Japanese Unexamined Patent Publication No. 2009-181921

Journal of Power Sources 240 (2013) 50-53

However, since the glass ceramics described in Patent Document 1 contain Ti as a constituent component, the reduction resistance at 25 ° C. is about 2.5 V (vs. Li), and when used as a solid electrolyte, Li ₄ Ti ₅ O There is a problem that a low-potential negative electrode active material such as ₁₂ or TiO _{2 cannot be used.}

According to the investigation by the present inventors, the solid electrolyte described in Non-Patent Document 1 has improved the reduction resistance at 25 ° C. to 1.5 V (vs. Li) or less, and the lithium ion conductivity at 25 ° C. It is also reported that it is relatively high at ^{0.7 × 10 -4} S · cm ^-1. However, in the production of this solid electrolyte, in order to suppress grain growth, it is necessary to use an energization sintering method (SPS) in which energization is performed at a high temperature and firing is performed while applying pressure, which poses a big problem industrially. Further, since this is a solid electrolyte containing almost no amorphous material (mostly in the crystalline phase), it is difficult to obtain glass ceramics. When this solid electrolyte is produced without using the energization sintering method, the grain growth is remarkable and the grain boundary resistance (resistance of ionic conduction generated at the contact interface between particles) of the obtained solid electrolyte becomes high. , Lithium ion conductivity is less than ^{1.0 × 10 -5} S · cm ^-1.

Therefore, the present invention can be used as a solid electrolyte such as an all-solid secondary battery, which can be easily manufactured industrially, has low grain boundary resistance, and has high lithium ion conductivity and high reduction resistance. An object of the present invention is to provide a lithium ion conductive glass ceramic.

In order to solve the above problems, the present inventor has diligently studied, and based on the oxide standard molar%, the Li ₂ O component is 10 to 20%, the P ₂ O ₅ component is 30 to 40%, and the ZrO ₂ component is 40 to 40. It contains 50%, Y ₂ O ₃ component 0 to 4%, Al ₂ O ₃ component 0 to 3%, and GeO ₂ component 0 to 2%, and contains a crystal phase having a rhombohedral-based NASICON type structure. Lithium ion conductive glass ceramics having an a-axis lattice constant of 8.872 Å or more of this crystal phase identified by X-ray diffraction and Rietbelt analysis can be industrially easily produced and can be easily produced. The present invention has been completed by finding that it is a glass ceramic having low grain boundary resistance, high lithium ion conductivity and high reduction resistance.

That is, the present invention is the following (1) to (10).
(1) In terms of oxide-based mol%, Li ₂ O component is 10 to 20%, P ₂ O ₅ component is 30 to 40%, ZrO ₂ component is 40 to 50%, and Y ₂ O ₃ component is 0 to 4. %, Al ₂ O ₃ component 0 to 3%, GeO ₂ component 0 to 2%, including a crystal phase having a rhombic crystal system NASICON type structure, identified by X-ray diffraction and Rietveld analysis. Lithium ion conductive glass ceramics having an a-axis lattice constant of the crystal phase of 8.872 Å or more.
(2) The lithium ion conductive glass ceramic according to (1), wherein the mass sum of the crystal components identified by X-ray diffraction and Rietveld analysis is less than 97% by mass with respect to the total mass.
(3) The lithium ion conductive glass ceramic according to (1) or (2), wherein the crystal phase contains a LiZr ₂ (PO ₄ ) ₃ α phase and / or a LiZr ₂ (PO ₄ ) _{3 α'phase.
(4) The lithium ion conductivity according to (3), wherein the mass of the LiZr 2} (PO ₄ ) ₃ α phase identified by X-ray diffraction and Rietveld analysis is 80% by mass or more with respect to the total mass. Glass ceramics.
(5) The lithium ion conductive glass ceramic according to any one of (1) to (4), which contains 0.1 to 5 _{% of a SiO 2} component in mol% based on an oxide.
(6) The lithium ion conductive glass ceramic according to (5), wherein the valence of Si contained in the entire lithium ion conductive glass ceramic is 3.5 valence or more and less than 3.9 valence.
(7) The lithium ion conductive glass ceramic according to (5) or (6), wherein the coordination number of O to Si contained in the entire lithium ion conductive glass ceramic is 5 or more and 7 or less.
(8) The lithium ion conductive glass ceramic according to any one of (1) to (7), which is a powder having a lithium ion conductivity of 1.0 × 10 ^-5 S · cm ^{-1 or more at 25 ° C.} ..
(9) The lithium according to any one of (1) to (7), which is a substrate having a lithium ion conductivity of 7.0 × 10 ^-5 S · cm ^{-1 or more and a thickness of 300 μm or less at 25 ° C.} Ion conductive glass ceramics.
(10) The lithium ion conductive glass ceramic according to (9), wherein the maximum particle size of the particles in the outermost particle layer of the substrate is 30 μm or less, and the average particle size of the particles is 15 μm or less.

According to the present invention, it is possible to obtain lithium ion conductive glass ceramics which can be easily manufactured industrially, have low grain boundary resistance, and have high lithium ion conductivity and high reduction resistance. Due to its high reduction resistance, this lithium ion conductive glass ceramic becomes a solid electrolyte having a wide potential window on the low potential side, whereby a low potential electrode active material and a high potential electrode active material are used in combination. As a result, a high-voltage battery can be obtained, so that it can be suitably used as a solid electrolyte for an all-solid secondary battery, a seawater battery, or the like.

It is a graph which shows the impedance measurement result (Cole-Cole Plot) of the solid electrolyte pellet of Example 2. It is a secondary electron image of the outermost particle layer in the solid electrolyte pellet of Comparative Example 1 (drawing substitute photograph). It is a secondary electron image of the outermost particle layer in the solid electrolyte pellet of Example 1 (photograph substitute for drawing). It is a secondary electron image of the outermost particle layer in the solid electrolyte pellet of Example 2 (drawing substitute photograph). _{LiZr 2} (PO ₄ ) ₃ α-phase a-axis lattice constant and lithium ion conductivity (S) in the solid electrolyte pellets of Example 1 (initial product, 1 year later) and Example 2 (initial product, 1 year later)・ It is a graph which shows the relationship with ^{cm -1).} It is a cyclic voltammogram of the solid electrolyte pellet (dotted line) of Example 1 and the solid electrolyte pellet (solid line) of Example 2. It is the XAFS analysis result of the solid electrolyte powder and SiO _{2 powder of Example 2.} It is K-end EXAFS data (standardized by absorption edge strength) of Si in the solid electrolyte powder and SiO _{2 powder of Example 2.} It is the K-end EXAFS vibration function data (k ² weighted) of Si in the solid electrolyte powder and SiO _{2 powder of Example 2.} ³ is a graph showing the effect of the ratio of Y on the lithium ion conductivity (S · cm ^-1 , solid line) and density (g / cm 3, dotted line) in the solid electrolyte pellets of Examples 2 to 6. ³ is a graph showing the effect of ^{the ratio of Si on the lithium ion conductivity (S · cm -1} , solid line) and density (g / cm 3, dotted line) in the solid electrolyte pellets of Examples 1 to 2 and 7 to 9. .. _{LiZr 2} (PO ₄ ) ₃ changed by fluctuations in the ratio of Y or the ratio of Si in the solid electrolyte pellets of Examples 3 to 6 (white plot) and Examples 1 to 2 and 7 to 9 (black-painted plot). It is a graph which shows the influence which the a-axis lattice constant (Å) of α phase has on the lithium ion conductivity (Scm ^-1). It is a graph which shows the impedance measurement result (Cole-Cole Plot) of the solid electrolyte substrate of Example 11. It is a secondary electron image of the outermost particle layer in the solid electrolyte substrate of Example 11 (photograph substitute for drawing). It is a state in which the particle size of the particles is measured from the secondary electron image of the outermost particle layer in the solid electrolyte substrate of Example 11 (drawing substitute photograph).

