KR102250408B1 - Lithium ion battery anode active material and lithium ion battery comprising the same - Google Patents

Abstract

Disclosed is a composite (0.1<=x<2) including: porous particles including reduced titania represented by TiO_2-x; and a titania (TiO_2) nanostructure formed at least partially on the surface of the porous particles. The composite provided in one aspect of the present invention can improve the performance and lifespan of a lithium ion battery by being used as a negative electrode active material of the lithium ion battery, and has a wide potential window.

Description

Lithium ion battery anode active material and lithium ion battery comprising the same

The present invention relates to a lithium ion battery negative electrode active material and a lithium ion battery including the same.

The increasing power demand in the green energy market, including plug-in hybrid electric vehicles and portable devices, has a higher energy density, no memory effect, has a long cycle, and has adjustable characteristics compared to other battery systems. LIB) system. Therefore, most of the research is mainly focused on optimizing the major components of LIB such as anode, cathode, binder and electrolyte.

Titanium dioxide (TiO ₂ ) is attracting a lot of attention as a new negative active material material of LIB due to its excellent material stability, cost efficiency, and non-toxicity. TiO ₂ has various traits, and representative traits include anatase, rutile, brookite, and bronze B, and the anatase TiO ₂ is being studied most actively. Li ⁺ intercalation by the comparison with the conventional LIB using graphite negative active material generated from a low electric potential ^{(0.2 V vs. Li / Li +} ), anatase TiO ₂ negative active material of Li ⁺ intercalation has a higher potential (1.75V vs. Li Occurs at /Li ^{+ ).} This high intercalation potential _{means that the working potential of the TiO 2} negative active material is high, which avoids lithium plating, decomposition of the electrolyte, and of the solid electrolyte interface (SEI) layer (generally occurring at ^{0.1 V vs. Li/Li + ).} It is advantageous to prevent formation. However, _{reversible lithium storage in the anatase TiO 2} structure is limited, and the actual capacity (~ 180 mA·h·g ^-1 ) is lower than the maximum theoretical capacity (335 mAh·g ^{-1 ).} The difference between the theoretical capacity and the actual capacity of the anatase TiO ₂ can be understood in consideration of the electrochemical lithium storage process, which can be expressed by the following half-cell reaction equation.

_{^{TiO 2 + xLi + + xe -}} → Li x TiO 2 (x = 0 ~ 0.55)

This process involves a _{phase transition from tetragonal anatase TiO 2} to Li _0.55 TiO ₂ , which is (i) slow Li ⁺ ion diffusion, (ii) low electrical conductivity (about 10 ^-12 S·cm ^{-1 for} _{anatase TiO 2)} ) And (iii) narrow reversible half-cell working potential window (1-3 V). Thus, the lithium content (x) is dynamically limited to 0.55, which corresponds to a charging capacity of ^{~ 180 mA·h·g −1.} However, in the nano-sized (<7 nm) anatase TiO ₂ crystal, an ^{additional phase transition was observed during the Li +} intercalation (lithiation) reaction, and a Li _0.98 TiO ₂ phase was generated, bringing the Li ⁺ ions closer to the theoretical limit in the range of 1-3 V. You can store up to (x = 0.98). In the case of large anatase TiO ₂ , a wider operating potential window (0.01-3 V) is required for ^{such Li +} ion intercalation. In this case, the formation of an _{irreversible LiTiO 2} ^{layer due to excessive Li +} ion intercalation. It can be a direct cause of poor Li ⁺ ion diffusion and subsequent kinetic limitations, resulting in inefficient electrochemical performance of ^{Li +} ions in the _{TiO 2 electrode and limiting power and energy density.}

In order to overcome these drawbacks, ^{numerous studies have been conducted to improve the limitation of Li +} _{ion storage kinetics in the TiO 2} negative active material, but the inherent low energy density caused by the narrow working potential window (1-3 V) is still It remains a major disadvantage of the _{TiO 2 negative active material.} Therefore, it would be highly desirable to develop a new titanium oxide based platform that has a wide working potential window for high energy density and can have fast Li ^{+ ion transport and electrical conductivity.}

It is an object of the present invention to provide a composite comprising reduced titania and titania, and to provide a lithium ion battery with improved performance by using it as a negative electrode active material for a lithium ion battery.

In order to achieve the above object, in one aspect of the present invention

Porous particles containing reduced titania represented by TiO _2-x; And

_{A titania (TiO 2} ) nanostructure formed at least partially on the surface of the porous particle;

A composite containing (0.1≦x<2) is provided.

In addition, in another aspect of the present invention

A lithium ion battery negative active material including the composite is provided.

In addition, in another aspect of the present invention

Preparing a mixture by mixing titania (TiO ₂ ) and a metal (step 1);

Heat-treating the mixture to _{prepare reduced titania (TiO 2-x} ) (Step 2);

Etching the metal by treating the reduced titania with an acid (step 3); And

Forming titania on at least a portion of the surface of the reduced titania by acid treating the reduced titania on which the metal has been etched (step 4);

A method for manufacturing a composite comprising a (0.1≦x<2) is provided.

Furthermore, in another aspect of the present invention

A lithium ion battery including the negative active material is provided.

The composite provided in one aspect of the present invention is used as an anode active material for a lithium ion battery, thereby improving the performance and life of a lithium ion battery, and has an effect of having a wide potential window compared to a _{commercial titanium oxide (TiO 2) anode active material.} .

1 schematically shows a composite manufacturing process according to an embodiment of the present invention,
2A is TiO of Preparation Example 1 of the present invention, and FIG. 2B shows a STEM image of the composite of Example 1,
2C is an enlarged HR-TEM showing the titania nanostructure of the composite of Example 1 of the present invention,
2D shows a selected area diffraction pattern of Example 1 of the present invention,
Figure 2e is a-TiO ₂ , shows the XRD pattern of Preparation Example 1, Example 1,
2F and 2G show the L-edge spectra of Ti (FIG. 2F) and K-edge spectra of O (FIG. 2G) of Preparation Example 1 and Example 1 according to electron energy loss spectroscopy, respectively,
3A is a view showing the XPS depth profile spectra according to the etching time in the Ti 2p region of Example 1 in Preparation Example 1 of the present invention, and FIG.
4A to 4F show cyclic voltammetry and galvanostatic charge discharge profiles for Preparation Examples and Examples of the present invention, and FIGS. 4A and 4B are a-TiO ₂ , FIG. 4C And FIG. 4D shows Preparation Example 1, and FIGS. 4E and 4F are for Example 1,
Figure 4g shows a-TiO ₂ , Preparation Example 1, Example 1 of 40 mA g ^-1 (0.2 C) to 4 A g ^-1 (20 C) shows the speed characteristics at the current density,
Figure 4h shows the long-term cycle characteristics at 0.2 C and 20 C for _{a-TiO 2} of the present invention, Preparation Example 1 and Example 1,
5A and 5B show CV curves for Preparation Examples 1 and 1 of the present invention at a scan speed ^{of 0.1 mVs -1} to 1 0.1 mVs ^-1,
Figure 5c shows a curve for determining the b value from the cathode peak current for Preparation Example 1 and Example 1 of the present invention,
5D shows the cathode peak current according to the scan speed for Preparation Examples 1 and 1 of the present invention,
5E and 5F are graphs showing the effect of pseudocapacitive and intercalation according to the scan speed for Preparation Examples 1 and 1 of the present invention,
6A and 6B show ex-situ XRD spectra in the charging and discharging process for Preparation Examples 1 and 1 of the present invention,
6C to 6H show ex-situ XPS depth profiles in the charging and discharging process for Preparation Examples 1 and 1 of the present invention,
7 is a graph showing electrical conductivity, specific capacity, and specific surface area for Preparation Examples and Examples of the present invention,
8A to 8D show mapping images for O, Mg, and Ti elements with respect to TiO-MgO in a preparation example of the present invention,
8E shows an XRD pattern for TiO-MgO in a Preparation Example of the present invention,
9A to 9F are STEM images and XRD patterns for Preparation Examples of the present invention,
10A to 10H show STEM images and XRD patterns for embodiments of the present invention,
11 shows the relative fractions of _{TiO and r-TiO 2} for Preparation Examples and Examples of the present invention,
_{12 shows a visual image for a-TiO 2} , Preparation Examples and Examples in the present invention,
13A and 13B show molecular orbitals for _{TiO 2 and TiO,
14 shows a profile of a-TiO 2} in the present invention, a Raman spectral analysis performed for Preparation Examples and Examples,
15A and 15B show _{UV-Vis spectra and Tauc plots for a-TiO 2} , Preparation Example 1 and Example 1 of the present invention, respectively,
16 is a graph showing electrical conductivity for _{a-TiO 2} , Preparation Examples and Examples in the present invention,
Figure 17a shows a-TiO _{2 and N 2} adsorption and desorption curves for Preparation Examples in the present invention, Figure 17b shows a pore size distribution (PSD) curve,
Figure 17c shows a-TiO _{2 in the present invention, N 2} adsorption and desorption curves for Preparation Examples and Examples, Figure 17d shows a pore size distribution (PSD) curve,
18A to 18F show cyclic voltammetry and galvanostatic charge discharge profiles for the preparation examples of the present invention,
19A to 19F show cyclic voltammetry and galvanostatic charge discharge profiles for embodiments of the present invention,
20A to 20F show cyclic voltammetry and galvanostatic charge discharge profiles of _{a-TiO 2} , r-TiO _{2, and TiO,}
21A and 21B show XRD patterns and N ₂ adsorption and desorption curves and pore size distribution curves of monoclinic TiO, respectively,
Figure 22 shows a graph of the electrochemical impedance spectroscopy method for _{a-TiO 2} , Preparation Examples and Examples in the present invention,
23A and 23B show cyclic voltammetry and galvanostatic charge discharge profiles of Comparative Example 1, respectively,
24 shows a table summarizing the surface properties for _{a-TiO 2} , Preparation Examples and Examples in the present invention,
_{25 is a table summarizing a-TiO 2} , TiO and TiO ₂ fraction calculation process and results for the preparation examples and examples in the present invention,
26 is a table summarizing the surface properties of monoclinic TiO,
27 is a table showing the fraction by type of Ti ions from the ex-situ XPS result of FIG. 6,
28 is a table showing impedance parameters for _{a-TiO 2} , Preparation Examples and Examples in the present invention,
29 is a table summarized by evaluating the performance as a lithium ion battery for the previously reported titanium oxides,
30 schematically shows a composite according to an embodiment of the present invention.

