KR20240117711A - Nanofibers for Li-Se batteries comprising reducted graphene oxides, Manufactuaring method of the same, Cathodes composition for Li-Se batteries comprising the same and Li-Se batteries comprising the same - Google Patents

Abstract

Nanofibers for lithium-selenium secondary batteries of the present invention include reduced-graphene oxide (r-GO) and carbon structures; And selenium (Se) may be included on the surface of the structure.
The method for producing nanofibers for lithium-selenium secondary batteries of the present invention includes preparing a mixture by mixing graphene oxide nanosheets, a carbon source, and thermally decomposable polymer particles; Producing fiber by electrospinning the mixture; Carbonizing the fiber; etching the carbonized fiber; And it may include loading selenium into the etched fiber.
The cathode material composition for a lithium-selenium secondary battery of the present invention may include the nanofibers.
The lithium-selenium secondary battery of the present invention may include the nanofibers.

Description

Nanofibers for lithium-selenium secondary batteries containing reduced-graphene oxide, method for manufacturing same, cathode material composition for lithium-selenium secondary batteries containing same, and lithium-selenium secondary batteries containing same oxides, Manufacturing method of the same, Cathodes composition for Li-Se batteries comprising the same and Li-Se batteries comprising the same}

The present invention relates to nanofibers for lithium-selenium secondary batteries containing reduced-graphene oxide, a method of manufacturing the same, a cathode material composition for lithium-selenium secondary batteries containing the same, and a lithium-selenium secondary battery containing the same.

Lithium-ion battery (LIB) technology has made tremendous progress over the past few decades, making it possible to utilize LIBs to achieve their full capacity in terms of energy density. But for commercial applications that require much higher energy densities, such as large-scale grid storage or transportation systems, new battery systems are needed.

Accordingly, lithium-sulfur batteries (LSBs) have emerged with high energy density and abundant resources, but there are some fundamental problems, such as the insulating properties of elemental sulfur and discharge products, severe volume changes, and unstable Li cathode.

Selenium has attracted the attention of researchers because it is similar to sulfur in terms of electrochemical properties, has the advantages of high capacity of Se-based anodes, low rate of side reactions due to low reactivity of Se, and easier control during electrochemical processes. . However, like LSB, Li-Se batteries also have the problem of polyselenide dissolving during the electrochemical process. Additionally, low availability of bulk Se along with poor electronic and ion transport properties results in sub-expected electrochemical performance.

The purpose of the present invention is to provide a nanofiber for a lithium-selenium secondary battery with a novel structure, a method for manufacturing the same, a cathode material composition for a lithium-selenium secondary battery containing the same, and a lithium-selenium secondary battery containing the same.

In order to achieve the purpose of the present invention as described above,

Nanofibers for lithium-selenium secondary batteries of the present invention include reduced-graphene oxide (r-GO) and carbon structures; And selenium (Se) may be included on the surface of the structure.

The method for producing nanofibers for lithium-selenium secondary batteries of the present invention includes preparing a mixture by mixing graphene oxide nanosheets, a carbon source, and thermally decomposable polymer particles; Producing fiber by electrospinning the mixture; Carbonizing the fiber; etching the carbonized fiber; And it may include loading selenium into the etched fiber.

The cathode material composition for a lithium-selenium secondary battery of the present invention may include the nanofibers.

The lithium-selenium secondary battery of the present invention may include the nanofibers.

In the nanofibers of the present invention, the highly conductive reduced graphene oxide matrix acts as a self-supporting framework that improves structural integrity while providing a conductive path for rapid charge transfer, and the interconnected chain-like mesopores ensure efficient electrolyte penetration. The micropores provide impregnation with active materials, resulting in excellent electrochemical performance, including excellent rate performance and overwhelming long-term cycling stability when used as an anode for lithium-selenium secondary batteries.

Therefore, nanofibers can be used in various fields such as secondary batteries, medical devices, and catalysts. Preferably, nanofibers according to various embodiments of the present invention can be used as a cathode material for a lithium-selenium secondary battery.

Figure 1 is a schematic diagram for explaining a nanofiber and a method of manufacturing the same according to an embodiment of the present invention.
Figure 2 is a schematic diagram to explain P-rGO NF and P-rGO@Se NF without micropores.
Figure 3 shows FE-SEM images, TEM images, HR-TEM images, XRD patterns, N ₂ adsorption-desorption isotherms, BJH pore size distribution curves, and element point mapping images of P-rGO NFs and Bi-P-rGO NFs.
Figure 4 shows FE-SEM images, TEM images, HR-TEM images, SAED patterns, XRD patterns, N ₂ adsorption-desorption isotherms, and BJH pore size distribution curves of P-rGO@Se NF and Bi-P-rGO@Se NF. and element point mapping image.
Figure 5 is a graph showing the XPS results of Bi-P-rGO@Se NF, TG curves, and Raman spectra of Bi-P-rGO@Se NF and P-rGO@Se NF.
Figure 6 shows the results of measuring the electrochemical performance of Example 1 and Comparative Example 1 of the present invention.
Figure 7 shows the CV curve of the first cycle of Example 1 and Comparative Example 1 of the present invention and charge/discharge profiles at various cycle numbers.
Figure 8 is a schematic diagram of the results and reaction of Experimental Example 5 of the present invention.
Figure 9 is a graph and FE-SEM image showing the results of Experimental Example 6 of the present invention.
Figure 10 is a FE-SEM image of the anode of Comparative Example 1 after 100 ^th cycle at 0.5 C in Experimental Example 6 of the present invention.
Figure 11 is a digital image of the surface of a glass fiber paired with the anode of Example 1 and Comparative Example 1 in Experimental Example 6 of the present invention.

