鋰離子電池是目前廣泛應用的儲能裝置，具有高工作電壓、長循環壽命和低自放電等優點。層狀富鋰錳基正極材料，例如xLi2MnO3·(1-x)LiTMO2 (TM=Ni、Co、Mn)，則具有高能量密度、超過250 mAhg-1的放電容量和低成本等優勢。然而，由於其表面陰離子的氧化還原反應導致氧氣釋放，導致結構劣化，進而引起不可逆的容量損失、低首圈庫倫效率和長期循環下的電壓衰減。因此，表面包覆和元素摻雜改質被視為有效抑制結構劣化的策略之一。
本文係採用共沉澱法合成前驅物並加入鋰源後高溫燒結成Li1.4Mn0.6Ni0.2Co0.2O2.4正極材料，利用LiNbO3和LiNi0.5Mn1.5O4尖晶石塗層與Mg、F和Mg*F進行摻雜。從電化學性能可知，LRM@Nb、首次庫倫效率和2C放電容量為86.9% 和 149 mAhg-1；LRM@1NM為90.2%和146 mAhg-1； LRM-Mg為88.0%和172 mAhg-1；LRM-F為93.2%和173 mAhg-1；LRM-Mg*F為88.6%和155 mAhg-1。在相同條件下，LRM-Bare首次庫倫效率和2C放電容量僅為85.5%和126 mAhg-1。Nb和Mg元素改善電化學性能是歸因於更強的Nb-O和Mg-O鍵形成，穩定晶格結構，有效的抑制充/放電循環間的結構重排，使鋰離子具高可逆性。此外，擴大的Li層間距，有利於鋰離子的擴散。F元素的摻雜則可抑制高電壓下Li2MnO3的氧氣釋放，同時保持陰離子之氧化還原反應。然而，LRM@1NM尖晶石結構提供三維Li離子擴散路徑，形成完全開放的表面，使Li離子在電極/電解質介面處實現快速轉移。
長期的充放循環會使富鋰正極材料中過渡金屬遷移至鋰層導致結構由層狀相變為尖晶石結構，且過渡金屬元素也會溶解至電解質，使循環效能大幅下降。LRM-Bare、LRM-Mg和LRM-F在2 – 4.6 V截止電壓下，以0.2 C充電/0.5 C放電循環170次，循環保持率分別為81、91和84 %。由結果可知Mg改質後，由於鹼金屬離子一般佔據鋰層四面體位置，而Mg離子半徑接近Li離子，但大於Ni、Co和Mn的離子半徑，因此可減少過渡金屬元素佔據Li位置，降低陽離子混排。而F摻雜會取代O的位置，由於F−的最高佔據分子軌道低於O2−，因此可以加強陰陽離子的鍵合，抑制過渡金屬離子的遷移，減少氧的損失，進而提高結構的穩定性。

Lithium-ion batteries have the advantages of high working voltage, long cycle life and small self-discharge, and are currently widely used energy storage devices. Layered lithium-rich manganese-based cathode materials, namely χLi2MnO3˙(1-χ)LiTMO2 (TM=Ni, Co, Mn), have the advantages of high energy density, discharge capacity exceeding 250 mAhg-1, and low cost. However, because the anion redox on the surface leads to the release of oxygen, which degrades the structure, resulting in large irreversible capacitance, low first-cycle Coulombic efficiency, and voltage decay under long-term cycling, particle surface coating and element doping modification are considered to inhibit the structure. Effective strategy for degradation.
In this paper, the co-precipitation method is used to synthesize the precursor and sinter at high temperature after adding lithium source to Li1.4Mn0.6Ni0.2Co0.2O2.4 cathode material, doping with Mg, F and Mg*F using LiNbO3 and LiNi0.5Mn1.5O4 spinel coatings..From the electrochemical performance, the initial Coulombic efficiency and 2C discharge capacity of LRM-Nb, LRM@1NM, LRM@2.5NM, LRM-Mg, LRM-F and LRM-Mg*F are 86.92% and 149 mAhg-1, respectively, 90.28 % and 146 mAhg-1, 86.61% and 131 mAhg-1, 88.04% and 172 mAhg-1, 93.29% and 173 mAhg-1, 88.66% and 155 mAhg-1. Under the same conditions, the initial Coulombic efficiency and 2C discharge capacity of LRM-Bare are only 85.51% and 126 mAhg-1. The improved electrochemical performance of Nb and Mg elements is attributed to stronger Nb-O and Mg-O bond formation, stable lattice structure, effective suppression of structural rearrangement between charge/discharge cycles, and high reversibility of lithium ions. . In addition, the enlarged Li interlayer spacing facilitates the diffusion of Li ions. The doping of F elements can suppress the oxygen release of Li2MnO3 under high voltage, while maintaining the redox reaction of anions. However, the LRM@1NM and LRM@2.5NM spinel structures provide three-dimensional Li ion diffusion pathways, forming completely open surfaces, enabling rapid Li ion transfer at the electrode/electrolyte interface.
In lithium-rich cathode materials, during the long-term cycle process, the transition metal migrates to the lithium layer, causing the structure to change from a layered phase to a spinel structure, and the transition metal element dissolves into the electrolyte, which greatly reduces the cycle efficiency. LRM-Bare, LRM-Mg, and LRM-F were cycled 170 times at 0.2C charge and 0.5C discharge at a cut-off voltage of 2 – 4.6 V, and the cycle retention rates were 81.97%, 91.10% and 84.87%, respectively. It was observed from the results that after Mg modification, the alkali metal ions generally occupy the tetrahedral position of the lithium layer, while the Mg ion radius is close to that of the Li ion, but larger than the ionic radius of Ni, Co and Mn, so the transition metal elements cannot occupy the Li position. , to reduce cation mixing. In addition, after F doping, this improvement is attributed to the structural stabilization effect of F substituting O. Since the highest occupied molecular orbital of F− is lower than that of O2−, it can strengthen the bonding of anions and cations and inhibit the migration of transition metal ions. Reduce the loss of oxygen, thereby improving the stability of the structure.