Unless otherwise defined, all technical and scientific terms used herein have the same meaning as commonly understood by an expert skilled in the art to which the present invention belongs. In general, the nomenclature used in this specification is well known and commonly used in the art.

In one aspect of the present invention

Porous particles containing reduced titania represented by TiO _2-x; And

_{A titania (TiO 2} ) nanostructure formed at least partially on the surface of the porous particle;

Complex containing

(0.1≦x<2) is provided.

Hereinafter, the composite provided in one aspect of the present invention will be described in detail for each configuration.

First, the composite provided in one aspect of the present invention includes porous particles containing reduced titania.

The reduced titania is _{represented by TiO 2-x} , and 0.1≦x<2.

In one embodiment, the reduced titania may be TiO.

The porous particles may have a size of 500 nm to 5 µm, preferably 1 µm to 3 µm. When the size of the porous particles exceeds 5 µm, the growth of the titania nanostructure may be delayed due to the excessive increase of the reduced titania porous particles, and thus the reproducibility of the change in the ratio of the composite may be degraded compared to the existing studies. If it is less than 500 nm, there may be a problem that only pure titania can be produced, not a complex of reduced titania and titania. That is, in more detail, TiO having small particles is rapidly dissolved and recrystallized in a heated acid solution, so that _{only pure r-TiO 2} is likely to be produced.

The pore size of the porous particles may be 1 nm to 1000 nm, preferably 10 nm to 500 nm, more preferably 30 nm to 400 nm.

Next, the composite provided in one aspect of the present invention includes a titania (TiO ₂ ) nanostructure.

The titania nanostructure may be formed at least partially on the surface of the porous particle.

The titania nanostructure may be a needle-shaped nanostructure. In the case of acicular, it is preferable in that its surface area can be increased.

The titania nanostructure may form a core-shell structure with the porous particle. That is, a titania nanostructure surrounding the porous particle core may be formed on the surface.

A plurality of titania nanostructures may be formed on the porous particles, and may have a needle-shaped shape extending radially from the porous particles (see FIG. 30).

The width of the titania nanostructure may be 1 nm to 10 nm, preferably 2 nm to 8 nm, more preferably 3 nm to 5 nm.

In addition, the length of the titania nanostructure may be 10 nm to 200 nm, preferably 20 nm to 150 nm, more preferably 30 nm to 100 nm.

It may be difficult to grow the size of the titania nanostructure beyond the above-described range, and the ^{diffusion length of Li +} ions in the corresponding range may be shortened, which may be advantageous as a negative electrode active material.

The titania nanostructure may be in the form of a multi-layered needle-shaped titania bundle.

The titania may be anatase or rutile crystalline. Preferably, it may be a rutile phase.

In addition, in another aspect of the present invention

A lithium ion battery negative active material including the composite is provided.

The composite has been described in detail above, and it is not described in duplicate.

When the composite is used as a negative active material for a lithium ion battery, the efficiency and life of the battery can be improved in that both can contribute to the storage of lithium due to the pseudocapacitive and intercalation reactions of the reduced titania and titania nanostructures of the porous particles. .

The complex is composed of TiO ₂ together with titania nanostructures (preferably r-TiO ₂₎ and the surface 2 (Ti ²⁺⁾ 4 that has a large surface area consisting of a (Ti ⁴⁺⁾ as titania composite having a different oxidation number It may be formed in the form of a porous particle (preferably TiO) core containing reduced titania represented by _-x. When each complex ^{reacts electrochemically with Li +} ions, intercalation reactions occur in titania nanostructures, and pseudocapacitive reactions occur in porous particles containing reduced titania, thereby allowing Li ⁺ ions to be stored in different ways. Therefore, by synthesizing a job-sharing composite having a different charging tendency, it is possible to have higher storage capacity and durability.

In addition, in another aspect of the present invention

Preparing a mixture by mixing titania (TiO ₂ ) and a metal (step 1);

Heat-treating the mixture to _{prepare reduced titania (TiO 2-x} ) (Step 2);

Etching the metal by treating the reduced titania with an acid (step 3); And

Forming titania on at least a portion of the surface of the reduced titania by acid treating the reduced titania on which the metal has been etched (step 4);

Composite manufacturing method comprising a

(0.1≦x<2) is provided.

Hereinafter, a method of manufacturing a composite provided in another aspect of the present invention will be described in detail for each step.

First, a method for manufacturing a composite provided in another aspect of the present invention includes a step (step 1) of preparing a mixture by mixing _{titania (TiO 2) and a metal.}

In step 1, titania and metal may be mixed in a weight ratio of 1:2 to 3:1. When the weight ratio of titania and metal is less than 1:2, there is a problem that a pure TiO phase may not be obtained, and when the weight ratio of titania and metal exceeds 3:1, unstable titania oxide (Mg _x Ti _y O _z , There is a problem that TiO _x (X <0.7)) or pure titanium metal (Ti) may be generated.

The metal may be one or more selected from the group consisting of Mg, Na, Li, Ca, Al, Zr, Ti, V, K, Rb, and Cs. In one embodiment, the metal may be Mg.

The metal may be a metal powder.

The titania may be an anatase or rutile phase. In one embodiment, it may be an anatase phase.

Through the step 1, a composite in which a metal is located on the titania particles may be formed.

Next, the composite manufacturing method provided in another aspect of the present invention may include a step (step 2) of preparing _{reduced titania (TiO 2-x) by heat-treating the mixture.}

In the reduced titania, 0.1≦x<2.

The heat treatment may be performed in an atmosphere supplied with an inert gas, but is not limited thereto. The inert gas is not limited as long as it is a gas commonly used in heat treatment, and may be, for example, H ₂ /Ar or H ₂ /N ₂ gas, but is not limited thereto. The reason for performing the heat treatment in an atmosphere supplied with the inert gas may be to prevent other reactions other than the reduction reaction of titania from occurring in the heat treatment process, but is not limited thereto.