Hereinafter, various embodiments of this document are described with reference to the attached drawings. The examples and terms used herein are not intended to limit the technology described in this document to specific embodiments, and should be understood to include various modifications, equivalents, and/or substitutes for the examples.

Nanofibers according to various embodiments of the present invention are nanofibers with a novel structure that can be applied to lithium-selenium secondary batteries. The nanofibers of the present invention include reduced-graphene oxide (r-GO) and carbon structures; And selenium (Se) may be included on the surface of the structure.

The reduced-graphene oxide (r-GO) and carbon structure may be porous, and the reduced-graphene oxide (r-GO) and carbon structure include pores with a relatively large diameter and pores with a relatively small diameter. It may have bimodal porosity.

The selenium (Se) may be chain-type selenium (c-Se) located in the relatively small diameter pores.

The relatively large diameter pores may have an average diameter in the range of 30 nm to 60 nm, and the relatively small diameter pores may have an average diameter in the range of 1 nm to 3 nm, but are not limited thereto.

Hereinafter, a method for manufacturing nanofibers for lithium-selenium secondary batteries according to various embodiments of the present invention will be described.

A method of manufacturing nanofibers for lithium-selenium secondary batteries according to various embodiments of the present invention includes preparing a mixture by mixing graphene oxide nanosheets, a carbon source, and thermally decomposable polymer particles; Producing fiber by electrospinning the mixture; Carbonizing the fiber; etching the carbonized fiber; And it may include loading selenium into the etched fiber.

Before preparing the mixture, a step of dispersing the graphene oxide nanosheets may be further included.

The carbon source includes sucrose, dextrin, citric acid, ethylen glycol, poly ethylen glycol, PVP (Polyvinylpyrrolidone), PEDOT (Polyethylenedioxythiophene), PAN (Polyacrylonitrile), PAA (polyacrylic acid), PVA (polyvinylalcohol), PMMA (polymethyl methacrylate), PVDF (polyvinylidene fluoride), PVac (polyvinylacetate), PS (polystyrene), PVC (polyvinylchloride), PEI (polyetherimide), PBI (polybenzimidasol), PEO (polyethyleneoxide), PCL (poly e-caprolactone), PA-6 (polyamide- 6), it may include at least one selected from the group consisting of PTT (polytrimethylenetetraphthalate), PDLA (poly D,L-lactic acid), polycarbonate, polydioxanone, polyglicolide, and dextran.

The thermally decomposable polymer particles include polystyrene, polyvinyl acetate, polymethyl methacrylate, polyvinylpyrrolidone, polyvinyl alcohol, polycarbonate, polyester, polyetherimide, polyethylene, polyethylene oxide, polyurethane, and polyvinyl chloride. It may include at least one selected from the group consisting of These thermally decomposable polymer particles may decompose into gas during heat treatment to form voids. By forming a space, gas diffusion can occur smoothly in the next heat treatment step, and nanofibers with the desired structure can be obtained.

Before the etching step, a step of mixing with potassium hydroxide may be additionally included.

In the present invention, nanofibers with a novel structure can be manufactured in a simple manner through electrospinning and etching. In addition, nanofibers with this novel structure can be mass-produced, making them highly likely to be used industrially.

The cathode material for a lithium-selenium secondary battery according to various embodiments of the present invention may include the nanofibers.

Lithium-selenium secondary batteries according to various embodiments of the present invention may include the nanofibers.

Hereinafter, embodiments of the present invention will be described in detail with reference to the attached drawings.

1 is a schematic diagram showing the nanofiber synthesis mechanism of the present invention.

First, referring to ① in Figure 1, graphene oxide nanosheets and size-controlled PS nanobeads (average diameter = 100 nm) are homogeneously mixed with PVA polymer. Through the carbonization process at 600°C, the PVA polymer changes into amorphous carbon (AC), while the PS nanobeads are completely decomposed into gas products, forming chain-shaped mesopores with an average size of 42 nm (② in Figure 1). The presence of mesopores not only facilitates effective electrolyte penetration but also increases the active surface area in contact with the electrolyte. Additionally, GO nanosheets are converted into a highly conductive reduced graphene oxide (rGO) matrix, which acts as a self-supporting framework to improve the structural integrity of the nanofibers. The nanofibers obtained in this way were heat treated at 800°C, chemically activated using KOH, and then washed repeatedly with H ₂ O/HCl (1:1 vol%) (③ in FIG. 1). Through this process, micropores are formed along with mesopores. Finally, bimodal porous nanofibers can be obtained through a typical melt diffusion process in which selenium (Se) element is mixed in the obtained nanofiber and then heat treated to fill the micropores with Se (④ in Figure 1).