The present invention will be described.
_{In the present invention, the Li 2} O component is 10 to 20%, the P ₂ O ₅ component is 30 to 40%, the ZrO ₂ component is 40 to 50%, and the Y ₂ O ₃ component is 0 to 0 to the molar% based on the oxide. It contains 4%, 0 to 3% of _{Al 2} O _{3 component, 0 to 2} % of GeO 2 component, contains a crystal phase having a rhombic crystal system NASICON type structure, and is identified by X-ray diffraction and Rietveld analysis. It is a lithium ion conductive glass ceramic having an a-axis lattice constant of 8.872 Å or more in the crystal phase of octopus.
In the following, this may be referred to as "glass-ceramics of the present invention".

Here, in the present invention, the "glass ceramics" is obtained by precipitating a crystal phase by heat-treating the raw material glass, and the crystal phase and the amorphous phase (amorphous phase) formed by the heat treatment are referred to. include. That is, it is a mixture of ceramics and glass.
Unless otherwise specified, the content of each component contained in the glass ceramics of the present invention is expressed in mol% based on oxides. The composition represented by "mol% based on oxides" is based on the assumption that the oxides, composite salts, metal fluorides, etc. used as raw materials for the glass ceramics of the present invention are all decomposed into oxides at the time of melting. In this case, the total number of moles of the produced oxide is 100 mol%, and each component contained in the glass ceramic of the present invention is described.

<Components>
First, each component constituting the glass ceramic of the present invention will be described in detail.

The Li ₂ O component is an essential component that imparts lithium ion conductivity to the glass ceramics of the present invention. Therefore, the _{lower limit of the content of the Li 2} O component is 10%, preferably 12%, and more preferably 14%. _{On the other hand, the content of the Li 2} O component is 20%, preferably 18%, because the chemical durability of the glass ceramics of the present invention can be enhanced to enhance the morphological stability when formed into a solid electrolyte layer or the like. , More preferably 16% is the upper limit.

The P ₂ O ₅ component is an essential component necessary for forming a crystal phase having a rhombohedral NASICON type structure in the glass ceramics of the present invention. Therefore, the _{lower limit of the content of the P 2} O ₅ component is 30%, preferably 33%, and more preferably 35%. On the other hand, since it is possible to suppress the formation of other crystal phases and make it difficult to reduce the lithium ion conductivity of the formed crystal phase, _{the content of the P 2} O ₅ component is 40%, preferably 39%, more. The upper limit is preferably 38%.

The ZrO ₂ component is an essential component capable of making it difficult for decomposition by reduction to occur in the crystal phase having lithium ion conductivity in the glass ceramics of the present invention. Therefore, the _{lower limit of the content of the ZrO 2} component is 40%, preferably 42%, and more preferably 44%. _{On the other hand, the content of the ZrO 2} component is set to 50%, preferably 48%, more preferably 47%, because it is possible to easily form a crystal phase having a rhombohedral-based NASICON type structure.

Incidentally, more preferably the molar ratio of P ₂ O ₅ component (P ₂ O ₅ component / Li ₂ O component) is 2.0-3.0 with respect to Li ₂ O component in the glass ceramic of the present invention, also, the molar ratio of ZrO ₂ component to Li ₂ O component (ZrO ₂ component / Li ₂ O component) is more preferably a 2.5-3.5.

The Y ₂ O ₃ component is an optional component capable of adjusting the lithium ion conductivity of the crystal phase in the glass ceramic of the present invention and adjusting the mechanical strength and size of the crystal phase. Therefore, _{the content of the Y 2} O ₃ component is preferably 0.1%, more preferably 0.5%, still more preferably 1%, still more preferably 1.2% as the lower limit. On the other hand, since it is possible to suppress the formation of other crystal phases and make it difficult to reduce the lithium ion conductivity of the formed crystal phase, _{the content of the Y 2} O ₃ component is limited to 4%, more preferably 3. It is preferable that the upper limit is 5.5%, more preferably 3.2%, still more preferably 3%, still more preferably 2.8%, still more preferably 2.5%.
The molar ratio of Y ₂ O ₃ component to Li ₂ O component in the glass ceramic of the present invention (Y ₂ O ₃ component / Li ₂ O component), more preferable to be 0.05 to 0.26 0 It is more preferably 0.05 to 0.20.

_{Like the Y 2} O ₃ component, the Al ₂ O ₃ component can also adjust the lithium ion conductivity of the crystal phase in the glass ceramics of the present invention, and can also adjust the mechanical strength and size of this crystal phase. It is an ingredient. Therefore, _{the content of the Al 2} O ₃ component is preferably 0.1%, more preferably 0.5%, still more preferably 1%, still more preferably 1.2% as the lower limit. Further, as in the case of the Y ₂ O ₃ component, the formation of other crystal phases can be suppressed and the lithium ion conductivity of the formed crystal phase can hardly be lowered. Therefore, the content of the _{Al 2} O _{3 component is high.} The upper limit is 3%, more preferably 2.8%, and even more preferably 2.5%.

The GeO ₂ component is an optional component that promotes the crystallization of the glass ceramics of the present invention. Therefore, _{the content of the GeO 2} component is preferably 0.1%, more preferably 0.5%, and even more preferably 1% as the lower limit. On the other hand _{, the GeO 2} component is composed of the GeO 2 component because it is difficult to reduce the lithium ion conductivity of the crystal phase formed in the same manner as _{the Y 2} O ₃ component and the Al ₂ O ₃ component, and it is easy to coexist with the Zr component. The upper limit of the content is 2%, preferably 1.5%.

The SiO ₂ component is an optional component capable of increasing the mechanical strength of the glass ceramic of the present invention and improving the lithium ion conductivity of the glass ceramic of the present invention by partially substituting the _{P 2} O _{5 component.} be. Therefore, _{the content of the SiO 2} component is preferably 0.1%, more preferably 0.5%, and even more preferably 0.7% as the lower limit. _{On the other hand, the content of the SiO 2} component is preferably 5% because it is possible to easily form a desired crystal phase and the crystals are likely to be adjacent to each other and the decrease in lithium ion conductivity can be suppressed. The upper limit is more preferably 4%, further preferably 3%, still more preferably 2.5%.
The molar ratio of SiO ₂ component to Li ₂ O component in the glass ceramic of the present invention (SiO ₂ component / Li ₂ O component), more preferable to be 0.007 to 0.17, 0.03 to 0 It is more preferably .17.

The CaO component and the MgO component are optional components that can increase the lithium ion conductivity by including a large amount of Li in the crystal phase due to the balance of valences. Therefore, both the content of the CaO component and the content of the MgO component are preferably 0.5%, more preferably 1%, and further preferably 2% as the lower limit. On the other hand, since the decrease in lithium ion conductivity of the glass ceramics of the present invention can be suppressed, the content of the CaO component and the content of the MgO component are both preferably 5%, more preferably 4%, still more preferable. Is preferably up to 3%.

Similar to the Y ₂ O ₃ component and the Al ₂ O ₃ component, the Sc ₂ O ₃ component and the Ga ₂ O ₃ component can adjust the lithium ion conductivity of the crystal phase in the glass ceramics of the present invention, and can adjust the lithium ion conductivity of the crystal phase. It is an optional component whose size and the like can be adjusted. Therefore, the _{lower limit of the content of the Sc 2} O ₃ component and the content of the Ga ₂ O ₃ component is preferably 0.1%, more preferably 0.5%, and further preferably 1%. Is. Further, _{similarly to the Y 2} O ₃ component and the Al ₂ O ₃ component, the formation of other crystal phases can be suppressed and the lithium ion conductivity of the formed crystal phase can be hardly lowered. Therefore, Sc ₂ O _The upper limit of the content of the three components and the content of the Ga ₂ O ₃ component is preferably 2%, more preferably 1.5%.

The SnO ₂ component, _{like the GeO 2} component, is an optional component that promotes the crystallization of the glass ceramics of the present invention. Therefore, _{the content of the SnO 2} component is preferably 0.1%, more preferably 0.5%, and even more preferably 1% as the lower limit. On the other hand, since the lithium ion conductivity of the formed crystal phase can be hardly lowered, _{the content of the SnO 2} component is preferably 2%, more preferably 1.5% as the upper limit.

Furthermore, the glass ceramics of the present invention may contain an inorganic component containing boron (B) or fluorine (F).

On the other hand, in the glass ceramics of the present invention, the content of titanium (Ti) is preferably reduced as much as possible (for example, less than 1%, further less than 0.1%, etc.), and more preferably no Ti is contained. This is because the reduction of the reduction resistance can be suppressed by reducing the Ti component. Further, not only Ti but also transition metal components such as niobium (Nb), vanadium (V), and nickel (Ni) are preferably reduced as much as possible, and more preferably not contained.