The heat treatment may be performed at about 500° C. to about 1,000° C., but is not limited thereto. For example, the heat treatment is about 500°C to about 1,000°C, about 600°C to about 1,000°C, about 700°C to about 1,000°C, about 800°C to about 1,000°C, about 900°C to about 1,000°C, about 500°C It may be performed at about 900°C, about 500°C to about 800°C, about 500°C to about 700°C, or about 500°C to about 600°C, but is not limited thereto. When the heat treatment is performed at less than 500°C, the reduction of titania may not be performed smoothly, and when the heat treatment is performed at more than 1,000°C, the reduction proceeds too severely, making it difficult to efficiently and systematically control the surface characteristics of the reduced titania. However, the problem of wasting unnecessary energy may occur, but is not limited thereto.

The heat treatment may be performed for about 3 hours to about 10 hours, but is not limited thereto. For example, the heat treatment is about 3 hours to about 10 hours, about 4 hours to about 10 hours, about 5 hours to about 10 hours, about 6 hours to about 10 hours, about 7 hours to about 10 hours, about 8 hours To about 10 hours, about 9 hours to about 10 hours, about 3 hours to about 9 hours, about 3 hours to about 8 hours, about 3 hours to about 7 hours, about 3 hours to about 6 hours, about 3 hours to about It may be performed for 5 hours, or about 3 hours to about 4 hours, but is not limited thereto. When the heat treatment is performed for less than 3 hours, reduction of the titania may not be performed smoothly, and when the heat treatment is performed for more than 10 hours, the reduction proceeds too severely, making it difficult to efficiently and systematically control the surface properties of the reduced titania, A problem of wasting unnecessary energy may occur, but is not limited thereto.

Next, a method for manufacturing a composite provided in another aspect of the present invention includes a step of etching the metal by treating the reduced titania with an acid (step 3).

The step 3 is a step of removing the metal component formed in the process of reducing _{the titania (TiO 2} ) by etching the mixture that has been heat-treated in step 2 with an acid, through which the metal component is not included. It is possible to prepare unreduced titania (TiO _2-x).

Preferably, the metal component removed in the above step may be present in the form of a metal oxide.

The step of treating with an acid may be performed by stirring for 24 hours in a 2 M HCl solution, but is not limited thereto.

_{The reduced titania (TiO 2-x} ) prepared by the above manufacturing method may have a cubic structure symmetry without an anatase and/or rutile phase.

Next, a method for manufacturing a composite provided in another aspect of the present invention includes a step (step 4) of forming titania on at least a portion of the surface of the reduced titania by acid-treating the reduced titania on which the metal has been etched.

Through the above steps, at least a portion of the reduced titania phase is converted into titania.

The titania may be an anatase phase or a rutile phase, but may be a rutile phase in an embodiment.

The step may be carried out for 10 to 50 hours at a temperature of 50 ℃ to 150 ℃. Preferably, it can be carried out for 15 to 23 hours at a temperature of 60 ℃ to 120 ℃.

When performed for less than 10 hours at a temperature of less than 50°C, there is a problem that the _{reduced titania (eg, TiO) may not be sufficiently converted to titania (eg, acicular r-TiO 2 ),} When the operation is performed at a temperature exceeding 150° C. for 50 hours or more, there is a problem in that the reduced titania is excessively converted to titania, thereby reducing the efficiency of the battery.

The step 4 may be continuously performed without any change except for the temperature change from step 3 above.

Therefore, the composite manufacturing method provided in another aspect of the present invention has the advantage of being able to form a composite of reduced titania (eg, TiO) and titania (eg, acicular r-TiO _{2) through a simple process.} There is this.

Furthermore, in another aspect of the present invention

A lithium ion battery including the negative active material is provided.

Hereinafter, a lithium ion battery provided in another aspect of the present invention will be described in detail for each configuration.

The lithium ion battery provided in another aspect of the present invention may include a positive electrode, a negative electrode, a separator, and an electrolyte.

The positive electrode may include a positive electrode active material, a conductive material, a binder, and a solvent.

As the positive electrode active material, any lithium-containing metal oxide may be used as long as it is commonly used in the art.

For example, LiCoO2, LiMn _x O _2x (x=1, 2), LiNi _1-x Mn _x O ₂ (0<x<1) or LiNi _1-xy Co _x Mn _y O ₂ (0≤x≤0.5 , 0≤y≤0.5), etc. Specifically, it is a compound capable of occluding/release of lithium such _{as LiMn 2} O ₄ , LiCoO ₂ , LiNiO ₂ , LiFeO ₂ , V ₂ O _{5, TiS or MoS.}

Carbon black and graphite fine particles may be used as the conductive material, and vinylidene fluoride/hexafluoropropylene copolymer, polyvinylidene fluoride (PVDF), polyacrylonitrile, polymethyl methacrylate, poly Tetrafluoroethylene and a mixture thereof or a styrene butadiene rubber-based polymer may be used, and as a solvent, N-methylpyrrolidone, acetone, or water may be used. The contents of the positive electrode active material, the conductive material, the binder, and the solvent may be levels commonly used in lithium batteries.

The negative electrode may include a negative active material, and the negative active material may include porous particles including reduced titania represented by _{TiO 2-x; And a titania (TiO 2} ) nanostructure formed at least partially on the surface of the porous particle; and a composite (0.1≦x<2) including.

The composite and the negative active material have been described above, and thus, overlapping descriptions will not be made.

A separator may be included between the anode and the cathode.

Any of the separators can be used as long as they are commonly used in lithium batteries. It is preferable that the electrolyte has a low resistance to ion migration and has excellent electrolyte moisture absorption capacity. For example, as selected from glass fiber, polyester, Teflon, polyethylene, polypropylene, polytetrafluoroethylene (PTFE), or a combination thereof, it may be in the form of a non-woven fabric or a woven fabric. Specifically, a rollable separator such as polyethylene or polypropylene is used for a lithium ion battery, and a separator having excellent organic electrolyte impregnation ability is used for a lithium ion polymer battery, and such a separator may be manufactured according to the following method. After the separation membrane composition is prepared by mixing a polymer resin, a filler, and a solvent, the separation membrane composition is directly coated and dried on an electrode to form a separation membrane, or the separation membrane composition is cast and dried on a support, and then peeled off from the support. The separator may be formed by laminating on the electrode.

The polymer resin is not particularly limited, and all materials used for the bonding material of the electrode plate may be used. For example, vinylidene fluoride/hexafluoropropylene copolymer, polyvinylidene fluoride, polyacrylonitrile, polymethyl methacrylate, or mixtures thereof may be used.

In addition, an electrolyte may be included between the positive electrode and the negative electrode.

The electrolyte is propylene carbonate, ethylene carbonate, fluoroethylene carbonate, diethyl carbonate, methylethyl carbonate, methylpropyl carbonate, butylene carbonate, benzonitrile, acetonitrile, tetrahydrofuran, 2-methyltetrahydrofuran, γ-buty Lolactone, dioxolane, 4-methyldioxolane, N,N-dimethylformamide, dimethylacetamide, dimethyl sulfoxide, dioxane, 1,2-dimethoxyethane, sulfolane, dichloroethane, chlorobenzene, _{LiPF 6} , LiBF ₄ , LiSbF ₆ , LiAsF ₆ , LiClO in a solvent such as nitrobenzene, dimethyl carbonate, methyl isopropyl carbonate, ethylpropyl carbonate, dipropyl carbonate, dibutyl carbonate, diethylene glycol, dimethyl ether, or a mixture thereof ₄ , LiCF ₃ SO ₃ , Li(CF ₃ SO ₂ ) ₂ N, LiC ₄ F ₉ SO ₃ , LiAlO ₂ , LiAlCl ₄ , LiN(C _x F _2x+1 SO ₂ )(C _y F _2y+1 SO ₂ ) (Where x and y are natural waters), LiCl, LiI, or a lithium salt such as a mixture thereof may be dissolved and used.