In the absence of micropores, Se accumulates in the form of bulky precipitates on mesopores, as can be seen in Figure 2. The presence of bulky Se causes insufficient electrolyte penetration and hinders the activity of highly active substances. Therefore, the synergistic effect of the structural advantages of the bimodal porous structure enables excellent electrochemical performance.

Single-pore porous reduced graphene oxide nanofibers (abbreviated as “P-rGO NF”) were obtained from composite fibers stabilized in N ₂ atmosphere at 600 °C for 3 hours, as shown in Figure 3 (a)-(g). got it The FE-SEM image in Figure 3 (a) shows a fiber shape with an average diameter of 400 nm. The fiber contains well-formed mesopores throughout the structure, as shown in the inset of Figure 3 (a). The TEM image in Figure 3 (b) confirms the FE-SEM results in the presence of chain-shaped mesopores (average diameter = 42 nm) uniformly formed in the fiber structure due to decomposition of PS nanobeads during carbonization. The HR-TEM image in Figure 3(c) confirms the presence of graphitic carbon in the fibers that make up the rGO matrix. During heat treatment on (002), (100), and (110) crystal planes, GO nanosheets were converted into rGO matrix. The well-formed rGO matrix acts as a self-supporting framework that increases the structural integrity of the nanofibers. Additionally, the XRD pattern in (d) of Figure 3 shows a broad peak centered at 2θ = 24°, which corresponds to the carbonaceous product of the nanofiber. The carbon content of the nanofibers comes from PVA polymer and rGO nanosheets after carbonization. The N ₂ adsorption-desorption isotherm of P-rGO NF in Figure 3(e) shows type IV behavior, indicating the presence of mesopores due to decomposition of PS nanobeads during the carbonization step. The presence of mesopores resulted in a-BET surface area of 344 m ² /g (pore volume: 0.59 cm ³ /g). In Figure 3(f), mesopores (average pore diameter 42 nm) of the nanofiber were also confirmed in the Barrett-Joyner-Halenda (BJH) pore size distribution curve. Additionally, the sharp peak that appears at 3.8 nm is due to the tensile strength effect of N ₂ desorption. The element-mapping image shown in (g) of Figure 3 confirms the uniform dispersion of C throughout the fiber length without any other elements being detected, indicating the carbonaceous structure of the fiber and Figure 3 This is consistent with the XRD result in (d).

Double porous rGO nanofibers (abbreviated as “Bi-P-rGO NF”) were obtained from P-rGO NF through a chemical activation process using KOH, followed by heat treatment at 800 °C in N2 atmosphere for 3 hours. The FE-SEM micrograph in Figure 3(h) shows that the fiber shape with an average diameter of ~300 nm is maintained even after the activation process. The decrease in fiber diameter is due to heat shrinkage at a high temperature of 800℃. Additionally, the mesopores of the fiber structure can be clearly seen in the high-magnification FE-SEM image in the inset of Figure 3 (h). The TEM image shown in Figure 3(i) clearly verifies the FE-SEM results. The presence of graphitized carbon due to the rGO matrix is also confirmed in the HR-TEM image (Figure 3(j)). The XRD pattern in (k) of Figure 3 shows a broad peak at 2θ = 25˚, indicating the carbonaceous characteristics of the prepared nanofibers even after KOH activation. To confirm the presence of micropores created by KOH activation, the N ₂ adsorption-desorption isotherm ((l) in Figure 3) and BJH pore size distribution ((m) in Figure 3) were confirmed. The N ₂ adsorption-desorption isotherm in (l) of Figure 3 shows type I isotherm characteristics, indicating the presence of micropores in the nanofiber. During the KOH activation process, potassium was inserted into the carbon matrix and then etched to form micropores. The coexistence of mesopores/micropores resulted in a BET surface area of 1178 m ² /g (pore volume: 1.30 cm ³ /g), which was much higher than that of P-rGO NF, which had a BET surface area of 344 m ² /g (pore volume: 0.59 cm ³ /g). caused the result. The BJH pore size distribution curve in (m) of Figure 3 shows a sharp peak at less than 3 nm and a wide peak at 40 nm, confirming that both mesopores and micropores exist in the nanofiber. The elemental dot mapping image shown in (n) of Figure 3 again confirms that carbon is evenly distributed in the porous nanofibers and that impurity species such as K were not detected, confirming that the same impurities were completely removed during the etching and cleaning process. there is.