Further, in the glass ceramics of the present invention, the content of the sulfur (S) component is preferably reduced as much as possible, and more preferably not contained. This is because the possibility of generating harmful gases such as hydrogen sulfide can be reduced in an all-solid-state secondary battery or the like by reducing the S component. Further, in order to avoid a decrease in lithium ion conductivity, it is preferable to reduce alkali metal (Na, K, etc.) components other than Li as much as possible, and it is more preferable not to contain them.

<Crystal phase, etc.>
Next, the composition of the crystal phase contained in the glass ceramic of the present invention and the composition of the phase other than this crystal phase will be described in detail.

The glass ceramic of the present invention contains a predetermined amount of each of the above-mentioned components, and further contains a crystal phase having a rhombohedral-based NASICON-type structure composed of at least a part of each of these components. The crystal phases contained in the glass ceramics of the present invention preferably have a rhombohedral NASICON type structure, but other lithium ion conductive crystal phases (for example, LISION type, perovskite type, garnet type, etc.) ) May be included in part. Even in this case, among all the crystal phases contained in the glass ceramics of the present invention, the crystal phase having a rhombohedral NASICON type structure is preferably 80% by mass or more, and 90% by mass or more. More preferably, it is 95% by mass or more.

The glass ceramics of the present invention containing the crystal phase having the above-mentioned rhombohedral NASICON type structure have the a-axis lattice constant of this crystal phase identified by X-ray diffraction (XRD) and Rietveld analysis. It is 872 Å or more. When the a-axis lattice constant of this crystal phase is extended more than a predetermined value, the glass ceramic has high lithium ion conductivity. The a-axis lattice constant is more preferably 8.874 Å or more, further preferably 8.876 Å or more, and further preferably 8.878 Å or more. The upper limit thereof is preferably 8.892 Å or less, more preferably 8.889 Å or less, and further preferably 8.888 Å or less.

Further, the glass ceramics of the present invention containing the crystal phase having the above-mentioned rhombic crystal system NASICON type structure is the sum of the masses of the crystal components (crystal phase components) identified by X-ray diffraction (XRD) and Rietveld analysis. However, it is preferably less than 97% by weight based on the total mass of the glass ceramics of the present invention. In other words, the components of the unspecified phase (phases other than the crystal phase) calculated by the difference between the total mass of the glass ceramics of the present invention and the mass sum of the crystal components identified by X-ray diffraction (XRD) and Rietbelt analysis. Is preferably 3% by mass or more with respect to the total mass of the glass ceramics of the present invention. This is because the increase in grain boundary resistance due to overgrowth of particles can be suppressed, and the elongation of the a-axis lattice constant of the crystal phase can be stabilized. The unspecified phase mainly includes an amorphous phase. The sum of the masses of the crystal components is more preferably less than 95% by mass with respect to the total mass of the glass ceramics of the present invention, that is, the mass of the components of the unspecified phase is more preferably 5% by mass or more. .. Further, the lower limit of the mass sum of the crystal components is not limited, but is preferably 80% by weight or more, more preferably 85% by weight or more, based on the total mass of the glass ceramics of the present invention. Suitable.

In the above-mentioned X-ray diffraction (XRD), ZnO having the same mol% amount as the initial sample amount is mixed as a standard sample by a powder X-ray diffractometer (D8 DISCOVER manufactured by Bruker), and CuKα ray is used. And measure.

Further, the Rietveld analysis described above is performed by the Rietveld analysis software "Z-Rieveld code" using the XRD data measured under the above conditions.

The crystal phase having a rhombic crystal-based NASICON type structure contained in the glass ceramics of the present invention has higher lithium ion conductivity at room temperature. Therefore, Li _{1 + x + y} Y _x Zr _{2- It preferably contains a crystalline phase of x S y} P _3-y O ₁₂ (0 ≤ x <2, 0 ≤ y <3), LiZr ₂ (PO ₄ ) ₃ α phase (high temperature phase) and / or LiZr ₂ ( PO ₄ ) It _{is more preferable to include 3} α'phase (low temperature phase). The crystal phase having a NASICON type structure of this rhombic crystal system is composed of LiZr ₂ (PO ₄ ) ₃ α phase (high temperature phase) and / or LiZr ₂ (PO ₄ ) ₃ α'phase (low temperature phase). However, in addition to the α phase and the α'phase _{, other phases of LiZr 2} (PO ₄ ) ₃ (β phase, γ phase, etc.) may be included. Incidentally, LiZr ₂ (PО ₄₎ is a-axis lattice constant is within the range described above ₃ alpha phase, 80 mass% or more with respect to the total mass of the lithium ion conductive glass-ceramics of the present invention, further more than 85 wt% If it is contained, it is very suitable because the stability over time becomes more excellent. Here, the alpha phase and [alpha] 'phase, CSD number 97658 Cambridge Structural Database _{CSD-System (α-LiZr 2} (PO 4) 3), CSD number _{89456 (α'-LiZr 2 (PO} 4) 3 ).

Further, _{when the glass ceramic of the present invention contains a SiO 2} component, the valence of Si contained in the entire glass ceramic of the present invention (crystal phase and unspecified phase) is 3.5 valence or more and less than 3.9 valence. It is more preferable to have it, and it is more preferable that it has a valence of 3.6 or more and a valence of 3.8 or less. This is because the stability of lithium ion conductivity and the mechanical strength of the crystal structure are better.

Further, _{when the glass ceramic of the present invention contains a SiO 2} component, the coordination number of O to the Si contained in the entire glass ceramic of the present invention (crystal phase and unspecified phase) is 5 or more and 7 or less. Is more preferable, and it is further preferable that the coordination number is 6. This is also because the stability of lithium ion conductivity and the mechanical strength of the crystal structure are better.

The valence of Si described above is measured by X-ray absorption fine structure analysis (XAFS: X-ray Absorption Fine Structure).
Further, the coordination number of O to Si described above is measured by wide area X-ray absorption fine structure analysis (EXAFS: Extended X-ray Absorption Fine Structure).

<Form>
Next, the form of the glass-ceramics of the present invention will be described in detail.

The form of the glass-ceramics of the present invention is not particularly limited, but is a powder or a substrate when used for a solid electrolyte layer or an electrode layer of an all-solid secondary battery, an electrode layer or a partition wall (separator) of a seawater cell, or the like. Is preferable. In particular, when the glass ceramics of the present invention and the electrode active material are mixed to form an electrode layer, the form of the glass ceramics of the present invention is preferably powder.

The glass ceramics of the present invention have high lithium ion conductivity and high reduction resistance. When the glass ceramics of the present invention are powders, for example, the lithium ion conductivity at 25 ° C. is preferably 1. 0 × 10 ^-5 S · cm ^-1 or more, more preferably 3.0 × 10 ^-5 S · cm ^-1 or more, still more preferably 6.0 × 10 ^-5 S · cm ^-1 or more, still more preferably 7 .0 × 10 ^-5 S · cm ^-1 or more, and the reduction resistance at 25 ° C. is preferably 1.5 V (vs. Li: potential for lithium) or less, more preferably 1.0 V (vs. Li) or less. Is preferable. When the lithium ion conductivity at 25 ° C. is ^{less than 1.0 × 10 -10} S · cm ^-1 , the lithium ion conduction does not substantially occur. Further, although not limited, the lithium ion conductivity of this powder at 25 ° C. may be 1.0 × 10 ⁻² S · cm ^-1 or less, and the reduction resistance at 25 ° C. is 0.3 V. It may be (vs. Li) or more. Further, the particles contained in this powder preferably have a maximum particle size of 30 μm or less and an average particle size of 15 μm or less.