The lithium ion battery structure including the positive electrode, the negative electrode, and a separator positioned between the positive electrode and the negative electrode is wound or folded to be accommodated in a cylindrical battery case or a prismatic battery case, and then, when the electrolyte is injected, a lithium ion battery may be completed. After the battery structure is stacked in a bi-cell structure, it is impregnated with an organic electrolyte, and the resulting product is accommodated in a pouch and sealed, thereby completing a lithium ion polymer battery.

In another embodiment of the present invention, a method for manufacturing a negative active material for a lithium ion battery including the method for manufacturing the composite described above may be provided.

In yet another embodiment of the present invention, a method for manufacturing a lithium ion battery including the method for manufacturing a negative active material for a lithium ion battery described above may be provided.

Hereinafter, the present invention will be described in more detail through examples and experimental examples. The scope of the present invention is not limited to specific embodiments, and should be construed by the appended claims. In addition, those who have acquired ordinary knowledge in this technical field should understand that many modifications and variations are possible without departing from the scope of the present invention.

<Preparation Example 1> Preparation of porous TiO (T-1)

A homogeneous mixture of anatase TiO ₂ (a-TiO ₂ ) powder (4 g) and Mg powder (1.21 g) (molar ratio 1: 1) was placed in a quartz tube furnace and heated at 650° C. under a flow of _{5% H 2 /Ar.} . After the magnesium heat reduction treatment, the heated sample was stirred with a 2 M HCl solution for 4 hours to selectively etched the MgO phase. As a result, a highly porous reduced TiO (T-1) without MgO was obtained.

<Preparation Example 2> Preparation of porous TiO (T-0.6)

Prepared in the same manner as in Preparation Example 1, but in order to investigate the effect of the degree of heat reduction of magnesium on the degree of phase change, magnesium was added to the anatase powder (4 g) at a molar ratio of 0.6 (0.726 g) and mixed. I did. The final sample obtained therefrom was named T-0.6.

<Preparation Example 3> Preparation of porous TiO (T-0.75)

It was prepared in the same manner as in Preparation Example 1, but with respect to the anatase powder (4 g), magnesium was added and mixed at a molar ratio of 0.75 (0.908 g). The final sample obtained therefrom was named T-0.75.

<Example 1> TiO part of the needle-shaped rutile TiO ₂ Convert in-situ (in-situ) to (TR-y) 1

Preparation Example 1 (reduced TiO without MgO, T-1) was continuously stirred in HCl solution at 80° C. for 20 hours. Thereafter, centrifugation and freeze-drying were sequentially performed _{to produce a final product in which TiO was gradually transformed into a unique r-TiO 2} nanoneedle on the surface of the porous TiO, which was designated as TR-20.

<Example 2> Part of TiO is a needle-shaped rutile TiO ₂ Convert in-situ (in-situ) to (TR-y) 2

Prepared in the same manner as in Example 1, except that Preparation Example 1 was stirred in HCl solution at 80° C. for 12 hours, which was designated as TR-12.

<Example 3> Part of TiO is a needle-shaped rutile TiO ₂ Convert in-situ (in-situ) to (TR-y) 3

It was prepared in the same manner as in Example 1, except that Preparation Example 1 was stirred in HCl solution at 80° C. for 24 hours, and this was named TR-24.

<Example 4> A part of TiO was used as a needle-shaped rutile TiO ₂ Convert in-situ (in-situ) to (TR-y) 4

It was prepared in the same manner as in Example 1, but Preparation Example 1 was stirred in HCl solution at 80°C for 36 hours, which was designated as TR-36.

<Comparative Example 1> TiO and r-TiO ₂ Simple blending

The TiO/r-TiO ₂ ratio was quantitatively mixed at a ratio of 6:4, and matched with the ratio of Example 1.

<Experimental Example 1> Confirmation of TiO formation

The phase structure of the sample was confirmed by powder X-ray diffraction (XRD, Rigaku Smartlab, 40 kV, 15 mA, 4° min ^-1 , Cu-Kα radiation, λ = 0.15406 nm).

As in Preparation Example 1, when a-TiO ₂ and Mg are mixed and then _{annealed at 650° C. under H 2} /Ar (5% H ₂ ), as can be seen from the XRD results of FIG. 8e, a-TiO ₂ peak Disappears, and a new MgO peak and a reduced TiO peak can be confirmed.

As shown in the scanning transmission electron microscope (STEM) and energy dispersive X-ray spectroscopy (EDX) elemental mapping images of FIG. 8, the MgO phase grows into the TiO framework during the magnesium heat reduction process.

During annealing at 650° C., Mg (melting point: 650° C.) starts melting and small a-TiO ₂ nanoparticles (NP) fuse together to grow into large micron-sized particles. Mg, a strong reducing agent, _{takes oxygen in the vicinity of the TiO 2} framework to form MgO according to Formula 1, while the fused a-TiO ₂ is converted into an oxygen-deficient cubic TiO phase in the form of a much larger particle size. a-TiO _{2 may also be reduced by H 2} according to the following Formula 2 to _{generate H 2} O as well as the TiO phase, but the Gibbs free energy (solid, -596.4 kJ mol ^-1 ) of _{MgO is H 2} O ( Much lower than that of gas, -228.6 kJ mol ^-1 ), MgO formation takes precedence during the magnesium thermal reduction process under _{H 2 /Ar flow.}

TiO ₂ + Mg ↔ TiO + MgO (Chemical Formula 1)

TiO ₂ + H ₂ ↔ TiO + H ₂ O (Chemical Formula 2)

<Experimental Example 2> Confirmation of TiO generation according to the content of Mg

The morphology of the samples was characterized by scanning transmission electron microscopy (STEM) and high resolution transmission electron microscopy (HR-TEM, Hitachi HF-3300) operated at 200 kV in TEM mode.

After annealing as in Preparation Example 1, subsequent HCl etching was performed to remove MgO grown in the TiO phase, and as a result, porous MgO-free having macropores in the range of 60 nm to 200 nm as shown in Figs. TiO(T-1) is formed.

In order to confirm the role of Mg in the reduction of a-TiO ₂ , based on Preparation Examples 1 to 3 during magnesium heat reduction, examples in which Mg and a-TiO ₂ have different molar ratios (0.6, 0.75 and 1.0) Was observed.

For different molar ratios as confirmed by the STEM image of FIG. 9, a similar porous structure was observed in Preparation Examples 1 to 3, but the XRD pattern showed that a pure TiO phase was not obtained until Preparation Example 1 having a molar ratio of 1 was reached. Confirmed by (Fig. 9f). Other Mg ₂ TiO ₄ and MgTiO ₃ impurities were also observed at a low molar ratio (Figs. 9b and 9d). Interestingly, T-1 exhibits an XRD pattern consistent with the cubic β-TiO phase centered on 37.23°, 43.25°, 62.83°, 75.35° and 79.34°, completely different _{from a-TiO 2 (FIGS. 2d and 2e)} .

<Experimental Example 3> r-TiO ₂ Creation confirmation

Subsequently, Preparation Example 1 was redispersed in a 2 M HCl solution as in Examples 1 to 4, and stirred at 80° C. for 12, 20, 24 and 36 hours. Depending on the treatment time, HCl treatment results in the formation of newly grown micron-sized long nanoneedle bundles throughout the TiO2. These hairy needle-like bundles on the TiO surface become more dense over time and _{are identified as square r-TiO 2} by STEM and XRD analysis (FIG. 10 ). This means that as the treatment time increases, the cubic β-TiO phase gradually changes _{from the HCl solution to the r-TiO 2 phase.} For Example 4 (TR-36), a sample treated for 36 hours, all TiO was completely transformed into hairy nano needles, resulting in a pure r-TiO ₂ phase (FIGS. 10g, 10h).

11 shows the quantitative ratio of _{TiO and r-TiO 2} determined by the reference intensity ratio (RIR) method of the XRD pattern for the samples of Examples 1 to 4 (TR-y). As the treatment time increases, the TiO phase decreases, while the r-TiO ₂ phase increases and becomes dominant.