Elemental Se impregnated Bi-P-rGO nanofibers (abbreviated as “Bi-P-rGO@Se NF”) and elemental Se-treated P-rGO nanofibers (abbreviated as “P-rGO@Se NF”) were melt-diffused. Obtained using the process. Both samples were heat treated by maintaining the mass ratio of sample:Se at 3:7. The physical properties of Se-infiltrated nanofibers are shown in Figure 4. Figure 4 (a) is a FE-SEM image that confirms the fiber shape of P-rGO@Se NF. In addition, the presence of bulky Se precipitates on the mesopores can be clearly confirmed through the inset images of (a) and (b) of Figure 4. The incomplete impregnation of Se is mainly due to the absence of micropores in the nanofibers. The TEM image in Figure 4(b) confirms the FE-SEM results with the clear presence of bulky Se nanoparticles (average 60 nm) that aggregate in the mesopores of P-rGO NF during the melt diffusion process. The HR-TEM image in Figure 4(c) shows clear lattice fringes corresponding to the (100) plane of elemental Se and the (002) plane of rGO, separated by 0.38 and 0.34 nm, respectively. The SAED pattern in Figure 4(d) also shows diffraction planes corresponding to single crystal Se nanoparticles. Additionally, in Figure 4(e), a sharp crystal peak can be seen in the XRD pattern, which clearly suggests the presence of Se element as a separate entity deposited on the mesopore. The N ₂ adsorption-desorption isotherm of P-rGO@Se NF ((f) in Figure 4) shows that bulky Se was largely deposited on the mesopores, resulting in a BET surface area of 18 m ² /g (pore volume 0.10 cm ³ /g). The elemental mapping image in Figure 4(g) shows that Se nanoparticles are present throughout the fiber surface with minimal impregnation.

The physical properties of Bi-P-rGO@Se NF are shown in (h) to (n) of Figure 4. The FE-SEM image in Figure 4(h) along with the inset image confirms that the porous one-dimensional fiber structure is well maintained even after Se impregnation. The TEM image in Figure 2 (i) shows a carbon framework (gray area) and chain-shaped interconnected mesopores (bright area). Additionally, unlike P-rGO@Se NF, bulky Se nanoparticles were not observed (Figure 4(i)), indicating that nano-sized Se penetrated well into the micropores (<3 nm). The HR-TEM image in Figure 4(j) mainly shows graphitized carbon within the rGO matrix, and the SAED and XRD patterns in Figure 4(k) and (l) indicate the presence of carbonaceous phase, nano-sized Se. This means that it has penetrated well into the porous carbon skeleton. The N ₂ adsorption-desorption isotherm in (m) of Figure 4 also confirms that nano-sized Se fills the inside of the porous carbon framework, which reduces the BET surface area to 67 m ² /g (pore volume 0.14 cm ³ /g). It is proven to have been done. The element dot mapping image in (n) of FIG. 4 shows that Se is uniformly distributed inside the porous nanofiber. Ultimately, the FE-SEM, TEM, XRD, and BET results of Bi-P-rGO@Se NF demonstrate that nano-sized Se was successfully impregnated inside the porous nanostructure compared with P-rGO@Se NF.

The chemical states and molecular environments of various elements of Bi-P-rGO@Se NF were confirmed through XPS (X-ray photoelectron spectroscopy) analysis, as shown in Figures 5 (a) to (c). The survey spectrum in Figure 5(a) shows photoelectron signals corresponding to C 1s, Se 3p and Se 3d orbitals. The ^high -resolution C ^1s It shows the peak. The very intense sp ² -hybridized C=C peak confirms the presence of carbonaceous products in the nanofibers in the form of graphitized carbon from the rGO matrix. The separated Se 3d spectrum in Figure 5(c) shows two main peaks attributed to Se 3d _5/2 (55.9 eV) and Se 3d _3/2 (56.7 eV) states, confirming the presence of infiltrated Se. . The Se content of the nanofibers was quantified using TG analysis (Figure 5(d)). Bi-P-rGO@Se NF and P-rGO@Se NF contain nearly 50.7 and 43.3 wt% of Se element, respectively. It can again be confirmed that the higher Se amount for Bi-P-rGO@Se NFs enables the use of high micropore volume during the impregnation process. Additionally, the rapid weight loss at higher temperature (375 °C) of Bi-P-rGO@Se compared to P-rGO@Se (~343 °C) is due to the chain type Se() filling the micropores of Bi-P-rGO@Se. c-Se) content. That is, there is a relatively strong chemical interaction between c-Se and carbonaceous species, and higher energy is required to break it.