When the glass ceramic of the present invention is a substrate, it is preferable that the substrate has a lithium ion conductivity of 7.0 × 10 ^-5 S · cm ^-1 or more and a thickness of 300 μm or less at 25 ° C. The thickness of this substrate is more preferably 200 μm or less, and further, but not limited to, the thickness of this substrate can be easily reduced when a solid electrolyte layer is used to reduce a short circuit between the positive electrode and the negative electrode. Is preferably 0.5 μm or more. Further, since the lithium ion conductivity is further enhanced, it is very preferable that the maximum particle size of the particles in the outermost particle layer of this substrate is 30 μm or less and the average particle size of the particles is 15 μm or less. be. Further, the reduction resistance of this substrate is preferably 1.5 V (vs. Li) or less at 25 ° C., more preferably 1.0 V (vs. Li) or less, as in the case of powder.

Here, in the present invention, the "outermost particle layer" of the substrate means a particle layer formed by particles exposed on the surface of the substrate in the glass ceramic of the present invention which is the substrate. The same applies to forms other than the substrate.
Further, in the present invention, the "maximum particle size" and the "average particle size" of a particle are completely included in the field of view of 24 μm × 19 μm by observation with a scanning electron microscope (SEM). It is the maximum value when all the particle diameters (the length of the longest diagonal line) are measured, and the average value thereof.

The glass-ceramics of the present invention having the above-mentioned structure have a composition in which the particles are densified without coarsening during firing, and the crystal structure does not easily change even when cooled in a furnace from a high temperature. A solid electrolyte that can be easily manufactured industrially, has high lithium ion conductivity and high reduction resistance, and has a wide potential window on the low potential side due to this high reduction resistance. Therefore, the low-potential electrode active material and the high-potential electrode active material can be used in combination, and as a result, a high-voltage battery can be obtained. Further, even if charging / discharging is performed at a high voltage, decomposition due to reduction is unlikely to occur, and lithium ion conductivity is stable for a long period of time. Therefore, it can be suitably used as a solid electrolyte for all-solid-state secondary batteries and seawater batteries.

For example, when the glass ceramics of the present invention are used as the solid electrolyte of the all-solid secondary battery, it can be used for at least one selected from the solid electrolyte layer and the electrode layer (positive electrode layer, negative electrode layer). In particular, since the glass ceramics of the present invention have high reduction resistance, they are preferably used for the solid electrolyte layer or the negative electrode layer. These will be described in detail below.

<Solid electrolyte layer>
A solid electrolyte layer for an all-solid secondary battery can be formed by mixing the glass ceramics of the present invention with an inorganic binder or the like, if necessary, and then sintering the glass ceramics. In particular, the solid electrolyte layer preferably contains 80% by mass or more of the glass ceramic of the present invention, more preferably 90% by mass or more, from the viewpoint of further enhancing the lithium ion conductivity, from the glass ceramic of the present invention. Is more preferable. In the case of a plate-shaped solid electrolyte layer, it is preferable that the thickness thereof is 0.5 μm or more and 300 μm or less as in the case of the above-mentioned substrate.

In the case of using an inorganic binder, it is preferable to use a lithium ion conductive inorganic binder Examples of the inorganic binder, LiPO ₃ of amorphous or polycrystalline, 70LiPO ₃ -30Li ₃ PO ₄ , Li ₂ O-SiO ₂ , Li ₂ O-SiO ₂ -P ₂ O _5- B ₂ O _5- BaO and the like. In particular, Li ₂ O-P ₂ O ₅ type glass, Li ₂ O-P ₂ O ₅ -M ₂ O ₃ type glass (including those in which P is replaced with Si, M is Al or B), And one or more selected from those obtained by quenching _{LiPO 3 after melting to make it amorphous are preferable.} A particularly preferable embodiment of LiPO ₃ which has been melted and then rapidly cooled to make it amorphous has a low glass transition temperature (Tg) of about 280 ° C. and is unlikely to undergo crystallization. Therefore, this inorganic binder and the glass ceramics of the present invention are used. By mixing and heating to 600 ° C., a solid electrolyte layer having high lithium ion conductivity and high reduction resistance can be formed. The content of the inorganic binder contained in the solid electrolyte layer is preferably 20% by mass or less, more preferably 10% by mass or less, and 5% by mass or less, based on the total mass of the solid electrolyte layer. Is more preferable.

<Electrode layer>
For all-solid-state secondary batteries, the glass ceramics of the present invention are mixed with an electrode active material (positive electrode active material or negative electrode active material) and, if necessary, a conductive auxiliary agent, an inorganic binder, etc., and then sintered. The electrode layer can be formed. In particular, this electrode layer preferably contains the glass ceramics of the present invention in an amount of 20% by mass or more, more preferably 40% by mass or more, and further preferably 50% by mass or more. This is because the movement path of lithium ions can be easily secured in the electrode layer, so that the charge / discharge characteristics and the battery capacity of the battery can be easily improved. On the other hand, in this electrode layer, the glass ceramic of the present invention is preferably 80% by mass or less, more preferably 70% by mass or less. This is because the filling amount of the electrode active material is secured and the battery capacity can be increased.

As the positive electrode active material, for example, NASICON type LiV ₂ (PO ₄ ) ₃ and olivine type Li _x J _y MtPO ₄ (however, J is at least one selected from Al, Mg, and W, and Mt is One or more selected from Ni, Co, Fe, and Mn, x satisfies 0.9 ≦ x ≦ 1.5, y satisfies 0 ≦ y ≦ 0.2), layered oxide, or spinel-type oxide. And so on. In particular, LiMtO ₂ and / or LiMt ₂ O ₄ (where Mt is one or more selected from Fe, Ni, Co and Mn) is preferable. This is because the glass ceramics of the present invention and the positive electrode active material can be easily bonded to each other, and lithium ions can be easily transferred between them, so that the charge / discharge characteristics of the all-solid-state secondary battery can be further improved. .. Specific examples of the positive electrode active material include LiCoPO ₄ , LiCoO ₂ , LiMn ₂ O _{4, and the} like. Further, by adding Mg as a trace component, thermal decomposition of the positive electrode active material can be suppressed and the discharge capacity can be improved.

The negative electrode active material is selected from, for example, NASICON-type, olivine-type, spinel-type crystal-containing oxides, rutile-type oxides, anatase-type oxides, amorphous metal oxides, metal alloys, and the like. At least one or more. Further, low potential negative electrode active materials such as _{Li 4} Ti ₅ O ₁₂ and TiO _{2 can also be used.} In particular, Li _{1 + x + z} Al _x Ti _2-x Si _z P _3-z O ₁₂ (where x satisfies 0 ≦ x ≦ 0.8 and z satisfies 0 ≦ z ≦ 0.6), Li ₄ More preferably, it is at least one selected from Ti ₅ O ₁₂ and _{Ti O 2.} This is because the glass ceramics of the present invention and the negative electrode active material can be easily bonded to each other, and lithium ions can be easily transferred between them, so that the charge / discharge characteristics of the all-solid-state secondary battery can be further improved. .. Specific examples of the negative electrode active material include, for example, Li ₂ V ₂ (PO ₄ ) ₃ , Li ₃ Fe ₂ (PO ₄ ) ₃ , LiFePO ₄ , Li ₄ Ti ₅ O ₁₂ , SiO _x (0.25 ≤ x ≤ 2). ), Cu ₆ Sn _{5 and the} like.

The content of these electrode active materials is preferably 10% by mass or more and 50% by mass or less in the electrode layer. In particular, by setting this content to 10% by mass or more, the battery capacity of the all-solid-state secondary battery can be further increased. Then, more preferably, the lower limit is 20% by mass. On the other hand, by setting this content to 50% by mass or less, it is possible to easily secure the ionic conductivity of the electrode layer. Then, it is more preferable that the upper limit is 40% by mass, more preferably 30% by mass.

The conductive auxiliary agent used for the electrode layer includes carbon compounds such as graphite, activated carbon, and carbon nanotubes, and metals composed of at least one selected from Ni, Fe, Mn, Co, Mo, Cr, Ag, and Cu. Examples include alloys, metals such as titanium, stainless steel, and aluminum, and precious metals such as platinum, gold, ruthenium, and rhodium. By using such a material having high electron conductivity as a conductive auxiliary agent, the amount of current that can be conducted through the narrow electron conduction path formed in the electrode layer is increased, so that the charge / discharge characteristics of the all-solid-state secondary battery are improved. be able to.