The visual color of the sample treated as shown in FIG. 12 is also clearly different. The color darkened as the content of Mg increased during the reduction process. It was almost black for the highest TiO phase content. On the other hand, when the HCl treatment was used, as more r-TiO ₂ was generated, the color of the sample was sequentially changed from black to white (FIG. 12 ). 25 shows the relative fractions of _{TiO and r-TiO 2} in various prepared samples using a reference intensity ratio (RIR) method based on a powder XRD pattern.

It can be seen that the HCl treatment has a dual role as (1) an etching process in which MgO is dissolved in a TiO-MgO composite to generate MgO-free porous TiO, and (2) a reaction medium as a supply means. TiO is transformed into r-TiO ₂ nano needles by in-situ transformation. The r-TiO ₂ nanoneedle in-situ (in-situ) growth process can occur due to oxidation of various titanium species formed by the thermal reduction reaction of magnesium.

In addition, X-ray photoelectron spectroscopy (XPS) was performed using an XPS system (ESCALAB 250) with a monochromatic Al Kα (150W) source. The energy scale was aligned using the Fermi level (4.10 eV vs. absolute vacuum value) of the XPS instrument.

As shown in Figure 3a, Ti 2p deconvoluted high-resolution X-ray photoelectron spectroscopy (XPS) depth profile, Preparation Example 1 (MgO-free T-1) is composed ^{of Ti 2+} , Ti ³⁺ and Ti ⁴⁺ can confirm. More specifically, the surface is ^{composed of Ti 3+} and Ti ⁴⁺ , and the core is mainly composed of Ti ²⁺ . These various oxidation states of titanium are gradually oxidized to _{r-TiO 2} recrystallized in-situ on the surface of TiO in aqueous HCl solution. 2B and 2C show STEM and HR-TEM (high-resolution transmission electron microscope) images, respectively. The size of the r-TiO ₂ nanoneedle is 5-7 nm in width and is uniformly formed over the entire TiO surface.

In the presence of HCl solution, Ti ³⁺ and Ti ⁴⁺ are released into the reaction solution. Ti ³⁺ produces ^{TiOH 2+} by oxidative hydrolysis of Ti ^3+. TiOH ^{2+ is oxidized to Ti 4+} by reaction with dissolved oxygen to form r-TiO _2. Such in-situ r-TiO ₂ conversion may be represented by the following Chemical Formulas 3 and 4.

Ti ³⁺ + H ₂ O → TiOH ²⁺ + H ⁺ (Chemical Formula 3)

TiOH ²⁺ + O ₂ → Ti(IV) ^{-oxide + O 2-} → TiO ₂ (Formula 4)

Ti ²⁺ is also dissolved in HCl solution, contributing to the formation of _{r-TiO 2 nano needles.} As a result, all titanium composite ions are used as growth units, and the _{mechanism of formation of r-TiO 2} nanoneedle can be explained as follows.

In the case of r-TiO _{2 , the TiO 6} octahedron is first formed by a bond between a Ti atom and 6 oxygen atoms. Thereafter, the TiO ₆ octahedron shares a pair of opposite edges with the next octahedron to form a chain-like structure. Since the transformation speed of the different crystal planes depends on the number of corners and edges of the available coordination polyhedron, the transformation speed of the rutile plane follows the order (110) <(100) <(101) <(001). Due to this mechanism, r-TiO ₂ nanoneedles growing along the [001] direction are formed.

In order to investigate the change of the cationic Ti valency state on the surfaces of Preparation Example 1 (T-1) and Example 1 (TR-20), electron energy loss spectroscopy (EELS) scans were performed (FIGS. 2F and 2G). Two features were identified. First, Ti L _{2 and 3} of the sample of Example 1 clearly show a peak followed by a chemical split corresponding to a ^{pure Ti 4+} valency state that is much less sharp in the sample of Preparation Example 1 (FIG. 2F). Peak split can be explained by the molecular orbital (MO) theory for oxides. The molecular orbitals for all titanium oxides with different oxidation states are similar, but the occupancy of each orbital is different depending on the oxidation state and atomic coordination. In the case of _{TiO 2 , MO is occupied up to the t 1g} level, and the 2t _2g level is empty (Fig. 13). In EELS, electrons at the Ti 2p core level _{can be excited to bin 2t 2g} and 3e _g levels, and as a result, it can be observed that the peaks at the edges of _{L 2 and 3 are split.} In the case of Preparation Example 1 (T-1), the 2t _2g level is already filled with electrons, and the first orbit available for the excited electrons is the 3e _g level, so the L ₂ -t _2g peak disappears from the EELS spectrum. In the case of the O edge, the same reason applies to the observed split peak.

_{In order to confirm the r-TiO 2} conversion in the TiO core, high-resolution X-ray photoelectron spectroscopy (XPS) depth profiling was performed, and as a result, ^{the distribution of other Ti x+} species present in the sample can be confirmed (FIG. 3). In Figure 3a, the surface of Preparation Example 1 (T-1) is mainly composed of Ti ⁴⁺ and Ti ³⁺ , Ti ⁴⁺ 2p _3/2 , Ti ⁴⁺ 2p _1/2 , Ti ³⁺ 2p _3/2 And peaks at 459.4, 465, 458.3 and 463.2 eV due to ^{Ti 3+} 2p _1/2. The lower shoulder peaks at 455.7 and 461.3 eV indicate Ti ²⁺ 2p _3/2 and Ti ²⁺ 2p _1/2 , respectively. Ti ³⁺ exists as a major Ti species on the surface of T-1, and the peak area decreases in the order of ^{Ti 3+} > Ti ⁴⁺ > Ti ^2+. However, after sputtering the sample with an Ar laser to expose the inner part of the sample, it was confirmed that ^{the distribution of Ti x+} was changed, and Ti ²⁺ became the dominant species inside. After sputtering for 3000 seconds, ^{the surface area ratio of Ti 2+} : Ti ³⁺ : Ti ⁴⁺ to T-1 was confirmed to be 55:30:15 from the XPS results.

In the case of Example 1 (TR-20), the Ti ⁴⁺ peaks at 459.4 and 465 eV correspond to Ti ⁴⁺ 2p _3/2 and Ti ⁴⁺ 2p _1/2 , respectively (FIG. 3B). That is, Ti ⁴⁺ is dominant due to the recrystallized r-TiO ₂ nano needles on the TR-20 surface. The sputtered TR-20 also showed a distribution similar to that of sputtered T-1 for ^{different Ti x+ species, with a peak area ratio of Ti 2+} : Ti ³⁺ : Ti ⁴⁺ at 54:28:18 after 3000 seconds sputtering ( Fig. 2b). This means that while the surface of TR-20 is Ti ⁴⁺ dominant r-TiO ₂ , the core is more likely to be composed ^{of Ti 2+ dominant TiO.}

Formation of r-TiO ₂ on the TiO surface was confirmed using Raman spectroscopy (FIG. 14). The surface was analyzed with a Raman spectrometer (Thermo Scientific, NICOLET ALMECA XR) and excited using a 532 nm laser beam. For a-TiO ₂ , the main identifiable peaks are the typical peak values _{for the anatase TiO 2} ^{144 cm -1} (E _g ), 197 cm ^-1 (E _g ), 397 cm ^-1 (B _{1 g} ), 518 cm ^-1 ( A _1g + B _1g ) and 640 cm ^-1 . However, cubic TiO(T-1) of the rock salt structure is inactive in Raman spectroscopy. Examples 1 to 4 (TR-y) Due to the formation of nano needles on the samples, 144 cm ^-1 (B _{1 g} ), 419 cm ^-1 (E _g ) and 613 cm ^-1 (A _{1 g) in the Raman spectrum} ), a distinct r-TiO ₂ peak is formed (the weak band corresponding to the B _2g ^{mode at 827 cm -1 is not shown).} A broad peak at 243 cm ^-1 occurs in the second process.

<Experimental Example 4> Confirmation of band gap and electrical conductivity characteristics

Ultraviolet-visible (UV-vis) absorption spectrum was observed using a CARY5000 UV-Vis spectrophotometer (Agilent Technology).