The structural properties of Bi-P-rGO@Se NF and P-rGO@Se NF were investigated through Raman spectroscopy, and the results are shown in Figures 5 (e) and (f). The high-frequency Raman spectrum in (e) of Figure 5 shows the presence of D-band (1360 cm ^-1 ) and G-band (1600 cm ^-1 ) for both samples. The G-band originates from the presence of sp ² -hybridized carbon species in rGO, while the D-band corresponds to disordered amorphous carbon (AC) formed by decomposition of PVA. Additionally, the crystallinity of the carbonaceous product was determined using the relative intensity ratio ( I _D /I _G ) of the two bands. The I _D /I _G ratio was observed to be 1.03 for Bi-P-rGO@Se and 0.92 for P-rGO@Se, respectively, which may occur during the chemical activation process using KOH in Bi-P-rGO@Se. Indicates the presence of more defects. The low-wavelength Raman spectrum for Bi-P-rGO@Se NF in (f) of Figure 5 is triangular (t-Se, 237 cm ^-1 ), chain (c-Se, 255 cm ^-1 ), and ring (r- Se, 267 cm ^-1 ) shows three peaks corresponding to the type. In contrast, P-rGO@Se NF shows only a high intensity band centered at 237 cm ^-1 corresponding to t-Se due to the bulky Se filled in the mesopores. The highest peak intensity of c-Se at 255 cm ^-1 mainly arises from the presence of micropores that favor short-range alignment in Bi-P-rGO@Se. Therefore, the Raman spectrum results suggest that the presence of c-Se in Bi-P-rGO@Se provides superior electrochemical performance compared to P-rGO@Se due to high Se availability.

Hereinafter, the present invention will be explained in detail by examples. However, the following examples are only for illustrating the present invention and the present invention is not limited by the following examples.

Example 1

0.5 g of graphene oxide (GO) nanosheets were dispersed in 10 mL of ethanol by ultrasound for 3 hours to obtain a uniform suspension. Afterwards, 2.0 g of polyvinyl alcohol (PVA 2000, Kanto Chemical Co.) and 1.5 g of PS nanobeads (average diameter = 100 nm) were added to the suspension while stirring vigorously overnight to prepare an electrospinning solution.

The prepared electrospinning solution was placed in a plastic syringe pump (12 mL) equipped with a 25 gauge stainless steel needle. The electrospinning solution was jetted at a rate of 1.0 mL/h, and the fibers were collected on a rotating drum (180 rpm) covered with Al foil 15 cm away from the needle tip. The voltage applied between the collector and the needle was fixed at 20 kV.

The spun composite fibers were stabilized at 200 °C for 12 hours in air, and the stabilized fibers were carbonized at 600 °C for 3 hours at a heating rate of 5 °C/min under N ₂ atmosphere. The fiber was then mixed with potassium hydroxide (KOH, Samchun Co.). Fiber and potassium hydroxide were mixed at a ratio of 1:2 using a mortar and pestle, and the mixture was activated by heat treatment at 800°C for 3 hours at a temperature increase rate of 5°C/min under N ₂ atmosphere. The fiber was then etched with H ₂ O/HCL (1:1 vol%) for 24 hours, washed with distilled water several times, and dried.

The obtained etched fibers were mixed with elemental Se powder (99.99%, Sigma-Aldrich) at a mass ratio of 3:7 and subjected to a two-step heat treatment at 260 °C for 12 h, followed by heating at a heating rate of 5 °C/min under N ₂ atmosphere. Bi-P-rGO@Se NF was obtained by heat treatment at 350 °C for 1 hour.

To measure the electrochemical performance of the obtained Bi-P-rGO@Se NF, a CR2032 type coin cell was manufactured. The electrode was made by mixing 70% by weight of Bi-P-rGO@Se NF obtained as an active material, 20% by weight of carbon black (Super-P) as a conductive material, and 10% by weight of sodium carboxymethyl cellulose (CMC) as a binder in deionized water and then covered with Al foil. It was manufactured by coating on . The coated slurry was dried overnight at 60° C. in a hot air oven using a doctor blade. Afterwards, circular electrodes (diameter = 14 mm) with an active material loading of ~ 0.3 mg/cm ² were punched and transferred inside the glove box. Li metal and microporous polypropylene film were used as the counter electrode and separator, respectively. The electrolyte was prepared by dissolving 1 M LiPF ₆ in a 1:1 (v/v) mixture of ethylene carbonate and diethyl carbonate (EC/DEC). Coin cell assembly was performed inside a glove box filled with high-purity Ar gas to produce a Li-Se secondary battery.

Comparative Example 1

P-rGO@Se NF was manufactured in the same manner as in Example 1, except that it was not subjected to potassium hydroxide treatment or etching, and a Li-Se secondary battery containing it as an active material was manufactured in the same manner as in Example 1.

Experimental Example 1 - CV Curve

The cells assembled in Example 1 and Comparative Example 1 were subjected to CV testing at a scan rate of 0.1 mV/s in a voltage range of 1.0 to 3.0 V for the initial five cycles. The first CV scan in Figure 6(a) for Example 1 shows a low intensity peak at 2.31 V (indicated by a red arrow) corresponding to the r-Se to c-Se conversion. However, the peak at 2.31 V disappears during the subsequent reduction cycle due to complete conversion during the first cycle. Additionally, the sharp decrease peak located at 1.76 V is attributed to the formation of Li ₂ Se from c-Se. During the oxidation scan, a single peak is observed at 2.06 V where Li ₂ Se is converted back to Se, completing the 2Li + Se ↔ Li ₂ Se redox cycle. It should be noted that the conversion reaction is a solid-solid process due to the insolubility of lithium polyselenide in carbonate-based electrolytes. Additionally, the oxidation peak remains stable during subsequent cycles (Figure 7(a)).