The content of this conductive auxiliary agent is 1% by mass or more with respect to the entire electrode material (that is, positive electrode active material or negative electrode active material) contained in the electrode layer in consideration of the balance between the battery capacity and the electron conductivity of the electrode layer. It is preferably 20% by mass or less, more preferably 2% by mass or more and 15% by mass or less, and most preferably 4% by mass or more and 10% by mass or less.

Further, as the inorganic binder used for the electrode layer, the same one as the above-mentioned solid electrolyte layer can be used. Further, a current collector may be provided in at least one of the electrode layers, that is, the positive electrode layer and the negative electrode layer. This is because electricity can be easily taken out through the current collector, so that the all-solid-state secondary battery can be easily charged and the all-solid-state secondary battery can be easily discharged. The current collector may be one in which a thin metal layer is laminated or bonded to a positive electrode layer and / or a negative electrode layer, and is fired after laminating a metal layer or a precursor of a conductor on a raw material composition. It may be the one that has been used. If the electron conductivity of the electrode layer itself is high, this current collector may not be provided.

An all-solid secondary battery can be formed by using the solid electrolyte layer and the electrode layer as described above, but one or two of the solid electrolyte layer, the positive electrode layer and the negative electrode layer, which are conventionally known, can be used. Alternatively, such a conventionally known one may be combined with a solid electrolyte layer containing the glass ceramics of the present invention or an electrode layer. However, by including the glass ceramics of the present invention in all of the solid electrolyte layer, the positive electrode layer, and the negative electrode layer, an all-solid secondary battery having higher charge / discharge characteristics can be formed.

The contents of the glass ceramics, the electrode active material, and the conductive auxiliary agent of the present invention and their compositions were mounted on a field emission transmission electron microscope (FE-TEM) by carving out a solid electrolyte layer or an electrode layer. It can be identified by using an energy loss analyzer or an X-ray analyzer, or an X-ray analyzer mounted on a field emission scanning microscope (FE-SEM). Here, when the X-ray analyzer is used, the Li ₂ O component cannot be analyzed directly, but _{the content of the Li 2} O component can be estimated by calculating the electric charge from other constituent components.

<Manufacturing method of glass ceramics of the present invention>
Next, the method for producing the glass ceramics of the present invention will be described in detail.
The glass ceramics of the present invention are characterized in that general methods for producing inorganic materials such as firing, melting, and mixed firing of inorganic materials can be used, and industrial production is easy. Then, although not limited to, a mixing step of mixing the raw material compositions and firing of the mixed raw material composition or molding of the raw material composition into a desired shape and then firing are desired. It is preferable to produce the glass ceramics of the present invention by a solid phase method including a firing step of producing the crystal phase of the above.

For example, as an example of the method for producing glass ceramics of the present invention, there is a raw material glass preparation step of melting and vitrifying a lithium ion conductive raw material containing an inorganic substance containing at least Li and / or P to obtain a raw glass. An amorphous precursor preparation step of crushing the raw material glass and then mixing an amorphous inorganic substance containing Zr to obtain a powdery amorphous precursor, and firing this amorphous precursor at 1100-1300 ° C. to form crystals. Precipitated, in terms of oxide-based mol%, Li ₂ O component is 10 to 20%, P ₂ O ₅ component is 30 to 40%, ZrO ₂ component is 40 to 50%, and Y ₂ O ₃ component is 0 to 4. %, Al ₂ O ₃ component 0 to 3%, and GeO ₂ component 0 to 2%, including a crystal phase having a rhombohedral-based NASICON type structure, identified by X-ray diffraction and Rietbelt analysis. Examples thereof include a manufacturing method including a firing step for obtaining lithium ion conductive glass ceramics having an a-axis lattice constant of 8.872 Å or more in the crystal phase of the octopus. The above-mentioned firing step may be performed in two or more steps, but it is preferably performed in one step from the viewpoint of manufacturing efficiency and the like.

Then, in the above-mentioned example of the manufacturing method, before the firing step, a molding step of molding the powdery amorphous precursor into a substrate having a target size and thickness is included, and the amorphous obtained by this molding step is included. By firing the molded body of the precursor by the firing step described above, a substrate-like lithium ion conductive glass ceramic (the glass ceramic of the present invention which is a substrate) can be obtained. Even in the case of producing the glass ceramics of the present invention in a form other than the substrate, the powdery amorphous precursor may be molded into the desired shape by the molding step in the same manner.

In the case of producing the glass ceramics of the present invention in a form other than powder, a method of firing a powdery amorphous precursor by the firing step described above and then molding the fired powder into a desired shape ( The glass ceramics of the present invention having low grain boundary resistance, high lithium ion conductivity and high reduction resistance cannot be obtained by the method in which the firing step is not performed after molding and the method in which molding is refired after firing). The same applies when the sheet is molded by mixing with a polymer or the like.

The embodiments described above are merely examples for facilitating the understanding of the present invention, and do not limit the present invention. That is, the components and the like described above can be changed and improved without departing from the spirit of the present invention, and it goes without saying that the present invention includes equivalents thereof.

Hereinafter, examples of the present invention will be described, but the present invention is not limited to the following examples, and various modifications can be made within the technical idea of the present invention.

<Preparation of solid electrolyte (Comparative Example 1) by solid phase method using a precursor having no amorphous composition>
Using zirconia oxide (ZrO ₂ ), lithium carbonate (Li ₂ CO ₃ ), ammonium dihydrogen phosphate (NH ₄ H ₂ PO ₄ ), and yttrium oxide (Y ₂ O ₃ ) as raw material powders, Li _1.15 Y _0.15 Zr _{After being prepared so as to have a quantitative ratio of 1.85} P ₃ O ₁₂ , the mixture was placed in a 600 ml PP cup and mixed with a YTZ ball having an outer diameter of 10 mm for 3 hours in a pot mill. Then, the obtained mixed powder was placed in a platinum pot and calcined at 900 ° C. for 3 hours to obtain a precursor. The rate of temperature rise at this time was set to 300 ° C./h.
Further, the obtained precursor was pulverized using an alumina mortar and pestle to obtain a 106 μm mesh pass. Then, it was crushed with a planetary port mill until the cumulative 90% particle size (D90) became 1 μm or less, and dried. Then, the obtained dry powder was placed in a platinum pot, fired in a cantal furnace at 1200 ° C. for 20 hours, and cooled in a furnace to obtain a solid electrolyte powder of Comparative Example 1.

Further, after crushing the dry powder before firing in an agate mortar, 1.5 g of a sample that passed 0.5 mm mesh was tablet-molded with a force of 20 kN using a mold having an outer diameter of 20 mm, and the obtained molded product was platinum. The solid electrolyte pellets of Comparative Example 1 were also obtained by placing them on a plate and calcining them in a cantal furnace at 1200 ° C. for 20 hours and cooling them.

The cumulative 90% particle size (D90) described above means a cumulative 90% by volume particle size from the smaller particle size in the particle size distribution, and is based on JIS R 1629 "Laser diffraction / scattering method for fine ceramic raw materials". It was measured and calculated by a particle size distribution measuring device (manufactured by Spectris, Mastersizer 3000) by the laser diffraction / scattering method of "Particle size distribution measuring method".

<Preparation of solid electrolyte (Examples 1 to 10) by solid phase method using a precursor having an amorphous composition>
The present invention is obtained by pulverizing and mixing an amorphous zirconium-containing inorganic raw material and a raw material glass obtained by vitrifying other raw materials, drying the mixture to obtain a precursor having an amorphous composition, and firing the precursor after firing or tablet molding. Solid electrolytes of Examples 1 to 10 which are glass ceramics were obtained. In the drawings described later, of these examples, Example 2 may be referred to as "LYZSP12" or "LYZSP".
The details of this method will be described below step by step.

(Making raw glass)
First, lithium metaphosphate (LiPO ₃ ), trilithium phosphate (Li ₃ PO ₄ ), yttrium oxide (Y ₂ O ₃ ), or silicon dioxide (SiO ₂ ) was prepared in the amounts shown in Table 1 below or Table 2 below. .. Each of the prepared samples was placed in a platinum pot, melted and vitrified with good stirring at 1100 ° C. or higher, and cast on a metal cast plate. The yield of the recovered raw material glass including the raw material glass adhering to the platinum pot was 99% or more by weight in each of Examples 1 to 10.