UV-Vis absorption spectra of a-TiO ₂ , Preparation Example 1 (T-1) and Example 1 (TR-20) are visible light despite similar absorption properties in the UV region (200-400 nm). In the region (400-800 nm), they exhibit significantly different absorption properties. In Fig. 15a. _{Unlike a-TiO 2} , Preparation Example 1 and Example 1 also absorb visible and infrared wavelengths. Bandgap values obtained from Tauc plots based on UV-diffuse reflectance spectroscopy (UV-DRS) are a-TiO ₂ , 3.25 eV, 1.65 eV and 2.71 eV for Preparation Example 1 and Example 1, respectively. The narrow bandgap energy of T-1 is due _{to the phase shift of the pure anatase TiO 2 to} the cubic TiO phase and ^{formation of Ti 3+.} The increasing bandgap energy in TR-20 is due to the partial transformation of _{TiO to r-TiO 2.} Reduced bandgap energy helps to improve electrical conductivity, which can improve storage kinetics in high-speed test performance. Four-probe electrical conductivity measurements of titanium oxide samples support the correlation between bandgap energy reduction and electrical conductivity (FIG. 16).

A four-probe configuration cell was used to measure the change in electrical conductivity under pressure, and Keithley Model 6220 and Model 2182A were used as a low direct current (DC) source and a voltmeter, respectively.

T-1 is 5.19 S·cm ^-1 at 20 MPa, indicating the highest electrical conductivity value compared to a-TiO ₂ (8.87 × 10 ^-5 S·cm ^{-1 ).} However, as the HCl treatment time increases, the conductive TiO in the TR-y sample _{is converted to r-TiO 2} having low conductivity, so that the electrical conductivity decreases.

<Experimental Example 5> Confirmation of pore characteristics

_{An N 2} adsorption-desorption isotherm was obtained using a Brunauer-Emmett-Teller (BET) surface analyzer at a test temperature of 77 K. Surface area, pore volume and pore size distribution were determined using BET and Barret-Joyner-Halenda (BJH) analysis. BET surface areas were calculated using the experimental points in the relative pressure P / P _O = 0.05-0.25. The pore size distribution was derived from the adsorption branch using the BJH method. The total pore volume was estimated by the amount of _{N 2} adsorbed at a relative pressure (P/P _{0) of 0.995.}

While the phase was changed from a-TiO ₂ to Tx and TR-y by sequential treatment, the porous surface properties were measured (FIGS. 17 and 24). Among the various titanium oxides, a-TiO ₂ exhibits the highest surface area and pore volume. The surface area and pore volume of T-1 ^{are calculated as 11.14 m 2} /g and 0.022 cm ³ /g, respectively, which are higher than T-0.6 and T-0.75. This phenomenon is caused by thermal reduction of magnesium in which _{a-TiO 2} NPs are fused to form larger agglomerated TiO particles by molten Mg at 650°C. In fused large reduced samples with large macropores, the surface area decreases sharply. Among the prepared samples, TR-20 exhibited the largest surface area (49.47 m ² /g) and total pore volume (0.055 cm ³ /g). As the portion of r-TiO ₂ increased, the surface area and pore volume continued to increase up to TR-20. However, after 20 hours of deformation, the surface area decreases due to the aggregation of _{r-TiO 2 nano needles blocking the inner channel (FIGS. 17C and 24 ).} Due to these unique characteristics, the phase transformation-optimized TR-20 can be a very promising negative active material for lithium ion batteries (LIBs). The pore size distribution also follows a similar trend to the surface area change, which means that the average pore size starts at 13.48 nm at T-1, but peaks at TR-20 (24.52 nm) and then decreases. These microporous and mesoporous distributions are mainly due to surface properties. In summary, TR-20 allows more surface exposure to the electrolyte due to its excess surface properties ^{and shortens the Li +} ion diffusion path. 24 shows the surface properties of various samples determined by the _{N 2 isotherm.}

<Experimental Example 6> Evaluation of electrochemical properties for each sample

Pure commercial titanium oxide and titanium oxide samples prepared from Preparation and Examples were tested in half-cells for a potential range of 0.01-3 V.

In the following examples, for the production of the anode electrode, a titanium oxide active material was used as a solvent using acetylene black (as a conductivity improver) and a binder (carboxymethylcellulose (CMC) and polyacrylic acid (PAA) in a 2: 1 w/w ratio). Distilled water was used as the mixture in a weight ratio of 70:15:15. A thin film having a thickness of 20 μm was prepared by casting the resulting uniform slurry on the etched copper foil using a doctor blade. The prepared thin film was dried at room temperature overnight and then dried at 80° C. overnight in a vacuum oven. Subsequently, the electrode was roll-pressed to improve interparticle contact. The loading density of the active material was ^{2.2-2.4 mg/cm 2.} Coin-shaped half-cells (CR 2032, Hohsen Co.) are anode electrodes (working electrodes), lithium foil (99.9% purity and 150 μm thick, countertop _{) in an argon-filled glove box (< 0.1 ppm of H 2} O and O _{2 ).} And a reference electrode) and a separator (Celgard 2320) were assembled. An electrolyte was prepared using _{1M LiPF 6} in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC; 1:1 v/v, Soulbrain Pte. Ltd.).

The electrochemical performance is 0.01 V-3.0 V vs. ambient temperature. It was tested by constant current charge-discharge measurement with a multi-channel battery test system (CTS, BaSyTec) in the voltage range of Li/Li ^{+ and a current density of 40-4000 mA/g.} Cyclic voltammetry (CV) experiments were performed using an electrochemical workstation (VSP-1, Bio Logic Science Instruments) in a potential range of 0.01V-3.0 V at a scan rate of 0.2 mV/s. Electrochemical impedance spectroscopy (EIS) measurements were performed with an amplitude of an overlapping alternating current (AC) signal of 0.01 V in the frequency range of 100 kHz-100 mHz.

In the case of a-TiO ₂ ^{, the capacity rapidly decreases, recording 64 mAh·g −1} within the initial 10 cycles at an appropriate charge/discharge rate of 0.2 C (FIGS. 4A and 4B ). However, although T-1 initially has a small capacity value _{, the capacity is well maintained compared to a-TiO 2} , the dynamics are fast, and the electrochemical stability is excellent (FIGS. 4C and 4D ).

From the point of view of the lithiation-discharge cycling operation, Preparation Example 1 (T-1) was 145-190 mA·h·g at 0.2 C without noticeable reduction (kept 97% capacity within 100 cycles). ^It has a specific capacity of -1, has a high reversible lithium storage capacity in long-term cycling, and operates within a potential window of 0.01-3 V. According to the results of cyclic voltammetry (CV) and constant current charging and discharging, the storage of electric charges occurs through Faradaic charge transfer on or around the T-1 surface. Both T-0.6 and T-0.75 have lower specific capacity than T-1 due to their low surface area (Fig. 18).

Under the same evaluation conditions, Example 1 (TR-20) showed ^{a surprising reversible initial capacity of 346 mA·h·g −1} at a charge/discharge rate of 0.2 C, which is more than twice the T-1 capacity (Fig. 4f). . During the first cycle at CV, the first reduction peak was observed at 1.75 V due to the intercalation reaction of ^{Li +} ions in the _{r-TiO 2 crystal structure.} The 0.8 V reduction peak that follows means that a solid electrolyte interface (SEI) layer is formed on the surface of the TR-20 electrode. Li ⁺ ion deintercalation reaction is also observed at the oxidation peak of 2.0 V (Fig. 4e). According to the second cycle, TR-20 exhibits a specific capacity ^{of ~ 347 mA·h·g −1} _{, which exceeds the theoretical capacity of r-TiO 2} (335 mA·h·g ⁻¹ ). This high dose is presumed to be due to the additional pseudocapacitive effect delivered by the contribution of the TiO dose.