However, the CV curve of Comparative Example 1 (FIGS. 6(a) and 7(b)) shows slightly different characteristics from the CV curve of Example 1. The irreversible peaks observed at 2.35 V and 2.20 V (marked by blue arrows in Figure 6(a)) correspond to the conversion of r-Se to c-Se and disappear during the subsequent reduction cycle as shown in Figure 7(b). . Additionally, two oxidation peaks at 1.9 V and 2.3 V are observed during the charging process, which is due to the stepwise conversion of Li ₂ Se to lithium selenide and then to elemental Se, respectively. Additionally, there is a noticeable difference in the shape and intensity of the redox peaks, which means that the redox reaction kinetics inside the cell are poor.

This phenomenon becomes more pronounced when considering the polarization potential (ΔV) of the nanofiber. The cell of Example 1 shows lower separation between redox peaks (ΔV=300 mV) compared to the cell of Comparative Example 1 (ΔV=390 mV). After the first cycle, the decrease peaks of both samples ((a) and (b) in Figure 7) shift significantly toward higher potentials, mainly due to the deformation of the anode during the lithiation process. In addition, the smaller the voltage shift of Example 1 compared to Comparative Example 1, the better electrochemical activity was shown in the sample of Example 1. The new shoulder peak at 1.7 V (marked with '*' in Figure 7(a)) corresponds to the reaction of lithium with two Se molecules with different chain lengths in the reduction scan after 2 ^nd cycle for the sample of Example 1. appears in

The resulting well-overlapping and symmetric CV curves indicate a highly reversible redox process. Additionally, the CV curve of the anode of Example 1 shows a higher current density than that of the anode of Comparative Example 1, which clearly indicates improved redox reaction kinetics inside the cell compared to the anode of Comparative Example 1.

Experimental Example 2 - Initial charge/discharge profile

The CV curve was further verified using the charge-discharge profile obtained at C-rate 0.5, as shown in (b) of Figure 6. The charge-discharge profile shows a voltage plateau that coincides with the CV curve. For example, the anode of Example 1 exhibits a long discharge voltage plateau at about 1.9 V, corresponding to the conversion of elemental Se to Li ₂ Se. Similarly, at 1.95 V, this charge plateau coincides with the conversion of Li ₂ Se to Se. In contrast, the anode of Comparative Example 1 exhibited a short charge-discharge voltage plateau at 1.81 (Li ₂ Se → Se) and 1.60 V (Se → Li ₂ Se), indicating that the discharge capacity value was lower than that of the anode of Example 1. it means. Additionally, the poor discharge plateau of the anode of Comparative Example 1 at about 2.25 V is due to the conversion of r-Se to c-Se. The positive electrode of Example 1 and the positive electrode of Comparative Example 1 showed initial discharge capacities of 603 mA·h/g and 408 mA·h/g, respectively. Additionally, a relatively high coulombic efficiency of 86% was observed for the anode of Example 1 compared to 47% for the anode of Comparative Example 1. The large initial capacity loss or low initial coulombic efficiency of the positive electrode of Comparative Example 1 is due to the presence of bulky Se, which is then dissolved to form lithium polyselenide. However, the anode of Example 1 with chain-type Se loaded in micropores alleviates polyselenide dissolution, limiting the multi-step reaction between Li and Se to a single-step reaction.

The high discharge capacity of the positive electrode of Example 1 can be attributed to the high availability of chain-type Se as the active material in the micropores. In contrast, the low discharge capacity of the positive electrode of Comparative Example 1 may be due to the presence of bulky Se deposits that remain unused during the redox process, as evidenced by the short charge-discharge plateau. Overall, the CV and initial charge-discharge profiles demonstrate that the synergistic effect of the highly porous and conductive framework in the anode of Example 1 improves electrochemical performance.

Experimental Example 3 - Cycle characteristics

At 0.5 C-rate, the positive electrode of Example 1 shows an initial discharge capacity of 603 mA·h/g compared to the positive electrode of Comparative Example 1 (408 mA·h/g) (FIG. 6(c)). As cycling progresses, the positive electrode of Example 1 exhibits a discharge capacity of 441 mA·h/g (73% capacity retention) at the end of 800 ^th continuous charge-discharge cycle, and in contrast, the positive electrode of Comparative Example 1 exhibits a discharge capacity of 441 mA·h/g (73% capacity retention) throughout cycling. shows low discharge capacity. The positive electrode of Comparative Example 1 showed a discharge capacity of 149 mA·h/g (37% capacity maintained) after 560 cycles. In addition, Comparative Example 1 shows an average capacity decrease of 0.11% per cycle, but Example 1 shows an average capacity decrease of 0.03% per cycle, further verifying the structural advantage.