(Preparation of glass ceramic solid electrolyte)
After crushing each of the above raw material glasses until a 106 μm mesh pass is obtained, in Examples 1 to 9, amorphous zirconium phosphate ((ZrO) ₂ (HPO ₄ ) ₂ ) was added in the amount shown in Table 1 below, respectively, in Example 10 Yttrium oxide (Y ₂ O ₃ ) and amorphous zirconium phosphate ((ZrO) ₂ (HPO ₄ ) ₂ ) were added in the amounts shown in Table 2 below, and 1-propanol was used as a dispersion medium and pulverized by a planetary ball mill. .. The crushed medium at this time was YTZ beads (manufactured by Nikkato Corporation) having an outer diameter of 2 mm. Then, the slurry after crushing is dried, and this dry powder is put into a platinum pot as an amorphous precursor, fired in a cantal furnace for 20 hours at 1200 ° C., and cooled in a furnace to obtain the solid electrolyte powders of Examples 1 to 10. Obtained.
Further, after each of the dry powders before firing was crushed in an agate mortar, 1.5 g of a sample that passed 0.5 mm mesh was tablet-molded with a force of 20 kN using a mold having an outer diameter of 20 mm, and the obtained molded product was obtained. It was placed on a platinum plate and calcined and cooled in a cantal furnace at 1200 ° C. for 20 hours to obtain solid electrolyte pellets of Examples 1 to 10.

<Density and lithium ion conductivity measurement>
The surfaces of the solid electrolyte pellets of Comparative Examples 1, 1 and 2 were polished with # 800 and # 2000 water-resistant abrasive paper and 1-propanol, and then calipers, a micrometer, and an electronic balance were used. , The diameter, thickness, and weight were measured, respectively, and the density was calculated. The results are shown in Table 3 below.

Further, gold electrodes were formed on both sides of each solid electrolyte pellet as blocking electrodes by a magnetron sputtering apparatus (SC-701HMC manufactured by Sanyu Electronics Co., Ltd.). Then, the impedance was measured by an electrochemical evaluation device (manufactured by Biologic, SP300) at 25 ° C. under the conditions of a frequency of 0.1 Hz to 7 MHz, an amplitude voltage of 10 mV, and an open circuit voltage, and the lithium ion conductivity was calculated. This result is also shown in Table 3 below. As a reference, the Core-Cole Plot obtained by the impedance measurement of Example 2 is shown in FIG. The above-mentioned lithium ion conductivity was calculated based on the resistance, thickness, and electrode area of the bending point on the low frequency side (the portion indicated by the arrow in FIG. 1 in Example 2).
It was shown that Example 1 and Example 2 are solid electrolytes having a lithium ion conductivity of one order or more higher than that of Comparative Example 1.

<Powder X-ray diffraction and Rietveld analysis>
Powder X-ray diffraction and Rietveld analysis were performed to confirm the crystal structure of the solid electrolyte. First, the solid electrolyte powders of Comparative Example 1, Example 1, and Example 2 were crushed in an alumina mortar, and then 1 g of the crushed sample and 0.1 g of zinc oxide (ZnO, manufactured by High Purity Chemical Laboratory Co., Ltd.) as a reference. , And mixed for 5 minutes using a Menou mortar and pestle, and powder X-ray diffraction measurement was performed. For powder X-ray diffraction measurement, D8 DISCOVER manufactured by Bruker Co., Ltd. is used, and for weight measurement, an electronic balance capable of weighing up to 0.1 mg is used, and an ionizer is used so that the weight does not change before and after the measurement due to static electricity or the like. Weighed while removing static electricity. ZnO was calibrated using α-Al ₂ O ₃ (STANDARD REFERANCE MATERIAL 674, US DEPARTMENT OF COMMERCE NATIONAL BUREAU OF STANDARDS). Incidentally, the calibration of ZnO, and ZnO 0.181 g, were mixed using an agate mortar and α-Al ₂ O ₃ 0.122g described above, was carried out by the powder X-ray diffraction measurement using the same method as described above.

In the Rietveld analysis, the powder X-ray diffraction measurement result described above was analyzed by the profile function "Split Pseudo Voigt function" by the Rietveld analysis software "Z-Rieveld code". The initial parameter values of crystal structure, alpha-phase in the CSD numbers _{(LiZr 2 (PО 4) 3} α) is 97 658, ZnO is twenty-six thousand one hundred seventy, [alpha] 'phase _{(LiZr 2 (PО 4) 3} α') is 89456, YPO ₄ was 201113 and ZrO ₂ was 80046. In addition, the space group was fixed as it was for all phases, and the lattice constant was variable for all the values allowed in the space group for all phases.

The results of this Rietveld analysis are shown in Table 4 below. Further, the structural parameters of the α phase of Comparative Example 1 obtained by this Rietbelt analysis are shown in Table 5 below, the structural parameters of the α phase of Example 1 are shown in Table 6 below, and the structural parameters of the α phase of Example 2 are shown in Table 6 below. Is shown in Table 7 below. The phases other than the crystalline phase (mainly unspecified phase and amorphous phase) were calculated by the following formula.

(Calculation formula for phases other than crystalline phase)
Mass ratio of phases other than crystal phase (standardized by ZnO) = (mass ratio of evaluation sample with mixed ZnO as 1-mass ratio of each crystal phase other than ZnO with ZnO as 1 obtained by Rietbelt analysis ÷ Mass ratio of the evaluation sample with the mixed ZnO as 1 (calculation formula of the mass ratio excluding ZnO of phases other than the crystal phase)
Mass ratio of phases other than crystal phase (standardized by ZnO) / Total mass ratio of each crystal phase other than ZnO x 100%
The correction coefficient of ZnO calibrated by Rietveld analysis was 0.879 times.

As can be seen from this result, the solid electrolyte powders of Comparative Example 1, Example 1, and Example 2 all contain a crystal phase of _{LiZr 2} (PO ₄ ) _{3 α phase, which is a rhombic crystal-based NASICON type structure.} However, while Comparative Example 1 has almost no phase other than the crystalline phase, in Examples 1 and 2, the phase other than the crystalline phase, which is mainly an amorphous phase, exceeds 3% by mass (that is, the mass of the crystal component). The sum was less than 97% by weight based on the total mass). Further, it was shown that the a-axis lattice constants of these α phases were less than 8.872 Å in Comparative Example 1 and more than 8.882 Å in Examples 1 and 2.

<Observation of secondary electron image>
Since the state of the particles greatly affects the lithium ion conductivity, the secondary electron image of the outermost particle layer was also observed for the solid electrolyte pellets of Comparative Example 1, Example 1, and Example 2. For the secondary electron image observation, S-3000N manufactured by Hitachi High-Technologies Corporation was used. Regarding this result, the secondary electron image of Comparative Example 1 is shown in FIG. 2, the secondary electron image of Example 1 is shown in FIG. 3, and the secondary electron image of Example 2 is shown in FIG. As for the observation conditions, for Comparative Example 1, the working distance (WD; working distance (distance from the center of the objective lens to the sample surface)) was 16.5 mm, the acceleration voltage was 15 kV, and the magnification was 500 times. In Example 1, the WD was 15.9 mm, the acceleration voltage was 15 kV, and the magnification was 500 times. In Example 2, the WD was 11.5 mm, the acceleration voltage was 15 kV, and the magnification was 500 times.
Even under the same firing conditions, the solid electrolyte of Comparative Example 1 obtained from the precursor having no amorphous composition showed remarkable grain growth (FIG. 2). On the other hand, in the solid electrolytes of Examples 1 and 2 obtained from the precursor having an amorphous composition, the formation of a clear neck was confirmed, and the grain boundaries were densified (FIGS. 3 and 4).

<Preservation test>
The solid electrolyte is mainly used for battery applications, and since it is necessary for the lithium ion conductivity to be stable for a long period of time in battery applications, a storage test was conducted. As an evaluation method, the solid electrolyte pellets of Comparative Example 1, Example 1, and Example 2 were stored in an atmosphere of 25 ° C., powder X-ray diffraction and Rietveld analysis were performed by the same method as described above, and density and density and Ion conductivity was measured. The storage period was one year.