Considering that the rate characteristics determine the output power, Example 1 (TR-20) was investigated at different charging and discharging rates (Fig. 4G). Even at a high rate of 20 C, the TR-20 can still ^{provide a specific capacity of 224 mA·h·g −1} , which is much higher than the capacities of T-1 and a-TiO _2. That is, the charging or discharging process can be completed within 3.4 minutes (202 seconds) while achieving a relatively higher capacity than _{a-TiO 2.} By switching the current density back to 0.2 C, a stable high capacity of 353 mA·h·g ^-1 can be obtained. As shown in FIG. 4h, Example 1 of the _{optimized TiO to r-TiO 2} ratio significantly improved the cycle performance compared to _{a-TiO 2.} Under the test conditions of 20° C. for 1000 cycles, the capacity _{decreased by 85% in the case of a-TiO 2} , whereas Example 1 showed a retention rate of 97%, which was the highest specific capacity and the best cycling performance among the reported titanium oxides. Is to have. In the case of other TR-y samples such as Examples 2 to 4, the initial capacity increases as the _{modified rutile (r-TiO 2) content increases under the same conditions.} However, the cycle performance is less than that of Example 1 (Fig. 19).

<Experimental Example 7> Evaluation of the contribution of pseudocapacitive reaction and intercalation reaction

In order to understand the electrochemical behavior of the prepared titanium oxides, CV data at various scan rates were used to distinguish the contribution of intercalation and pseudocapacitive reactions in ^{Li + storage.} Analysis was performed in relation to the operation of the peak current by assuming that the current i obeys the relationship between the scan speed and the power law (FIGS. 5A and 5B).

i = av ^b (Equation 1)

In Equation 1, a is a constant and v is a scan rate, and a value of b may vary from 0.5 (half-infinite diffusion process) to 1 (capacitive process).

5C shows a plot of log(|i|) versus log(v), where the slope of the line corresponds to the value of b. The b value was determined to be ~ 0.23 V (vs. Li/Li ⁺ ) in the case of Preparation Example 1 (T-1) and ~ 2.0 V in the case of Example 1 (TR-20), using an anodic peak. . A value of b = 0.5 indicates that the current is controlled by the semi-infinite diffusion behavior and b = 1 indicates the capacitive behavior. In the case of T-1, the b value is 0.924, indicating a capacitor-like dynamics model, which is one of the characteristics of pseudocapacitance. In contrast, values in the range of b = 0.8-0.9 of TR-20 mean that a combination of pseudocapacitive and intercalation reactions occurs.

In order to quantitatively distinguish the contribution of the surface-limited pseudocapacitive reaction and the diffusion-controlled intercalation reaction, the following Equation 2 was used based on the dependence between the peak current and the scan rate.

i = k ₁ v + k ₂ v ^1/2 (Equation 2)

In Equation 2, i is the current at a given potential, v is the scan rate, and k ₁ and k ₂ are constants. Here, k ₁ v and k ₂ v ^1/2 represent pseudocapacitive and intercalation reactions, respectively. Dividing both sides of Equation 2 by the square root of the scan rate, ^{a plot of i/v 1/2} versus v ^1/2 creates a straight line with a y-intercept constant k ₂ and a slope constant k ₁ , which is a pseudocapacitive reaction and intercalation. It provides quantitative information on the contribution of the reaction (Figure 5d). According to this quantitative analysis, it can be confirmed that the pseudocapacitive reaction of T-1 is superior to that of TR-20, and that the intercalation reaction _{of the r-TiO 2 nanoneedle grown on the TiO structure of TR-20 contributes relatively more (} Figures 5e and 5f).

To investigate the tendency of diffusion-controlled intercalation and surface-limited pseudocapacitive reactions, electrochemical measurements were performed with TR-y samples of Examples 1-4. From FIG. 19, it can be seen that the process of the intercalation reaction in which lithium diffusion is controlled increases sequentially as the _{r-TiO 2 content increases, whereas the surface-limited pseudocapacitive reaction occurring on the surface decreases proportionally.} As shown in Fig. 19h, as the rutile (r-TiO ₂ ) content increases, the stability of the performance decreases rapidly.

The same electrochemical measurements were performed for commercial r-TiO ₂ (square) and monoclinic TiO. Commercial rutile (r-TiO ₂ ) shows distinct cathode (1.75 V) and anode (2.25 V) peaks at CV (FIGS. 20C and 20D ). In the first charging process (0.8 V vs. Li/Li ⁺ ), the SEI layer is formed on the _{r-TiO 2} , and the charge inserted along the c-axis to form _{LiTiO 2} ^{by one Li +} ion is 0.01-3.0 V It is caused by the electrochemical mechanism of the rutile (r-TiO _{2) negative active material operating in the potential range of (FIG. 20C).} Like anatase (a-TiO ₂ ), r-TiO ₂ also undergoes an irreversible excess process during the charge-discharge process within the operating potential range of 0.01-3.0 V (FIG. 20D). After the first charge to form LiTiO ₂ , it can be estimated that the specific capacity decreases because reversible ions less than ^{1 Li + are extracted during discharge.} Unlike conventional a-TiO ₂ , r-TiO ₂ is reliable in an extended operating potential range (0.01-3 V) compared to commercial graphite (372 mA·h·g ^{-1) and has a competitive specific capacity (313 mA).} H·g ^-1 ). However, due to the reversible capacity loss in the first cycle, the specific capacity decreases sharply in the next cycle. Commercial monoclinic TiO exhibits a pseudocapacitive reaction due to the low surface area of the micron-scale particle size, and exhibits a much lower specific capacity compared to the cubic TiO (T-1) prepared in Preparation Example (FIGS. 20e and 20f). 21 shows the crystal phase and surface properties of commercial monoclinic TiO.

In the case of the TR-20, one of the explanations for the increased storage capacity may be a'job-sharing' mechanism. Li ⁺ ions are _{stored in interstitial sites on the oxide side of the TiO 2} crystal interface, while electrons are stored in the carbon additive or SEI layer. Similarly, TR-20 is suitable as an ultra-fast charging negative active material due to the synergistic effect of the r-TiO ₂ ^{nano-needle that stores Li +} ions and the conductive porous TiO phase that stores electrons, and is an addition for ^{Li + ions by utilizing charge separation.} Can provide storage capacity. Therefore, the hybrid titanium oxide composite, in which two different phases are interconnected, exhibits a pseudocapacitive effect of improving energy density without affecting high power or cycle life compared to conventional titanium oxide. While the capacity reduction can be minimized, it is possible to maintain high durability and specific capacity characteristics at various C speeds and cycling performances as shown in FIGS. 4G and 4H.

A constant current of 40 mA·g ^-1 (0.2 C) was applied across the electrode. Fig. 6 shows the ex-situ XPS depth profiles of the T-1 (Figs. 6c, 6d and 6e) and TR-20 (Figs. 6f, 6g and 6h) electrodes. Data were obtained after 10 cycles of clean, fully charged (0.01 V) and discharged (3 V). Each electrode was prepared independently and sputtered by Ar laser for 1000 seconds.

The ex-situ XRD patterns of Preparation Example 1 (T-1) and Example 1 (TR-20) electrodes were obtained in the charging and discharging states in the original first cycle to determine the reaction mechanism in FIGS. 6A and 6B. lost. The T-1 electrode does not undergo a phase change after being fully charged at 0.01 V (FIG. 6A). The (111) and (200) diffraction peaks for TiO are slightly shifted to 2θ (smaller d-spacing) values in the charged state, which may be due to the Coulomb interaction between ^{TiO and Li + ions.} Even after full discharge at 3 V, the diffraction can still be indexed into the (111) and (200) peaks of the parent structure, and the lattice constant increases slightly compared to the initial state electrode. Therefore, these structural data are evidence that electrochemically induced phase transition is suppressed in T-1, indicating that the porous cubic TiO exhibits high structural stability. That is, at 1.5 V, an irreversible peak is not observed at the CV of T-1, but can be verified by the fact that a strong irreversible peak is observed at _{a-TiO 2.} It has been reported as a feature of pseudocapacitance to minimize the structural changes resulting from the inhibition of lithium intercalation and phase transition.

In the case of TR-20, characteristic rutile peaks of (110) and (101) appear at the baseline of the pattern after the first charging, but the peak does not appear again after discharging (FIG. 6B). Due to the unique properties of r-TiO ₂ ^{, some Li +} ions remain trapped in the TR-20 structure, as shown by the above-described structural phase transition. In theory, the bulk Li _0.5 TiO ₂ and LiTiO ₂ phases have distinct diffraction peaks, so an extra peak should be detected. However, it is difficult to distinguish between _{nano-sized Li 0.5} TiO ₂ and LiTiO ₂ due to excessive peak expansion of nano-sized materials. These results are consistent with the TR-20's constant current charge-discharge curve, which can explain why the ^{charge capacity is 706 mA·h·g -1} , while the discharge capacity is ~347 mA·h·g ^{-1 (} Fig. 4f).