The high coulombic efficiency of 99% throughout the cycling process further explains the high reversibility of the redox processes inside the cell. These results are even more evident when considering charge-discharge profiles at various cycle numbers. The high discharge capacity of the positive electrode of Example 1 is mainly due to the long voltage plateau ((c) in Figure 7) that is better formed than the voltage plateau ((d) in Figure 7) of the positive electrode in Comparative Example 1. The cycling results are compared The high active material utilization can again be seen in the anode of Example 1 compared to the anode of Example 1, where the low capacity means a high proportion of unutilized active material (large Se deposits), which results in slowing down inside the cell. Resulting in reaction kinetics.

Experimental Example 4 - Rate characteristics

The anode of Example 1 exhibits discharge capacities of 613, 454, 428, 407, 385, 357, and 326 mA·h/g at C-rates of 0.5, 1.0, 3.0, 5.0, 7.0, 10.0, and 20.0, respectively (Figure (d)) of 6. In addition, when the current is restored to 0.5 C, it shows a stable discharge capacity of 440 mA·h/g. In contrast, the positive electrode of Comparative Example 1 shows low discharge capacities of 400, 169, 143, 117, 101, 85, and 50 mA·h/g at the corresponding C-rate mentioned above. The positive electrode of Example 1 exhibits high discharge capacity even at a high C-rate of 20.0 C, and its bimodal porous structure including mesopores and micropores ensures high active material utilization in addition to efficient electrolyte penetration or wettability of the electrode. Additionally, the highly conductive rGO matrix not only increases the structural integrity of the electrode but also provides numerous conductive channels that support fast charge transfer redox kinetics.

Experimental Example 5 - Li ion diffusion coefficient

To authenticate the excellent redox dynamics inside the Li-Se cell, _the Li ion diffusion coefficient ( DLi ⁺ ) was examined by examining the CV data in the voltage range from 1.0 to 3.0 V and various scan rates from 0.05 to 1.0 mV/s, as shown in Figure 8. shown in CV curves obtained at various voltage scan rates show typical Li-Se redox peaks, suggesting electrochemical processes occurring between elemental Se and Li ₂ Se through solid-solid processes inside the cell. Even at a high voltage scan rate of 1.0 mV/s, a higher current intensity value is observed at the anode of Example 1 (Figure 8(a)) than the anode of Comparative Example 1 (Figure 8(c)), which is due to It refers to a fast electrochemical process. The Randles-Sevcik equation was additionally applied to the D _Li ⁺ measurements of all cells, and is shown in <Equation 1> below.

<Equation 1>

Here, I _p is the redox peak current, n is the number of electrons involved in the redox reaction (n=2), A is the electrode surface area (cm ² ), C _Li is the Li-ion concentration (mol/L), υ is the voltage refresh rate (V/s).

I _p vs. obtained for the anode of Example 1 and the anode of Comparative Example 1. The υ ^0.5 plot is shown in (b) and (d) of Figures 8, and the D _Li ⁺ values calculated for all Li-Se cells are shown in Table 1 below.

Sample D _Li ⁺ (cm ² /s) × 10 ^-9 Average D _Li ⁺ (cm ² /s) × 10 ^-9 Peak Ⅰ Peak Ⅱ Peak Ⅲ Example 1 8.06 3.40 6.50 5.99 Comparative Example 1 0.61 0.38 0.65 0.55

The Li-Se battery using the positive electrode of Example 1 had a higher lithium ion diffusion coefficient value (5.99 × 10 ^-9 cm ² /s) than the battery using the positive electrode of Comparative Example 1 (0.55 × 10 ^-9 cm ² /s). s), indicating excellent diffusion dynamics inside the cell. The high diffusion coefficient results confirm that the unique one-dimensional design strategy with chain-like Se filled in the micropores results in kinetically favorable redox dynamics, enhancing the overall improvement of electrochemical performance (Figure 8). (e)-①). In contrast, the high proportion of unused active material (a large amount of Se precipitate) in the anode of Comparative Example 1 results in slow reaction kinetics inside the cell, resulting in poor electrochemical properties ((e)-② in Figure 8). .

Experimental Example 6 - Electrochemical Impedance Spectroscopy (EIS)