Table 8 below shows the results of Rietveld analysis before the start of the storage test (initial product) and after storage for 1 year (1 year later), and Table 9 below shows the density and lithium ion conductivity measurement results. In the solid electrolyte of Example 1 containing no Si component, the α phase decreased from 96.6% by mass to 28.8% by mass after storage for 1 year, and decomposed into _{α'phase, YPO 4} , ZrO _{2, and the like.} On the other hand, the decomposition of the α phase (decrease in mass ratio) in the solid electrolyte of Example 2 containing the Si component was suppressed to less than 10% by mass. The solid electrolyte of Example 1 containing no Si component has a lithium ion conductivity reduced by about 46% after storage for one year, whereas the solid electrolyte of Example 2 containing a Si component has a lithium ion conductivity. The decrease was about 28%.

In addition, the results of this storage test and the literature values of similar products (Li _1.15 Y _0.15 Zr _1.85 P ₃ O ₁₂ (JCPDS 01-083-4639), Li _1.2 Ca _0.1 Zr _1.9 P ₃ O ₁₂ (RSC Advances 2011 (1)) ) 1728-1731), LiZr ₂ P ₃ O ₁₂ rhombic crystal system (JCPDS 01-084-0998)) with lattice constants (a-axis (a / Å), c-axis (c / Å)) Table 10 shows a summary of lithium ion conductivity, and FIG. 5 shows a plot of the initial values of Examples 1 and 2 and the lithium ion conductivity and a-axis lattice constant one year later. The a-axis lattice constants of the α phase of Examples 1 and 2 prepared from the precursor having an amorphous composition are the a-axis lattices of the literature values of Comparative Example 1 and similar products prepared from the precursor having no amorphous composition. It was larger than the constant, and in particular, Example 2 containing the Si component was the largest. From the relationship between the lithium ion conductivity and the a-axis lattice constant, it can be confirmed that the lithium ion conductivity tends to increase as the a-axis lattice constant increases. It is presumed that this is because lithium ions can move more freely in the lattice of the crystal phase.
That is, it was confirmed that the solid electrolyte of Example 2 containing the Si component was particularly excellent in stability at room temperature (25 ° C.), and by replacing a part of P with Si, the a-axis direction of the lattice was further extended. It was confirmed that the lithium ion conductivity was further increased, and the stability of the crystal structure over time and the stability of the lithium ion conductivity were further enhanced.

<Confirmation of redox potential by cyclic voltammetry>
The reduction resistance of the solid electrolytes of Examples 1 and 2 was evaluated by cyclic voltammetry (CV) with lithium as the counter electrode.

A Li metal foil (electrode) and a polymer electrolyte are attached to a copper foil which is a current collector, and a solid electrolyte pellet of Example 1 or Example 2 in which an Au electrode has been sputtered is placed on the working electrode side. The current was collected with aluminum foil. This evaluation sample was vacuum-packed with an aluminum laminate pack, and only the copper foil and aluminum foil of the current collector could be taken out with a tab film. As a pretreatment of the hermetically sealed evaluation sample, it was held at 60 ° C. for 2 hours or more to promote the bonding between the polymer electrolyte and the solid electrolyte, and then the evaluation was performed.
The measurement temperature was 25 ° C., the starting voltage was an open circuit voltage, the scanning speed was 0.2 mV / s, and scanning was performed from 5 V to 0.2 V. The obtained cyclic voltammetry (current-potential curve: CV curve) is shown in FIG. It was confirmed that the solid electrolyte of Example 2 containing the Si component had a small change in current and was more excellent in reduction resistance.

<Confirmation of effective valence of Si by X-ray absorption fine structure analysis (XAFS)>
From the above results, it was confirmed that the solid electrolyte of Example 2 has high lithium ion conductivity and high reduction resistance. However, since the effect of the Si component and the state of Si in the solid electrolyte of Example 2 were unknown, the valence of Si was confirmed by X-ray absorption fine structure analysis (XAFS) for the solid electrolyte pellet of Example 2. .. The X-ray absorption fine structure analysis was performed at BL6N1 of the Aichi Synchrotron Optical Center. The result is shown in FIG.

The results in FIG. 7 showed good agreement with the XAFS spectrum (Photon Factory Activity Report 2012 # 30 (2013) B) actually measured and calculated for _{the SiO 6 coordination.} From this, it was suggested that Si is likely to be in the SiO ₆ coordination in the solid electrolyte, and is likely to be present in the ZrO ₆ site in the crystal or in the amorphous portion in the solid electrolyte.

The arrow in FIG. 7 indicates a portion where the difference from _{SiO 2 is remarkable.} In particular, from the new peak generated at about 1845 eV indicated by the arrow on the left side, it can be interpreted that the solid electrolyte of Example 2 is contaminated with Si having a valence lower than tetravalent. Since the difference in absorption edge energy between Si equivalent to 0 valence and SiO ₂ (equivalent to 4 valence) is 9 eV, Si having a valence of about 3.7 valence is obtained from the above-mentioned peak energy difference (about 3 eV). It was confirmed that it was included.

<Confirmation of the coordination number of O to Si by wide area X-ray absorption fine structure analysis (EXAFS)>
Furthermore, the coordination number of O to Si in the solid electrolyte of Example 2 was also confirmed by wide area X-ray absorption fine structure analysis (EXAFS). For the wide area X-ray absorption fine structure analysis, the synchrotron radiation data measured by BL6N1 of the Aichi Synchrotron Center used in the above-mentioned XAFS measurement was used. The analysis conditions are as follows.

The K end of Si is E = 1847 eV, while the LIII, LII, and LI ends of Y are E = 2080 eV, 2156 eV, 2373 eV, and Zr LIII, LII, and LI ends are E = 2222 eV, 2307 eV, and 2532 eV, respectively. Yes, it exists in the K-end EXAFS region of Si. The data in FIG. 8 shows the contamination of the LIII and LII ends of Y. Therefore, in the EXAFS analysis of the solid electrolyte of Example 2, it is necessary to limit the E range (wavenumber k range) used in the analysis in order to avoid mixing these. Therefore, in order to improve the accuracy of analysis as much as possible even within this limitation, the following measures were taken.

The peripheral atoms are limited to the closest O so that the experimental data can be reproduced with as few variables as possible and the closest O coordination number can be determined. As a result, the fitting variable scaling factor S ₀ ² O coordination number (Nо), Si-O distance (r _Si-O), the absorption edge (E _0), became Debye-Waller factor (σ ₀ ^2).
Then, the fitting variables were sequentially determined in the following flow with the limit E range: 1897 to 2061 eVkmax = 7.4, the wave number weight order: secondary, and the fitting space: wave number space.
(1) The reference SiO ₂ is analyzed in a wide E range to determine _{σ 0} ^2.
(2) The reference SiO ₂ is analyzed in the limit E range (σ ₀ ² is fixed to the value of (1)), and S ₀ ² is determined by this data that can be fixed to No = 4.
(3) Example 2 (LYZSP12) was analyzed in the restricted E range (σ ₀ ² was fixed to the value of (1) and S ₀ ² was fixed to the value of (2)), and Nо of the target crystal phase was evaluated.

The analysis results are shown in Table 11 and FIG. 9 below. FIG. 9 shows a “radial structure function in real space” extracted by the Fourier transform. Here, the R range and the k (or q) range used in the extraction are shown as a window function. From the fact that the function extracted from this data and the function of the fitting result matched and the R factor in Table 11, it was judged that this analysis was performed appropriately.
Then, the O-coordination number for Si in the solid electrolyte of Example 2 was estimated to be 6 instead of 4. (Table 11). From this result, the inclusion of the Si component in the solid electrolyte strengthens the covalent bond in the crystal phase (particularly in the α phase), and Si is arranged at the Zr site to energetically stabilize the lattice. It was suggested that the amorphous phase is likely to be stabilized by 6-coordinated Si.