In order to reveal the electrochemical reversibility of the two materials of Preparation Example 1 (T-1) and Example 1 (TR-20), ex for the initial state electrode and the electrode after the 10th charge (0.01 V) and discharge (3 V) cycles. -situ XPS depth profile was investigated (Fig. 6). ^{Observation of the XPS spectrum reveals} the presence of Ti 4+ due to charging and discharging in both T-1 and TR-20. The first T-1 electrode had ^{a 10% fraction of Ti 4+} , while the first TR-20 electrode had ^{a Ti 4+} fraction of 27% due to the contribution of the _{r-TiO 2 nanoneedle.} Looking at the depth profile of the electrode in which the initial state electrode 10th charging has been performed, the ratio of ^{Ti 4+} increases while the relative ratio of ^{Ti 2+ decreases.} However, as the discharge proceeds to 3 V, the relative fraction changes completely differently compared to the charging process. This phenomenon can be presumed to be due to the fact that ^{Ti 2+} is covered by the reduction of ^{Li +} ions due to the pseudocapacitive reaction of the TiO core surface, and thus has a relatively low sensitivity during surface analysis. Since Ti ³⁺ does not participate in the reaction, the change in the fraction ^{of Ti 3+ is less pronounced.} Figure 27 shows the relative fraction of ^{Ti x+} measured from ex-situ XPS depth profiles for T-1 and TR-20 electrodes during initial, charging and discharging. This reversible phenomenon during charging and discharging provides a viable mechanism for obtaining stable capacity in the TR-20. As a result, _{TR-20 characteristics that exceed the theoretical a-TiO 2} capacity (a-TiO ₂ = 335 mA·h·g ^-1 ) and stability are reversible intercalation and deintercalation with TiO showing rapid kinetics by pseudocapacitive charge reaction. This is due to the synergistic effect of the r-TiO _{2 nano acicular structure.}

<Experimental Example 8> Confirmation of resistance characteristics for various samples

In order to confirm the relationship between the internal electronic resistance and the various phase changes of titanium oxide, electrochemical impedance spectroscopy (EIS) was performed, which is shown in FIG. 22. The Nyquist plot was recorded at the open circuit voltage of a new cell showing a single semicircle in the high frequency region and a slanted straight line in the low frequency region, corresponding to the charge-transfer resistance (R _ct ) and ^{diffusion of Li + ions.} T-1 with a high TiO ratio represents _{the lowest R s of 2.04 Ω.} When conductive TiO is gradually transformed from Example 2 (TR-12) to Example 4 (TR-36) into r-TiO ₂ nanoneedles, the R _s value gradually increases. This result indicates that TiO promotes rapid charge transfer, suggesting that T-1 may have the highest pseudocapacitive performance. However, TR-20 has a larger R _ct (6.99 Ω) than T-1, and the high surface area (49.47 m ² /g) of TR-20 causes a high contact area between the electrode and the electrolyte.

The inclined straight line in the low frequency region represents the ion diffusion parameter of the electrolyte. Example 4 (TR-36), in which 100% was _{transformed into r-TiO 2} ^{nano-needle, exhibited the lowest Li +} ion diffusion because it had the lowest electrical conductivity. 28 is a table showing the impedance parameters determined for various samples. As a result, the excellence of Example 1 (TR-20) among various titanium oxide samples in which the phase ratio is controlled can be explained for the following reasons. First, compared to Preparation Example 1 (T-1), Example 1 having an optimum ratio of r-TiO ₂ ^{nanoneedles storing Li +} ions facilitates electrolyte-electrode contact, and the diffusion length of ^{Li + ions} Is shortened, and thus ^{, it is possible to well accommodate an enormous load caused by volume change due to Li +} ion intercalation and deintercalation, which can ensure stable specific capacity performance in a high-speed charge/discharge environment. Second, the TiO core effectively improves electrical conductivity (interface charge transfer), so that an electrochemical reaction can occur quickly. However, Example 4 (TR-36) having a fully modified r-TiO ₂ nanoneedle acts as a barrier to improper diffusion and migration of ^{Li +} ions with higher charge transfer resistance, resulting in more ^{Li + ions.} By making it irreversible, capacity and charging and discharging performance are limited.

7 shows the relationship between electrical conductivity (at 20 MPa), BET surface area, and specific capacity according to the quantitative ratio of TiO/r-TiO _2. The specific capacity of the prepared sample was calculated in the 10th and 100th cycles. Although the electrical conductivity is much lower than that of Preparation Example 1 (T-1) and Example 2 (TR-12), Example 1 (TR-20) with the largest surface area exhibits the highest total specific capacity. The high surface area will contribute to the high specific capacity of Example 1. Thus, based on the trade-off between electrical conductivity, surface properties, and _{the quantitative ratio of TiO/r-TiO 2} nanoneedles, Example 1 has the best electrochemical performance among several samples. This performance is close to the theoretical capacity of graphite (372 mA·h·g ⁻¹ ), and can be an important achievement for an in-situ synthesized titanium oxide anode active material. As a result, the job-sharing TR-20 composite exhibits superior capacity and excellent stability retention compared to other titanium oxide materials, and can overcome the low chemical diffusion and energy density limitations of conventional titanium oxide materials ( Fig. 29).

<Experimental Example 9> Comparison of Example 1 and Comparative Example 1

In addition, TR-20 of Example 1 was compared with Comparative Example 1 in _{which TiO and r-TiO 2 were simply mixed.} The performance for Comparative Example 1 is shown in FIG. 23.

Referring to FIG. 23, in the case of Comparative Example 1, _{intercalation and deintercalation reactions occurring in r-TiO 2} did not occur during charging and discharging, and showed a tendency to have the same pseudocapacitive behavior as in Preparation Example 1 (T-1). The first cycle specific capacity of Comparative Example 1 was 43.3 mAh·g ^-1 , and the 100th cycle had a specific capacity of ^{46.6 mAh·g -1.}

That is, it can be seen that when TiO and r-TiO ₂ are simply physically mixed, the same effect as in Example 1 cannot be obtained.

10 Reduced titania
20 titania nanostructures
100 complex

Claims (12)

Porous particles containing reduced titania represented by TiO _2-x; And
_{A titania (TiO 2} ) nanostructure formed at least partially on the surface of the porous particle;
Complex containing
(0.1≦x<2).
The method of claim 1,
The reduced titania is a composite, characterized in that TiO.
The method of claim 1,
The titania nanostructure is a composite, characterized in that the needle-shaped nanostructure.
The method of claim 1,
The titania is a complex, characterized in that the rutile (rutile) phase.
A lithium ion battery negative active material comprising the composite of claim 1.
Mixing titania (TiO ₂ ) and a metal to prepare a mixture (step 1);
Heat-treating the mixture to _{prepare reduced titania (TiO 2-x} ) (Step 2);
Etching the metal by treating the reduced titania with an acid (step 3); And
Forming titania on at least a portion of the surface of the reduced titania by acid-treating the reduced titania on which the metal has been etched (step 4);
Composite manufacturing method comprising a
(0.1≦x<2).
The method of claim 6,
In step 1, titania and metal are mixed in a weight ratio of 1:2 to 3:1.
The method of claim 6,
The method of manufacturing a composite, wherein the metal is at least one selected from the group consisting of Mg, Na, Li, Ca, Al, Zr, Ti, V, K, Rb and Cs.
The method of claim 6,
The method of manufacturing a composite, characterized in that the metal is a metal powder.
The method of claim 6,
The reduced titania is TiO, characterized in that the composite manufacturing method.
The method of claim 6,
The acid treatment of step 4 is a composite manufacturing method, characterized in that the 10 to 50 hours at a temperature of 50 ℃ to 150 ℃.
A lithium ion battery comprising the negative active material of claim 1.

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