The results of confirming the reaction dynamics inside the Li-Se cell through EIS are shown in Figure 9. The impedance of all Li-Se cells was measured in the fresh and uncycled state and also in the fully charged state at different cycle numbers while cycling at 0.5 C. Compared to the Li-Se battery using the anode of Comparative Example 1, the Li-Se battery using the anode of Example 1 exhibits the lowest charge transfer resistance (R _ct ) throughout the cycling process. For example, a fresh cell using the anode of Example 1 exhibited an R _ct value of 172 Ω (Figure 9(a)), which rapidly decreased to 54 Ω after 1 ^st cycle (Figure 9(b)). ) It mostly remains stable. After 100 ^th cycle, the Li-Se cell using the anode of Comparative Example 1 shows a higher R _ct value (75 Ω) compared to the Li-Se cell using the anode of Example 1 (FIG. 9(c)). . These results confirm that the structural advantage of Example 1 not only ensures a kinetically favorable redox process inside the Li-Se cell, but also improves the structural integrity of the Se electrode during long-term cycling. The FE-SEM photograph of the anode of Example 1 (FIG. 9(d)) and the FE-SEM photograph of the anode of Comparative Example 1 (FIG. 10) obtained after 100 ^th cycle at 0.5 C clearly prove the above claim. . The one-dimensional fibrous morphology of the positive electrode of Example 1 remained intact compared to Comparative Example 1, illustrating the excellent structural robustness of the as-prepared positive electrode, which not only allows for higher active material impregnation but also allows for lithium polyselenium during the electrochemical process. Accommodates large volume changes with cargo securing. Digital images of the glass fiber surfaces facing the lithium cathode and paired with the anodes of Example 1 and Comparative Example 1, respectively, are shown in Figure 11 and further analyzed for polyselenide immobilization. Referring to FIG. 11, compared to the dark color of the glass fiber used with the anode of Comparative Example 1 ((b) in FIG. 11), the white color of the glass fiber paired with the anode of Example 1 ((b) of FIG. 11) a)) clearly indicates better capture and reuse of polyselenide.

To obtain a high-performance Bi-P-rGO@Se anode, utilizing one-dimensional nanofibers composed of bimodal pores and a highly conductive rGO matrix as the active material host results in improved electrochemical performance, including excellent rate performance and overwhelming long-term cycling stability. This has been proven.

So far, the present invention has been examined focusing on its preferred embodiments. A person skilled in the art to which the present invention pertains will understand that the present invention may be implemented in a modified form without departing from the essential characteristics of the present invention. Therefore, the disclosed embodiments should be considered from an illustrative rather than a restrictive perspective. The scope of the present invention is indicated in the claims rather than the foregoing description, and all differences within the equivalent scope should be construed as being included in the present invention.

Claims (13)

reduced-graphene oxide (r-GO) and carbon structures; and
A nanofiber for a lithium-selenium secondary battery, characterized in that it contains selenium (Se) on the surface of the structure. According to paragraph 1,
A nanofiber for a lithium-selenium secondary battery, wherein the reduced-graphene oxide (r-GO) and carbon structure are porous. According to paragraph 2,
The reduced-graphene oxide (r-GO) and carbon structures are nanofibers for lithium-selenium secondary batteries, characterized in that they have bimodal porosity including pores of relatively large diameter and pores of relatively small diameter. According to paragraph 3,
The selenium (Se) is a lithium-selenium secondary battery nanofiber, characterized in that chain-type selenium (c-Se) located in the relatively small diameter pores. According to paragraph 3,
A nanofiber for a lithium-selenium secondary battery, wherein the relatively large diameter pores have an average diameter in the range of 30 nm to 60 nm. According to paragraph 3,
A nanofiber for a lithium-selenium secondary battery, wherein the relatively small diameter pores have an average diameter in the range of 1 nm to 3 nm. Preparing a mixture by mixing graphene oxide nanosheets, a carbon source, and pyrolytic polymer particles;
Producing fiber by electrospinning the mixture;
Carbonizing the fiber;
etching the carbonized fiber; and
Loading selenium into the etched fiber;
Method for producing nanofibers for lithium-selenium secondary batteries comprising. In clause 7,
Before preparing the mixture,
A method of producing nanofibers for lithium-selenium secondary batteries, characterized in that it further comprises the step of dispersing the graphene oxide nanosheets. In clause 7,
The carbon source includes sucrose, dextrin, citric acid, ethylen glycol, poly ethylen glycol, PVP (Polyvinylpyrrolidone), PEDOT (Polyethylenedioxythiophene), PAN (Polyacrylonitrile), PAA (polyacrylic acid), PVA (polyvinylalcohol), PMMA (polymethyl methacrylate), PVDF (polyvinylidene fluoride), PVac (polyvinylacetate), PS (polystyrene), PVC (polyvinylchloride), PEI (polyetherimide), PBI (polybenzimidasol), PEO (polyethyleneoxide), PCL (poly e-caprolactone), PA-6 (polyamide- 6), Manufacturing of nanofibers for lithium-selenium secondary batteries, characterized in that they contain at least one selected from the group consisting of PTT (polytrimethylenetetraphthalate), PDLA (poly D,L-lactic acid), polycarbonate, polydioxanone, polyglicolide, and dextran. method. In clause 7,
The thermally decomposable polymer particles include polystyrene, polyvinyl acetate, polymethyl methacrylate, polyvinylpyrrolidone, polyvinyl alcohol, polycarbonate, polyester, polyetherimide, polyethylene, polyethylene oxide, polyurethane, and polyvinyl chloride. A method for producing nanofibers for lithium-selenium secondary batteries, comprising at least one selected from the group consisting of: In clause 7,
A method of producing nanofibers for lithium-selenium secondary batteries, characterized in that it further comprises the step of mixing with potassium hydroxide before the etching step. A cathode material composition for a lithium-selenium secondary battery comprising the nanofibers according to any one of claims 1 to 6. A lithium-selenium secondary battery comprising the nanofibers according to any one of claims 1 to 6.