<Confirmation of the effect of the ratio of Y and Si on lithium ion conductivity and density>
_{Based on Example 2, the ratio of Y of Li 1.05 + x} Y _x Zr _2-x Si _0.05 P _2.95 O ₁₂ (the value of x in the formula) in the solid electrolyte pellets of Examples 3 to 6 is the lithium ion conductivity. The results of confirming the effects of density and α-phase on the a-axis lattice constant are shown in Table 12 and FIG. 10 below.
_{Further, based on Examples 1 and 2, the ratio of Si of Li 1.15 + y} Y _0.15 Zr _1.85 Si _y P _3-y O ₁₂ in the solid electrolyte pellets of Examples 7 to 9 (value of y in the formula). The effects of) on the lithium ion conductivity, density and α-phase a-axis lattice constant are shown in Tables 13 and 11.
Furthermore, _{the ratio of Y of Li 1.05 + x} Y _x Zr _2-x Si _0.05 P _2.95 O ₁₂ (value of x in the formula) and Si of _{Li 1.15 + y} Y _0.15 Zr _{1.85 Si y} P _3-y O ₁₂ FIG. 12 shows the result of confirming the influence of the a-axis lattice constant of the α phase changed by the fluctuation of the ratio (the value of y in the equation) on the lithium ion conductivity.
_{Furthermore, Li 1 + x + y Y} x Zr 2-x Si y P 3-y O 12 in the ratio of Si to (value of y in the formula) was _{0, Li 1 + x Y x} Zr 2-x Table 14 shows the results of confirming the effects of the Y ratio of P ₃ O ₁₂ (the value of x in the equation) on the lithium ion conductivity, density, and the a-axis lattice constant of the α phase.

The lithium ion conductivity, density, and α-phase a-axis lattice constant of each solid electrolyte pellet were all determined by the same method as described above. In addition, the lithium ion conductivity and density were measured twice for each example (Tables 12 to 14, FIGS. 10 to 12).

As a result, when _{the ratio of Y in Li 1.05 + x} Y _x Zr _2-x Si _0.05 P _2.95 O ₁₂ is 0.10 to 0.17, and when Li _{1.15 + y} Y _0.15 Zr _1.85 Si _y P _{3- It was shown that when the ratio of Si in y} O ₁₂ was 0.00 to 0.07, the lithium ion conductivity was higher.
The a-axis lattice constant of the α phase is the ratio of Y _{in Li 1.05 + x} Y _x Zr _2-x Si _0.05 P _2.95 O ₁₂ or Si in _{Li 1.15 + y} Y _0.15 Zr _{1.85 Si y} P _3-y O ₁₂ . By changing the ratio of 8.8880 Å to 8.8892 Å, it was confirmed that the value of the a-axis lattice constant of the α phase and the value of the lithium ion conductivity are correlated. That is, the lithium ion conductivity increases as the a-axis lattice constant of the α phase increases from 8.880 Å, and the lithium ion conductivity is highest when the a-axis lattice constant of the α phase is around 8.885 Å. It was confirmed that when the a-axis lattice constant of the phase exceeds 8.892 Å, the lithium ion conductivity tends to decrease.
Further, even when _{the ratio of Y in Li 1 + x + y} Y _x Zr _2-x Si _y P _3-y O ₁₂ exceeds 0.25 and 0.2, if the ratio of Si is 0. It was confirmed that the high lithium ion conductivity was relatively maintained as compared with the case where the ratio of Si was 0.05. The α-phase a-axis lattice constant of Example 10 was 8.888 Å, which is considered to be due to the fact that it was 8.892 Å or less, which is the threshold value of the α-phase a-axis lattice constant described above.

<Manufacturing of sintered body by sheet molding method>
For the solid electrolyte of Example 2 in which high lithium ion conductivity and stable performance were confirmed in pellets and powders, a sheet-like sintered body was produced by green sheet molding.
Specifically, 1-propanol, a binder, and a dispersant were added as a solvent to the powder obtained by pulverizing the dry powder of the amorphous precursor having the composition of Example 2 with D90 to a thickness of 1 μm or less, and the thickness was 20 μm. After forming the sheet, 12 sheets were laminated, vacuum packed and then subjected to hot water hydrostatic pressure pressing, and fired at 1200 ° C. The obtained sheet-like sintered body (sheet-like substrate) was designated as Example 11.

<Measurement of density and lithium ion conductivity of sheet-shaped substrate>
The sheet-shaped substrate of Example 11 was measured in diameter, thickness, and weight with a caliper, a micrometer, and an electronic balance, respectively, and the density was calculated. The density was 2.6-2.85 g / cm ³ . Further, a gold electrode was formed on both sides of the sheet-shaped substrate of Example 11 as a blocking electrode by a magnetron sputtering apparatus (SC-701HMC manufactured by Sanyu Electronics Co., Ltd.). Then, this was placed in a constant temperature bath at 25 ° C., and impedance was measured with an electrochemical evaluation device (Biologic, SP300) at a frequency of 7 MHz to 0.1 Hz and an amplitude voltage of 10 mV to measure lithium ion conductivity. .. This Core-Cole Plot is shown in FIG. The lithium ion conductivity at 25 ° C. is 8.5 to 9.3 × 10 ^-5 S · cm ^-1 , which is a solid electrolyte synthesized by SPS (Li _1.15 Y _0.15 Zr _1.85 P ₃ O ₁₂ ). It was a value higher than the lithium ion conductivity (literature value) of.

<Results of secondary electron image observation>
The sheet-shaped substrate of Example 11 was observed at a WD of 15.0 mm and an acceleration voltage of 15.1 kV using the same apparatus as described above. Then, the magnification was set to 5000 times, and the secondary electron image of the outermost particle layer was observed. The results are shown in FIGS. 14 and 15. From this result, it was confirmed that the grain interfaces were densely joined, although the existence of grain boundaries was recognized.
As for the particles existing in the outermost particle layer, 27 particles in which the particles were completely contained in the field of view of 24 μm × 19 μm were counted. The results are shown in Table 15 below. The particle size was the length of the longest diagonal line (maximum value of the diagonal line) (FIG. 15). From this result, it was confirmed that sintering was possible with a small particle size such as a maximum particle size of 10.2 μm, a minimum particle size of 0.7 μm, and an average particle size of 3.7 μm.

This application claims priority based on Japanese application Japanese Patent Application No. 2020-076009 filed on April 22, 2020, and incorporates all of its disclosures herein.

Claims (10)

Oxide-based mol%,
Li ₂ O component 10-20%,
30-40% of P ₂ O _{5 ingredients,}
40-50% of ZrO _{2 component,}
Y ₂ O ₃ component 0-4%,
Al ₂ O ₃ component 0-3%,
0-2% of GeO _{2 component}
Contains,
It contains a crystal phase that is a rhombohedral NASICON type structure.
The a-axis lattice constant of the crystal phase identified by X-ray diffraction and Rietveld analysis is 8.872 Å or more.
Lithium ion conductive glass ceramics. The lithium ion conductive glass ceramic according to claim 1, wherein the mass sum of the crystal components identified by X-ray diffraction and Rietveld analysis is less than 97% by mass with respect to the total mass. The lithium ion conductive glass ceramic according to claim 1 or 2, wherein the crystal phase includes a LiZr ₂ (PO ₄ ) ₃ α phase and / or a LiZr ₂ (PO ₄ ) _{3 α'phase. The lithium ion conductive glass ceramic according to claim 3, wherein the mass of the LiZr 2} (PO ₄ ) ₃ α phase identified by X-ray diffraction and Rietveld analysis is 80% by mass or more with respect to the total mass. Oxide-based mol%,
0.1 to 5% of SiO _{2 component}
The lithium ion conductive glass ceramic according to any one of claims 1 to 4, which is contained. The lithium ion conductive glass ceramic according to claim 5, wherein the valence of Si contained in the entire lithium ion conductive glass ceramic is 3.5 valence or more and less than 3.9 valence. The lithium ion conductive glass ceramic according to claim 5 or 6, wherein the coordination number of O to Si contained in the entire lithium ion conductive glass ceramic is 5 or more and 7 or less. The lithium ion conductive glass ceramic according to any one of claims 1 to 7, which is a powder having a lithium ion conductivity of 1.0 × 10 ^-5 S · cm ^{-1 or more at 25 ° C.} The lithium ion conductive glass ceramic according to any one of claims 1 to 7, which is a substrate having a lithium ion conductivity of 7.0 × 10 ^-5 S · cm ^{-1 or more and a thickness of 300 μm or less at 25 ° C.} .. The lithium ion conductive glass ceramic according to claim 9, wherein the maximum particle size of the particles in the outermost particle layer of the substrate is 30 μm or less, and the average particle size of the particles is 15 μm or less.

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