摘要...i
Abstract...ii
誌謝...iv
目錄...v
表目錄...viii
圖目錄...ix
第一章 緒論...1
1.1 前言...1
1.2研究動機...1
第二章 文獻回顧...3
2.1 鋰離子電池介紹...3
2.1.1 鋰離子工作原理介紹...4
2.1.2 鋰離子電池的特性...4
2.2 正極材料...5
2.2.1 LiCoO2 (LCO)...6
2.2.2 LiNiO2 (LNO)...7
2.2.3 LiMn2O4 (LMO)...8
2.2.4 LiFePO4 (LFP)...8
2.2.5 LiNi1/3Co1/3Mn1/3O2 (NCM111)...9
2.3 負極材料...10
2.3.1 石墨碳...10
2.3.2 鋰鈦氧化物(LTO)...10
2.3.3 二氧化鈦(TiO2)...10
2.3.4 合金材料...11
2.4 隔離膜...11
2.5 電解質...11
2.6 正極材料合成方法...12
2.6.1 水熱法(hydrothermal methods)...12
2.6.2 噴霧造粒法(Spray Pyrolysis Method)...13
2.6.3 固相合成法(Solid-state Method)...14
2.6.4 溶膠-凝膠法(Sol-gel Method)...15
2.6.5 化學共沉澱法...16
2.7 富鋰正極材料...16
2.7.1 富鋰正極材料結構...17
2.7.2 F元素表面摻雜改質...18
2.7.3 Mg元素摻雜改質...22
2.7.4 Nb元素摻雜改質...24
2.7.5 Ti元素摻雜改質...28
2.7.6 Al和B元素共摻雜...32
2.7.7 Na和F元素共摻雜...35
2.7.8 Zr元素包覆...40
第三章 實驗方法...44
3.1 實驗材料與設備...44
3.1.1 實驗材料...44
3.1.2 實驗設備...44
3.2 實驗流程...48
3.2.1 以共沉澱法合成富鋰正極材料...49
3.2.2 富鋰正極材料表面元素摻雜...49
3.2.3 富鋰正極材料表面包覆改質...50
3.2.4 富鋰正極材料漿料製作...50
3.2.5 富鋰正極材料塗佈...51
3.2.6 富鋰正極材料極片加工...51
3.2.7. 鈕扣型半電池組裝...51
3.3 分析儀器...52
3.3.1 充/放電機(承德和新威)...52
3.3.2 X光繞射儀 (X-ray Diffraction)...54
3.3.3 場發射電子顯微鏡...55
3.3.4 雷射粒徑分析儀...55
3.3.5 恆電位儀...56
3.3.6 X射線光電子能譜儀...57
第四章 結果與討論...58
4.1 表面包覆與離子摻雜之Li1.4 Mn0.6Ni0.2Co0.2O2.4性能...59
4.1.1 XRD 晶體結構分析...59
4.1.2 SEM 粒子形貌分析...61
4.1.3 倍率性能(C-rate)測試...68
4.1.4 快速充電倍率性能(Quick charging C-rate)測試...75
4.1.5 循環性能(Cycle)測試...79
4.1.6 電壓衰減研究...81
4.1.7 微分容量/電壓曲線(dQ/dV)測試...85
4.1.8 X-ray photoelectron spectroscopy (XPS) 分析...87
4.1.9 交流阻抗(AC Impendence)測試...91
4.2 Mg/F共摻雜Li1.4Ni0.2Co0.2Mn0.6O2.4...93
4.2.1 XRD 晶體結構分析...93
4.1.2 SEM 粒子形貌分析...95
4.2.3 倍率性能(C-rate)測試...100
4.2.4 高溫循環性能(Cycle)測試...104
4.2.5 高溫電壓衰減研究...107
4.2.6 微分容量/電壓曲線(dQ/dV)分析...109
參考文獻...115
Extended Abstract...124

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